《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Fabrication and Characterization of an Ultra-High-Temperature

J.Am. Ceran.Soe,90o3320-3322(2007) journal Fabrication and Characterization of an Ultra-High-Temperature Carbon Fiber-Reinforced ZrB2-SiC Matrix Composite Sufang Tang, Jingyi Deng, Shijun Wang, and Wenchuan Liu Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China ed ZrBr-SiC matrix composite der having an average particle size of 1.5 um. Alternatively was fabricated infiltration of si were built up, followed by needle punching at the Z-direction a and the two successive weftless plies were oriented at an angle of of 148 MPa. a ess of 5.6 MPa m", and a good 900. The fiber and ZrB2 volume contents were 24.9% and 3.9% oxidation and al respectively. Second, the fiber-powder preform was clamped in a graphite clamp in order to wipe off the fiber size and maintain hape. The processing temperature was increased to L. Introduction 1200C and maintained at this temperature for 2 h in an argon ZaRcicatio ns toride is an m at iale w pa stila r interesta tos. dated by a small amount of pyrocarbon in order to produce such as engine cowl inlets, wing leading edges, and nosecaps, obtaining from pyrolysis of natural gas by 10 h at 1000CA because its high melting point coupled with the ability to form an isothermal CVI apparatus. Finally, the green composite was refractory oxide scale gives the material the capability to with- tand temperatures in the 1900%-2500oC range. 'Addition of rapidly densified by the Sic matrix through a novel HCVI tech- Sic has been proved to enhance oxidation resistance and limits nique, described elsewhere. Density of the C/ZrB2-SiC com- diboride grain growth. Although the material ZrBr-SiC is be- posite was measured by the Archimedes method. Flexural ing developed, it is less mature at this time and is still facing tests were carried out according to a three-point bending meth- od, using 80 mm x 10 mm x 6 mm-size-specimens with a span many barriers. First, the poor sinterability and bulk-forming equal to 70 mm and loaded at I mm/min. Young's modulus techniques limit a feasible scale-up of dense components with complex shapes at affordable manufacturing conditions because was obtained from the slope of the flexural curve at the initial of its strongly covalent bond. linear domain. Fracture toughness was measured on samples of In addition, the low fracture toughness and the poor thermal 45 mm x 10 mm x 4.8 mm through a single-edge notch beam shock resistance are still of concern although carbon is used (SENB)test using a 40 mm span, a 0.2 mm notch width, and a 5 mm notch depth with a cross-head speed of 0.05 mm/min that the zrb-sic material will benefit from the incorporation of Five and three samples were measured for the flexural and the a fiber reinforcement phase in order to improve its fracture toughness, impart an acceptable level of thermal shock, and low C/OrB -SiC er its density. Carbon fiber is an obvious and attractive candidate mm×7 mm were carried out at I000°,1200°,andl400°cin for this role, because of its very high specific strength and stiff air in an electrical furnace. The samples were exposed to 5-min ness, its thermodynamic compatibility with Zr.-SiC under 3000 xidation at the selected temperature, and then taken out K, and its high-temperature endurance beyond 2000 K the furnace and kept at room temperature for 5 min. Then the To the best of the authors'knowledge, no open literature mples were placed into the furnace again for the second about carbon fiber-reinforced ZrBr-SiC ceramic composites has thermal cycle. Seven circles were carried out and the masses of been published. In this paper, a novel processing approach, in the samples were recorded before and after each circle. Three samples for each temperature were used. Ablation was per which carbon fibers were incorporated into the ZrB2 matrix us- formed by an arc-heated wind