S. Zhu et al./ Composites Science and Technology 59(1999)833-857 that of SiC/SiC at room temperature. The interface reduce the thermally induced stress. Since CVI proces- liding resistance decreases with increasing temperature. sing temperature is about 1000C, the thermally induced Therefore, the fatigue fracture morphology of Sic/Sic stress at 1000oC will be greatly decreased at 1000 C is similar to that of carbon/SiC at room tem- The observation of cracks on two interrupted speci perature. This demonstrates that the interface sliding mens(at 96 and at 187 MPa in tension at room tem- resistance of SiC/SiC markedly influences the evolution perature)showed that only small cracks initiated at the of fatigue damage at high temperature large pores at 96 MPa and long cracks formed in 90 bundles at 187 MPa, which is above the fatigue limit at 3.4. Cyclic fatigue mechanisms room temperature. Therefore, at room temperature cyclic fatigue is controlled by crack propagation. At the Recently, Evans et al. [67] reviewed fatigue of ceramic- stress below the fatigue limit, the driving force is not matrix composites at room temperature. Two possible enough to cause crack propagation. However, at mechanisms for CMCs were proposed [67]:(1)changes 1000 C cracks can propagate since the driving force is in the interface sliding resistance during cycling;(2) increased by reduction of fiber-bridging stress degradation of the strength of the fibers by cyclic sliding DiCarlo [40] pointed out that creep onset of Nica along the interface by means of an abrasion mechanism, lonM fibers was at 900 C. Therefore, fiber creep is an which introduces flaws in the fibers important factor for considering fatigue mechanism at The specimens fatigued for 107 cycles at the peak 1000oC. Henager and Jones[75-78] have shown that the ress of 160 MPa at room temperature have the same creep-induced stress relaxation of crack-bridging Nica- strength as the original specimens. The reason is that lonM fibers can account for the crack growth rate in the development of many matrix cracks and fiber/ CVI SiC reinforced with Nicalon TM fibers at 1100 C matrix debonding effectively reduce the number of stress However, creep data of Nicalon fibers at 1000C are particular crack extension. Fron gy associated with any very limited, not enough to be used for a prediction of concentration points and the ener this point, the 2D SiC/ creep relaxation of fibers in SiC/SiC Sic composite is considered to be a fatigue-resistant At room temperature, the fatigue resistance of SiC/ material at room temperature Sic or other CMCs is good with a fatigue limit of about At 1000 C, however, the residual strength of the spe- 80% UTS. However, at high temperatures, even if the cimens fatigued for 10 cycles at the peak stress of 75 temperature is not high enough to cause creep of fibers, MPa is lower than that of original specimens, although fatigue resistance of the composites was decreased [20] the difference is not great. This means that damage The present results show a similar phenomenon: fatigue mechanism during cyclic fatigue at 1000C is different resistance of SiC/Sic was much decreased at 1000C. As rom that at room temperature. At high temperature, discussed in the last section, fiber creep and decreased effects of oxidation, creep and sliding resistance of the sliding resistance of interface were causing the decreased interface should be considered. The creep rate of the Sic fatigue resistance. The lower sliding resistance of inter matrix at 1000C is very low and little oxidation occurs face is good for the failure strain and UTS under in argon. Therefore, the cyclic fatigue mechanisms of monotonic loading, but inferior to cyclic fatigue resis Sic/SiC at 1000 C will be explored only by the effects of tance. This is why UTS and failure strain at 1000C are fiber creep and the interfacial sliding resistance in the higher those at room temperature and the fatigue limit at 1000.c is lower than that at room temperature The friction stress at the interface is given by [55] Therefore, a balance should be made for monotonic properties and for cyclic fatigue resistance when ≈ designing or modifying 2D-woven CMCs where u is the coefficient of sliding friction, or is the 3.5. Fatigue of enhanced SiC/Sic and Hi-Nicalon/ radial thermally induced stress, og is the radial stress caused by roughness of interface, op is the radial stress arising from the difference in Poissons ratios between The cyclic fatigue life versus the maximum stress of fiber and matrix Hi-Nicalon/SiC at 1300.C in air is almost the same The thermally induced stress caused by thermal as that of enhanced Sic/SiC in air [57, 58], but longer expansion mismatch has an important role in both the than that of standard SiC/Sic in air and argon at formation and the effectiveness of fiber bridging. The 1300C [56](Fig 9) thermal expansion coefficient (3. 1x10-6 K-)of Fig. 10 shows the evolution of the stress strain fiber is lower than that (4.8x10-6))of the esis loops. The slope decreases and the width of the matrix [74]. As a result, a compressive residual increases with cycles. The former indicates the de ( fiber-clamping stress) is produced after processing at of the modulus and the latter means the decrease of the high temperature Increasing temperature is expected to interfacial sliding resistance. The hysteresis loops movethat of SiC/SiC at room temperature. The interface sliding resistance decreases with increasing temperature. Therefore, the fatigue fracture morphology of SiC/SiC at 1000C is similar to that of carbon/SiC at room tem￾perature. This demonstrates that the interface sliding resistance of SiC/SiC markedly in¯uences the evolution of fatigue damage at high temperature. 