Joumal of the American Ceramic Sociery-Morscher et al. Vol. 87. No. I (1) The increased strain to failure of outside-debonding composi can be attributed to the lower t of the bn-cvi sic interface over that of the fiber/BN interface. However, if T decreased and global load sharing exists, one would expect the ultimate strength properties to decrease. "4 The converse has been observed for strongly bonded interfaces compared with weakly bonded inter Ined Debonding faces in CVI SiC matrix composites where the higher T interface composites exhibit higher ultimate strengths. In this study, the lower T composites did not lose strength and in some cases were stronger than comparable high-T composites. Two explanations 0 can account for this. First, it is possible that high-T composites TIme, hr exhibit local load sharing and the lower-- composites exhibit global load sharing. If the high-T composites exhibit local loac 8. Stress-rupture of as-received and precracked SYL and SYL-iBn sharing, stress concentrations would develop for load-bearing IC composites at 815C in air. fibers surrounding individual or groups of broken fibers in a matrix crack. This would result in lower composite ultimate strengths than expected based on global load sharing.2 Second, global load sharing may occur for both low-and high- composites, and as Xi ome fracture surfaces were examined from burner-nig-exposed and Curtin" have theorized, fibers with an adhered coating would SYL and SYL-ibN composites. The composites exhibited long be effectively stronger than fibers without a coating because the allout lengths similar to as-produced specimens(Fig. 4(a)and a flaws on the fiber surface are constrained to some degree by the Sio2-containing layer was often observed on the surface of the Bn coating, even for low-modulus coatings such as C or BN. In fact in between the BN and the matrix throughout the cross section their model would predict approximately the same composite (Fig. 10). Evidently, oxidation occurred through the BN/CVI Sic strength for outside-debonding composites with a T= 10 MPa and interface region at the exposed cut edge into the interior of the inside-debonding composites with a T =70 MPa composite. HNS composites exhibited a flat fracture surface and strong bonding of fibers as has been reported for other systems The improved intermediate-temperature rupt with carbon layers that exist at the fiber surface. 3 bonding composites occurs in the manner put forward in Fi 1(b)(Fig. 7). Even after 100 h at 815C in a bridged matrix crack, a significant portion of the bn remained as a barier between the (3)Analysis of the BN-CVI SiC interface oxidation reaction product and the fiber surface for the majority of aes depth profiles were conducted on specimen surfaces that fiber circumference. However, thinner regions of BN separating were fractured in the aES chamber for several specimens exhib- nearest-neighbor fibers were oxidized and appear to have led to the iting inside and outside debonding. Depth profiles through BN time-dependent strength degradation. This would explain why the layers adhered to the CVi SiC matrix were performed(not shown). Sylramic composites with outside debonding are poorer in A mild enrichment of C appeared to exist at the BN-CVI Sic stress-rupture than the Sylramic- iBN composites( Fig. 7).The interface for both inside-and outside-debonding specimens. How former consists of many fibers nearly contacting another fiber with er, no difference in the amount of C enrichment at the bn-cVI little or no bn interphase in between, whereas the latter possesses SiC interface for inside-and outside-debonding composites could the in situ BN layers on the fiber surface that enable greater be discerned given the error in the aEs measurement (-10%). protection of the fibers as well as some degree of fiber separation Representative TEM micrographs of inside- and outside- Finally, for typical inside-debonding composites that fail at sig bonding specimens are shown in Fi of the nificantly shorter lives and lower stresses at the same temperature ame region are also shown. There does appear to be some c no bn was detectable in the oxidized nrichment at the BN-CVi SiC interface for the outside-debonding (see, e.g, Refs. 1 and 27). In other words the entire matrix and composites and little if any C enrichment for the inside-debonding interphase region of inside-debonding composites would be com mposites.However, this cannot be considered conclusive evi- pletely oxidized, with the fibers strongly bonded to the matrix dence of a C-layer (2)Why Outside Debonding Two potential mechanisms are considered for outside debond ng for these MI composite systems:(1)a weaker BN-CVI Sic 450 nterface than BN-fiber interface and(2)sufficient residual stress 400 at the interface to cause debonding of the weak interface probabl as-produced on cooling after infiltration of molten Si composites possess a lower T than inside-debonding composites interface is also lower than the debond energy of the fiber/BN nterface as well. Residual compression exists in the matrix( Fig 2), which presumably forces the fibers into residual tension. Thi 100 is due to free Si. The volume expansion of Si from the liquid to solid state is -9%. Therefore, expansion of the Si phase takes 50 place during cooling of the composite from its fabrication temper ature for MI(1400C depending on the additives to the Si). this 1020304050.6 places the Si in compression. Si also has a lower thermal expansion coefficient than SiC, -3 X 10rC compared with -4. x Strain. 