ournal JAm. Geren so,82112134-500(1999 Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites and Its Effect on Fiber Strength Jose Perez-Rigueiro, Jose Antonio Celemin, and Javier LLorca Departamento de Ciencia de Materiales, Universidad Politecnica de Madrid, EtS de Ingenieros de Caminos, 28040 Madrid, Spain Pilar herrero Instituto de ciencia de materiales deMadrid,CSIC,Cantoblanco,28049Madrid,Spain Microstructural changes at the interface were analyzed in degrading the fiber properties, and the oxidation resista two Nicalon-fiber ceramic-matrix composites with a dual BN is significantly better than that of carbon. 10 In addition, BN BN/SiC coating on the fibers after thermal exposure at dif. coatings exhibit a hexagonal layered s n coating outer trating cture with turbostratic which pre ferent environments(air and argon). The outer SiC coating motes fiber/matrix decohesion, whereas the acted as a barrier to oxygen, which penetrated into the protects the Bn from chemical attack during matrix infiltration omposite via pipeline diffusion along the BN/fiber inter- and limits the access of oxygen to the interface faces. Oxygen penetration led to the formation of an Sic These duplex BN/SiC duce. but do not eliminate layer by oxidation of the fiber surfaces. The in situ fiber oxidation problems at high temperature. For instance, the bI strength at different temperatures, as determined from the Nicalon-fiber interface can be oxidized at 1200C, leading to radIus of th e mIrror on on the fiber fracture surface the nucleation and growth of defects on the fiber surface, an ndicated that this SiO, layer severely degraded the fiber effect that significantly reduces the in situ fiber strength and strength. Oxidation was highly dependent on the nature of thus, the overall composite strength and toughness. 3 Other the BN/fiber interface. The presence of a thin carbon-rich investigations 4, s also present evidence of oxidation at the nterlayer, which burned out rapidly at high temperature fiber/BN interface at lower temperatures(6000-850oC)in com- favored the entry of oxygen and accelerated oxidation of posites with duplex BN/SiC coatings, although the effect of he fibers oxidation on the fiber and interface properties has not been measured. Determining the suitability of duplex BN/SIC coat . Introduction ings for high-temperature applications requires a better knowl- edge of the factors that control chemical stability at the inter F BER-REINFORCED ceramics with a weak fiber/matrix inter face and of the effect of oxidation on the fiber properties. The face exhibit damage-tolerant, tough behavior. This weal present investigation is aimed at analyzing these questions in nterface appears spontaneously in the form of a thin carbon- two fiber-reinforced ceramics in the temperature range 800 rich layer when polymer-derived Si-C-O fibers are incorp ceramic matrices at high temperature or in- troduced by coating the fibers with one or several carbon layers IL. Materials rior to matrix infiltration. 3, 4 However, these composites often experience severe embrittlement when they are tested at high Two different fiber-reinforced ceramics(Lanxide Corp temperature in oxidizing atmospheres. -7 Under such condi newark, DE) tions, the carbonaceous coating of burned-out carbon is re were reinforced with 37 vol% of ceramic-grade Nicalon(Nip- placed by an amorphous silicate, formed from the reaction pon Carbon Co., Tokyo, Japan) SiC fibers. The fiber preform between the oxidized fiber surface and the matrix and which was manufactured by stacking several layers of bidirectional bonds the fiber strongly to the matrix. -lc 0o-90o)Nicalon 8 harness satin-weave fabric. The fibers in Such results have spurred the development of new coatings the lay-up were coated by CVD with a thin layer of BN(-100- vith relatively low interfacial shear strength and good oxida- 300 nm)and afterward with a thicker layer of SiC (3 um) tion resistance at high temperature(21100C). Very promising onto the bn. The Al,O matrix in the first composite was results have been obtained to date using dual BN/SiC fiber infiltrated using a direct metal-oxidatiss jumin\ s jd th ocess. The preform coatings in Al,O, /Nicalon and barium magnesium alumino- was brought into contact with molten silicate(BMAS)glass-ceramic/Nicalon/2 