tunnel test. During the ablation ing a fiber-powder molding technique and then the fiber-powder preform densified by heaterless chemical vapor infiltration the specimen was exposed to the flame with high temperature (HCVD) for infiltration of Sic matrix, was employed to fabri (over 2000C)and hypersonic speed (3.5 Mach number) air. cate the C/zrBr Sic composite. The composite is characterized The total enthalpy was 10.3-14.3 MJ/kg and the ablated period ccording to its microstructure and its mechanica was 650s. more details are described elsewhere. Microstructure nd ablation properties and morphology of the C/Zrbx-Sic composite before and after oxidation were observed by a scanning electron microscope SEM) Il. Experimental Procedure In this work, a four-stage process was used to fabricate the C/ II. Results and discussion ZrBxSiC composite First, as-received 12-k tows of Toray T7 (Tokyo, Japan) carbon fibers were composited with Zr B2 pow Figure I displays the typical backscattered electron microscopy mages(BSEM) of the polished cross sections of the as-pro- cessed C/ZrBr-SiC composite It can be observed from Fig. 1(a) Contributing editor that the cross-section area includes 90 long fibers SiC matrix short fibers-zrB, particles-SiC matrix, and 0 long fibers-SiC matrix, and these layers alternately occur in the composite ch 10, 2006; approved May 7, 2007 The bright-white ZrB, particles can be clearly identified and om correspondence should be addressed. e-mail: jydengaimr accn then mainly distribute in the interlayer regions between the 3320

Fabrication and Characterization of an Ultra-High-Temperature Carbon Fiber-Reinforced ZrB2–SiC Matrix Composite Sufang Tang, Jingyi Deng,w Shijun Wang, and Wenchuan Liu Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China A novel carbon fiber-reinforced ZrB2–SiC matrix composite was fabricated by heaterless chemical vapor infiltration through infiltration of SiC matrix into a carbon fiber-ZrB2 powder pre￾form. The C/ZrB2–SiC composite presented a flexural strength of 148 MPa, a fracture toughness of 5.6 MPa . m1/2, and a good oxidation and ablation resistance. I. Introduction ZIRCONIUM diboride is a material of particular interest for applications to hypersonic vehicles with sharp aerosurfaces, such as engine cowl inlets, wing leading edges, and nosecaps, because its high melting point coupled with the ability to form refractory oxide scale gives the material the capability to with￾stand temperatures in the 19001–25001C range.1–3 Addition of SiC has been proved to enhance oxidation resistance and limits diboride grain growth.4 Although the material ZrB2–SiC is be￾ing developed, it is less mature at this time and is still facing many barriers. First, the poor sinterability and bulk-forming techniques limit a feasible scale-up of dense components with complex shapes at affordable manufacturing conditions because of its strongly covalent bond.5 In addition, the low fracture toughness and the poor thermal shock resistance are still of concern although carbon is used as an additive to improve the thermal stress resistance.1 Thus, it is likely that the ZrB2–SiC material will benefit from the incorporation of a fiber reinforcement phase in order to improve its fracture toughness, impart an acceptable level of thermal shock, and low￾er its density. Carbon fiber is an obvious and attractive candidate for this role, because of its very high specific strength and stiff￾ness, its thermodynamic compatibility with ZrB2–SiC under 3000 K, and its high-temperature endurance beyond 2000 K. To the best of the authors’ knowledge, no open literature about carbon fiber-reinforced ZrB2–SiC ceramic composites has been published. In this paper, a novel processing approach, in which carbon fibers were incorporated into the ZrB2 matrix us￾ing a fiber-powder molding technique and then the fiber-powder preform densified by heaterless chemical vapor infiltration (HCVI) for