10/C, respectively. Therefore, on further cooling, the Si placed in further compression. The crack closure effect( Fig. 2) side debonding. The HNS stress-strain curves are offset on the strain axis fiber/interphase bundles taken together are in residual compression for clarity necessitating residual tension in the fibers. Locally, the residual110 Journal of the American Ceramic Society--Morscher et al. Vol. 87, No. 1 6 170 1 150 4 I 0 100 200 300 400 lime, hr Fig. 8. Stress-rupture of as-received and precracked SYL and SYL-iBN SiClSiC composites at 815°C in air. Some fracture surfaces were examined from burner-rig-exposed SYL and SYL-iBN composites. The composites exhibited long pullout lengths similar to as-produced specimens (Fig. 4(a)) and a SO,-containing layer was often observed on the surface of the BN in between the BN and the matrix throughout the cross section (Fig. 10). Evidently, oxidation occurred through the BN/CVI Sic interface region at the exposed cut edge into the interior of the composite. HNS composites exhibited a flat fracture surface and strong bonding of fibers as has been reported for other systems with carbon layers that exist at the fiber surface.23 (3) Analysis of the BN-CVZ Sic Interface AES depth profiles were conducted on specimen surfaces that were fractured in the AES chamber for several specimens exhib￾iting inside and outside debonding. Depth profiles through BN layers adhered to the CVI Sic matrix were performed (not shown). A mild enrichment of C appeared to exist at the BN-CVI Sic interface for both inside- and outside-debonding specimens. How￾ever, no difference in the amount of C enrichment at the BN-CVI Sic interface for inside- and outside-debonding composites could be discerned given the error in the AES measurement (-10%). Representative TEM micrographs of inside- and outside￾debonding specimens are shown in Fig. 11. Carbon maps of the same region are also shown. There does appear to be some C enrichment at the BN-CVI Sic interface for the outside-debonding composites and little if any C enrichment for the inside-debonding composites. However, this cannot be considered conclusive evi￾dence of a C-layer. 500 1 450 400 2 350 B 300 # 250 150 100 50 0 4 200 I 0 0.1 0.2 0.3 0.4 0.5 0.6 Strain, % Fig. 9. Room-temperature tensile stress-strain curves of as-received and burner-rig-exposed SYL-iBN and HNS SiC/SiC specimens showing out￾side debonding. The HNS stress-strain curves are offset on the strain axis for clarity. IV. Discussion (1) Zmproved Mechanical Properties The increased strain to failure of outside-debondjng composites can be attributed to the lower 7 of the BN-CVI Sic interface over that of the fiber/BN interface. However, if 7 decreased and global load sharing exists, one would expect the ultimate strength properties to decrease.24 The converse has been observed for strongly bonded interfaces compared with weakly bonded inter￾faces in CVI Sic matrix composites where the higher 7 interface composites exhibit higher ultimate strengths5 In *his study, the lower T composites did not lose strength and in some cases were stronger than comparable high-? composites. Two explanations can account for this. First, it is possible that high-? composites exhibit local load sharing and the lower-? composites exhibit global load sharing. If the high-? composites exhibit local load sharing, stress concentrations would develop for load-bearing fibers surrounding individual or groups of broken fi6ers in a matrix crack. This would result in lower composite ultimate strengths than expected based on global load sharing.25 Second, global load sharing may occur for both low- and high-? composites, and as Xia and CurtinZ6 have theorized, fibers with an adhered coating would be effectively stronger than fibers without a coating because the flaws on the fiber surface are constrained to some degree by the coating, even for low-modulus coatings such as C or BN. In fact, their model would predict approximately the same composite strength for outside-debonding composites with a 7 = 10 MPa and inside-debonding composites with a T = 70 MPa. The improved intermediate-temperature rupture life of outside￾debonding composites occurs in the manner put forward in Fig. l(b) (Fig. 7). Even after 100 h at 815°C in a bridged matrix crack, a significant portion of the BN remained as a barrier between the oxidation reaction product and the fiber surface for the majority of fiber circumference. However, thinner regions of BN separating nearest-neighbor fibers were oxidized and appear to have led to the time-dependent strength degradation. This would explain why the Sylramic@ composites with outside debonding are poorer in stress-rupture than the Sylramic-iBN composites (Fig. 7). The former consists of many fibers nearly contacting another fiber with little or no BN interphase in between, whereas the latter possesses the in situ BN layers on the fiber surface that enable greater protection of the fibers as well as some degree of fiber separation. Finally, for typical inside-debonding composites that fail at sig￾nificantly shorter lives and lower stresses at the same temperature, no BN was detectable in the oxidized portion of a fracture surface (see, e.g., Refs. 1 and 27). In other words, the entire matrix and interphase region of inside-debonding composites would be com￾pletely oxidized, with the fibers strongly bonded to the matrix. (2) Why Outside Debonding? Two potential mechanisms are considered for outside debond￾ing for these MI composite systems: (1) a weaker BN-CVI Sic interface than BN-fiber interface and (2) sufficient residual stress at the interface to cause debonding of the weak interface probably on cooling after infiltration of molten Si. The outside-debonding composites possess a lower 7 than inside-debonding composites (Table I); presumably, the debond energy of the BN/CVI-Sic interface is also lower than the debond energy of the fiberlSN interface as well. Residual compression exists in the matrix (Fig. 2), which presumably forces the fibers into residual tension. This is due to free Si. The volume expansion of Si from the liquid to solid state is -9%. Therefore, expansion of the Si phase takes place during cooling of the composite from its fabrication temper￾ature for MI (- 1400°C depending on the additives to the Si). This places the Si in compression. Si also has a lower thermal expansion coefficient than Sic, -3 X 10-6/oC compared with -4.5 X 10-6/oC, respectively. Therefore, on further cooling, the Si is placed in further compression. The crack closure effect (Fig. 2) reflects the global residual stress state of the entire matrix; i.e., the free Si, the particulate Sic, the CVI Sic, and the unbridged 90" fibedinterphase bundles taken together are in residual compression necessitating residual tension in the fibers. Locally, the residual