January 1999 Fiber Effects on Minicomposite Mechanical Properties for Several SiC/cv1-SiC Matrix Systems d measured at the minicomposite failure stress, respectively. Similarly, EA(o)and EAF(out are the cumulative AE ener- gies for the same sample measured at the peak hysteresis-loop Sylramic-PBN stress of interest and measured at the minicomposite failure stress, respectively. Equation(6) was used to best fit the stress- shear stresses of 10 and 60 MPa best fit the HN-pbn and Syl-PBN data, respectively, for the estimated number of cracks 5 30 ormed during the experiment. These values were very similar to those found by using the hysteresis-loop technique for the HN-PBN (15+ 10 MPa over the entire stress range)and Syl- PBN (65 MPa at low stresses) minicomposites Hi-NiC-PBN x Also plotted in Fig, 10 are predicted stress-strain data for value greater than three times the best-fit values. The Sy Pbn predicted curves change much less, absolutely and rela- 400 tively, compared to the HN- pbn predicted curves, as t in- Hysteresis L。opP。 ak Stres$,MPa creases. This observation implies that the estimation of T from such a model is more accurate for the HN-PBN minicomposite. Fig, II of peak applied stress on the interfacial shear stress Conversely, estimating T or predicting the stress-strain curve same incomposite for the Syl-PBN minicomposite using the above-mentioned ncreasing peak-stress hysteresis loops. approach, is not very good. It will be argued later in this paper that the stress-strain behavior of the Syl-PBN minicomposite cannot be properly modeled with a constant T, because T resistance due to fiber roughness, 2 122(ii)a large interface hanges with o for the Syl-PBN minicomposite debond energy, 19,23 or(ii) overlapping sliding zones. 18,24 If a fiber-roughness effect was responsible for the increase in the 3) Dependence of T on Composite Stress apparent""T, a real increase in sliding resistance would have T was determined from the hysteresis loops just prior to occurred and T would not be constant along the sliding length. failure. For the Syl-PBN minicomposite, a relatively low-peak A large debond energy or overlapping sliding zones both as- stress hysteresis loop(o 330 MPa)resulted in a lower T sume a constant T value. The measured increase in T for these (65 MPa) compared to that for a relatively high-peak-stress two situations would be due to sliding lengths that do not hysteresis loop(T=170 MPa for p =460 MPa). There was increase with stress as expected at a given crack because the ny difference in T values determined for different debond lengths are hindered from attaining the assumed sliding ak-stress hysteresis loops of HN-PBN or Nic-3MBN mini- lengths(Eq (5), because of the high interfacial fracture energy omposites. To determine if this was a variation in stress or or interaction with another debonded zone, respectively ample-to-sample variation, T was determined for lower-peak An increase in the measured sliding resistance due to over- stress hysteresis loops of the sample that failed at 460 MPa lapping sliding zones is unlikely because most of the cracking The crack spacing was estimated from the AE energy, as de occurred near failure(the highest 5% of stress). The load on the cribed previously and in the Appendix. Figure 1l shows T alues that have been determined from several hysteresis loop were bridging cracks; therefore, the sliding lengths could only for several Syl-PBN and HN-PBN minicomposites. There is be restricted over a very small stress range. The net reduction ittle difference in T for the HN-PBN minicomposite, as a func- in loop width(increased T)would be minimal for the case of a tion of stress, whereas the Syl-PBN minicomposite shows an restricted sliding length. In addition, the le average fiber pull-out almost-linear dependence on the minicomposite peak stress lengths are almost an order of magnitude less than the smallest Several possible explanations exist for the stress-dependent matrix-crack spacings, which indicates a very short sliding dis T behavior of Syl-PBN minicomposites:(i)increased sliding tance. It is possible for short pull-out lengths to be the result of per Weib Weibull modulus for as-produced Sylramic fiber was -5.25 which is a very low value. This result is compared to a Weibull Syl-PBN Predicted Syl modulus of-8 for as-produced Hi-Nicalon fiber. 25 If the fiber/matrix interface is characterized by a large debond energy, the slope of the hysteresis loop would be linear edicted hn at the end of unloading or reloading portion of the loop. 9 If the ■t=30MPa Syl-PBN minicomposite strain scale was increased, there does not appear to be a linear region for the hysteresis loop. How- ever, because these loops are depicted at smaller strains, the Predicted HN sensitivity of the displacement measurement may be inad- equate to distinguish between a linear or parabolic-shaped hys- HN-PBN teresis loop at the end of the unloading or reloading portion the hysteresis loop. This mechanism cannot be excluded for ∧ two reasons. First, it is known that the fiber begins to decom- pose as the carbon-rich Hi-Nicalon fiber is subjected to tem- peratures >1300.C, resulting in a very thin carbon layer on the Model surface of the fiber. 