dez, G N. Morscher /Journal of the Europ amic Society 20 (2000)2627-2636 HN As-produced HN minicomposites at 1200C still could be due to a (o RT, 2800 MPa) slower recession of the bn interface 4.2. Dependence of the stress-rupture behavior on precrack load The load-time curve of Fig. I indicates that cracks are being created in the minicomposite up to failure. It clear then that the increase of the precrack load from ° C Experiments9so" CExperiments1200°c 280 to 295 MPa implies an increase in the number of 000 35000 40000 matrix cracks. Based on the direct relationship betwee the number of cracks and AE events(Table Al)and the Fig.8. Plot of the stress on fibers if fully loaded vS LM parameter. number of events recorded for several samples as Previous data included. See text for further discussion function of load(Fig. 2), there would be one or two cracks in the hot zone after the 119n precrack load For the 126 n precrack load, between two and four kp is 10-m2s-,and approximately 6 mm of carbon cracks would be estimated in the hot zone. A larger interphase length would be lost in about I h. This would number of cracks cause faster oxidation of the entire correspond to the entire hot zone gage length for a phase in the hot zone. This results in a larger length minicomposite with one crack in the hot zone region. of the fiber holding the maximum load (at least for short The fiber length exposed to the maximum load when the tests), more environmental degradation of the fibers, interphase was completely oxidized was about 25 times and more locations for stress concentration associated larger than the fiber length if there was no oxidation of with SiOz formation and the strong bonding between the interface(assuming one crack, as in the Introduc- fiber and matrix. all these factors would cause a decrease ion). This difference in load-bearing length will result inin the stress-rupture survival time. During the precrack decrease in strength of 47 to 28% for a Weibull mod- loading step it was rather common(approximately 25% ulus from 5 to 10. For more than one crack in the hot of the cases) for composites to fail between 290 to 295 zone, a smaller reduction in strength would occur. This MPa. This indicates that, in addition to an increase in the gage-length effect best explains the decrease of strength crack density, a greater occurrence of fiber failure would of C-Hn compared with BN-Hn since the rupture be occurring in minicomposites precracked at 295 MPa strength of the C-Hn minicomposites are approxi than at 280 MPa, and could also contribute to the poorer mately 25% reduced compared to the shorter gage stress-rupture behavior compared to samples precracked length loaded BN-HN minicomposites at 280 MPa(Fig. 6). At 1200oC, failure is dominated by The presence of fibers with no pullout on the fracture the creep-rupture behavior of the fibers, explaining the surface was observed for some of the fibers on the smaller dependence of the survival time with precrack minicomposite fracture surfaces tested at 700C and load for this temperature occurred for all of the fibers on the minicomposite frac ture surfaces tested at 950C. This indicates that strong 43. Fatigue bonding of the fibers to the matrix was a factor at these temperatures. When one or more fibers failed at or The results from fatigue experiments(Fig. 5)clearly away from the matrix crack, the added load to the indicate that it is at intermediate temperatures when the neighboring fibers in the matrix crack would be relative movement of matrix and fiber, due to fatigue, enhanced due to strong bonding resulting in local stress- affect the survival time most dramatically. This is pre- concentrations. This causes most of the fibers to pre- sumably due to the local concentration of stress during Terentially fail in the plane of the matrix crack. The fatigue on the sites where the fiber and matrix are bon longer recession distance with C interphases not only ded because of the Sio2 formation. At 1200 C, the increases the fiber gage length, it also allows fibers to minicomposite failed about 2 cm from the center of the bond more quickly compared to BN interphases. The furnace. In these regions the temperature is approxi- BN interphase only recesses a few microns separating mately 900@C. This shows that the resistance to fatigue the fiber from the matrix. For C interphase mini- is worse at intermediate temperatures than at 1200 C, omposites, after total C interphase removal, the fibers even though a greater amount of oxide reaction produc are free to move towards and contact the matrix is formed at 1200 C in the matrix cracks than at inter. At 1200 C, SiO2 formation is extensive. The stress- mediate temperatures. SiO2 at 1200C does flow to some upture properties appear to be controlled by the creep extent and may relieve some of the stress-concentrations of the fibers and are similar for both minicomposite produced at the fiber-SiOr-matrix bond compared to ystems. The differences between C and BN interphase lower temperatures where