《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_j.1151-2916.1998.tb02621.x[1]

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ournal 1. Am Ceram Soc, s1 [9) 2269-77(1998) Creep and Fatigue Behavior in an Enhanced Sic/SiC Composite at High Temperature Shijie Zhu, tt Mineo Mizuno, ".T Yasuo Nagano, Jianwu Cao, .f Yutaka Kagawa, ".f and Hiroshi Kava Japan Fine Ceramics Center, Nagoya 456, Japan, Institute of Industrial Sciences, University of Tokyo, Tokyo 106, apan, and Petroleum Energy Center, Tokyo 106, Japan The monotonic tension, creep, and fatigue behavior of an enhanced SiC/SiC composite will become glass and flow to enhanced SiC/SiC composite was investigated at a tempera- seal the cracks at high temperatures. As a result, the glass phase ture of 1300C in air and argon. The improved creep and prevents further oxidation and heals the cracks, which may fatigue resistances were determined and compared to those improve the creep and fatigue resistance of the composite at of the standard SiC/SiC composite. The effects of additives high temperature glass-forming, boron-based particulates) in the matrix on The cyclic-fatigue behavior of CMCs at high temperatures is the creep and environmental resistance of the enhanced not well understood. Elements such as environmental factors, SiC/SiC composite were discussed. Crack propagation in creep of constituents, thermally induced stresses at interfaces, the matrix of the enhanced SiC/SiC composite was different and interfacial sliding resistance can cause the reduction of from that in the standard SiC/SiC composite. The filling of fatigue life at high temperatures. 8, 18-20 If the temperature is he glassy phases in the cracks prohibited the diffusion of lower than that for the onset of creep of the fibers and the oxygen from the environment. As a result, creep and fa- matrix, the decrease of sliding resistance due to the relaxation tigue properties in the enhanced siC/SiC composite in air at of the residual stress is attributed to the dominant fatigue high temperatures was improved. dation of the fibers(such as creep), the subject becomes complicated and is a priority for further re- search. 2 Fatigue and creep damage mechanisms can operate N THE recent decade, creep and fatigue of continuous-fiber simultaneously under high-temperature cyclic loading. Fatigu reinforced ceramic-matrix composites(CMCs) have been hot loading at high temperatures results in creep-fatigue interac- topics, -20 because these properties are very important for the tion, which causes a reduction in the number of cycles to fail application of CMCs. To obtain high fracture toughness and ure. Creep of the fiber and degradation of the interfacial sliding hermal shock resistance, CMCs are designed with weak inter- resistance have been considered to be the reasons for decreased faces between the fibers and the matrix; e. g, the interface in a fatigue resistance at high temperatures in standard SiC/SiC SiC/SiC composite is coated with carbon or BN. The weak con interface can cause cracks to deflect along the interfaces which Creep and fatigue tests at 1300@C were both conducted in air omits intact fibers to bridge matrix-crack faces. However, at the same maximum stresses, to compare the time-dependent although the use of weak interfaces can increase fracture re- deterioration with cyclic-dependent damage. The creep and fa- sistance, it is not compatible with creep and fatigue resistance tigue data of the enhanced SiC/SiC composite were compare at high temperature, which demands strong interfaces that re- with those of the standard SiC/SiC composite, to investigate sist the nucleation and growth of cavities 2, 3 the effects of the additives in the matrix on the mechanical The carbon layer in a SiC/SiC composite leads to low oxi- roperties. The creep and fatigue crack-propagation paths and dation resistance at high temperatures. 21-4Glass-forming, bo- mechanisms were studied. To understand the interaction of ron-based particulates that react with oxygen to produce a seal environmental creep, creep tests in pure argon were also con- ant glass that inhibits oxidation of the carbon layer can be ducted at the same temperature added to the matrix. This technology has been applied to SiC/SiC composites. SiC/SiC composites modified in this way are called enhanced SiC/SiC composites. 5 Il. Materials and Experimental Procedures Because the creep resistance of SiC fibers(NicalonTM, Nip- The composites used in this investigation were processed by pon Carbon, Tokyo, Japan)is lower than that of the SiC matrix, using chemical vapor infiltration(CVI) of Sic into plane matrix cracking is the dominant damage mechanism. 7 This woven 0/90 SiC-fiber preforms(made by DuPont Lanxide condition is undesirable for creep and environmental resistance Composites, Wilmington, DE). Before infiltration, the pre of the composites, because the increase in matrix stress via forms were coated with carbon by using chemical vapor del stress redistribution causes the development of periodic matrix sition( CVD), to decrease the interface bonding between the cracking. The glass-forming ulates in the matrix of the fibers and the matrix, thereby increasing the toughness. The size of the as-received composite panels was 200 mm x 200 mm with a thickness of 3.2 mm; the composite contained 40 vol% SiC fibers and had a porosity of 9.7%. The diameter of R. Naslain--contributing editor the SiC (NicalonTM) fiber was-12 um, and each bundle con- sisted of 500 fibers The tensile specimens were machined from the panels usil diamond cutting tools. The shape and dimensions of the sp mens for the monotonic tension, creep, and cyclic-fatigue tests inducted by the pet ole tm Entrgy tenter (pic as taphan development programs have been described by Mizuno et a1.&The surfaces of the ecimens were polished by an 800-grit grinding wheel before Lnstitute of Industrial Sciences, University of Tokyo testing; thus, the specimens were unprotected by a final CVI PEtroleum Energy Center run after machining

Creep and Fatigue Behavior in an Enhanced SiC/SiC Composite at High Temperature Shijie Zhu,†,‡ Mineo Mizuno,*,† Yasuo Nagano,*,† Jianwu Cao,*,† Yutaka Kagawa,*,‡ and Hiroshi Kaya§ Japan Fine Ceramics Center, Nagoya 456, Japan, Institute of Industrial Sciences, University of Tokyo, Tokyo 106, Japan, and Petroleum Energy Center, Tokyo 106, Japan The monotonic tension, creep, and fatigue behavior of an enhanced SiC/SiC composite was investigated at a temperature of 1300°C in air and argon. The improved creep and fatigue resistances were determined and compared to those of the standard SiC/SiC composite. The effects of additives (glass-forming, boron-based particulates) in the matrix on the creep and environmental resistance of the enhanced SiC/SiC composite were discussed. Crack propagation in the matrix of the enhanced SiC/SiC composite was different from that in the standard SiC/SiC composite. The filling of the glassy phases in the cracks prohibited the diffusion of oxygen from the environment. As a result, creep and fatigue properties in the enhanced SiC/SiC composite in air at high temperatures was improved. I. Introduction I N THE recent decade, creep and fatigue of continuous-fiberreinforced ceramic-matrix composites (CMCs) have been hot topics,1–20 because these properties are very important for the application of CMCs. To obtain high fracture toughness and thermal shock resistance, CMCs are designed with weak interfaces between the fibers and the matrix; e.g., the interface in a SiC/SiC composite is coated with carbon or BN. The weak interface can cause cracks to deflect along the interfaces, which permits intact fibers to bridge matrix-crack faces.1 However, although the use of weak interfaces can increase fracture resistance, it is not compatible with creep and fatigue resistance at high temperature, which demands strong interfaces that resist the nucleation and growth of cavities.2,3 The carbon layer in a SiC/SiC composite leads to low oxidation resistance at high temperatures.21–24 Glass-forming, boron-based particulates that react with oxygen to produce a sealant glass that inhibits oxidation of the carbon layer can be added to the matrix.6 This technology has been applied to SiC/SiC composites. SiC/SiC composites modified in this way are called enhanced SiC/SiC composites.5 Because the creep resistance of SiC fibers (Nicalon™, Nippon Carbon, Tokyo, Japan) is lower than that of the SiC matrix, matrix cracking is the dominant damage mechanism.7 This condition is undesirable for creep and environmental resistance of the composites, because the increase in matrix stress via stress redistribution causes the development of periodic matrix cracking.3 The glass-forming particulates in the matrix of the enhanced SiC/SiC composite will become glass and flow to seal the cracks at high temperatures. As a result, the glass phase prevents further oxidation and heals the cracks, which may improve the creep and fatigue resistance of the composite at high temperature. The cyclic-fatigue behavior of CMCs at high temperatures is not well understood. Elements such as environmental factors, creep of constituents, thermally induced stresses at interfaces, and interfacial sliding resistance can cause the reduction of fatigue life at high temperatures.8,18–20 If the temperature is lower than that for the onset of creep of the fibers and the matrix, the decrease of sliding resistance due to the relaxation of the residual stress is attributed to the dominant fatigue mechanism.8,19 When the temperature is high enough to produce strength degradation of the fibers (such as creep), the subject becomes complicated and is a priority for further research.12 Fatigue and creep damage mechanisms can operate simultaneously under high-temperature cyclic loading. Fatigue loading at high temperatures results in creep–fatigue interaction, which causes a reduction in the number of cycles to failure. Creep of the fiber and degradation of the interfacial sliding resistance have been considered to be the reasons for decreased fatigue resistance at high temperatures in standard SiC/SiC composites.8,19 Creep and fatigue tests at 1300°C were both conducted in air at the same maximum stresses, to compare the time-dependent deterioration with cyclic-dependent damage. The creep and fatigue data of the enhanced SiC/SiC composite were compared with those of the standard SiC/SiC composite, to investigate the effects of the additives in the matrix on the mechanical properties. The creep and fatigue crack-propagation paths and mechanisms were studied. To understand the interaction of environmental creep, creep tests in pure argon were also conducted at the same temperature. II. Materials and Experimental Procedures The composites used in this investigation were processed by using chemical vapor infiltration (CVI) of SiC into planewoven 0°/90° SiC-fiber preforms (made by DuPont Lanxide Composites, Wilmington, DE). Before infiltration, the preforms were coated with carbon by using chemical vapor deposition (CVD), to decrease the interface bonding between the fibers and the matrix, thereby increasing the toughness. The size of the as-received composite panels was 200 mm × 200 mm with a thickness of 3.2 mm; the composite contained 40 vol% SiC fibers and had a porosity of 9.7%. The diameter of the SiC (Nicalon™) fiber was ∼12 mm, and each bundle consisted of 500 fibers. The tensile specimens were machined from the panels using diamond cutting tools. The shape and dimensions of the specimens for the monotonic tension, creep, and cyclic-fatigue tests have been described by Mizuno et al.8 The surfaces of the specimens were polished by an 800-grit grinding wheel before testing; thus, the specimens were unprotected by a final CVI run after machining. R. Naslain—contributing editor Manuscript No. 191494. Received August 26, 1997; approved December 16, 1997. This work is a part of the automotive ceramic gas turbine development programs conducted by the Petroleum Energy Center (PEC) in Japan. *Member, American Ceramic Society. † Japan Fine Ceramics Center. ‡ Institute of Industrial Sciences, University of Tokyo. § Petroleum Energy Center. J. Am. Ceram. Soc., 81 [9] 2269–77 (1998) Journal 2269