composites, which -1000C. The aluminum reacted with the oxygen to maintain their ambient-temperature mechanical properties ul matrix of porous Al,O3, which grew into the preform, and the to 1200@C. Duplex BN/SiC layers can be applied sequentially residual aluminum was removed afterward from the al, O,ma- onto the fibers by chemical vapor deposition(CVD)without where 1 The composite was received in the form of rectangu x. More details of the processing route are found else- lar plates of 3 mm nominal thickness. The porosity of the composites was -8% R. Naslain--contributing editor The matrix in the second cor te was introduced user polymer infiltration and pyrolysis. The precursor thermoset mer(polyureasilazane) was infiltrated into the fiber pre- form using a vacuum bagging approach. It then was pyrolyzed in an argon atmosphere at temperatures <1200C, leading to Member, American Ceramic Society. Mdrl de Madrid, and NATO an Si-C-N amorphous matrix. The infiltration-pyrolysis cycle was repeated several times until the required density was 3494
Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites and Its Effect on Fiber Strength Jose´ Pe´rez-Rigueiro, Jose´ Antonio Celemín, and Javier LLorca* Departamento de Ciencia de Materiales, Universidad Polite´cnica de Madrid, ETS de Ingenieros de Caminos, 28040 Madrid, Spain Pilar Herrero Instituto de Ciencia de Materiales de Madrid, CSIC, Cantoblanco, 28049 Madrid, Spain Microstructural changes at the interface were analyzed in two Nicalon-fiber ceramic-matrix composites with a dual BN/SiC coating on the fibers after thermal exposure at different temperatures (in the range 800°–1400°C) and in different environments (air and argon). The outer SiC coating acted as a barrier to oxygen, which penetrated into the composite via pipeline diffusion along the BN/fiber interfaces. Oxygen penetration led to the formation of an SiO2 layer by oxidation of the fiber surfaces. The in situ fiber strength at different temperatures, as determined from the radius of the mirror region on the fiber fracture surface, indicated that this SiO2 layer severely degraded the fiber strength. Oxidation was highly dependent on the nature of the BN/fiber interface. The presence of a thin carbon-rich interlayer, which burned out rapidly at high temperature, favored the entry of oxygen and accelerated oxidation of the fibers. I. Introduction FIBER-REINFORCED ceramics with a weak fiber/matrix interface exhibit damage-tolerant, tough behavior. This weak interface appears spontaneously in the form of a thin carbonrich layer when polymer-derived Si–C–O fibers are incorporated into glass-ceramic matrices at high temperature1,2 or introduced by coating the fibers with one or several carbon layers prior to matrix infiltration.3,4 However, these composites often experience severe embrittlement when they are tested at high temperature in oxidizing atmospheres.5–7 Under such conditions, the carbonaceous coating of burned-out carbon is replaced by an amorphous silicate, formed from the reaction between the oxidized fiber surface and the matrix, and which bonds the fiber strongly to the matrix.8–10 Such results have spurred the development of new coatings with relatively low interfacial shear strength and good oxidation resistance at high temperature ($1100°C). Very promising results have been obtained to date using dual BN/SiC fiber coatings in Al2O3/Nicalon11 and barium magnesium aluminosilicate (BMAS) glass-ceramic/Nicalon12 composites, which maintain their ambient-temperature mechanical properties up to 1200°C. Duplex BN/SiC layers can be applied sequentially onto the fibers by chemical vapor deposition (CVD) without degrading the fiber properties, and the oxidation resistance of BN is significantly better than that of carbon.10 In addition, BN coatings exhibit a hexagonal layered structure with turbostratic disorder, analogous to that of carbon-rich coatings, which promotes fiber/matrix decohesion, whereas the SiC outer coating protects the BN from chemical attack during matrix infiltration and limits the access of oxygen to the interface. These duplex BN/SiC coatings reduce, but do not eliminate, oxidation problems at high temperature. For instance, the BN/ Nicalon-fiber interface can be oxidized at 1200°C, leading to the nucleation