December 1999 Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites Table 1. Fiber Fracture Surfaces at Table I. In Situ Fiber Strength of Different Temperatures Fracture surface ype I Type Ill alues listed are percentages of fiber. arranged in ascending order, and a corresponding failure prob. wide distribution of fiber-surface defects, likely introduced ability, F=(i-0.5)N, is assigned to each strength, based on during composite processing. This wide distribution was unal- experimental d where i is the rank and N the total number of tered by the oxidation processes on the fiber surface at high nk statistics ta. The lowest strength is attributed to the fi temperature, and the Weibull modulus remained almost con- bers with a specular fracture surface(type Ill), following the stant or increased slightly. Second, the characteristic strength, analysis of Eckel and Bradt. 6 The highest strength is assigned fe, was practically equivalent in the as-received material and in to the fibers with a very short mirror radius(type I). The results the specimens tested at 800 C; thus, the modifications in the are shown in Fig. 6. According to the Weibull statistics, the BN/fiber interlayer region at 800 C were confined to this re fiber fracture probability is given by gion and did not nucleate defects on the fiber. The microstruc- tural characterization of this region, presented above(Figs 81-9-( (2) I(A)and 2(A)), showed no growth in the interlayer thickness after I h at 800C. The only change was the removal of carbon from this layer. This carbon removal should not have affected Here, the parameter of is obtained from Fig. 6, as the stress he fiber strength, a result in agreement with the fiber strength that gives a fracture probability of 63%, and m* is computed by obtained from the fracture mirrors the least-squares fitting of Eq. (2)to the experimental results The interfacial region did grow thicker after I h at 1000%and plotted in Fig. 6. The values of o* and m* derived from the 1200oC, and energy-dispersive spectroscopy indicated that the racture mirror data are not, in general, identical to the true in region was made of SiO,. This layer nucleated fiber-surfac situ fiber characteristic strength, oe, and Weibull modulus, m. of the fibers The values must be corrected to account for the screening of demonstrated earlier on isolated Nicalon fibers tested in air at flaws over a finite fiber length on either side of a fiber failure similar temperatures.23-25,27 The process, and not just the tem- Following the procedure developed by Curtin 17 for the case degradation. This conclusion is supported by a previous inves of multiple matrix cracks, the corrected values of o and were estimated in the present study and are presented in Tabl depth of the fiber defects on polished sections of samples in the lI for each temperature. The values are comparable to thos as-received condition and after 1 h at 1200%C. The stud reported in the literature for Nicalon SiC fibers23-25, 27 and showed that the average defect depth increased by -30%with clearly show the degradation of fiber strength in the composites heat treatment. in agreement with the reduction of in situ fiber tested at1000°and1200°C strength reported above. In addition, flexure tests conducted at Several conclusions were drawn from the present results ambient temperature on samples in the as-received condition First, the Weibull modulus of the fibers in the as-received and after I h at 1200C showed a significant decrease in composite was extremely low. This low value indicates a very strength(from 450 to 320 MPa), indicative of permanent fiber damage during high-temperature exposure VI. Concluding Remarks This present investigation showed that dual BN/SiC fiber coatings applied by chemical vapor deposition(CvD)retained a weak fiber/matrix interface. even when the fiber-reinforced composite was exposed at 1400C for I h in an oxidizing atmosphere. Interface characterization revealed no significant changes in the structure of either BN or SiC coatings or at the BN/SiC interface. This lack of change indicated that the external SiC coating acted as a barrier, actually limiting the diffusion of oxygen from the environment, and that the BN/SiC interface was thermodynamically inert in the absence of oxygen In contrast, the BN/fiber interface was modified by thermal Temperature treatments at high temperature, but the degree of change de- 252 pended on the nature of the BN/fiber interface and the tem- perature. A carbon-rich interlayer -10 nm thick was found at 520 the BN/fiber interface in the as-received Al,O3-matrix com- -1000 posite. The carbon in this interlayer burned out rapidly and disappeared at high temperature in air. Carbon removal was the only effect that oxygen had in the samples treated at 800C for 1000 1500 2000 200 I h, and the fiber strength was unaffected. However, at higher temperatures, oxygen from the environment led to the oxida- FIber Strength, S( MPa) tion of the fiber surfaces, and an SiO, layer formed, severely degrading the fiber strength. Fig. 6. Cumulative fracture probability of the fibers, F, at different Oxidation at the BN/fiber interface was very different in the Si-C-N-matrix composite, which did not present a carbon-richarranged in ascending order, and a corresponding failure prob￾ability, F 4 (i − 0.5)/N, is assigned to each strength, based on rank statistics, where i is the rank and N the total number