2270 Journal of the American Ceramic Society-Zhu et al. Vol 8l. No 9 all th d were performed using a servo- To compare the creep and fatigue properties of the enhanced hydraulic testing system del mrS 810, MTS System Corp SiC/SiC composite with those of the standard SiC/SiC com- Eden Prairie, MN)at a temperature of 1300.C The monotonic posite, creep and fatigue tests of the standard SiC/SiC com- tensile tests were conducted in air, under a constant displace posite were conducted under the same conditions as those for ment rate of 0.5 mm/min. The specimens were allowed to soak the enhanced SiC/SiC composite in air. The creep and fatigue for -30 min at 1300%C before the tensile tests were started. The data of the standard SiC/SiC composite in argon were cited alignment between the upper and lower grips of the load unit from Zhu et al. 7 Mizuno et al. and Zhu et al. 10 for comparison was veri ified using the steel dummy specimen for verificatio After fracture, the specimens were examined by using optical that was supplied by MTS Corp. to allow a bending strain of microscopy and scanning electron microscopy (SEM) 5%. in accordance with ASTM Standard E 1012-89. Ana lytical and empirical analysis studies have concluded that, for lL. Results and discussion negligible effects on the estimates of the strength-distribution parameters(for example, the Weibull modulus and character- istic strength) of monolithic advanced ceramics, the allowable 1 Microstructures and Monotonic Tension percent bending, as defined in ASTM Practice E 1012, should Micrographs of the enhanced and standard SiC/SiC compos- ites in their original states are shown in Fig. 1. The differend tensile strength distributions of continuous-fiber-reinforced between them are that there are glassy phases in the matrix CMCs do not exist. ASTM Practice C 1275-94 has adopted the the enhanced SiC/SiC composite(the gray phases in the matrix recommendations for the tensile testing of monolithic advanced in Fig. I(c). The thickness of the carbon layer at the interfaces ceramics. Because CMCs have inelastic deformation, which of the enhanced SiC/SiC composite(0.5-0.6 um) is larger than can redistribute the stress state and sometimes lead to notch that of the standard SiC/SiC composite(0. 1-0. 2 um). More- insensitivity, a bending strain of 5% should not affect the over, there are more pores in the fiber bundles of the enhanced strength distribution SiC/SiC composite than in those of the standard SiC/SiC The fatigue tests were performed with a sinusoidal loading composite. frequency of 20 Hz in air. The stress ratio, which is defined The tensile stress-versus-strain relation of the enhanced and the ratio of minimum stress to maximum stress, was 0. 1 for the standard SiC/SiC composites at 1300%C is shown in Fig. 2. The testing temperature. Creep tests were conducted under a con curves of the enhanced SiC/SiC composite indicate linear elas- tant load in air and in an argon atmosphere. Creep strain was tic behavior up to the proportional limit of -70 MPa, and this measured directly from the gauge length of the specimen using a contact extensometer(Model MTS 632.53-F71,M MPa). The UTS of the enhanced SiC/SiC composite is almost System Corp )that had a measuring range of +2.5 mm over its the same as that of the standard SiC/SiC composite; however, gauge length of 25 mm. Periodically, partial unloading the strain at the UTS of the enhanced SiC/SiC composite is reloading was applied, to measure the modulus change durin much higher than that of the standard SiC/SiC composite. It is the creep tests. The specimens were allowed to soak for >30 evident that the addition of glass-forming particulates in the matrix increases the ductility of the composite. A possible rea A controlled-atmosphere furnace(Model MTS 659. Mrs son is the decrease of creep resistance of the matrix in the System Corp was used for the creep tests in the argon atmo- sphere. For the tests in argon, the chamber was first allowed to The Youngs modulus that is obtained from the linear ump down to <13.3 Pa(100 mtorr) and then the chamber was tion of the curve in the enhanced SiC/SiC composite is-89 backfilled with high-purity argon gas. These steps were re- GPa, which is lower than that at 1000C(127 GPa)and much ated three times to ensure a thorough purge. The argon gas lower than that of the standard SiC/SiC composite at 1300C was flowed through the chamber enough to equal five times the (200 GPa). The Young's modulus of the enhanced SiC/Sic chamber volume. The volume percentage of oxygen in the ity argon gas was <I ppm 1000C to 1300 C, whereas the Youngs modulus of the stan dard SiC/SiC composite decreases -23%. It was reported that the Youngs modulus of the ceramic-grade NicalonTM fibers decreased 21% as the temperature increased from 1000C to American Society for Testing and Materials, Philadelphia, PA 1300oC, which is similar to the change of the Youngs modu- 20um Fig 1. Microstructures of the standard and enhanced SiC/SiC composites(a)standard SiC/SiC composite, (b)enhanced SiC/SiC composite, and (c) glassy phases of the matrix of the enhanced SiC/SiC composite