and growth of defects on the fiber surface, an effect that significantly reduces the in situ fiber strength and, thus, the overall composite strength and toughness.13 Other investigations14,15 also present evidence of oxidation at the fiber/BN interface at lower temperatures (600°–850°C) in composites with duplex BN/SiC coatings, although the effect of oxidation on the fiber and interface properties has not been measured. Determining the suitability of duplex BN/SiC coatings for high-temperature applications requires a better knowledge of the factors that control chemical stability at the interface and of the effect of oxidation on the fiber properties. The present investigation is aimed at analyzing these questions in two fiber-reinforced ceramics in the temperature range 800°– 1400°C. II. Materials Two different fiber-reinforced ceramics (Lanxide Corp., Newark, DE) were used in the present study. Both composites were reinforced with 37 vol% of ceramic-grade Nicalon (Nippon Carbon Co., Tokyo, Japan) SiC fibers. The fiber preform was manufactured by stacking several layers of bidirectional (0°–90°) Nicalon 8 harness satin-weave fabric. The fibers in the lay-up were coated by CVD with a thin layer of BN (∼100– 300 nm) and afterward with a thicker layer of SiC (∼3 mm) onto the BN. The Al2O3 matrix in the first composite was infiltrated using a direct metal-oxidation process. The preform was brought into contact with molten aluminum in air at ∼1000°C. The aluminum reacted with the oxygen to form a matrix of porous Al2O3, which grew into the preform, and the residual aluminum was removed afterward from the Al2O3 matrix. More details of the processing route are found elsewhere.11 The composite was received in the form of rectangular plates of 3 mm nominal thickness. The porosity of the composites was ∼8%. The matrix in the second composite was introduced using polymer infiltration and pyrolysis. The precursor thermoset polymer (polyureasilazane) was infiltrated into the fiber preform using a vacuum bagging approach. It then was pyrolyzed in an argon atmosphere at temperatures <1200°C, leading to an Si–C–N amorphous matrix. The infiltration–pyrolysis cycle was repeated several times until the required density was R. Naslain—contributing editor Manuscript No 189770. Received October 30, 1998; approved June 24, 1999. Supported by CICYT (Spain), Universidad Polite´cnica de Madrid, and NATO through Grants No. MAT 95-787, A9708, and CRG-941033, respectively. *Member, American Ceramic Society. J. Am. Ceram. Soc., 82 [12] 3494–500 (1999) Journal 3494
December 1999 Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites 4 achieved. The final composite was received in the form of and heat-treated at 130.C for 5 min, followed by another 9 rectangular plates of 3 mm nominal thickness. The porosity of min at 180C. The samples were cut with a diamond saw, the composite was in the range 5%10% mechanically polished up to 100 um thick, and dimpled until hey perforated IlL. Experimental Techniques The in situ fiber strength at different temperatures was de termined for the Al,,-matrix composites from the size of the Parallelepiped-shaped samples(15 mm x 15 mm x 3 mm) mirror region on the fiber fracture surfaces. 13, 16, 7 To this ene were cut from the center of the composite plates and heat- tensile dog-bone-shaped samples were machined from th treated I h in air in an electric-resistance furnace. The heat plates. The central portion of each sample had a uniform width reatment temperatures for the AlO-matrix composite were of 4 mm and a length of 15 mm. The samples were tested in 800°,10001200°,andl400°C. In addition, one sample was tension, In aIr,at25°,800°,1000,and1200°C. All of the heat-treated at 1200.C for I h in an argon atmosphere, to samples were held at the test temperature for l h before tes ucidate the influence of the environment on the oxidation more details about the mechanical tests are found elsewhere processes. The ShC-N-matrix composite samples were heat The fracture surfaces of the broken samples were examined by treated at800°and1200° for 1 h in ai scanning electron microscopy(SEM; Model No 6300, JEOL All of the samples for the transmission electron microscopy and the fracture surfaces of approximately 130 fibers, selected (TEM) studies were prepared