of experimental data. The lowest strength is attributed to the fi￾bers with a specular fracture surface (type III), following the analysis of Eckel and Bradt.16 The highest strength is assigned to the fibers with a very short mirror radius (type I). The results are shown in Fig. 6. According to the Weibull statistics, the fiber fracture probability is given by F~S! = 1 − exp F−S S s*c D m* G (2) Here, the parameter s*c is obtained from Fig. 6, as the stress that gives a fracture probability of 63%, and m* is computed by the least-squares fitting of Eq. (2) to the experimental results plotted in Fig. 6. The values of s*c and m* derived from the fracture mirror data are not, in general, identical to the true in situ fiber characteristic strength, sc, and Weibull modulus, m. The values must be corrected to account for the screening of flaws over a finite fiber length on either side of a fiber failure site. Following the procedure developed by Curtin 17 for the case of multiple matrix cracks, the corrected values of sc and m were estimated in the present study and are presented in Table II for each temperature. The values are comparable to those reported in the literature for Nicalon SiC fibers23–25,27 and clearly show the degradation of fiber strength in the composites tested at 1000° and 1200°C. Several conclusions were drawn from the present results. First, the Weibull modulus of the fibers in the as-received composite was extremely low. This low value indicates a very wide distribution of fiber-surface defects, likely introduced during composite processing. This wide distribution was unal￾tered by the oxidation processes on the fiber surface at high temperature, and the Weibull modulus remained almost con￾stant or increased slightly. Second, the characteristic strength, sc, was practically equivalent in the as-received material and in the specimens tested at 800°C; thus, the modifications in the BN/fiber interlayer region at 800°C were confined to this re￾gion and did not nucleate defects on the fiber. The microstruc￾tural characterization of this region, presented above (Figs. 1(A) and 2(A)), showed no growth in the interlayer thickness after 1 h at 800°C. The only change was the removal of carbon from this layer. This carbon removal should not have affected the fiber strength, a result in agreement with the fiber strength obtained from the fracture mirrors. The interfacial region did grow thicker after 1 h at 1000° and 1200°C, and energy-dispersive spectroscopy indicated that the region was made of SiO2. This layer nucleated fiber-surface defects, which greatly reduced the strength of the fibers, as demonstrated earlier on isolated Nicalon fibers tested in air at similar temperatures.23–25,27 The process, and not just the tem￾perature effect, seemed to be the main cause of fiber-strength degradation. This conclusion is supported by a previous inves￾tigation, using quantitative microscopy, that measured the depth of the fiber defects on polished sections of samples in the as-received condition and after 1 h at 1200°C. The study13 showed that the average defect depth increased by ∼30% with heat treatment, in agreement with the reduction of in situ fiber strength reported above. In addition, flexure tests conducted at ambient temperature on samples in the as-received condition and after 1 h at 1200°C showed a significant decrease in strength (from 450 to 320 MPa), indicative of permanent fiber damage during high-temperature exposure.28 VI. Concluding Remarks This present investigation showed that dual BN/SiC fiber coatings applied by chemical vapor deposition (CVD) retained a weak fiber/matrix interface, even when the fiber-reinforced composite was exposed at 1400°C for 1 h in an oxidizing atmosphere. Interface characterization revealed no significant changes in the structure of either BN or SiC coatings or at the BN/SiC interface. This lack of change indicated that the external SiC coating acted as a barrier, actually limiting the diffusion of oxygen from the environment, and that the BN/SiC interface was thermodynamically inert in the absence of oxygen. In contrast, the BN/fiber interface was modified by thermal treatments at high temperature, but the degree of change de￾pended on the nature of the BN/fiber interface and the tem￾perature. A carbon-rich interlayer ∼10 nm thick was found at the BN/fiber interface in the as-received Al2O3-matrix com￾posite. The carbon in this interlayer burned out rapidly and disappeared at high temperature in air. Carbon removal was the only effect that oxygen had in the samples treated at 800°C for 1 h, and the fiber strength was unaffected. However, at higher temperatures, oxygen from the environment led to the oxida￾tion of the fiber surfaces, and an SiO2 layer formed, severely degrading the fiber strength. Oxidation at the BN/fiber interface was very different in the Si–C–N-matrix composite, which did not present a carbon-rich Table I. Fiber Fracture Surfaces at Different Temperatures† Fracture surface Temperature (°C) 25 800 1000 1200 Type I 22 20 24 31 Type II 66 63 69 58 Type III 12 17 7 11 † Values listed are percentages of fiber. Fig. 6. Cumulative fracture probability of the fibers, F, at different temperatures. Table II. In Situ Fiber Strength of Ceramic-Matrix Composites Temperature (°C) sc (GPa) m 25 1.82 2.0 800 1.78 2.3 1000 1.44 2.9 1200 1.36 2.6 December 1999 Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites 3499