January 1999 Fiber Effects on Minicomposite Mechanical Properties for Several SiC/cvl-SiC Matrix Syster 3日-M白Y This observation would mean that for every crack formed, the amount of matrix around a crack that is incapable of further cracking is much greater for the Nic-3MBN and HN-PBN minicomposites than for the Syl-PBN minicomposites(from Eq(5) By combining these two factors, the local nature of cracking in Syl-PBN minicomposites can be deduced. Initial cracking in all three minicomposites occurs at the same stress(Fig. 3), because of the large flaws in the matrix. Most likely, there is some variabl ity in the matrix volume fraction along the length, and initial fracture occurs at regions with greater matrix frac- tions(this was observed for single-fiber microcomposites) Cracking eventually stops in these weak matrix regions for the Nic-3MBN and HN-PBN minicomposites, because the low in- terfacial shear stress blocks"large"regions of matrix from greater amounts of stress. However, cracking occurs in other egions of the matrix with smaller flaws as the matrix stress is HNAP 5 gapm increased. For the Syl-PBN minicomposite, the high interfacial shear stress enabled a larger portion of the weak matrix regions to be stressed, so cracking was concentrated more locally along the length of the minicomposite Matrix stresses that were high enough to crack/saturate other portions of the matrix were The local nature of cracking for the Syl-PBN minicomposite as well as the lower strain to failure of this fiber type, explains he lower strain to failure of this minicomposite system. Be- cause only a fraction of the matrix is actually cracked over the entire gauge length, and the sliding length is relatively small at the ultimate failure stress and strain. To obtain larger ded at each crack, a relatively small length of fiber is fully trains with this composite, lower interfacial shear st lower matrix cracking stresses, and/or higher fiber sti must be achieved o The minicomposite test has been proven to be very effective evaluating the tensile stress-strain behavior of several dif- SYLAP 5.0 kV x6.00K 5.00H ferent types of SiC fiber/CVI-SiC matrix systems. The ultimate strength, ultimate strain, and interfacial properties could be directly correlated with the fiber strengths, fiber moduli, and Fig. 14. SEM micrographs of fiber surfaces for(a)Hi-Nicalon and fiber roughnesses for the three systems studied. It was evident (b) Sylramic as-produced fibers that the lower strain to failure and large interfacial shear stresses measured for the Sylramic minicomposites were re- spectively due to the high modulus and rough surface of this composites, because of the fiber surface and/or the fiber. This result may have good and bad implications for com- lack of a carbon layer at th iber interface. A constant osite use. On the negative side higher strain to failures would interfacial shear stress co account for the stress be desired for most composite applications that require com- dependent T behavior and ot adequately model Syl- ant behavior On the positive side the smaller crack spacing ramic minicomposite behavior (and presumably smaller crack openings) could potentially have benefits for high-temperature applications if the access of (4) Matrix-Crack Spacing the environment to the interphase region was hindered, because Two factors account for cracking behavior in composites smaller matrix cracks"seal"'faster(compared to larger crack First, the matrix has a matrix-flaw distribution that is quite openings)from reactions between the matrix and the environ- large and accounts for the large stress range over which crack ment. In addition, the higher-T, higher-modulus Sylramic fiber ing occurs(Fig. 3 and Morscher et al. 1). Figure 3(a)shows system offers a higher stiffness material after damage, which that it is reasonable to assume that the matrix -flaw distribution may be important for some applications. These implications in the Syl-PBN and HN-PBN minicomposites is approximately will have to be considered and studied when implementing the same, because the AE activity of both minicomposites f these fiber types in a specific composite system for a specific lows the same trend with composite stress(matrix stress). The Nic-3MBN minicomposite may have a different matrix-flaw distribution. The coated-Nicalon tow already had a Sic layer APPENDIX on top of the BN coating prior to CVI-SiC infiltration, which or matrix infiltration and a thicker SiC region Estimation of the number of cracks from the exterior of the minicomposite(Fig. 4). This thicker SiC Acoustic Emission ion most likely resulted in a greater density of larger fl: Unfortunately, matrix cracks were not visible via stereom- which resulted in more cracking at lower stresses. Second, the croscopy during the tensile test. To determine the value of T, sliding length determines the amount of matrix around a crack the number of cracks for a given hysteresis-loop peak stress that can still be cracked. 