from the center of the parallel at random, were analyzed in each sample. Previous studie do, such a way that the fiber bundles were perpendicular to the values of in situ fiber strength, enough to obtain reliable epiped-shaped samples, by cutting square slices(600 um thick are surface. The slices were polished to 180 um thick and mounted with wax on an alumina rod 3 mm in diameter. The amples were dimpled down to 20 um and, finally, thinned by IV. Interface Degradation enter was perforated. The milled samples were examined by ()Ay OrAtrix Composite TEM (Model No 2000FX Il, JEOL, Tokyo, Japan). Com As shown in previous studies, O, II the porous Al2O3 matrix sitional analysis of the microstructural features was performed of the Al2O3-matrix composite filled the space between the by energy-dispersive spectroscopy, using probe sizes of-10- fibers, which were surrounded by the thick polycrystalline SiC 20nm. coating deposited by CVD. The thin BN layer(100 nm)be In order to ascertain whether the silicon oxides found in tween the fiber and the Sic external coating could be resolved les were actually silica or silicates, an SiO2 stan- only at higher magnification, by TEM (Fig. I(A)). No reaction dard was prepared, mixing equal volumes of SiO2 powder(0.8 zone was detected at the SiC/BN interface, in agreement with um average grain size, 99.9% purity) with a thermosetting he results of previous investigations. 8, 9 In contrast, a very resin(Araldite AT, Ciba Specialty Chemicals Corp, East Lan- thin region, of-10 nm, with slightly different contrast was sing, MI). The mixture was compacted in an aluminum tube detected at the BN/fiber interface. The energy spectra, obtained (A) ms8 08 0 00 0.5 1.2 S8N Fig. 1. (A) TEM photograph showing the BN coating between the nicalon fiber and the Sic external coating(carbon-rich interlayer region at ACMNMNiLYYMYW BN/fiber interface is marked with arrows);(B) energy spectrum corresponding to BN coating, (C)energy spectrum for BN/fiber interlayer
achieved. The final composite was received in the form of rectangular plates of 3 mm nominal thickness. The porosity of the composite was in the range 5%–10%. III. Experimental Techniques Parallelepiped-shaped samples (15 mm × 15 mm × 3 mm) were cut from the center of the composite plates and heattreated 1 h in air in an electric-resistance furnace. The heattreatment temperatures for the Al2O3-matrix composite were 800°, 1000°, 1200°, and 1400°C. In addition, one sample was heat-treated at 1200°C for 1 h in an argon atmosphere, to elucidate the influence of the environment on the oxidation processes. The Si–C–N-matrix composite samples were heattreated at 800° and 1200°C for 1 h in air. All of the samples for the transmission electron microscopy (TEM) studies were prepared from the center of the parallelepiped-shaped samples, by cutting square slices (600 mm thick) in such a way that the fiber bundles were perpendicular to the square surface. The slices were polished to 180 mm thick and mounted with wax on an alumina rod 3 mm in diameter. The samples were dimpled down to 20 mm and, finally, thinned by ion milling (milling conditions: 5 kV, 2.5 mA, 15°) until the center was perforated. The milled samples were examined by TEM (Model No. 2000FX II, JEOL, Tokyo, Japan). Compositional analysis of the microstructural features was performed by energy-dispersive spectroscopy, using probe sizes of ∼10– 20 nm. In order to ascertain whether the silicon oxides found in several samples were actually silica or silicates, an SiO2 standard was prepared, mixing equal volumes of SiO2 powder (0.8 mm average grain size, 99.9% purity) with a thermosetting resin (Araldite AT, Ciba Specialty Chemicals Corp., East Lansing, MI). The mixture was compacted in an aluminum tube and heat-treated at 130°C for 5 min, followed by another 90 min at 180°C. The samples were cut with a diamond saw, mechanically polished up to 100 mm thick, and dimpled until they perforated. The in situ fiber strength at different temperatures was determined for the Al2O3-matrix composites from the size of the mirror region on the fiber fracture surfaces.13,16,17 To this end, tensile dog-bone-shaped samples were machined from the plates. The central portion of each sample had a uniform width of 4 mm and a length of 15 mm. The samples