no relaxation of the glasskp is 10ÿ8 m2 sÿ1 , 20 and approximately 6 mm of carbon interphase length would be lost in about 1 h. This would correspond to the entire hot zone gage length for a minicomposite with one crack in the hot zone region. The ®ber length exposed to the maximum load when the interphase was completely oxidized was about 25 times larger than the ®ber length if there was no oxidation of the interface (assuming one crack, as in the Introduc￾tion). This di€erence in load-bearing length will result in a decrease in strength of 47 to 28% for a Weibull mod￾ulus from 5 to 10. For more than one crack in the hot zone, a smaller reduction in strength would occur. This gage-length e€ect best explains the decrease of strength of C±HN compared with BN±HN since the rupture strength of the C±HN minicomposites are approxi￾mately 25% reduced compared to the shorter gage￾length loaded BN±HN minicomposites. The presence of ®bers with no pullout on the fracture surface was observed for some of the ®bers on the minicomposite fracture surfaces tested at 700C and occurred for all of the ®bers on the minicomposite frac￾ture surfaces tested at 950C. This indicates that strong bonding of the ®bers to the matrix was a factor at these temperatures. When one or more ®bers failed at or away from the matrix crack, the added load to the neighboring ®bers in the matrix crack would be enhanced due to strong bonding resulting in local stress￾concentrations. This causes most of the ®bers to pre￾ferentially fail in the plane of the matrix crack. The longer recession distance with C interphases not only increases the ®ber gage length, it also allows ®bers to bond more quickly compared to BN interphases. The BN interphase only recesses a few microns separating the ®ber from the matrix. For C interphase mini￾composites, after total C interphase removal, the ®bers are free to move towards and contact the matrix. At 1200C, SiO2 formation is extensive. The stress￾rupture properties appear to be controlled by the creep of the ®bers and are similar for both minicomposite systems. The di€erences between C and BN interphase HN minicomposites at 1200C still could be due to a slower recession of the BN interface. 4.2. Dependence of the stress±rupture behavior on precrack load The load±time curve of Fig. 1 indicates that cracks are being created in the minicomposite up to failure. It is clear then that the increase of the precrack load from 280 to 295 MPa implies an increase in the number of matrix cracks. Based on the direct relationship between the number of cracks and AE events (Table A1) and the number of events recorded for several samples as a function of load (Fig. 2), there would be one or two cracks in the hot zone after the 119 N precrack load. For the 126 N precrack load, between two and four cracks would be estimated in the hot zone. A larger number of cracks cause faster oxidation of the entire interphase in the hot zone. This results in a larger length of the ®ber holding the maximum load (at least for short tests), more environmental degradation of the ®bers, and more locations for stress concentration associated with SiO2 formation and the strong bonding between ®ber and matrix. All these factors would cause a decrease in the stress-rupture survival time. During the precrack loading step, it was rather common (approximately 25% of the cases) for composites to fail between 290 to 295 MPa. This indicates that, in addition to an increase in the crack density, a greater occurrence of ®ber failure would be occurring in minicomposites precracked at 295 MPa than at 280 MPa, and could also contribute to the poorer stress-rupture behavior compared to samples precracked at 280 MPa (Fig. 6). At 1200C, failure is dominated by the creep-rupture behavior of the ®bers, explaining the smaller dependence of the survival time with precrack load for this temperature. 4.3. Fatigue The results from fatigue experiments (Fig. 5) clearly indicate that it is at intermediate temperatures when the relative movement of matrix and ®ber, due to fatigue, a€ect the survival time most dramatically. This is pre￾sumably due to the local concentration of stress during fatigue on the sites where the ®ber and matrix are bon￾ded because of the SiO2 formation. At 1200C, the minicomposite failed about 2 cm from the center of the furnace. In these regions the temperature is approxi￾mately 900C. This shows that the resistance to fatigue is worse at intermediate temperatures than at 1200C, even though a greater amount of oxide reaction product is formed at 1200C in the matrix cracks than at inter￾mediate temperatures. SiO2 at 1200C does ¯ow to some extent and may relieve some of the stress-concentrations produced at the ®ber±SiO2±matrix bond compared to lower temperatures where no relaxation of the glass Fig. 8. Plot of the stress on ®bers if fully loaded vs LM parameter. Previous data included.1 See text for further discussion. J. Marti nez-FernaÂndez, G.N. Morscher / Journal of the European Ceramic Society 20 (2000) 2627±2636 2633