All the mechanical tests were performed using a servohydraulic testing system (Model MTS 810, MTS System Corp., Eden Prairie, MN) at a temperature of 1300°C. The monotonic tensile tests were conducted in air, under a constant displacement rate of 0.5 mm/min. The specimens were allowed to soak for ∼30 min at 1300°C before the tensile tests were started. The alignment between the upper and lower grips of the load unit was verified using the steel dummy specimen for verification that was supplied by MTS Corp. to allow a bending strain of 5%. Similar studies of the effect of bending on the tensile strength distributions of continuous-fiber-reinforced CMCs do not exist. ASTM Practice C 1275-94 has adopted the recommendations for the tensile testing of monolithic advanced ceramics. Because CMCs have inelastic deformation, which can redistribute the stress state and sometimes lead to notch insensitivity, a bending strain of 5% should not affect the strength distribution. The fatigue tests were performed with a sinusoidal loading frequency of 20 Hz in air. The stress ratio, which is defined as the ratio of minimum stress to maximum stress, was 0.1 for the testing temperature. Creep tests were conducted under a constant load in air and in an argon atmosphere. Creep strain was measured directly from the gauge length of the specimen by using a contact extensometer (Model MTS 632.53-F71, MTS System Corp.) that had a measuring range of ±2.5 mm over its gauge length of 25 mm. Periodically, partial unloading– reloading was applied, to measure the modulus change during the creep tests. The specimens were allowed to soak for >30 min at 1300°C before creep or cyclic-fatigue tests were started. A controlled-atmosphere furnace (Model MTS 659, MTS System Corp.) was used for the creep tests in the argon atmosphere. For the tests in argon, the chamber was first allowed to pump down to <13.3 Pa (100 mtorr) and then the chamber was backfilled with high-purity argon gas. These steps were repeated three times to ensure a thorough purge. The argon gas was flowed through the chamber enough to equal five times the chamber volume. The volume percentage of oxygen in the high-purity argon gas was <1 ppm. To compare the creep and fatigue properties of the enhanced SiC/SiC composite with those of the standard SiC/SiC composite, creep and fatigue tests of the standard SiC/SiC composite were conducted under the same conditions as those for the enhanced SiC/SiC composite in air. The creep and fatigue data of the standard SiC/SiC composite in argon were cited from Zhu et al.7 Mizuno et al.8 and Zhu et al.10 for comparison. After fracture, the specimens were examined by using optical microscopy and scanning electron microscopy (SEM). III. Results and Discussion (1) Microstructures and Monotonic Tension Micrographs of the enhanced and standard SiC/SiC composites in their original states are shown in Fig. 1. The differences between them are that there are glassy phases in the matrix of the enhanced SiC/SiC composite (the gray phases in the matrix in Fig. 1(c)). The thickness of the carbon layer at the interfaces of the enhanced SiC/SiC composite (0.5–0.6 mm) is larger than that of the standard SiC/SiC composite (0.1–0.2 mm). Moreover, there are more pores in the fiber bundles of the enhanced SiC/SiC composite than in those of the standard SiC/SiC composite. The tensile stress-versus-strain relation of the enhanced and standard SiC/SiC composites at 1300°C is shown in Fig. 2. The curves of the enhanced SiC/SiC composite indicate linear elastic behavior up to the proportional limit of ∼70 MPa, and this stress is ∼30% of the ultimate tensile strength (UTS) (225 MPa). The UTS of the enhanced SiC/SiC composite is almost the same as that of the standard SiC/SiC composite; however, the strain at the UTS of the enhanced SiC/SiC composite is much higher than that of the standard SiC/SiC composite. It is evident that the addition of glass-forming particulates in the matrix increases the ductility of the composite. A possible reason is the decrease of creep resistance of the matrix in the enhanced SiC/SiC composite. The Young’s modulus that is obtained from the linear portion of the curve in the enhanced SiC/SiC composite is ∼89 GPa, which is lower than that at 1000°C (127 GPa) and much lower than that of the standard SiC/SiC composite at 1300°C (200 GPa). The Young’s modulus of the enhanced SiC/SiC composite decreases 30% as the temperature increases from 1000°C to 1300°C, whereas the Young’s modulus of the standard SiC/SiC composite decreases ∼23%. It was reported that the Young’s modulus of the ceramic-grade Nicalon™ fibers decreased 21% as the temperature increased from 1000°C to 1300°C,25 which is similar to the change of the Young’s modu- ¶ American Society for Testing and Materials, Philadelphia, PA. Fig. 1. Microstructures of the standard and enhanced SiC/SiC composites ((a) standard SiC/SiC composite, (b) enhanced SiC/SiC composite, and (c) glassy phases of the matrix of the enhanced SiC/SiC composite). 2270 Journal of the American Ceramic Society—Zhu et al. Vol. 81, No. 9

September 1998 Creep and Fatigue Behavior in an Enhanced SiC/Sic Composite at High Temperature 2271 300 250 n200 150 1300ˇc 100 Enhanced, Air 50 Standard. ar 0.002 0.003 0.005 0.006 Strai Fig. 2. Monotonic tensile stress versus strain of the standard SiC/SiC composite in argon and the enhanced SiC/SiC composite in air, at 1300oC (the displacement rate was 0.5 mm/min). s modulus of the and moreso the matrix, contribute to the reduction of the modu- enhanced SiC matrix is-70 GPa at 1300oC. Because the cree lus of the enhanced SiC/SiC composite with temperature. Un- rate of the ceramic-grade Nicalon M fibers at 1300oC is suffi- ently hanced SiC matrix. The Young's modulus of the matrix can be within the time frame of their tests, as shown in Fig 3 calculated if the mixture law of modulus is assumed for the composite, which is (2) Creep and fatigue E=E4+Em(1--) Plots of the tensile creep strain versus time in the enhanced SiC/SiC composite at different maximum applied stresses in air where Ee, Es and Em are the Youngs moduli of the at 1300C are shown in Fig. 3. Only the transient creep stage fiber, and the matrix, respectively; Ve is the equivale exists at stresses >90 MPa, which is similar to the creep of fraction of fibers, and Ip is the volume fraction of or Nicalon TM fibers at 1300oC.28, 29 A long transient creep stage he standard SiC/SiC composite at room temperature, E and very short tertiary creep stage appear at 90 MPa. The 260 GPa. Er= 190 GPa 25 and Vn 10%. When v= 20% steady-state or minimum strain rates of cyclic creep are lower Em 320 GPa, which is similar to the measured Young's than those of static creep(Figs. 3 and 4 ). The steady-state of modulus of the CVI SiC matrix(300 GPa2). For the enhanced minimum strain rates of cyclic creep are calculated using the 0.007 0.014 0.006 1300C. 150 MPa. Air 0.005 0.01 0.004 0.008 Creep 0.003 E0.006 口 Fatigue 0.002 0,004 0.001 a Fatigue 0.002 03006009001200 5x1041×10 Time (9 Fig 3. Tensile creep strain versus time of the enhanced SiC/SiC composite under constant load(creep) and cyclic loading(fatigue)in air at 1300C loads of (a)150 and(b)90 MPa)

lus of the standard SiC/SiC composite. Therefore, the fibers, and moreso the matrix, contribute to the reduction of the modulus of the enhanced SiC/SiC composite with temperature. Unfortunately, there are no Young’s modulus data for the enhanced SiC matrix. The Young’s modulus of the matrix can be calculated if the mixture law of modulus is assumed for the composite, which is Ec 4 Ef Vf e + Em(1 − Vf e − Vp) (1) where Ec, Ef , and Em are the Young’s moduli of the composite, fiber, and the matrix, respectively; Vf e is the equivalent volume fraction of fibers, and Vp is the volume fraction of pores. For the standard SiC/SiC composite at room temperature, Ec 4 260 GPa,8 Ef 4 190 GPa,25 and Vp 4 10%. When Vf e 4 20%, Em ≈ 320 GPa, which is similar to the measured Young’s modulus of the CVI SiC matrix (300 GPa26). For the enhanced SiC/SiC composite, the calculated Young’s modulus of the enhanced SiC matrix is ∼70 GPa at 1300°C. Because the creep rate of the ceramic-grade Nicalon™ fibers at 1300°C is sufficiently high,25,27–29 the Young’s modulus is time dependent, within the time frame of their tests, as shown in Fig. 3. (2) Creep and Fatigue Plots of the tensile creep strain versus time in the enhanced SiC/SiC composite at different maximum applied stresses in air at 1300°C are shown in Fig. 3. Only the transient creep stage exists at stresses >90 MPa, which is similar to the creep of Nicalon™ fibers at 1300°C.28,29 A long transient creep stage and very short tertiary creep stage appear at 90 MPa. The steady-state or minimum strain rates of cyclic creep are lower than those of static creep (Figs. 3 and 4). The steady-state of minimum strain rates of cyclic creep are calculated using the Fig. 3. Tensile creep strain versus time of the enhanced SiC/SiC composite under constant load (creep) and cyclic loading (fatigue) in air at 1300°C (loads of (a) 150 and (b) 90 MPa). Fig. 2. Monotonic tensile stress versus strain of the standard SiC/SiC composite in argon and the enhanced SiC/SiC composite in air, at 1300°C (the displacement rate was 0.5 mm/min). September 1998 Creep and Fatigue Behavior in an Enhanced SiC/SiC Composite at High Temperature 2271

