CARBON 47(2009)I034-I04 1037 Table 2-Constituent parameters and values of C/sic and is about 4. x 10-6/K for the isotropic Sic matrix. Wher composites thermal misfit generated in the C/Sic, both axial and radial residual tensile stress in brittle Sic matrix on the fibers easily Parameter led to these two families of matrix cracks. Below the processing matrix Em GPa temperature of 1000C, the type I cracks mainly occurred be- cTe of Sic matrix 10-6/K cause the sic matrix encountered the axial tensile residual Fracture strength of Sic matrix mu 58 stress resulted from its much greater shrinkage than the axial Youngs modulus of C fiber fiber (in this case, partial interficial debonding also initiated C Aber radius R 3.5 Room along fiber/matrix interfaces due to the greater shrinkage of mperature 1273 the radial fibers than the matrix upon cooling from the process CTE of C fber axial ing temperature). Above the processing temperature of 10(radial) 1000C, the type Il cracks tend to form because the Sic matrix a Coefficient of thermal encountered the loop tensile residual stress because it has much less expansion than the radial fiber. The type I cracks are now those cracks oriented perpendicular to the loading cracks: type I cracks perpendicular to fiber axis and type I direction(0 fibers). As a consequence some type II cracks be- cracks parallel to flber axis, are clearly indicated in Fig. 2a. come type I cracks when a load is applied parallel to fiber direc- As we know, the carbon fibers display a significant anisotropy tion. Many previous researchers also mainly observed the type in the axial and radial directions whereas the CVI-Sic matrix I cracks in the as-fabricated C/SiCs [14-17. Upon the extra is generally considered as isotropy. For the long and continu- mechanical loading, these type I thermal cracks transversely ous carbon fiber, as listed in Table 2, the radial cte is much grow and propagate leading to progressive increase inresidual larger than its surrounding matrix one whereas the axial strain of the composite [18] and continuous TRS relief in addi- CtE is much smaller than the matrix one. That is tion to the processing-induced thermal load damage 》 In the current study, therefore, we are type I cracks perpendicular to loading direction(parallel to fi- radal and f al were the CTE of the fibers in the radial and ax- ber axis) and their effect on the axial TRS evolution whereas ial direction, which were well known to be about 10- and the radial TRS relief (i.e, partial interficial debonding) will 0x10-/K, respectively. zm denotes the CTE of the CVI-matrix be neglected b240 200 80 20 0.00.10.20.3040.5060.7 Strain(%) Strain(%) 240 3D C/SiC 200 2. 5D C/SiC Strain(%) Strain(%) Fig 3- Typical reloading/unloading hysteresis loop evolutions of (a)needled C/Sic, (b)2D C/sic, (c)2.5D C/Sic, and (d 3D C/siccracks: type I cracks perpendicular to fiber axis and type II cracks parallel to fiber axis, are clearly indicated in Fig. 2a. As we know, the carbon fibers display a significant anisotropy in the axial and radial directions whereas the CVI-SiC matrix is generally considered as isotropy. For the long and continu￾ous carbon fiber, as listed in Table 2, the radial CTE is much larger than its surrounding matrix one whereas the axial CTE is much smaller than the matrix one. That is aradial f  am  aaxial f : ð2Þ aradial f and aaxial f were the CTE of the fibers in the radial and ax￾ial direction, which were well known to be about 10 · 106 and 0 · 106 /K, respectively. am denotes the CTE of the CVI-matrix and is about 4.6 · 106 /K for the isotropic SiC matrix. When thermal misfit generated in the C/SiC, both axial and radial residual tensile stress in brittle SiC matrix on the fibers easily led to these two families of matrix cracks. Below the processing temperature of 1000 C, the type I cracks mainly occurred be￾cause the SiC matrix encountered the axial tensile residual stress resulted from its much greater shrinkage than the axial fiber (in this case, partial interficial debonding also initiated along fiber/matrix interfaces due to the greater shrinkage of the radial fibers than thematrix upon cooling from the process￾ing temperature). Above the processing temperature of 1000 C, the type II cracks tend to form because the SiC matrix encountered the loop tensile residual stress because it has much less expansion than the radial fiber. The type I cracks are now those cracks oriented perpendicular to the loading direction (0 fibers). As a consequence some type II cracks be￾come type I cracks when a load is applied parallel to fiber direc￾tion. Many previous researchers also mainly observed the type I cracks in the as-fabricated C/SiCs [14–17]. Upon the extra mechanical loading, these type I thermal cracks transversely grow and propagate leading to progressive increase in residual strain of the composite [18] and continuous TRS relief in addi￾tion to the processing-induced thermal load damage. In the current study, therefore, we are concerned with the type I cracks perpendicular to loading direction (parallel to fi- ber axis) and their effect on the axial TRS evolution whereas the radial TRS relief (i.e., partial interficial debonding) will be neglected. Table 2 – Constituent parameters and values of C/SiC composites. Parameter Symbol Value Units Young’s modulus of SiC matrix Em 350 GPa CTEa of SiC matrix am 4.6 106 /K Fracture strength of SiC matrix rmu 58 MPa Young’s modulus of C fiber Ef 230 GPa C fiber radius R 3.5 lm Room temperature T0 298 K Processing temperature Tp 1273 K CTE of C fiber af 0 (axial) 10 (radial) 106 /K a Coefficient of thermal expansion. Fig. 3 – Typical reloading/unloading hysteresis loop evolutions of (a) needled C/SiC, (b) 2D C/SiC, (c) 2.5D C/SiC, and (d) 3D C/SiC composite specimens. CARBON 47 (2009) 1034 – 1042 1037