3.4. Cyclic fatigue mechanisms Recently, Evans et al. [67] reviewed fatigue of ceramic￾matrix composites at room temperature. Two possible mechanisms for CMCs were proposed [67]: (1) changes in the interface sliding resistance during cycling; (2) degradation of the strength of the ®bers by cyclic sliding along the interface by means of an abrasion mechanism, which introduces ¯aws in the ®bers. The specimens fatigued for 107 cycles at the peak stress of 160 MPa at room temperature have the same strength as the original specimens. The reason is that the development of many matrix cracks and ®ber/ matrix debonding e€ectively reduce the number of stress concentration points and the energy associated with any particular crack extension. From this point, the 2D SiC/ SiC composite is considered to be a fatigue-resistant material at room temperature. At 1000C, however, the residual strength of the spe￾cimens fatigued for 107 cycles at the peak stress of 75 MPa is lower than that of original specimens, although the di€erence is not great. This means that damage mechanism during cyclic fatigue at 1000C is di€erent from that at room temperature. At high temperature, e€ects of oxidation, creep and sliding resistance of the interface should be considered. The creep rate of the SiC matrix at 1000C is very low and little oxidation occurs in argon. Therefore, the cyclic fatigue mechanisms of SiC/SiC at 1000C will be explored only by the e€ects of ®ber creep and the interfacial sliding resistance in the following. The friction stress at the interface is given by [55] i  …T ‡ R ‡ P† …2† where  is the coecient of sliding friction, T is the radial thermally induced stress, R is the radial stress caused by roughness of interface, P is the radial stress arising from the di€erence in Poisson's ratios between ®ber and matrix. The thermally induced stress caused by thermal expansion mismatch has an important role in both the formation and the e€ectiveness of ®ber bridging. The thermal expansion coecient (3.110ÿ6 Kÿ1 ) of SiC ®ber is lower than that (4.810ÿ6 Kÿ1 ) of the SiC matrix [74]. As a result, a compressive residual stress (®ber-clamping stress) is produced after processing at high temperature. Increasing temperature is expected to reduce the thermally induced stress. Since CVI proces￾sing temperature is about 1000C, the thermally induced stress at 1000C will be greatly decreased. The observation of cracks on two interrupted speci￾mens (at 96 and at 187 MPa in tension at room tem￾perature) showed that only small cracks initiated at the large pores at 96 MPa and long cracks formed in 90 bundles at 187 MPa, which is above the fatigue limit at room temperature. Therefore, at room temperature cyclic fatigue is controlled by crack propagation. At the stress below the fatigue limit, the driving force is not enough to cause crack propagation. However, at 1000C cracks can propagate since the driving force is increased by reduction of ®ber-bridging stress. DiCarlo [40] pointed out that creep onset of Nica￾lonTM ®bers was at 900C. Therefore, ®ber creep is an important factor for considering fatigue mechanism at 1000C. Henager and Jones [75±78] have shown that the creep-induced stress relaxation of crack-bridging Nica￾lonTM ®bers can account for the crack growth rate in CVI SiC reinforced with NicalonTM ®bers at 1100C. However, creep data of NicalonTM ®bers at 1000C are very limited, not enough to be used for a prediction of creep relaxation of ®bers in SiC/SiC. At room temperature, the fatigue resistance of SiC/ SiC or other CMCs is good with a fatigue limit of about 80% UTS. However, at high temperatures, even if the temperature is not high enough to cause creep of ®bers, fatigue resistance of the composites was decreased [20]. The present results show a similar phenomenon: fatigue resistance of SiC/SiC was much decreased at 1000C. As discussed in the last section, ®ber creep and decreased sliding resistance of interface were causing the decreased fatigue resistance. The lower sliding resistance of inter￾face is good for the failure strain and UTS under monotonic loading, but inferior to cyclic fatigue resis￾tance. This is why UTS and failure strain at 1000C are higher those at room temperature and the fatigue limit at 1000C is lower than that at room temperature. Therefore, a balance should be made for monotonic properties and for cyclic fatigue resistance when designing or modifying 2D-woven CMCs. 3.5. Fatigue of enhanced SiC/SiC and Hi-NicalonTM/ SiC The cyclic fatigue life versus the maximum stress of Hi-NicalonTM/SiC at 1300C in air is almost the same as that of enhanced SiC/SiC in air [57,58], but longer than that of standard SiC/SiC in air and argon at 1300C [56] (Fig. 9). Fig. 10 shows the evolution of the stress strain hyster￾esis loops. The slope decreases and the width of the loops increases with cycles. The former indicates the decrease of the modulus and the latter means the decrease of the interfacial sliding resistance. The hysteresis loops move S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851 841