2272 Journal of the American Ceramic Society-Zhu et al. Vol 81. No 9 eep,n=8 1300°c,Air 如d 后 6 1300“c Standard. air Maximum stress, MP Stress, MPa Fig. 4. Minimum creep strain rate versus stress of the enhanced ig. 6. Minimum creep strain rate versus stress of the standard and SiC/SiC composite under constant load (creep) and cyclic loading enhanced SiC/SiC composites under constant load in air and in argon (fatigue)in air at 1300C atl300°C naximum strain, i.e, without considering the strain of the un- loading cycle one hand, is much improved. On the other hand, this condition We assume that creep strain rates(e)of the composites can is partly because the creep rates of Sic fibers in argon are be described by the power law igher than those in air. 27 Figure 6 also shows that the creep rates of the enhanced SiC/SiC composite in argon are higher than those of the stan- dard SiC/SiC composite, because of the low creep resistance of he enhanced SiC matrix. However, in air, the creep rates in the where A is a constant, n the stress exponent fo 2 the enhanced SiC/SiC composite are substantially lower than those activation energy for creep, R the gas constar in the standard SiC/SiC composite olute temperature, The stress exponent for cycl 10, All the stress exponents(n) for the creep of the enhanced and which is slightly higher than that for static andard SiC/SiC composites in air and argon are much higher (Fg.4). han the stress exponent(n= 1-2.5) for the creep of Nicalon Although there is an obvious difference in creep rates be- fibers.27-29 However, in a severe-matrix-cracking AL,O,/SiC a given maximum stress are almost the same(Fig. 5). If a consistent with that of the Al2 O, fiber ip of the composite is In Nicalon TM-fiber- shorter time at the maximum stress under fatigue is considered, reinforced glass-ceramics, n for creep of the composite is also a longer time to rupture under fatigue should be expected, the same as that of NicalonTM fibers. The creep rate is con- ompared to that for creep. The same time to rupture under trolled by the creep of fibers in either severe- matrix-cracking fatigue and creep means that the real behavior of fatigue is Al2O,/SiC composites or glass-ceramic-matrix SiC/calcium more complicated, which will be discussed in the last section. aluminosilicate(SiC/CAS) composites, because of their weak Figure 6 shows the creep rates of the enhanced and standard matrix. In SiC/SiC composites, the Sic SiC/SiC composites in air and argon. For the standard SiC/Sic resistance. Therefore, the creep-rate-co of the composite, the creep rates in air are much higher than those in enhanced SiC/SiC composite should be argon, because of the oxidation effects on creep. Conversely, of the fibers constrained by the matrix, an creep of free for the enhanced SiC/SiC composite, which implies that the Figure 7 shows that the time to rupture in air is longer than oxidation resistance of the enhanced SiC/SiC composite that in argon at a given stress in the enhanced SiC/SiC com- 1300°c Cree ague Strength Standard. Ai 50 10210310410 102 TIme to Rupture (s) TIme to Rupture, s Fig. 5. under constant in air at I300°C

maximum strain, i.e., without considering the strain of the unloading cycle. We assume that creep strain rates (e . ) of the composites can be described by the power law e˙ = Asn expS− Q RTD (2) where A is a constant, n the stress exponent for creep, Q the activation energy for creep, R the gas constant, and T the absolute temperature. The stress exponent for cyclic creep is 10, which is slightly higher than that for static creep (n 4 8) (Fig. 4). Although there is an obvious difference in creep rates between creep and fatigue (Fig. 4), their time-to-rupture values at a given maximum stress are almost the same (Fig. 5). If a shorter time at the maximum stress under fatigue is considered, a longer time to rupture under fatigue should be expected, compared to that for creep. The same time to rupture under fatigue and creep means that the real behavior of fatigue is more complicated, which will be discussed in the last section. Figure 6 shows the creep rates of the enhanced and standard SiC/SiC composites in air and argon. For the standard SiC/SiC composite, the creep rates in air are much higher than those in argon, because of the oxidation effects on creep. Conversely, the creep rates in argon are evidently higher than those in air for the enhanced SiC/SiC composite, which implies that the oxidation resistance of the enhanced SiC/SiC composite, on one hand, is much improved. On the other hand, this condition is partly because the creep rates of SiC fibers in argon are higher than those in air.27 Figure 6 also shows that the creep rates of the enhanced SiC/SiC composite in argon are higher than those of the standard SiC/SiC composite, because of the low creep resistance of the enhanced SiC matrix. However, in air, the creep rates in the enhanced SiC/SiC composite are substantially lower than those in the standard SiC/SiC composite. All the stress exponents (n) for the creep of the enhanced and standard SiC/SiC composites in air and argon are much higher than the stress exponent (n 4 1–2.5) for the creep of Nicalon™ fibers.27–29 However, in a severe-matrix-cracking Al2O3/SiC composite, the stress exponent n for creep of the composite is consistent with that of the Al2O3 fiber.11 In Nicalon™-fiberreinforced glass-ceramics, n for creep of the composite is also the same as that of Nicalon™ fibers. The creep rate is controlled by the creep of fibers in either severe-matrix-cracking Al2O3/SiC composites or glass-ceramic-matrix SiC/calcium aluminosilicate (SiC/CAS) composites, because of their weak matrix. In SiC/SiC composites, the SiC matrix has high creep resistance. Therefore, the creep-rate-controlling process of the enhanced SiC/SiC composite should be considered to be creep of the fibers constrained by the matrix, rather than creep of free fibers. Figure 7 shows that the time to rupture in air is longer than that in argon at a given stress in the enhanced SiC/SiC comFig. 4. Minimum creep strain rate versus stress of the enhanced SiC/SiC composite under constant load (creep) and cyclic loading (fatigue) in air at 1300°C. Fig. 5. Maximum stress versus time to rupture of the enhanced SiC/ SiC composite under constant load (creep) and cyclic loading (fatigue) in air at 1300°C. Fig. 6. Minimum creep strain rate versus stress of the standard and enhanced SiC/SiC composites under constant load in air and in argon at 1300°C. Fig. 7. Stress versus time to rupture of the standard and enhanced SiC/SiC composites under constant load in air and in argon at 1300°C. 2272 Journal of the American Ceramic Society—Zhu et al. Vol. 81, No. 9