infiltration of SiC matrix, was employed to fabri￾cate the C/ZrB2–SiC composite. The composite is characterized according to its microstructure, and its mechanical, oxidation, and ablation properties. II. Experimental Procedure In this work, a four-stage process was used to fabricate the C/ ZrB2–SiC composite. First, as-received 12-k tows of Toray T700 (Tokyo, Japan) carbon fibers were composited with ZrB2 pow￾der having an average particle size of 1.5 mm. Alternatively stacked layers of powder, short-cut-fiber webs, and weftless plies were built up, followed by needle punching at the Z-direction; and the two successive weftless plies were oriented at an angle of 901. The fiber and ZrB2 volume contents were 24.9% and 3.9%, respectively. Second, the fiber-powder preform was clamped in a graphite clamp in order to wipe off the fiber size and maintain the preform shape. The processing temperature was increased to 12001C and maintained at this temperature for 2 h in an argon atmosphere. In the third stage, the treated preform was consol￾idated by a small amount of pyrocarbon in order to produce an interface and increase thermal conductivity of the preform, obtaining from pyrolysis of natural gas by 10 h at 10001C in an isothermal CVI apparatus. Finally, the green composite was rapidly densified by the SiC matrix through a novel HCVI tech￾nique, described elsewhere.6 Density of the C/ZrB2–SiC com￾posite was measured by the Archimedes method. Flexural tests were carried out according to a three-point bending meth￾od, using 80 mm
10 mm
6 mm-size-specimens with a span equal to 70 mm and loaded at 1 mm/min. Young’s modulus was obtained from the slope of the flexural curve at the initial linear domain. Fracture toughness was measured on samples of 45 mm
10 mm
4.8 mm through a single-edge notch beam (SENB) test using a 40 mm span, a 0.2 mm notch width, and a 5 mm notch depth with a cross-head speed of 0.05 mm/min. Five and three samples were measured for the flexural and the SENB tests, respectively. Isothermal oxidation tests of the C/ZrB2–SiC composite with a dimension of 10 mm
10 mm
7 mm were carried out at 10001, 12001, and 14001C in air in an electrical furnace. The samples were exposed to 5-min oxidation at the selected temperature, and then taken out of the furnace and kept at room temperature for 5 min. Then the samples were placed into the furnace again for the second thermal cycle. Seven circles were carried out and the masses of the samples were recorded before and after each circle. Three samples for each temperature were used. Ablation was per￾formed by an arc-heated wind tunnel test. During the ablation, the specimen was exposed to the flame with high temperature (over 20001C) and hypersonic speed (3.5 Mach number) air. The total enthalpy was 10.3–14.3 MJ/kg and the ablated period was 650 s. More details are described elsewhere.7 Microstructure and morphology of the C/ZrB2–SiC composite before and after oxidation were observed by a scanning electron microscope (SEM). III. Results and Discussion Figure 1 displays the typical backscattered electron microscopy images (BSEM) of the polished cross sections of the as-pro￾cessed C/ZrB2–SiC composite. It can be observed from Fig. 1(a) that the cross-section area includes 901 long fibers–SiC matrix, short fibers–ZrB2 particles–SiC matrix, and 01 long fibers–SiC matrix, and these layers alternately occur in the composite. The bright-white ZrB2 particles can be clearly identified and then mainly distribute in the interlayer regions between the two K. Kakimoto—contributing editor w Author to whom correspondence should be addressed. e-mail: jydeng@imr.ac.cn Manuscript No. 22984. Received March 10, 2006; approved May 7, 2007. Journal J. Am. Ceram. Soc., 90 [10] 3320–3322 (2007) DOI: 10.1111/j.1551-2916.2007.01876.x r 2007 The American Ceramic Society 3320

