2628 Martine. dez, G N. Morscher /Journal of the European Ceraic Society 20 (2000)2627-2636 the fiber fracture mirrors were observed to increase how much worse, if at all, is C interphase HN/ SiC in commensurate with the degree of minicomposite rup rupture than BN interphase HN/SiC? Therefore, the Ire strength loss. Therefore, it is evident that the impetus for to determine the rupture mechanisms causing NIC/SiC rupture with carbon behavior of minicomposites with C interphases interphases include a flaw growth or flaw creation to provide ison for HN/SiC minicomposites mechanism in addition to the two already mentioned with BN interphases as well as NIC/Sic minicomposites It is likely that the surface of the NIC fiber may have with C interphases b n modified during CVI SiC infiltration. Naslain0 describes four NIC/C/CVI SiC composite scenarios where the surface of the Si-C-O containing NIC fibers 2. Experimental procedure are altered after CVI SiC composite fabrication. All four scenarios have a complex carbon-rich layer on the Tows of 500 hn fibers were used to fabricate the fiber surface in between the Cvi deposited carbon layer minicomposites studied in this work. The tows were and the Si-C-O fiber. In some cases, SiO2 is present and mounted on graphite racks, coated with carbon and in others it is not depending on the CVI approach taken then composited with Sic by chemical vapor infiltration and/or a fiber pretreatment. Nevertheless, local carbon( Hyper-Therm Inc, Huntington Beach, CA). The fiber rich areas at the fiber surface are present that woul volume fraction(0.16+0.01)and minicomposite cross- oxidize during rupture testing. In other words, sectional area was determined based on the measured or mechanism due to surface recession or"pit formation estimated weights and densities of the minicomposite as a result of the oxidation of carbon layers or local constituents. 3 The observation of the minicomposite carbon rich regions on the fiber surface could lead to polished cross-sections indicated that the carbon inter rapid strength loss. This type of mechanism better phase was uniform with an average thickness of 0. 4 um explains the observed rapid rupture strength loss with The fracture surfaces of the minicomposites had a time for NIC/C/SiC composites. It is not known if the fibrous appearance. the Sic matrix being thicker on the more thermally stable Hi-NicalonM(HN)fibers withC outside but more uniformly distributed in the interior interphases would undergo similar degradation after than in the NIC fiber, C-interphase minicomposites CVI SiC processing previously studied. 2 In the earlier study, the minicomposite tensile stress Room temperature tensile testing was performed rupture properties of NIC reinforced SiC with a BN using an universal testing machine (Model 4502, interphase had superior elevated temperature stress- Instron, Canton, MA); the test set-up is described in rupture life in air compared to C interphase mini- detail elsewhere. 2 The minicomposites were mounted composites. This has also been shown for woven NIC/ onto cardboard tabs with epoxy. Monotonic loading Sic composites tested in flexure. It was suggested in tensile tests were performed to determine the ultimate Ref. 2 that the BN interphase NIC/SiC rupture beha- failure load and the optimum precrack load. Modal vior was superior to C interphase NIC/SiC because of acoustic emission(AE)was monitored with sensors the formation of a measurable oxide layer (0.5 H attached to the epoxy just above and below the mini- 700C after 12 h)on the fibers that had a bn inter- composite gage-section. The AE analyzer(Digital Wave phase. This was due to the enhanced oxidation of Sic Corporation, Englewood, CO) recorded and digitized when in contact with BN. This oxide layer, which was the true sound wave form for each event on both chan predominantly SiOz, protected the fibers from the nels(sensors). The number of events and location of strength degrading mechanism of the NIC fibers when each event could then be determined once the speed of directly exposed to the environment as was the case for sound was estimated The details of the ae analysis are C interphase NIC/SiC. The presence of thick oxide included in the Appendix A. scales on fibers in BN interphase NIC/SiC causing only Constant-load stress-rupture tests were run in a dead- minimal strength degradation (n-50) also is further weight load stress-rupture rig. A furnace was located at evidence against the oxide scale mechanism discussed the center of the minicomposite. The total length of the above for NIC/SiC with C interphases. It was also furnace was 35 mm with a hot zone of 12 mm. These hown in Ref 2 that the rupture behavior of Hn/BN/ tests were run at 700, 950, and 1200C. The mini- iC minicomposites were superior to NIC/BN/SiC composites were fully loaded before increasing the tem minicomposites in air minature at 100C/min up to the test temperature. The Unfortunately, in the earlier study, HN/SiC mini- minicomposites were precracked at room temperature composites with C interphases were not studied. Since with loads of 119 or 126 N, which corresponds to a HN is a more thermally stable fiber than NIC, it may be composite stress of 280 or 295 MPa, respectively expected that HN SiC with C interphases is not as sus- The fatigue tests were als o run at700,950,and ceptible to severe rupture-strength degradation as NIC/ 1200C with the universal testing machine and same SiC with C interphases. Thus, the question remains, furnace set-up. The load was cycled from a minimumthe ®ber fracture mirrors