were tested in tension, in air, at 25°, 800°, 1000°, and 1200°C. All of the samples were held at the test temperature for 1 h before testing. More details about the mechanical tests are found elsewhere.13 The fracture surfaces of the broken samples were examined by scanning electron microscopy (SEM; Model No. 6300, JEOL), and the fracture surfaces of approximately 130 fibers, selected at random, were analyzed in each sample. Previous studies demonstrated that this number is enough to obtain reliable values of in situ fiber strength.13,16 IV. Interface Degradation (1) Al2O3-Matrix Composite As shown in previous studies,10,11 the porous Al2O3 matrix of the Al2O3-matrix composite filled the space between the fibers, which were surrounded by the thick polycrystalline SiC coating deposited by CVD. The thin BN layer (∼100 nm) between the fiber and the SiC external coating could be resolved only at higher magnification, by TEM (Fig. 1(A)). No reaction zone was detected at the SiC/BN interface, in agreement with the results of previous investigations.18,19 In contrast, a very thin region, of ∼10 nm, with slightly different contrast was detected at the BN/fiber interface. The energy spectra, obtained Fig. 1. (A) TEM photograph showing the BN coating between the Nicalon fiber and the SiC external coating (carbon-rich interlayer region at BN/fiber interface is marked with arrows); (B) energy spectrum corresponding to BN coating; (C) energy spectrum for BN/fiber interlayer. December 1999 Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites 3495
Journal of the American Ceramic SociefyPeres-Rigueiro et al. Vol. 82, No. 12 by energy-dispersive X-ray microanalysis(EDAX), for the BN revealed that the carbon had disappeared after I h exposure at coating and the BN/fiber interlayer are plotted in Figs. I(B)and 800oC(Fig. 2(A)and that the main constituents were silicon (C), respectively. The first plot shows a significant amount of and oxygen. The thicker interlayer that formed after heat treat arbon and oxygen in the BN coating of the as-received com ment at 1200C allowed a more precise characterization by posite. Other authors 820 also detected oxygen and carbon-in EDAX. In Fig. 2(B), the energy spectrum is compared with that the range 5-10 wt%in BN coatings deposited by CVD. The obtained for the SiO2 standard, prepared as indicated in Section origin of the oxygen was unclear, but oxygen seemed to be Ill. Both spectra are practically identical, indicating that the introduced by contamination during the deposition process terlayer formed during high-temperature exposure was Sio and not any type of silicate The cuation associated with the application of the BN coating No oxidation of the bulk bn was observed. even in the samples treated at 1400C for 1 h. The energy spectrum of the interlayer was carbon rich. As indicated in the Introduction BN coating under this condition is shown in Fig. 3. The only arbon-rich layers appeared spontaneously on the surface of difference from the spectrum obtained in the as-received ma- alon fibers incorporated in glass-ceramic matrice terial is the absence of carbon after high-tem exposure, temperature. 1 The layers also were found when bn coating he peaks of nitrogen and oxygen are unaltered. This finding were applied at high temperature onto Nicalon fibers 8-20 and emphasizes the need to improve our knowledge of the oxida- were attributed to fiber decomposition at the surface, although tion processes in BN. The oxidation resistance of this coating there is no clear consensus on this point improves with crystallinity and purity whereas, on the contrary The effect on the BN/fiber interface of oxidation treatment the presence of water vapor leads to the rapid volatilization of for I h at800°and1200° C is shown in F he borosilicate glasses formed by oxidation. 21 22 ely The composite treated for I h at 1000C presented High-resolution analysis by tem (not shown here)indicated an intermediate appearance and is not plotted here, for the sake that this Bn coating exhibited a turbostratic structure and, as of brevity. The BN/SiC interface remained sharp and free from mentioned above, contained significant amounts of carbon and any reaction layer. The thin BN/fiber interlayer, in contrast oxygen(Fig. I(B). Although the present heat\"preatment atments were grew thicker with temperature, reaching -100 nm after heat performed in a laboratory atmosphere, with a rela treatment at 1200C. The energy spectrum of the interlayer of-50%, the Bn coating was unaffected by heat treatment, BN SiO2 Fiber tandard 吕06 M Energy(kev) Fig. 2. TEM photographs showing BN coating after I h at(A)800 and(B)1200C Energy spectra of the BN/fiber interlayer after I h at 800 and at 1200C are shown below the corresponding micrographs
