wwceramics. org/ACT SiC Fiber-Reinforced MI SiC Composites 157 mental degradation at elevated temperatures. This onset failure Table ID)multiplied by the volume fraction of ominimatrix, 95 MPa, was about the same as that found fibers in the loading direction(Table D). For the lower in morscher, 16 modulus fiber composites, slightly lower t values were When multiplied by the final matrix crack density, determined as would be expected for smoother surfaces Eq(4)can be used to derive a general relationship for and finer grain sizes of these fiber types. However, for composites. For example, based on Curtin and col- 23+5 MPa. In an earlier study for MI composites fab- strain can be described ricated with lower fiber volume fraction SA lbers the E=o/Ec +adpc/Er(o+oth) best-fit t was found to be between 40 and 50 MPa' For (4) similar Syl-iBN composites from a different vendor, t for pe>28 was found to be x 70 MPa. This difference in t is probably best explained by the thicker BN interphase where the stress-dependent crack density, Pe, can be thicknesses of this study(<1.0+0. 1 um)compared found from Eq (3)multiplied by the final crack density with the thinner and converted back into composite stress, the stress- the other studies S N interphases used(<0.5um)in ndent sliding length the minimatrix approach described above, it 6=r(+oth)/21 (5) follows that a simple relationship for stress-strain be- where r is the fiber radius, t is the interfacial shear stress havior and matrix cracking can be established for a wide range of 2D 0/90 MI woven composite systems. The approach assumes that the matrix porosity is low enough a=(1-fEm/fE (6 so that matrix cracks emanate from the 90% minicompos Em, the elastic modulus of the minimatrix, is assumed to ites in the 2D architecture and that this cracking mech- be everything in the composite other than the load nism is negligibly dependent on specimen width, length bearing fibers and can be determined from rule of mix- and thickness (volume effects). It is shown here that this understanding can also be applied to creep-rupture prop- cures knowing fo, Et, and Ee. The only parameter not erties, as discussed in the next section known is t which can be determined by best fitting Eq 5)to empirical Figure 1b(circles) shows predicte stress-strain urves using the best-fit t values listed in Table Ill, Elevated Temperature Creep-Rupture Behavior which also shows the measured matrix crack densities Tensile creep-rupture studies were performed in The maximum stress circle for the predicted stress- ambient air at 1200C for ZMI, SA, and Syl-iBN com- strain curves in Fig. 1b is the average fiber strength at posites and at 1315C for SA and Syl-iBN composites Table Ill. Matrix Crack Density and Interfacial Representative data on total strain versus time data are shown in Fig.6at1200°Cand1315°C. The creep curves are dominated by a decreasing strain rate. a clear Final measured Best fit steady state or constant creep-rate region was never crack density(#/mm) t(MPa) achieved in the creep tests used for most specimens in this study. For specimens that failed in rupture during SYLiBN-1 25 SYLiBN-2 the test, a region of increasing strain rate just before rupture was typically observed. For specimens that sur- tes SA 7.8 H stopped and either a fast-fracture test was performed Z-1 at the creep temperature or the specimen was cooled n, remove HNS-1* rig, and tested at room tem- similar unload-reload and ae monitor- HNS-2 ing as used for the as-produced specimens. A limited DiCarlo et al° number of specimens were tested for times >100 h tomental degradation at elevated temperatures. This onset sminimatrix, B95 MPa, was about the same as that found in Morscher.16 When multiplied by the final matrix crack density, Eq. (4) can be used to derive a general relationship for matrix cracking in these 2D woven composites and can then be used to model stress/strain behavior for these composites.7 For example, based on Curtin and col￾leagues,17,18 strain can be described as e ¼ s=Ec þ adrc=Efðs þ sthÞ; for r1 c > 2d ð4Þ where the stress-dependent crack density, rc, can be found from Eq. (3) multiplied by the final crack density and converted back into composite stress, the stress￾dependent sliding length d ¼ arðs þ sthÞ=2t ð5Þ where r is the fiber radius, t is the interfacial shear stress, and a ¼ ð1 f ÞEm=fEc ð6Þ Em, the elastic modulus of the minimatrix, is assumed to be everything in the composite other than the load￾bearing fibers and can be determined from rule of mix￾tures knowing fo, Ef, and Ec. The only parameter not known is t which can be determined by best fitting Eq. (5) to empirical stress–strain behavior. Figure 1b (circles) shows predicted stress–strain curves using the best-fit t values listed in Table III, which also shows the measured matrix crack densities. The maximum stress circle for the predicted stress– strain curves in Fig. 1b is the average fiber strength at failure (Table II) multiplied by the volume fraction of fibers in the loading direction (Table I). For the lower modulus fiber composites, slightly lower t values were determined as would be expected for smoother surfaces and finer grain sizes of these fiber types. However, for the composites reinforced with the polycrystalline fibers (Syl-iBN and SA), the t values were unexpectedly low at 2375 MPa. In an earlier study for MI composites fab￾ricated with lower fiber volume fraction SA fibers, the best-fit t was found to be between 40 and 50 MPa.5 For similar Syl-iBN composites from a different vendor,7 t was found to be B70 MPa. This difference in t is probably best explained by the thicker BN interphase thicknesses of this study (B1.070.1 mm) compared with the thinner BN interphases used (r0.5 mm) in the other studies.4,5 Using the minimatrix approach described above, it follows that a simple relationship for stress–strain be￾havior and matrix cracking can be established for a wide range of 2D 0/90 MI woven composite systems. The approach assumes that the matrix porosity is low enough so that matrix cracks emanate from the 901 minicompos￾ites in the 2D architecture and that this cracking mech￾anism is negligibly dependent on specimen width, length, and thickness (volume effects). It is shown here that this understanding can also be applied to creep-rupture prop￾erties, as discussed in the next section. Elevated Temperature Creep-Rupture Behavior Tensile creep-rupture studies were performed in ambient air at 12001C for ZMI, SA, and Syl-iBN com￾posites and at 13151C for SA and Syl-iBN composites. Representative data on total strain versus time data are shown in Fig. 6 at 12001C and 13151C. The creep curves are dominated by a decreasing strain rate. A clear steady state or constant creep-rate region was never achieved in the creep tests used for most specimens in this study. For specimens that failed in rupture during the test, a region of increasing strain rate just before rupture was typically observed. For specimens that sur￾vived the creep test, typically B100 h, the test was stopped and either a fast-fracture test was performed at the creep temperature or the specimen was cooled down, removed from the rig, and tested at room tem￾perature, using similar unload–reload and AE monitor￾ing as used for the as-produced specimens. A limited number of specimens were tested for times 4100 h to Table III. Matrix Crack Density and Interfacial Shear Stress Panel Final measured crack density (#/mm) Best fit s (MPa) SYLiBN-1 7.5 25 SYLiBN-2 8.1 18 SYLiBN-3 7.9 20 SA-1 7.8 28 HN 2 9 Z-1 4.1 20 HNS-1 7.0 36 HNS-2 6.4 20 DiCarlo et al. 6 www.ceramics.org/ACT SiC Fiber-Reinforced MI SiC Composites 157