were observed to increase commensurate with the degree of minicomposite rup￾ture strength loss. Therefore, it is evident that the mechanisms causing NIC/SiC rupture with carbon interphases include a ¯aw growth or ¯aw creation mechanism in addition to the two already mentioned. It is likely that the surface of the NIC ®ber may have been modi®ed during CVI SiC in®ltration. Naslain10 describes four NIC/C/CVI SiC composite scenarios where the surface of the Si±C±O containing NIC ®bers are altered after CVI SiC composite fabrication. All four scenarios have a complex carbon-rich layer on the ®ber surface in between the CVI deposited carbon layer and the Si±C±O ®ber. In some cases, SiO2 is present and in others it is not depending on the CVI approach taken and/or a ®ber pretreatment. Nevertheless, local carbon rich areas at the ®ber surface are present that would oxidize during rupture testing. In other words, a mechanism due to surface recession or ``pit formation'' as a result of the oxidation of carbon layers or local carbon rich regions on the ®ber surface could lead to rapid strength loss. This type of mechanism better explains the observed rapid rupture strength loss with time for NIC/C/SiC composites. It is not known if the more thermally stable Hi-NicalonTM (HN) ®bers with C interphases would undergo similar degradation after CVI SiC processing. In the earlier study,2 the minicomposite tensile stress rupture properties of NIC reinforced SiC with a BN interphase had superior elevated temperature stress￾rupture life in air compared to C interphase mini￾composites. This has also been shown for woven NIC/ SiC composites tested in ¯exure.11 It was suggested in Ref. 2 that the BN interphase NIC/SiC rupture beha￾vior was superior to C interphase NIC/SiC because of the formation of a measurable oxide layer (0.5 m at 700C after 12 h) on the ®bers that had a BN inter￾phase. This was due to the enhanced oxidation of SiC when in contact with BN.12 This oxide layer, which was predominantly SiO2, protected the ®bers from the strength degrading mechanism of the NIC ®bers when directly exposed to the environment as was the case for C interphase NIC/SiC. The presence of thick oxide scales on ®bers in BN interphase NIC/SiC causing only minimal strength degradation (n50) also is further evidence against the oxide scale mechanism discussed above for NIC/SiC with C interphases. It was also shown in Ref. 2 that the rupture behavior of HN/BN/ SiC minicomposites were superior to NIC/BN/SiC minicomposites in air. Unfortunately, in the earlier study, HN/SiC mini￾composites with C interphases were not studied. Since HN is a more thermally stable ®ber than NIC, it may be expected that HN/SiC with C interphases is not as sus￾ceptible to severe rupture-strength degradation as NIC/ SiC with C interphases. Thus, the question remains, how much worse, if at all, is C interphase HN/SiC in rupture than BN interphase HN/SiC? Therefore, the impetus for this work was to determine the rupture behavior of HN/SiC minicomposites with C interphases to provide a comparison for HN/SiC minicomposites with BN interphases as well as NIC/SiC minicomposites with C interphases. 2. Experimental procedure Tows of 500 HN ®bers were used to fabricate the minicomposites studied in this work. The tows were mounted on graphite racks, coated with carbon and then composited with SiC by chemical vapor in®ltration (Hyper-Therm Inc., Huntington Beach, CA). The ®ber volume fraction (0.160.01) and minicomposite cross￾sectional area was determined based on the measured or estimated weights and densities of the minicomposite constituents.13 The observation of the minicomposite polished cross-sections indicated that the carbon inter￾phase was uniform with an average thickness of 0.4 mm. The fracture surfaces of the minicomposites had a ®brous appearance, the SiC matrix being thicker on the outside, but more uniformly distributed in the interior than in the NIC ®ber, C-interphase minicomposites previously studied.2 Room temperature tensile testing was performed using an universal testing machine (Model 4502, Instron, Canton, MA); the test set-up is described in detail elsewhere.2 The minicomposites were mounted onto cardboard tabs with epoxy. Monotonic loading tensile tests were performed to determine the ultimate failure load and the optimum precrack load. Modal acoustic emission (AE) was monitored with sensors attached to the epoxy just above and below the mini￾composite gage-section. The AE analyzer (Digital Wave Corporation, Englewood, CO) recorded and digitized the true sound wave form for each event on both chan￾nels (sensors). The number of events and location of each event could then be determined once the speed of sound was estimated. The details of the AE analysis are included in the Appendix A. Constant-load stress-rupture tests were run in a dead￾weight load stress-rupture rig. A furnace was located at the center of the minicomposite. The total length of the furnace was 35 mm with a hot zone of 12 mm. These tests were run at 700, 950, and 1200C. The mini￾composites were fully loaded before increasing the tem￾perature at 100C/min up to the test temperature. The minicomposites were precracked at room temperature with loads of 119 or 126 N, which corresponds to a composite stress of 280 or 295 MPa, respectively. The fatigue tests were also run at 700, 950, and 1200C with the universal testing machine and same furnace set-up. The load was cycled from a minimum 2628 J. Marti nez-FernaÂndez, G.N. Morscher / Journal of the European Ceramic Society 20 (2000) 2627±2636