2 The interfacial shear stress of the must be known, because crack saturation for most of the hys- Syl-PBN minicomposite, although probably not constant over teresis-loop peak stresses had not been reached. If the cumu- g length, is significantly greater than the interfacial lative AE energy could be related to the cracks produced, M ss of the Nic-3MBN and HN-PBN minicomposites. (Eq(6))could be approximated for a given hycomposites, because of the rougher fiber surface and/or the lack of a carbon layer at the BN/fiber interface. A constant interfacial shear stress could not account for the stress￾dependent t behavior and would not adequately model Syl￾ramic minicomposite behavior. (4) Matrix-Crack Spacing Two factors account for cracking behavior in composites. First, the matrix has a matrix-flaw distribution that is quite large and accounts for the large stress range over which crack￾ing occurs (Fig. 3 and Morscher et al.11). Figure 3(a) shows that it is reasonable to assume that the matrix-flaw distribution in the Syl-PBN and HN-PBN minicomposites is approximately the same, because the AE activity of both minicomposites fol￾lows the same trend with composite stress (matrix stress). The Nic-3MBN minicomposite may have a different matrix-flaw distribution. The coated-Nicalon tow already had a SiC layer on top of the BN coating prior to CVI-SiC infiltration, which resulted in poor matrix infiltration and a thicker SiC region on the exterior of the minicomposite (Fig. 4). This thicker SiC region most likely resulted in a greater density of larger flaws, which resulted in more cracking at lower stresses. Second, the sliding length determines the amount of matrix around a crack that can still be cracked.24 The interfacial shear stress of the Syl-PBN minicomposite, although probably not constant over the sliding length, is significantly greater than the interfacial shear stress of the Nic-3MBN and HN-PBN minicomposites. This observation would mean that for every crack formed, the amount of matrix around a crack that is incapable of further cracking is much greater for the Nic-3MBN and HN-PBN minicomposites than for the Syl-PBN minicomposites (from Eq. (5)). By combining these two factors, the local nature of cracking in Syl-PBN minicomposites can be deduced. Initial cracking in all three minicomposites occurs at the same stress (Fig. 3), because of the large flaws in the matrix. Most likely, there is some variability in the matrix volume fraction along the length, and initial fracture occurs at regions with greater matrix frac￾tions (this was observed for single-fiber microcomposites29). Cracking eventually stops in these weak matrix regions for the Nic-3MBN and HN-PBN minicomposites, because the low in￾terfacial shear stress blocks ‘‘large’’ regions of matrix from greater amounts of stress. However, cracking occurs in other regions of the matrix with smaller flaws as the matrix stress is increased. For the Syl-PBN minicomposite, the high interfacial shear stress enabled a larger portion of the weak matrix regions to be stressed, so cracking was concentrated more locally along the length of the minicomposite. Matrix stresses that were high enough to crack/saturate other portions of the matrix were never achieved. The local nature of cracking for the Syl-PBN minicomposite, as well as the lower strain to failure of this fiber type, explains the lower strain to failure of this minicomposite system. Be￾cause only a fraction of the matrix is actually cracked over the entire gauge length, and the sliding length is relatively small at each crack, a relatively small length of fiber is fully loaded at the ultimate failure stress and strain. To obtain larger failure strains with this composite, lower interfacial shear stresses, lower matrix cracking stresses, and/or higher fiber strengths must be achieved. V. Conclusions The minicomposite test has been proven to be very effective for evaluating the tensile stress–strain behavior of several dif￾ferent types of SiC fiber/CVI-SiC matrix systems. The ultimate strength, ultimate strain, and interfacial properties could be directly correlated with the fiber strengths, fiber moduli, and fiber roughnesses for the three systems studied. It was evident that the lower strain to failure and large interfacial shear stresses measured for the Sylramic minicomposites were re￾spectively due to the high modulus and rough surface of this fiber. This result may have good and bad implications for com￾posite use. On the negative side, higher strain to failures would be desired for most composite applications that require com￾pliant behavior. On the positive side, the smaller crack spacing (and presumably smaller crack openings) could potentially have benefits for high-temperature applications if the access of the environment to the interphase region was hindered, because smaller matrix cracks ‘‘seal’’ faster (compared to larger crack openings) from reactions between the matrix and the environ￾ment. In addition, the higher-t, higher-modulus Sylramic fiber system offers a higher stiffness material after damage, which may be important for some applications. These implications will have to be considered and studied when implementing these fiber types in a specific composite system for a specific application. APPENDIX Estimation of the Number of Cracks from Acoustic Emission Unfortunately, matrix cracks were not visible via stereomi￾croscopy during the tensile test. To determine the value of t, the number of cracks for a given hysteresis-loop peak stress must be known, because crack saturation for most of the hys￾teresis-loop peak stresses had not been reached. If the cumu￾lative AE energy could be related to the cracks produced, N (Eq. (6)) could be approximated for a given hysteresis-loop Fig. 14. SEM micrographs of fiber surfaces for (a) Hi-Nicalon and (b) Sylramic as-produced fibers. January 1999 Fiber Effects on Minicomposite Mechanical Properties for Several SiC/CVI-SiC Matrix Systems 153