Journal of the American Ceramic Society-Zhu et al. Vol 81. No 9 15m Fig. 10. Micrograph depicting creep crack propagation paths in a specimen of the enhanced SiC/SiC composite crept in argon at 1300oC and a load of 90 MPa for 2 h SiC composite in air and argon at 1300 C are shown in Fig 8. temperatures. The borosilicate glass does not wet the surface At the same maximum stress the fatigue life of the enhanced and acts to seal the interior from oxidation 6 Although Fig. 12 SiC/SiC composite is much higher than that of the standard shows carbon interphase oxidation on the fracture surfaces, in most areas in the specimens, no oxidative attack exists and The results(Figs. 7 and 8)show that the addition of the carbon remains at the interfaces. Carbon interphase oxidation lass-forming, boron-based particulates in the SiC matrix will be inhibited as long as the glass is present creases creep and the fatigue life in SiC/SiC composites at high temperature in air. Observation of the crack-propagation (4 Creep and Fatigue Damage Evolution paths will give information and evidence to understand this Gradual modulus degradation during cyclic fatigue has been reported for the unidirectional and laminated ceramic compo ites at room temperature 35-39 and at elevated temper 3) Microscopic Damage and fracture tures.5, 16, 17 40 It has been shown that the gradual damage Creep and fatigue cracks are always found at the large pores growth accompanies modulus decrease in the CMCs under mong the fiber bundles(indicated by arrows in Fig. 9). When ue racks meet 0 fibers, debonding of interfaces between the To understand the damage evolution during fatigue fibers and the matrix occurs. The 0 fibers bridge crack faces enhanced SiC/SiC composite at 1300C, the Youngs and therefore decrease the driving force at the crack tip as a were measured. Figure 13 shows the evolution of the bridging component. 1-3,7-12 At 1300 C, the glassy phases be- strain hysteresis loops. The slope decreases and the width of come liquid and flow into cracks. At room temperature, they he loops increases as the number of cycles increases. The resolidify and become situated in the cracks(Fig 9(c).Around the large pores between the fiber bundles, the matrix is the last coating layer by the pure SiC(the light-colored layer), which the same as the matrix in the standard SiC/SiC composite Crack propagation is very straight in this layer but is deflected or discontinuous in the inner matrix( the darker-colored areas in Figs. 9 and 10). This observation means that creep propa- gation occurs in the matrix of the enhanced SiC/SiC composite Severe oxidation of the fibers and the matrix can be found occasionally in the specimens after long time tests in air( Fig 11). This type of oxidation surely decreases the failure time 星 The filling of the glassy phases in the cracks hinders the dif- fusion of oxygen along crack paths. As a result, the oxidation esistance in the enhanced SiC/SiC composite is improved Fracture surfaces under creep or fatigue in air show that the interfaces between the fibers and the matrix become hollow (Fig. 12(a), because the pyrolytic carbon layer at the interface reacts with oxygen,6 and forms gas that evaporates into the environment. There are many pores or holes in the matrix(Fig 12(b). Their size is the same as that of the glass-forming particulates(Fig. I(c). Therefore, it is thought that the par- Fig. I1. Micrograph showing the oxidation fibers and the matrix ticulates become glass and flow out to seal cracks at high after fatigue at a load of 90 MPa for 2. x 106 cycles

SiC composite in air and argon at 1300°C are shown in Fig. 8. At the same maximum stress, the fatigue life of the enhanced SiC/SiC composite is much higher than that of the standard SiC/SiC composite in air, and also is higher than that in argon. The results (Figs. 7 and 8) show that the addition of the glass-forming, boron-based particulates in the SiC matrix increases creep and the fatigue life in SiC/SiC composites at high temperature in air. Observation of the crack-propagation paths will give information and evidence to understand this phenomenon. (3) Microscopic Damage and Fracture Creep and fatigue cracks are always found at the large pores among the fiber bundles (indicated by arrows in Fig. 9). When cracks meet 0° fibers, debonding of interfaces between the fibers and the matrix occurs. The 0° fibers bridge crack faces and therefore decrease the driving force at the crack tip as a bridging component.1–3,7–12 At 1300°C, the glassy phases become liquid and flow into cracks. At room temperature, they resolidify and become situated in the cracks (Fig. 9(c)). Around the large pores between the fiber bundles, the matrix is the last coating layer by the pure SiC (the light-colored layer), which is the same as the matrix in the standard SiC/SiC composite. Crack propagation is very straight in this layer but is deflected or discontinuous in the inner matrix (the darker-colored areas in Figs. 9 and 10). This observation means that creep propagation occurs in the matrix of the enhanced SiC/SiC composite. Severe oxidation of the fibers and the matrix can be found occasionally in the specimens after long time tests in air (Fig. 11). This type of oxidation surely decreases the failure time. The filling of the glassy phases in the cracks hinders the diffusion of oxygen along crack paths. As a result, the oxidation resistance in the enhanced SiC/SiC composite is improved. Fracture surfaces under creep or fatigue in air show that the interfaces between the fibers and the matrix become hollow (Fig. 12(a)), because the pyrolytic carbon layer at the interfaces reacts with oxygen5,6 and forms gas that evaporates into the environment. There are many pores or holes in the matrix (Fig. 12(b)). Their size is the same as that of the glass-forming particulates (Fig. 1(c)). Therefore, it is thought that the particulates become glass and flow out to seal cracks at high temperatures. The borosilicate glass does not wet the surface and acts to seal the interior from oxidation.6 Although Fig. 12 shows carbon interphase oxidation on the fracture surfaces, in most areas in the specimens, no oxidative attack exists and carbon remains at the interfaces. Carbon interphase oxidation will be inhibited as long as the glass is present.6 (4) Creep and Fatigue Damage Evolution Gradual modulus degradation during cyclic fatigue has been reported for the unidirectional and laminated ceramic composites at room temperature35–39 and at elevated temperatures.5,16,17,40 It has been shown that the gradual damage growth accompanies modulus decrease in the CMCs under fatigue loading.37,39 To understand the damage evolution during fatigue of the enhanced SiC/SiC composite at 1300°C, the Young’s moduli were measured. Figure 13 shows the evolution of the stress– strain hysteresis loops. The slope decreases and the width of the loops increases as the number of cycles increases. The Fig. 10. Micrograph depicting creep crack propagation paths in a specimen of the enhanced SiC/SiC composite crept in argon at 1300°C and a load of 90 MPa for 2 h. Fig. 11. Micrograph showing the oxidation fibers and the matrix after fatigue at a load of 90 MPa for 2.8 × 106 cycles. 2274 Journal of the American Ceramic Society—Zhu et al. Vol. 81, No. 9

September 1998 Creep and Fatigue Behavior in an Enhanced SiC/Sic Composite at High Temperature 2275 Matrix Fig. 12. Fractures surfaces in a specimen of the enhanced SiC/SiC composite fatigued in air at 1300 C and a load of 120 MPa for 4. x 10 cycles ((a) general view and(b) former indicates a decrease of the modulus, and the latter further studied. Creep of the bridged fibers transfers stress to means a decrease of the interfacial sliding resistance. The he matrix and causes matrix cracking 30-33 As a result, the teresis loops move to the right along the strain axis, which is modulus decreases as the number of cycles increases known as ratchetting, because of time-dependent deformation Creep tests show similar results of the decrease of the modu- lus with creep time in the enhanced SiC/SiC composite. At 60 plot of the Young's modulus, normalized by the value MPa, the modulus remains constant during creep in air but from the linear portion during the first loading, versus the decreases as the time increases in argon( Fig. 15). Although the number of cycles is shown in Fig. 14. When the modulus applied stress for the creep test is 60 MPa below the matrix decreases to -80% of the original value. the frac- cracking stress(70 MPa), the damage led to degradation of the also matrix cra (E e2, where Er is the modulus of the fibers and V is the total mosphere. Observation of the specimen that was crept at 60 volume fraction of fibers in the composite. This observation MPa in argon shows that matrix cracks are initiated from the means that the matrix still contributes to the modulus of the large pores(Fig. 16); this condition may be due to the lower c At 120 and 90 MPa, the modulus remains constant, up to 10 of the fibers transfers the stress onto the matrix and causes and 10 cycles, respectively, and then decreases as the number matrix cracking. Moreover, the sealing of the cracks by the of cycles increases Matrix cracks formed by the first loading glass may be more effective in air than in argon, because there and during the early stage of fatigue are not sufficient to affect are more oxygen atoms available in air to react with the glas the modulus, which implies that there may exist a critical level forming particulates than in argon of damage for the decrease of modulus. which needs to be The limit modulus for fracture under creep tests in air and c 2x10e 0.8 0.6 40 150 MPa 04 --120 MPg 02 -A-90 MPa 0 3103 110102103104105108107 Strain Cycle Fig. 13. Evolution of the hysteresis l normalized by the value of the SiC/SiC composite fatigued in air at C and a load of 90 MPa ing(EEo), versus the number of cycles of the enhanced SiC/SiC composite under cyclic loading in air at 1300oC