26 The interphase processing temperature Matrix was 1400C for-I min, which could have resulted in a very 010203040.06070800 thin carbon layer between the fiber and the BN. Such a carbo Strain, layer would not form on Sylramic fibers, because these fibers are stoichiometric SiC and are processed at temperatures much higher than the interphase processing temperature. Therefore, IN-PBN minicomposites(the hyster. it is likely that the actual debonding interface for HN-PBN curve has been removed ) The T value was that minicomposites is in the carbon or at the carbon/SiC interface which was The actual debonding interface for Syl-PBN minicompositesand measured at the minicomposite failure stress, respectively. Similarly, EAE(s) and EAE(sult) are the cumulative AE ener￾gies for the same sample measured at the peak hysteresis-loop stress of interest and measured at the minicomposite failure stress, respectively. Equation (6) was used to best fit the stress– strain curve for the Syl-PBN and HN-PBN minicomposites for t. Figure 10 shows the predicted and tensile data. Interfacial shear stresses of 10 and 60 MPa best fit the HN-PBN and Syl-PBN data, respectively, for the estimated number of cracks formed during the experiment. These values were very similar to those found by using the hysteresis-loop technique for the HN-PBN (15 ± 10 MPa over the entire stress range) and Syl￾PBN (65 MPa at low stresses) minicomposites. Also plotted in Fig. 10 are predicted stress–strain data for a t value greater than three times the best-fit values. The Syl￾PBN predicted curves change much less, absolutely and rela￾tively, compared to the HN-PBN predicted curves, as t in￾creases. This observation implies that the estimation of t from such a model is more accurate for the HN-PBN minicomposite. Conversely, estimating t or predicting the stress–strain curve for the Syl-PBN minicomposite using the above-mentioned approach, is not very good. It will be argued later in this paper that the stress–strain behavior of the Syl-PBN minicomposite cannot be properly modeled with a constant t, because t changes with s for the Syl-PBN minicomposite. (3) Dependence of t on Composite Stress t was determined from the hysteresis loops just prior to failure. For the Syl-PBN minicomposite, a relatively low-peak￾stress hysteresis loop (sp 4 330 MPa) resulted in a lower t (∼65 MPa) compared to that for a relatively high-peak-stress hysteresis loop (t ≈ 170 MPa for sp 4 460 MPa). There was little if any difference in t values determined for different peak-stress hysteresis loops of HN-PBN or Nic-3MBN mini￾composites. To determine if this was a variation in stress or a sample-to-sample variation, t was determined for lower-peak￾stress hysteresis loops of the sample that failed at 460 MPa. The crack spacing was estimated from the AE energy, as de￾scribed previously and in the Appendix. Figure 11 shows t values that have been determined from several hysteresis loops for several Syl-PBN and HN-PBN minicomposites. There is little difference in t for the HN-PBN minicomposite, as a func￾tion of stress, whereas the Syl-PBN minicomposite shows an almost-linear dependence on the minicomposite peak stress. Several possible explanations exist for the stress-dependent t behavior of Syl-PBN minicomposites: (i) increased sliding resistance due to fiber roughness,21,22 (ii) a large interface debond energy,19,23 or (iii) overlapping sliding zones.18,24 If a fiber-roughness effect was responsible for the increase in the ‘‘apparent’’ t, a real increase in sliding resistance would have occurred and t would not be constant along the sliding length. A large debond energy or overlapping sliding zones both as￾sume a constant t value. The measured increase in t for these two situations would be due to sliding lengths that do not increase with stress as expected at a given crack because the debond lengths are hindered from attaining the assumed sliding lengths (Eq. (5)), because of the high interfacial fracture energy or interaction with another debonded zone, respectively. An increase in the measured sliding resistance due to over￾lapping sliding zones is unlikely because most of the cracking occurred near failure (the highest 5% of stress). The load on the fibers did not increase significantly for most of the fibers that were bridging cracks; therefore, the sliding lengths could only be restricted over a very small stress range. The net reduction in loop width (increased t) would be minimal for the case of a restricted sliding length. In addition, the average fiber pull-out lengths are almost an order of magnitude less than the smallest matrix-crack spacings, which indicates a very short sliding dis￾tance. It is possible for short pull-out lengths to be the result of a low t value and high fiber Weibull modulus. However, the Weibull modulus for as-produced Sylramic fiber was ∼5,25 which is a very low value. This result is compared to a Weibull modulus of ∼8 for as-produced Hi-Nicalon fiber.25 If the fiber/matrix interface is characterized by a large debond energy, the slope of the hysteresis loop would be linear at the end of unloading or reloading portion of the loop.19 If the Syl-PBN minicomposite strain scale was increased, there does not appear to be a linear region for the hysteresis loop. How￾ever, because these loops are depicted at smaller strains, the sensitivity of the displacement measurement may be inad￾equate to distinguish between a linear or parabolic-shaped hys￾teresis loop at the end of the unloading or reloading portion of the hysteresis loop. This mechanism cannot be excluded for two reasons. First, it is known that the fiber begins to decom￾pose as the carbon-rich Hi-Nicalon fiber is subjected to tem￾peratures >1300°C, resulting in a very thin carbon layer on the surface of the fiber.26 The interphase processing temperature was 1400°C for ∼1 min,16 which could have resulted in a very thin carbon layer between the fiber and the BN. Such a carbon layer would not form on Sylramic fibers, because these fibers are stoichiometric SiC and are processed at temperatures much higher than the interphase processing temperature.3 Therefore, it is likely that the actual debonding interface for HN-PBN minicomposites is in the carbon or at the carbon/SiC interface. The actual debonding interface for Syl-PBN minicomposites Fig. 10. Measured (solid lines) and modeled (data points) stress– strain behavior of Syl-PBN and HN-PBN minicomposites (the hyster￾esis portion of the curve has been removed). The t value was that which was used in the model. Fig. 11. Effect of peak applied stress on the interfacial shear stress. Each set of data points has been taken from the same minicomposite tensile test for increasing peak-stress hysteresis loops. January 1999 Fiber Effects on Minicomposite Mechanical Properties for Several SiC/CVI-SiC Matrix Systems 151