October 2007 Communications of the American Ceramic Society 3321 (a) 0° longfibers SiC matrix 200m Fig 1. Backscattered electron microscopy images of the cross section of the C/zrBr-SiC composite (a) The continuous and alternate distribution of 90 long fibers-SiC matrix, short fibers-ZrB2 particles-SiC matrix, and 0 long fibers-SiC matrix; (b)the distribution of carbon fibers, ZrB2 particles, and SiC matrix in the interlayer. successive weftless plies. Further observation of the interlayer presented in Fig. I(b) shows that the dense Sic matrix almost ompletely fills up the inter fiber pores and integrates the ZrB particles and the carbon fibers Table I summarizes the characteristics of the CzrBrSiC osite. For comparison, the properties of the C/SiC com- with a same fiber stacked-up mode and a fiber volum content of 30.6% are also listed, also prepared by the HCVI technique. The C/zrBy-Sic composite has a slightly lower 3458 ding 1% while it has a pronouncedly higher gs modulus of 100. 4%, although it possesses a lower content of 18.3% compared with the C/SiC composite. The main reasons for the different bending strengths should be at- ributed to the different fiber contents, while the ZrB2 particles 0.00 with a high modulus, occurring in the C/zrBr-SiC composite 203040.50.60.70.8091.0 ould mainly contribute to the discrepancy of Youngs moduli The fracture toughness value(5.6 MPa.m")of the C/ZrBr-SiC Fig. 2. Curves of load and fracture toughness composite is largely higher, compared with a value of 2.8 and 2.3 displacement of MPa.m" for monolithic ZrB2 and SiC, respectively. -Its the C/zrB-SiC composite fracture behavior, as shown in Fig. 2. becomes somewhat duc- the surface morphologies after the seventh cycle of oxidation. At tile, i.e., cracks gradually propagate in the composite and no 000C, a continuous and densified glass in the form of many catastrophic breaking occurs while the monolithic cerami particle-shaped aggregates is found for the C/ZrBxSiC compos- show a linear increase in deflection with increasing load and te while much powder-shaped Sic debris appears on the surface suddenly breaks(catastrophic breaking) of the C/Sic composite. At 1200 and 1400.C, the aggregates Figure 3 exhibits the loss curves of the C/zrBx-Sic continuously increase with the temperature increasing in the Cl composite after exposure to seven cycles of 5-min static furnace ZrBr-SiC composite, resulting in the formation of the more-uni- oxidation at 1000, 1200, and 1400.C. The interpretation of the form oxide scale. In the case of the C/SiC composite, many data needs to be tempered by the fact that the specimens were cut hemisphere-shaped granules with big and small bulks are formed from a larger plate without providing any oxidation protection at 1200C. indicating that a small amount of Sic matrix was for the cut edges, open pores, and microcracks. This subjects the fibers to rapid oxidation attack since the naked fibers are greatly activated and many microcracks and open pores supply the ox- ■一C/sc1000° ygen diffusion channels. As expected, the C/zrBr-SiC composite b-c/sic1200° possesses a relatively high mass loss rate while it still presents an ▲- C/SiC1400° bvious advantage compared with the c/sic composite. From v-C/ZrB2-SiC1000°c mass loss rates with oxidation. In addition, the m g trend in the Fig. 3, it can be clearly seen that there is a decreas CZrB2SiC1200°C ass loss rates cZrB2Sic1400°C are similar at 1200 and 1400.C and are obviously higher than that at 1000C for the C/zrBr-SiC composite. Figure 4 shows 20 Table I. Characteristics of C/ZrBx-SiC Composite C/ZrBr-SiC Fiber content(vol %) 24.9 Density(g/cm) Flexural strength(MPa) 148+12 Youngs modulus(GPa) 58.0+5.0 Fracture toughness(MPa."2)5.6+1.6 Ablated surface temperature(C) 1188+5 Fig 3. Curves of mass loss rate versus temperature of the C/zrBr-Sic and the C/SiC composites