former indicates a decrease of the modulus, and the latter means a decrease of the interfacial sliding resistance. The hysteresis loops move to the right along the strain axis, which is known as ratchetting, because of time-dependent deformation (creep). A plot of the Young’s modulus, normalized by the value from the linear portion during the first loading, versus the number of cycles is shown in Fig. 14. When the modulus decreases to ∼80% of the original value, the specimens fracture. Eighty percent of the original modulus is still higher than (Ef Vf )/2, where Ef is the modulus of the fibers and Vf is the total volume fraction of fibers in the composite. This observation means that the matrix still contributes to the modulus of the composites. At 120 and 90 MPa, the modulus remains constant, up to 10 and 104 cycles, respectively, and then decreases as the number of cycles increases. Matrix cracks formed by the first loading and during the early stage of fatigue are not sufficient to affect the modulus, which implies that there may exist a critical level of damage for the decrease of modulus, which needs to be further studied. Creep of the bridged fibers transfers stress to the matrix and causes matrix cracking.30–33 As a result, the modulus decreases as the number of cycles increases. Creep tests show similar results of the decrease of the modulus with creep time in the enhanced SiC/SiC composite. At 60 MPa, the modulus remains constant during creep in air but decreases as the time increases in argon (Fig. 15). Although the applied stress for the creep test is 60 MPa below the matrix cracking stress (70 MPa), the damage led to degradation of the Young’s modulus, and also matrix cracking, in the argon atmosphere. Observation of the specimen that was crept at 60 MPa in argon shows that matrix cracks are initiated from the large pores (Fig. 16); this condition may be due to the lower creep resistance of the fibers in argon than in air.27 The creep of the fibers transfers the stress onto the matrix and causes matrix cracking. Moreover, the sealing of the cracks by the glass may be more effective in air than in argon, because there are more oxygen atoms available in air to react with the glassforming particulates than in argon. The limit modulus for fracture under creep tests in air and Fig. 12. Fractures surfaces in a specimen of the enhanced SiC/SiC composite fatigued in air at 1300°C and a load of 120 MPa for 4.8 × 104 cycles ((a) general view and (b) cross-sectional view). Fig. 13. Evolution of the hysteresis loops in fatigue of the enhanced SiC/SiC composite fatigued in air at 1300°C and a load of 90 MPa in air. Fig. 14. Young’s modulus, normalized by the value of the first loading (E/E0), versus the number of cycles of the enhanced SiC/SiC composite under cyclic loading in air at 1300°C. September 1998 Creep and Fatigue Behavior in an Enhanced SiC/SiC Composite at High Temperature 2275

Journal of the American Ceramic Society-Zhu et al. Vol 81. No 9 ◆60MPa,A o 60 MPa, Alr 0.5 50 102 Time (s Fig. 16. Micrograph depicting the matrix crack in a specimen of the enhanced SiC/SiC composite crept at a load of 60 MPa in argon 1g modulus, normalized by the value of the first load- ne time to rupture of the enhanced SiC/SiC com nt load (60 MPa)in air and in argon at 1300oC the enhanced SiC/SiC argon is% of the original value, which is lower than that it is also higher thana o he at trd is several orders of magnitude SiC/SiC composite in air, and standard SiC/SiC composite in argon at1300°C lus between creep and fatigue has not yet been well understood (7) Creep ar ue cracks are always found at the large A possible interpretation is that the larger debonded length of pores among the fiber bundles ack propagation occurs the fibers(producing a larger strain in Fig. 3)under fatigue in the matrix of the enhanced composite. The crack an under creep leads to failure of the enhanced fibers at th propagation of the composite is controlled by both the bridged ame maximum stress. 2 Therefore. when the matrix cracks are fibers and the matrix similar in damage extent, fatigue failure occurs before creep Acknowledgment: The authors are grateful for the assistance of Mr S failure Ogawa in the mechanical tests. v. Conclusion References 1) The ultimate tensile strength (UTS)of the enhanced A.G. Evans, ""Perspective on the Development of High-Toughness Ceram- SiC/SiC composite is similar to that of the standard SiC/SiC ics,”J.Am. Cera.Soc,73{2]187-206(1990 composite; however, the strain at UTS of the enhanced SiC/SiC Goto and Y. Kagawa, ""Fracture Behavior and Toughness of a plain- composite is much higher than that of the standard SiC/SiC n SiC Fiber-Reinforced SiC Matrix Composite, Mater. Sci. Eng.A, A211,72-81 composite. The Youngs modulus that is calculated from the ndamental Research in Structural Ceramics for Service Near linear portion of the curve is -89 GPa, which is much lower 000°C,”J.Am. Ceram. soc,7692147-74(1993) than that of the standard SiC/SiC composite at 1300 C(200 A. G. Evans and Life Prediction Issues for High-Temperat ering Ceramics and Their Composites, ""Acta Mater., 45[1] 23-40 2) The stress exponent(n) for cyclic creep is 10, which is K and T. Dunyak, Fully-Reversed Cyclic higher than that for static creep(n= 8). Both stress exponents posites at Elevated Temperature, for creep are much higher than the stress exponent for the creep 6D. S. Fox, ""Oxidation Kinetics of Enhanced SiC/SiC, ""Ceram. Eng. Sci of SiC fibers(n=1-2.5).27-29 The creep-rate-controlling pro Proc,16|S87-84(1995 cess of the composite may be the creep of the fibers that are Behavior of a Sic-Fiber/SiC Composite at Elevated Temperatures,"Compos. by the matrix (3) For the enhanced SiC/SiC composite, the creep rates in Sc. Technol.,57,1629-37(1997) XM. Mizuno, S Zhu, Y. Nagano, Y Sakaida, Y Kagawa, and M. Watanabe argon are evidently higher than hoe that in argon at a given tures, J.Amz Ceram Soc. 79[12]3065-77(1996). stress. However, the creep resistance in air is in argon, because of the oxidation effects on creep for the standard SiC/SiC composite. The oxidation resistance of the Temperature, Mater. Sci. Eng, A, A220, 100-108(1996). enhanced SiC/SiC composite is much improved. and Fatique behavion o t sicaversi c com. sateant, ngh .mayer at resep (4) The creep rate of the enhanced SiC/SiC composite is F Lamouroux, M. Steen, and J. L. Valles,"Uniaxial Tensile and Creep the time to rupture of the enhanced SiC/SiC composite is much perimental evans,F W. Okrand Soc, 14, 529-37 1 Composite:LEX longer than that of the standard SiC/SiC composite in air, be cause of the improved oxidation resistance of the enhanced SiC/SiC composite 3. W. Holmes, Y. Park, and J Creep and Creep Recov- (5) Although the creep rate of the enhanced SiC/SiC avior of a SiC-Fiber Sia N, Matrix Composite, J. Am. Ceram Soc., 76 posite in argon is higher than that of the standard SiC/Si I4X. Wu and J. W. Holmes, "Tensile Creep and Creep-Strain Recovery Be- composite in argon, the time to rupture of the enhanced SiC/ havior of silicon Carbide fiber/calcium aly SiC composite is still longer than that of the standard SiC/SiC 可mmum可 the enhanced SiC/SiC composite posite, "J. Am. Ceram.个mcm olmes, ""Influence of Stress-Ratio on the Elevated Temperature Fatigue of a SiC Fiber-Reinforced Si3N4 Composite, J. Am. Ceram. Soc., 74 7]163945(1991) (6) At the same maximum stress, the cyclic-fatigue life of 7M. Elahi. K. Liao. J Lesko K. Reifsnider. and w. Stinchcomb. " Elevated

