S. Zhu et al. /Composites Science and Technology 59(1999)833-851 in the plat region and finally in the rough region in Creep at 1300C in Argon Fig. 18(b) 4.4. Creep mechanism Since little is known about the creep behavior of the CVI SiC matrix at the temperatures of 1000-13000C, it is difficult to say if the creep rate of the matrix is higher than that of the fiber in SiC/SiC although it was classi fied by a creep mismatch ratio <1 [31]. The minimum steady-state creep rate of CVD SiC is 9. 3x10-10 s-I in argon at 1400 C and 220 MPa [91], which is lower than that(10-7-10-8 s-l)of Nicalon M fiber in argon at 300C and 200 MPa extrapolating from the data at High stress 180 MPa high stresses [43, 44]. If the CVI SiC matrix in the com (fracture time: 1.9 min) posite is considered to be similar to CVD SiC, the SiC/ Sic composite should be classified by a creep mismatch Creep at 1300C in Argon Standard SiC/SiC Fracture Surface ratio >1. Therefore, a time-dependent stress redistribu tion will occur from the less creep-resistant fibers to the more creep resistant matrix after loading. In this case, the stress transfer to the matrix will cause matrix cracking. However, this is based on no matrix cracking during the initial loading and on the existence of a good stress transfer across the interfaces Since the thermal expansion coefficient of the Sic matrix is larger than that of Nicalon M fibers, the resi dual stress on the fibers is compressive when the tem perature is lower than the processing temperature, and relaxes with increasing temperature. Because the CVI processing temperature is about 1000%C, the residual ow stress: 45 MPa (fracture time; 6.6Hr stress on the fibers is close to zero at and above 1000%C The stress state at the interface and the weakness of the Creep at 1300C in Argon interface caused by carbon coating lead to a low stress Standard SiC/SIC transfer capability of the interface. Therefore, the extent of the stress redistribution according to CMr is very difficult to estimate As shown in Fig. 12, a constant tensile load produces an instantaneous strain response followed by a time dependent strain in SiC/SiC composite. The instanta neous strain includes elastic and inelastic strains depending on the stress level because both fiber creep and matrix cracking occur during the initial loading at high stresses. At high stresses, there is only a transient creep stage. This is similar to the creep behavior of NicalonTM fibers [43, 44]. Fiber creep seems to be the (c)Left: Fast fracture area (d)Right: Slow crack growth area rate controlling creep process for the composite in ten Low stress: 45 MPa (fracture time 6.6 Hr) sion. However, the creep strain of Sic/SiC composite is one order of magnitude lower than that of Nicalon TM Fig 18. Tensile creep fracture surfaces in argon.(a)1300C 180 MPa; fibers at a given rupture time [43-45]. This means that (b)1300 C, 45 MPa; ( c)left part in(b), and(d)right part in(b) creep of the fibers is constrained by either the matrix or the weave architecture. Moreover, the stress exponent stage is sometimes and only exhibits a minimum and the activation energy for creep of the composite are creep rate. The tertiary stage is evident and covers a inconsistent with those of Nicalon Tm fibers, as descri- large portion of the creep strain. The progressive bed later debonding of interfaces and the statistical rupture of At low stresses, there are three stages of creep: tran- fibers were identified as the mechanism of tertiary creep sient, steady-state and tertiary stages. The steady-state in an Al2O3 (D/CVI-SiC composite with a 2D wovenin the plat region and ®nally in the rough region in Fig. 18(b). 4.4. Creep mechanism Since little is known about the creep behavior of the CVI SiC matrix at the temperatures of 1000±1300C, it is dicult to say if the creep rate of the matrix is higher than that of the ®ber in SiC/SiC although it was classi- ®ed by a creep mismatch ratio <1 [31]. The minimum steady-state creep rate of CVD SiC is 9.310ÿ10 sÿ1 in argon at 1400C and 220 MPa [91], which is lower than that (10ÿ7 ÿ10ÿ8 sÿ1 ) of NicalonTM ®ber in argon at 1300C and 200 MPa extrapolating from the data at high stresses [43,44]. If the CVI SiC matrix in the com￾posite is considered to be similar to CVD SiC, the SiC/ SiC composite should be classi®ed by a creep mismatch ratio >1. Therefore, a time-dependent stress redistribu￾tion will occur from the less creep-resistant ®bers to the more creep resistant matrix after loading. In this case, the stress transfer to the matrix will cause matrix cracking. However, this is based on no matrix cracking during the initial loading and on the existence of a good stress transfer across the interfaces. Since the thermal expansion coecient of the SiC matrix is larger than that of NicalonTM ®bers, the resi￾dual stress on the ®bers is compressive when the tem￾perature is lower than the processing temperature, and relaxes with increasing temperature. Because the CVI processing temperature is about 1000C, the residual stress on the ®bers is close to zero at and above 1000C. The stress state at the interface and the weakness of the interface caused by carbon coating lead to a low stress transfer capability of the interface. Therefore, the extent of the stress redistribution according to CMR is very dicult to estimate. As shown in Fig. 12, a constant tensile load produces an instantaneous strain response followed by a time dependent strain in SiC/SiC composite. The instanta￾neous strain includes elastic and inelastic strains depending on the stress level because both ®ber creep and matrix cracking occur during the initial loading at high stresses. At high stresses, there is only a transient creep stage. This is similar to the creep behavior of NicalonTM ®bers [43,44]. Fiber creep seems to be the rate controlling creep process for the composite in ten￾sion. However, the creep strain of SiC/SiC composite is one order of magnitude lower than that of NicalonTM ®bers at a given rupture time [43±45]. This means that creep of the ®bers is constrained by either the matrix or the weave architecture. Moreover, the stress exponent and the activation energy for creep of the composite are inconsistent with those of NicalonTM ®bers, as descri￾bed later. At low stresses, there are three stages of creep: tran￾sient, steady-state and tertiary stages. The steady-state stage is sometimes short and only exhibits a minimum creep rate. The tertiary stage is evident and covers a large portion of the creep strain. The progressive debonding of interfaces and the statistical rupture of ®bers were identi®ed as the mechanism of tertiary creep in an Al2O3(f)/CVI-SiC composite with a 2D woven Fig. 18. Tensile creep fracture surfaces in argon. (a) 1300C 180 MPa; (b) 1300C, 45 MPa; (c) left part in (b), and (d) right part in (b). S. Zhu et al. / Composites Science and Technology 59 (1999) 833±851 845