successive weftless plies. Further observation of the interlayer presented in Fig. 1(b) shows that the dense SiC matrix almost completely fills up the inter fiber pores and integrates the ZrB2 particles and the carbon fibers. Table I summarizes the characteristics of the C/ZrB2–SiC composite. For comparison, the properties of the C/SiC com￾posite with a same fiber stacked-up mode and a fiber volume content of 30.6% are also listed, also prepared by the HCVI technique.6 The C/ZrB2–SiC composite has a slightly lower bending strength of 14.1% while it has a pronouncedly higher Young’s modulus of 100.4% , although it possesses a lower fiber content of 18.3% compared with the C/SiC composite. The main reasons for the different bending strengths should be at￾tributed to the different fiber contents, while the ZrB2 particles with a high modulus, occurring in the C/ZrB2–SiC composite, should mainly contribute to the discrepancy of Young’s moduli. The fracture toughness value (5.6 MPa m1/2) of the C/ZrB2–SiC composite is largely higher, compared with a value of 2.8 and 2.3 MPa m1/2 for monolithic ZrB2 and SiC, respectively.8–9 Its fracture behavior, as shown in Fig. 2, becomes somewhat duc￾tile, i.e., cracks gradually propagate in the composite and no catastrophic breaking occurs while the monolithic ceramics show a linear increase in deflection with increasing load and suddenly breaks (catastrophic breaking). Figure 3 exhibits the mass loss curves of the C/ZrB2–SiC composite after exposure to seven cycles of 5-min static furnace oxidation at 10001, 12001, and 14001C. The interpretation of the data needs to be tempered by the fact that the specimens were cut from a larger plate without providing any oxidation protection for the cut edges, open pores, and microcracks. This subjects the fibers to rapid oxidation attack since the naked fibers are greatly activated and many microcracks and open pores supply the ox￾ygen diffusion channels. As expected, the C/ZrB2–SiC composite possesses a relatively high mass loss rate while it still presents an obvious advantage compared with the C/SiC composite. From Fig. 3, it can be clearly seen that there is a decreasing trend in the mass loss rates with oxidation. In addition, the mass loss rates are similar at 12001 and 14001C and are obviously higher than that at 10001C for the C/ZrB2–SiC composite. Figure 4 shows the surface morphologies after the seventh cycle of oxidation. At 10001C, a continuous and densified glass in the form of many particle-shaped aggregates is found for the C/ZrB2–SiC compos￾ite while much powder-shaped SiC debris appears on the surface of the C/SiC composite. At 12001 and 14001C, the aggregates continuously increase with the temperature increasing in the C/ ZrB2–SiC composite, resulting in the formation of the more-uni￾form oxide scale. In the case of the C/SiC composite, many hemisphere-shaped granules with big and small bulks are formed at 12001C, indicating that a small amount of SiC matrix was 200µm (a) 10 µm 90° long fibers SiC matrix 0° longfibers SiC matrix shortfibers SiC matrix ZrB2 particle ZrB2 particle Carbon fiber SiC matrix (b) Fig. 1. Backscattered electron microscopy images of the cross section of the C/ZrB2–SiC composite. (a) The continuous and alternate distribution of 901 long fibers–SiC matrix, short fibers-ZrB2 particles–SiC matrix, and 01 long fibers–SiC matrix; (b) the distribution of carbon fibers, ZrB2 particles, and SiC matrix in the interlayer. Table I. Characteristics of C/ZrB2–SiC Composite Composition C/ZrB2–SiC C/SiC Fiber content (vol%) 24.9 30.6 Density (g/cm3 ) 2.470.1 2.270.1 Flexural strength (MPa) 148712 16378 Young’s modulus (GPa) 58.075.0 26.572.2 Fracture toughness (MPa m1/2) 5.671.6 6.571.2 Ablated surface temperature (1C) 118875 123473 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 300 350 200 250 100 150 0 50 Displacement (mm) Load (N) 0.00 1.15 2.30 3.45 4.60 5.75 6.90 8.05 1 2 3 Fracture toughness (MPam1/2) Fig. 2. Curves of load and fracture toughness versus displacement of the C/ZrB2–SiC composite. 5 10 15 20 25 30 35 1.2 1.6 2.0 2.4 2.8 3.2 Mass loss rate (%) Time (min) C/SiC 1000°C C/SiC 1200°C C/SiC 1400°C C/ZrB2-SiC 1000°C C/ZrB2-SiC 1200°C C/ZrB2-SiC 1400°C Fig. 3. Curves of mass loss rate versus temperature of the C/ZrB2–SiC and the C/SiC composites. October 2007 Communications of the American Ceramic Society 3321