argon is ∼40% of the original value, which is lower than that under fatigue. The reason for the difference of the limit modulus between creep and fatigue has not yet been well understood. A possible interpretation is that the larger debonded length of the fibers (producing a larger strain in Fig. 3) under fatigue than under creep leads to failure of the enhanced fibers at the same maximum stress.12 Therefore, when the matrix cracks are similar in damage extent, fatigue failure occurs before creep failure. IV. Conclusion (1) The ultimate tensile strength (UTS) of the enhanced SiC/SiC composite is similar to that of the standard SiC/SiC composite; however, the strain at UTS of the enhanced SiC/SiC composite is much higher than that of the standard SiC/SiC composite. The Young’s modulus that is calculated from the linear portion of the curve is ∼89 GPa, which is much lower than that of the standard SiC/SiC composite at 1300°C (200 GPa). (2) The stress exponent (n) for cyclic creep is 10, which is higher than that for static creep (n 4 8). Both stress exponents for creep are much higher than the stress exponent for the creep of SiC fibers (n 4 1–2.5).27–29 The creep-rate-controlling process of the composite may be the creep of the fibers that are constrained by the matrix. (3) For the enhanced SiC/SiC composite, the creep rates in argon are evidently higher than those in air; consequently, the time to rupture in air is longer than that in argon at a given stress. However, the creep resistance in air is much lower than in argon, because of the oxidation effects on creep for the standard SiC/SiC composite. The oxidation resistance of the enhanced SiC/SiC composite is much improved. (4) The creep rate of the enhanced SiC/SiC composite is much lower than that of the standard SiC/SiC composite, and the time to rupture of the enhanced SiC/SiC composite is much longer than that of the standard SiC/SiC composite in air, because of the improved oxidation resistance of the enhanced SiC/SiC composite. (5) Although the creep rate of the enhanced SiC/SiC composite in argon is higher than that of the standard SiC/SiC composite in argon, the time to rupture of the enhanced SiC/ SiC composite is still longer than that of the standard SiC/SiC composite. This observation implies that the addition of a glassy phase in the matrix of the enhanced SiC/SiC composite increases creep rates but much improves the total creep time to rupture. (6) At the same maximum stress, the cyclic-fatigue life of the enhanced SiC/SiC composite is several orders of magnitude higher than that of the standard SiC/SiC composite in air, and it is also higher than that of the standard SiC/SiC composite in argon at 1300°C. (7) Creep and fatigue cracks are always found at the large pores among the fiber bundles. Slow crack propagation occurs in the matrix of the enhanced SiC/SiC composite. The crack propagation of the composite is controlled by both the bridged fibers and the matrix. Acknowledgment: The authors are grateful for the assistance of Mr. S. Ogawa in the mechanical tests. References 1 A. G. Evans, ‘‘Perspective on the Development of High-Toughness Ceramics,’’ J. Am. Ceram. Soc., 73 [2] 187–206 (1990). 2 K. Goto and Y. Kagawa, ‘‘Fracture Behavior and Toughness of a PlainWoven SiC Fiber-Reinforced SiC Matrix Composite,’’ Mater. Sci. Eng., A, A211, 72–81 (1996). 3 R. Raj, ‘‘Fundamental Research in Structural Ceramics for Service Near 2000°C,’’ J. Am. Ceram. Soc., 76 [9] 2147–74 (1993). 4 A. G. Evans, ‘‘Design and Life Prediction Issues for High-Temperature Engineering Ceramics and Their Composites,’’ Acta Mater., 45 [1] 23–40 (1997). 5 M. Elahi, K. Liao, K. Reifsnider, and T. Dunyak, ‘‘Fully-Reversed Cyclic Fatigue Response of Ceramic Matrix Composites at Elevated Temperature,’’ Ceram. Eng. Sci. Proc., 16 [4] 75–85 (1995). 6 D. S. Fox, ‘‘Oxidation Kinetics of Enhanced SiC/SiC,’’ Ceram. Eng. Sci. Proc., 16 [5] 877–84 (1995). 7 S. Zhu, M. Mizuno, Y. Nagano, Y. Kagawa, and H. Kaya, ‘‘Tensile Creep Behavior of a SiC-Fiber/SiC Composite at Elevated Temperatures,’’ Compos. Sci. Technol., 57, 1629–37 (1997). 8 M. Mizuno, S. Zhu, Y. Nagano, Y. Sakaida, Y. Kagawa, and M. Watanabe, ‘‘Cyclic Fatigue Behavior of SiC/SiC Composite at Room and High Temperatures,’’ J. Am. Ceram. Soc., 79 [12] 3065–77 (1996). 9 S. Zhu, Y. Kagawa, M. Mizuno, S. Guo, Y. Nagano, and H. Kaya, ‘‘In Situ Observation of Fatigue Crack Propagation of SiC-Fiber/SiC Composite at Room Temperature,’’ Mater. Sci. Eng., A, A220, 100–108 (1996). 10S. Zhu, M. Mizuno, Y. Kagawa, J. Cao, Y. Nagano, and H. Kaya, ‘‘Creep and Fatigue Behavior of SiC-Fiber/SiC Composite at High Temperatures,’’ Mater. Sci. Eng., A, A225, 69–77 (1997). 11F. Lamouroux, M. Steen, and J. L. Valles, ‘‘Uniaxial Tensile and Creep Behaviour of an Alumina Fiber-Reinforced Ceramic Matrix Composite: I. Experimental Study,’’ J. Eur. Ceram. Soc., 14, 529–37 (1994). 12A. G. Evans, F. W. Zok, and R. M. McMeeking, ‘‘Fatigue of Ceramic Matrix Composites,’’ Acta Metall. Mater., 43 [3] 859–75 (1995). 13J. W. Holmes, Y. Park, and J. W. Jones, ‘‘Tensile Creep and Creep Recovery Behavior of a SiC-Fiber Si3N4 Matrix Composite,’’ J. Am. Ceram. Soc., 76 [5] 1281–93 (1993). 14X. Wu and J. W. Holmes, ‘‘Tensile Creep and Creep-Strain Recovery Behavior of Silicon Carbide Fiber/Calcium Aluminosilicate Matrix Ceramic Composites,’’ J. Am. Ceram. Soc., 76 [10] 2695–700 (1993). 15C. H. Weber, J. P. A. Lofvander, and A. G. Evans, ‘‘Creep Anisotropy of a Continuous-Fiber-Reinforced Silicon Carbide/Calcium Aluminosilicate Composite,’’ J. Am. Ceram. Soc., 77 [7] 1745–52 (1994). 16J. W. Holmes, ‘‘Influence of Stress-Ratio on the Elevated Temperature Fatigue of a SiC Fiber-Reinforced Si3N4 Composite,’’ J. Am. Ceram. Soc., 74 [7] 1639–45 (1991). 17M. Elahi, K. Liao, J. Lesko, K. Reifsnider, and W. Stinchcomb, ‘‘Elevated Fig. 16. Micrograph depicting the matrix crack in a specimen of the enhanced SiC/SiC composite crept at a load of 60 MPa in argon. Fig. 15. Young’s modulus, normalized by the value of the first loading (E/E0), versus the time to rupture of the enhanced SiC/SiC composite under constant load (60 MPa) in air and in argon at 1300°C. 2276 Journal of the American Ceramic Society—Zhu et al. Vol. 81, No. 9

September 1998 Creep and Fatigue Behavior in an Enhanced SiC/SiC Composite at High Temperature 2277 Temperature Cyclic Fatigue of Silicon Carbide Fiber Reinforced Silicon Car- 29R. Bodet, J. Lamon, N. Jia, and R. E. Tressler, " Microstru J C Stability and Creep Behavior of Si-C-O(Nicalon) Fibers in F W. Zok, and A G. Evans, "Oxidation Monoxide and Argon Environments, " J. A. Ceram. Soc., 79 [10J 2673-86 Embrittlement Probe for Ceramic-Matrix Composites, " J. Am. Ceram. Soc., 78 82097-100(1 PA. G. Evans and F. w. Zok, " The Physics and Mechanics of Brittle Matrix IP. Reynaud, D. Rouby, G. Fantozzi, F. Abbe, and P. Peres, "Cyclic Fatigue of Ceramic-Matrix Composites, pp 85-94 in Ce M.R. Begley, B N Cox, and R. M. MeMeeking, " Time Dependent Crack tre Ceramic-Matrix Ce opposites 1. De Growth in Ceramic Matrix Composites with Creeping Fibers, " Acta Metall. sign, Durability, and Performance. Edited by A G. Evans and R. Naslain Maer,43U]3927-36(1995) American Ceramic Society, Westerville, OH, 1995 M. M y o A. sites wnsh aneeR g MeM eekingech emps soptuse 43 Fatigue Mechanisms in Graphite/SiC Composites at Room and High Tem- 15]727-40(1995) perature, 'J.Am. Ceram Soc., 77[3]792-800(1994) 33A.G. Evans and C. Weber, ""Creep Damage in SiC/SiC Composites,' cing Ageing Effects on Mater. Sci. Eng, A, A208, 1-6(1996) 3-C R. Jones, C H. Henager Jr, and R. H, Jones, ""Crack Bridging by Sic 2L. Filipuzzi, G. Camus, R. Naslain, and J. Thebault, "" Oxidation Mecha Fibers during slow Crack Growth and the resultar sms and Kinetics of lD-SiC/SiC Composite Material, I: An Experimen tal SiC/SiC Composites, "Scr. Metall. Mater, 33[12] 2067-72(1995) pproach,J Am Ceram Soc 721.459-66(1994) on of Oxi- ns in Thermostructural Composites, "J. Anm Ceram Soc., 77 cracking and Elastic M. L. Bouchetou, T. Cutard, M. Huger, D. Fargeot, and C. Gault, "Effect Chou, Damage Development and R. Bodet, X. Bourrat, J. Lamon, and R. Naslain, Tensile Creep Behaviour elope of a Silicon Carbide-Based Fiber with a Low Oxygen Content, J. Mater. Sci. under Static Fatigue and Cyclic Fatigue, 'J. Am. Ceram. Soc., 76[7J1720-28 d mechaniam a. "CMC. Microcracking: Relationship with Microstructure L P. Zawada, L M. Butkus, and G. A. Hartman, Tensile and Fatigue ber-Reinforced Aluminosilicate Glass, J.Am Ceran.Soc,742851-58(1991) G. Simon and A R. Bunsell,""Creep Behaviour and Structural Character 3%S, F. Shuler, J. W. Holmes, and X. Wu, " Influence of Loading Frequency ization at High Temperatures of Nicalon SiC Fibres, J Mater. Sci, 19, 3658- on the Room-Temperature Fatigue of a Carbon-Fiber/SiC-Matrix Composite, 70(1984) JAm. Ceram.Soc,7619]2327-36(1993) N. Jia, R. Bodet, and R E. Tressler, "" Effects of Microstructural Instability 40M. Steen,"Fatigue of 2D Woven Ceramic Comp n the Creep Behavior of Sic-C-O(Nicalon) Fibers in Argon, J. Am. Ceram. tures", pp. 1687-92 in Fatigue 96, Vol. Ill. Edited by G. Ldtjering and H. Soc,762]3051-60(1993) Nowack. Elsevier Science. Oxford. U. K. 1996

