CARBON 47(2009)I034-I04 1039 between microcrack lips slightly removed from their initial 0 dled C/SIc 0.25 In contrast to the 2D, 2.SD and 3D C/Sic composite, the needled C/Sic exhibits the most extensive hysteresis loop 0.20 主 width in the last cycle and the largest inelastic residual strain accounting for considerable interficial sliding and crack clo sure impediment. On the one hand, only about 37.5% fibers in the needled C/Sic were arranged in the tensile load direc- tion. The few load-bearing fibers easily led to multiple matrix cracking transversely and extensive interfacial debonding upon loading. Also, due to too few reinforcing fibers in the load direction the applied stress induced damage strain (i.e 0 40 80 120 160 200 240 280 320 matrix cracking and interfacial debonding) became very diffi Peak applied stress(MPa) cult to recover upon unloading, resulting in rapid increase in residual strain and accumulation in damage. On the other hand, disorder of the large fibers in the other directions Needled C/Sic (non-load direction) enhanced the crack closure impediment 0.6 3D C/SiC manent strain and the damage accumulation of the needled C/Sic eventually resulted in the earlier fracture and lower 04 strength(Fig. Sa). Comparatively, equivalent 93% load-bearing fibers in the 3D C/Sic were regularly arranged along the load direction, which was very advantageous in recovery of the amage strain leading to much more elastic strain and less inelastic residual strain present in the composites after unloading. Consequently, it can be concluded from Fig. 5 that, 00 the more the effective load-bearing fibers parallel to tensile 04080120160200240280320 axis, the greater the ECFL, the slower the increase of inelastic Peak applied stress(MPa) residual strain and the smaller the damage factor associated with stiffness of the composites eedled C/SiC 3. 3. Axial TRS ●2Dc/siC A-2.5D C/SIC ←30csC According to the calculation methods reported in [20 0.5 through solving the Y-coordinates of the common interse on point of the regr lines of the reloading-l loops, the axial TRS could be measured out and then listed in Table 1. the measured TRs are: 91 MPa for the needled c Sic, 130 MPa for the 2D C/Sic, 127 MPa for the 2.5D C/Sic and 109 MPa for 3D C/SiC composites. The detailed calcula 00 tion procedure has been described in [20] for a specified case of 2D C/Sic material. It must be noted that the measured val 4080120160200240280320 ues herein were actual TRS present in these four composi plied stress(MPa) after partial relief of the maximum thermal misfit stress dur- Fig. 5-Comparison in key mechanical parameters of the ing the cooling from the processing temperature down to ture. Consider a composite with perfe ect inter. composites as a function of the peak applied stress: (a) face bond and non- cracked matrix, theoretical maximum val permanent strain,(b)total strain and (c)damage factor. ues of the axial TRS in matrix can be classically estimated as: zEvI+Em Vn (x-xm)(T。-Tp) unloading cycles(Fig. Sc) because the same level of load were where m and zf refer to the linear CTE of the matrix and fiber uniformly shared by the more reinforcing filbers. It is widely respectively. Tp and To are the processing temperature and accepted [7 that the increase of inelastic residual strain in operation temperature. vm and vi are the volume fraction of C/SiCs may be attributed to the interaction of several phe- the matrix and fiber. Em and Er are the Young moduli nomena: @)the release of axial residual stresses during the the matrix and fibers. It must be mentioned that Eq(6)is va- loading/unloading cycles; (n) partial irreversible sliding aris- lid along fiber direction only and thus the volume fraction of ing from the various energy dissipative frictional mecha- the fibers in this formulas was normalized by ECFl i.Note closure possibly related to fiber roughness and/or the contact assumed as the fully dense matrix modulus because of theunloading cycles (Fig. 5c) because the same level of load were uniformly shared by the more reinforcing fibers. It is widely accepted [7] that the increase of inelastic residual strain in C/SiCs may be attributed to the interaction of several phe￾nomena: (i) the release of axial residual stresses during the loading/unloading cycles; (ii) partial irreversible sliding aris￾ing from the various energy dissipative frictional mecha￾nisms; and (iii) a mechanical impediment of complete crack closure possibly related to fiber roughness and/or the contact between microcrack lips slightly removed from their initial positions. In contrast to the 2D, 2.5D and 3D C/SiC composite, the needled C/SiC exhibits the most extensive hysteresis loop width in the last cycle and the largest inelastic residual strain accounting for considerable interficial sliding and crack clo￾sure impediment. On the one hand, only about 37.5% fibers in the needled C/SiC were arranged in the tensile load direc￾tion. The few load-bearing fibers easily led to multiple matrix cracking transversely and extensive interfacial debonding upon loading. Also, due to too few reinforcing fibers in the load direction the applied stress induced damage strain (i.e., matrix cracking and interfacial debonding) became very diffi- cult to recover upon unloading, resulting in rapid increase in residual strain and accumulation in damage. On the other hand, disorder of the large fibers in the other directions (non-load direction) enhanced the crack closure impediment after unloading. As a consequence, rapid increase in the per￾manent strain and the damage accumulation of the needled C/SiC eventually resulted in the earlier fracture and lower strength (Fig. 5a). Comparatively, equivalent 93% load-bearing fibers in the 3D C/SiC were regularly arranged along the load direction, which was very advantageous in recovery of the damage strain leading to much more elastic strain and less inelastic residual strain present in the composites after unloading. Consequently, it can be concluded from Fig. 5 that, the more the effective load-bearing fibers parallel to tensile axis, the greater the ECFL, the slower the increase of inelastic residual strain and the smaller the damage factor associated with stiffness of the composites. 3.3. Axial TRS comparison According to the calculation methods reported in [20], through solving the Y-coordinates of the common intersec￾tion point of the regression lines of the reloading–unloading loops, the axial TRS could be measured out and then listed in Table 1. the measured TRS are: 91 MPa for the needled C/ SiC, 130 MPa for the 2D C/SiC, 127 MPa for the 2.5D C/SiC, and 109 MPa for 3D C/SiC composites. The detailed calcula￾tion procedure has been described in [20] for a specified case of 2D C/SiC material. It must be noted that the measured val￾ues herein were actual TRS present in these four composites after partial relief of the maximum thermal misfit stress dur￾ing the cooling from the processing temperature down to room temperature. Consider a composite with perfect inter￾face bond and non-cracked matrix, theoretical maximum val￾ues of the axial TRS in matrix can be classically estimated as: rm r ¼ E m kEfVf kEfVf þ EmVm ðaf  amÞðTo  TpÞ; ð6Þ where am and af refer to the linear CTE of the matrix and fiber, respectively. Tp and To are the processing temperature and operation temperature. Vm and Vf are the volume fraction of the matrix and fiber. Em and Ef are the Young moduli of the matrix and fibers. It must be mentioned that Eq. (6) is va￾lid along fiber direction only and thus the volume fraction of the fibers in this formulas was normalized by ECFL k. Note that the effective matrix modulus E m here can not be assumed as the fully dense matrix modulus because of the Fig. 5 – Comparison in key mechanical parameters of the needled C/SiC, 2D C/SiC, 2.5D C/SiC, and 3D C/SiC composites as a function of the peak applied stress: (a) permanent strain, (b) total strain and (c) damage factor. CARBON 47 (2009) 1034 – 1042 1039