3322 Communications of the American Ceramic Society VoL. 90. No. 10 1000°c 11200c (a2)1400°c 1000°c 可「1200c 2)1400°c 3 Fig 4. Surface morphologies after seven cycles of 5-min oxidation at 1000, 1200, and 1400C.(al, a2, a3)C/zrBr-SiC composite: (bl, b2. b3)C/SiC g fied Sio2 glass is formed at 1400C, illustrating the oxida tion of a large amount of SiC matrix. The CzrBrSiC composites were prepared by a HCV tech- idation, the following reaction can proceed for the nique through densification of a carbon fiber-zrB2 powder pre- ZrBr-Sic composite e one introduction to ZrBySiC ceramic significantly improved the zrB2(s)+5/2O2(g)=ZrO2(s)+B2O3( (1) fracture toughness of monolithic ceramic SiC and ZrB,, and on the other hand, ZrB, addition to the C/SiC composite greatly B2O3(=B2O3(g) enhanced the young s modulus and the oxidation and ablat resistance. From the material point of view, the C/zrB-Sic SiC(s)+3/202(g)=SiO2(0)+Co(g) (3) composites are promising as a novel ultra-high thermal prote tion material, and from the technological point of view, the fab- C(s)+1/2O2(g)=CO(g) (4) cation process offers the advantage of low-cost and short-time production of the C/ZrB--SiC composite with good pro ()+B2O3(g)= Borosilicate glass() References When the temperature is at 1000.C, the predominant glass is M. M. Opeka, I G. Talmy, and J. A. Zaykoski, " Oxidation-Based Materials B2O3, which has a melting point of 450%C, a high vapor pres- Selection for 2000.C+ Hypersonic Aerosurfaces: Theoretical Considerations and sure, and a low oxygen diffusion rate. The B 03 glass plays a"s R. Levine, E J Opia, M. C. Halbig. I D. Kiser, M Singh, and 1. A. Salem, “ evaluat rature Ceramics for Aeropropulsion Use cover and protect the naked carbon fibers, open pores, and mi- J.Er. Ceran.Soc.22.2757-67(20 crocracks At 1200C, only a small amount of SiO2 is FF. Monteverde and A Bellosi "The Resistance to Oxidation of an to form a borosilicate glass, supplying limited oxidation tion, while at the same time the major glass compor Aw. C. Tripp, HH. Davis, and H C. Graham, "Effect of a SiC Addition on th will greatly vaporize glass will be produced, which has high viscosity, low oxygen F. Moteverde and A. Bellosi, ""Microstructure and Properties of an HB,--SiCc diffusivity, and low vapor pressure, and hence, provides much 331-6(2004) 6S. F. Tang J. Y. Deng. S.J. Wang, and W. C. Liu, "Fabrication and Char- Hence, the mass loss rates at 1400%C are similar to those at Stoichiometric Matrix by heaterless chemical vapor infiltration, " Mater. Sci 1200.C although the oxidation temperature is higher. nie Ablation results show that no thickness loss appears after a and ablative behaviors go Cizrb-sic composites, and cisic cor 650-s arc-heated wind tunnel test. An orange-yellow oxide film is formed on the surface. It should be noted that the temperature F. Moteverde S Guicciardi. and A. Bellosi "Ad of the ablated surface is 46C lower than that of the C/sic Mechanical Properties of Zirconium Diboride Based Ceramics, Mater. Sci. Eng composite. This can be attributed to the higher thermal con- ductivity of the C/ZrBrSiC composite and the evaporation of