Temperature Cyclic Fatigue of Silicon Carbide Fiber Reinforced Silicon Carbide Matrix Composites,’’ Ceram. Eng. Sci. Proc., 15 [4] 3–12 (1994). 18F. E. Heredia, J. C. McNulty, F. W. Zok, and A. G. Evans, ‘‘Oxidation Embrittlement Probe for Ceramic-Matrix Composites,’’ J. Am. Ceram. Soc., 78 [8] 2097–100 (1995). 19P. Reynaud, D. Rouby, G. Fantozzi, F. Abbe, and P. Peres, ‘‘Cyclic Fatigue at High Temperatures of Ceramic-Matrix Composites’’; pp. 85–94 in Ceramic Transactions, Vol. 57, High-Temperature Ceramic-Matrix Composites I: Design, Durability, and Performance. Edited by A. G. Evans and R. Naslain. American Ceramic Society, Westerville, OH, 1995. 20W. L. Morris, B. N. Cox, D. B. Marshall, R. V. Inman, and M. R. James, ‘‘Fatigue Mechanisms in Graphite/SiC Composites at Room and High Temperature,’’ J. Am. Ceram. Soc., 77 [3] 792–800 (1994). 21N. Frety, R. Molins, and M. Boussuge, ‘‘Oxidizing Ageing Effects on SiC–SiC Composites,’’ J. Mater. Sci., 27, 5084–90 (1992). 22L. Filipuzzi, G. Camus, R. Naslain, and J. The´bault, ‘‘Oxidation Mechanisms and Kinetics of 1D-SiC/SiC Composite Material, I: An Experimental Approach,’’ J. Am. Ceram. Soc., 77 [2] 459–66 (1994). 23M. Huger, D. Fargeot, and C. Gault, ‘‘Ultrasonic Characterization of Oxidation Mechanisms in Thermostructural Composites,’’ J. Am. Ceram. Soc., 77 [10] 2554–60 (1994). 24M. L. Bouchetou, T. Cutard, M. Huger, D. Fargeot, and C. Gault, ‘‘Effect of the Environment on CMCs’’; see Ref. 19, pp. 181–90. 25R. Bodet, X. Bourrat, J. Lamon, and R. Naslain, ‘‘Tensile Creep Behaviour of a Silicon Carbide-Based Fiber with a Low Oxygen Content,’’ J. Mater. Sci., 30, 661–77 (1995). 26L. Guillaumat, ‘‘CMC Microcracking: Relationship with Microstructure and Mechanical Behavior’’; Ph.D. Thesis. Laboratoire des Composites Thermostructurales, Domaine Universitaire, Bordeaux, France, Feb. 1994. 27G. Simon and A. R. Bunsell, ‘‘Creep Behaviour and Structural Characterization at High Temperatures of Nicalon SiC Fibres,’’ J. Mater. Sci., 19, 3658– 70 (1984). 28N. Jia, R. Bodet, and R. E. Tressler, ‘‘Effects of Microstructural Instability on the Creep Behavior of SiC-C-O (Nicalon) Fibers in Argon,’’ J. Am. Ceram. Soc., 76 [12] 3051–60 (1993). 29R. Bodet, J. Lamon, N. Jia, and R. E. Tressler, ‘‘Microstructural Stability and Creep Behavior of Si-C-O (Nicalon) Fibers in Carbon Monoxide and Argon Environments,’’ J. Am. Ceram. Soc., 79 [10] 2673–86 (1996). 30A. G. Evans and F. W. Zok, ‘‘The Physics and Mechanics of Brittle Matrix Composites,’’ J. Mater. Sci., 29, 3857–96 (1994). 31M. R. Begley, B. N. Cox, and R. M. McMeeking, ‘‘Time Dependent Crack Growth in Ceramic Matrix Composites with Creeping Fibers,’’ Acta Metall. Mater., 43 [11] 3927–36 (1995). 32M. R. Begley, A. G. Evans, and R. M. McMeeking, ‘‘Creep Rupture in Ceramic Matrix Composites with Creeping Fibers,’’ J. Mech. Phys. Solids, 43 [5] 727–40 (1995). 33A. G. Evans and C. Weber, ‘‘Creep Damage in SiC/SiC Composites,’’ Mater. Sci. Eng., A, A208, 1–6 (1996). 34C. R. Jones, C. H. Henager Jr., and R. H. Jones, ‘‘Crack Bridging by SiC Fibers during Slow Crack Growth and the Resultant Fracture Toughness of SiC/SiCf Composites,’’ Scr. Metall. Mater., 33 [12] 2067–72 (1995). 35K. M. Prewo, ‘‘Fatigue and Stress Rupture of Silicon Carbide FiberReinforced Glass-Ceramics,’’ J. Mater. Sci., 22, 2695–701 (1987). 36P. G. Karandikar and T.-W. Chou, ‘‘Microcracking and Elastic Moduli Reductions in Unidirectional Nicalon-CAS Composite under Cyclic Fatigue Loading,’’ Ceram. Eng. Sci. Proc., 13 [9–10] 881–88 (1992). 37P. G. Karandikar and T.-W. Chou, ‘‘Damage Development and Moduli Reductions in Nicalon–Calcium Aluminosilicate Composites under Static Fatigue and Cyclic Fatigue,’’ J. Am. Ceram. Soc., 76 [7] 1720–28 (1993). 38L. P. Zawada, L. M. Butkus, and G. A. Hartman, ‘‘Tensile and Fatigue Behavior of Silicon Carbide Fiber-Reinforced Aluminosilicate Glass,’’ J. Am. Ceram. Soc., 74 [11] 2851–58 (1991). 39S. F. Shuler, J. W. Holmes, and X. Wu, ‘‘Influence of Loading Frequency on the Room-Temperature Fatigue of a Carbon-Fiber/SiC-Matrix Composite,’’ J. Am. Ceram. Soc., 76 [9] 2327–36 (1993). 40M. Steen, ‘‘Fatigue of 2D Woven Ceramic Composites at High Temperatures’’; pp. 1687–92 in Fatigue ’96, Vol. III. Edited by G. Lu¨tjering and H. Nowack. Elsevier Science, Oxford, U.K., 1996. h September 1998 Creep and Fatigue Behavior in an Enhanced SiC/SiC Composite at High Temperature 2277

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《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_j.1151-2916.1998.tb02621.x[1]

《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_j.1151-2916.1998.tb02621.x[1]

《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_j.1151-2916.1998.tb02621.x[1]