gradually oxidized to produce SiO2; however, a continuous and densified SiO2 glass is formed at 14001C, illustrating the oxida￾tion of a large amount of SiC matrix. During oxidation, the following reaction can proceed for the C/ZrB2–SiC composite: ZrB2ðsÞ þ 5=2O2ðgÞ ¼ ZrO2ðsÞ þ B2O3ðlÞ (1) B2O3ðlÞ ¼ B2O3ðgÞ (2) SiCðsÞ þ 3=2O2ðgÞ ¼ SiO2ðlÞ þ COðgÞ (3) CðsÞ þ 1=2O2ðgÞ ¼ COðgÞ (4) SiO2ðlÞ þ B2O3ðgÞ ¼ Borosilicate glassðlÞ (5) When the temperature is at 10001C, the predominant glass is B2O3, which has a melting point of 4501C, a high vapor pres￾sure, and a low oxygen diffusion rate. The B2O3 glass plays a positive role in oxidation resistance because the fluid glass can cover and protect the naked carbon fibers, open pores, and mi￾crocracks. At 12001C, only a small amount of SiO2 is produced to form a borosilicate glass, supplying limited oxidation protec￾tion, while at the same time the major glass component B2O3 will greatly vaporize. However, a significant amount of SiO2 glass will be produced, which has high viscosity, low oxygen diffusivity, and low vapor pressure, and hence, provides much more effective oxidation-protective capabilities at 14001C. Hence, the mass loss rates at 14001C are similar to those at 12001C although the oxidation temperature is higher. Ablation results show that no thickness loss appears after a 650-s arc-heated wind tunnel test. An orange-yellow oxide film is formed on the surface. It should be noted that the temperature of the ablated surface is 461C lower than that of the C/SiC composite. This can be attributed to the higher thermal con￾ductivity of the C/ZrB2–SiC composite and the evaporation of the B2O3. 7 IV. Conclusions The C/ZrB2–SiC composites were prepared by a HCVI tech￾nique through densification of a carbon fiber–ZrB2 powder pre￾form by the SiC matrix. On the one hand, the carbon fiber introduction to ZrB2–SiC ceramic significantly improved the fracture toughness of monolithic ceramic SiC and ZrB2, and on the other hand, ZrB2 addition to the C/SiC composite greatly enhanced the Young’s modulus and the oxidation and ablation resistance. From the material point of view, the C/ZrB2–SiC composites are promising as a novel ultra-high thermal protec￾tion material, and from the technological point of view, the fab￾rication process offers the advantage of low-cost and short-time production of the C/ZrB2–SiC composite with good properties. References 1 M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, ‘‘Oxidation-Based Materials Selection for 20001C1 Hypersonic Aerosurfaces: Theoretical Considerations and Historical Experience,’’ J. Mater. Sci., 39 [0] 5887–904 (2004). 2 S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem, ‘‘Evaluation of Ultra-High Temperature Ceramics for Aeropropulsion Use,’’ J. Eur. Ceram. Soc., 22, 2757–67 (2002). 3 F. Monteverde and A. Bellosi, ‘‘The Resistance to Oxidation of an HfB2–SiC Composite,’’ J. Eur. Ceram. Soc., 25, 1025–31 (2005). 4 W. C. Tripp, H. H. Davis, and H. C. Graham, ‘‘Effect of a SiC Addition on the Oxidation of ZrB2,’’ Ceram. Bull., 52 [8] 612–6 (1973). 5 F. Moteverde and A. Bellosi, ‘‘Microstructure and Properties of an HfB2–-SiC Composite for Ultra High Temperature Applications,’’ Adv. Eng. Mater., 6 [5] 331–6 (2004). 6 S. F. Tang, J. Y. Deng, S. J. Wang, and W. C. Liu, ‘‘Fabrication and Char￾acterization of C/SiC Composites With Large Thickness, High Density and Near￾Stoichiometric Matrix by Heaterless Chemical Vapor Infiltration,’’ Mater. Sci. Eng. A., 465 [1-2] 290–4 (2007). 7 S. F. Tang, J. Y. Deng, S. J. Wang, and W. C. Liu, ‘‘Comparison of Thermal and Ablative Behaviors of C/ZrB2–SiC Composites and C/SiC Composites,’’ submitted to Compos. Sci. Technol. 8 F. Moteverde, S. Guicciardi, and A. Bellosi, ‘‘Advances in Microstructure and Mechanical Properties of Zirconium Diboride Based Ceramics,’’ Mater. Sci. Eng. A, 346 [0] 310–9 (2003). 9 Q. W. Huang and L. H. Zhu, ‘‘High-Temperature Strength and Toughness Behaviors for Reaction-Bonded SiC Ceramics Below 14001C,’’ Mater. Lett., 59 [1415] 1732–5 (2005). & 1000 °C 1200°C 1400°C 1000°C 1200°C 1400°C (a1) (a2) (a3) (b1) (b2) (b3) 20 µm 20 µm Fig. 4. Surface morphologies after seven cycles of 5-min oxidation at 10001, 12001, and 14001C. (a1, a2, a3) C/ZrB2–SiC composite; (b1, b2, b3) C/SiC composite. 3322 Communications of the American Ceramic Society Vol. 90, No. 10

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