by energy-dispersive X-ray microanalysis (EDAX), for the BN coating and the BN/fiber interlayer are plotted in Figs. 1(B) and 1(C), respectively. The first plot shows a significant amount of carbon and oxygen in the BN coating of the as-received composite. Other authors18,20 also detected oxygen and carbon—in the range 5–10 wt%—in BN coatings deposited by CVD. The origin of the oxygen was unclear, but oxygen seemed to be introduced by contamination during the deposition process. Carbon added to BN reportedly15 prevents fiber-strength degradation associated with the application of the BN coating. The second plot, Fig. 1(C), shows that the present BN/fiber interlayer was carbon rich. As indicated in the Introduction, carbon-rich layers appeared spontaneously on the surface of Nicalon fibers incorporated in glass-ceramic matrices at high temperature.1,2 The layers also were found when BN coatings were applied at high temperature onto Nicalon fibers18–20 and were attributed to fiber decomposition at the surface, although there is no clear consensus on this point. The effect on the BN/fiber interface of oxidation treatment for 1 h at 800° and 1200°C is shown in Figs. 2(A) and (B), respectively. The composite treated for 1 h at 1000°C presented an intermediate appearance and is not plotted here, for the sake of brevity. The BN/SiC interface remained sharp and free from any reaction layer. The thin BN/fiber interlayer, in contrast, grew thicker with temperature, reaching ∼100 nm after heat treatment at 1200°C. The energy spectrum of the interlayer revealed that the carbon had disappeared after 1 h exposure at 800°C (Fig. 2(A)) and that the main constituents were silicon and oxygen. The thicker interlayer that formed after heat treatment at 1200°C allowed a more precise characterization by EDAX. In Fig. 2(B), the energy spectrum is compared with that obtained for the SiO2 standard, prepared as indicated in Section III. Both spectra are practically identical, indicating that the interlayer formed during high-temperature exposure was SiO2 and not any type of silicate. No oxidation of the bulk BN was observed, even in the samples treated at 1400°C for 1 h. The energy spectrum of the BN coating under this condition is shown in Fig. 3. The only difference from the spectrum obtained in the as-received material is the absence of carbon after high-temperature exposure; the peaks of nitrogen and oxygen are unaltered. This finding emphasizes the need to improve our knowledge of the oxidation processes in BN. The oxidation resistance of this coating improves with crystallinity and purity whereas, on the contrary, the presence of water vapor leads to the rapid volatilization of the borosilicate glasses formed by oxidation.21,22 High-resolution analysis by TEM (not shown here) indicated that this BN coating exhibited a turbostratic structure and, as mentioned above, contained significant amounts of carbon and oxygen (Fig. 1(B)). Although the present heat treatments were performed in a laboratory atmosphere, with a relative humidity of ∼50%, the BN coating was unaffected by heat treatment, Fig. 2. TEM photographs showing BN coating after 1 h at (A) 800° and (B) 1200°C. Energy spectra of the BN/fiber interlayer after 1 h at 800° and at 1200°C are shown below the corresponding micrographs. 3496 Journal of the American Ceramic Society—Pe´rez-Rigueiro et al. Vol. 82, No. 12
December 1999 Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites 08 0 Energy(kev) Fig 3. Energy spectrum of BN coating after I h at 1400C even at 1400.C. The stability of the BN coating and the pres- ence of SiO,(and not silicates) at the BN/fiber interface sug gest that the observed changes resulted from oxidation of th Nicalon fiber surface. Ample experimental evidence showed the formation of an amorphous SiO, layer on the fiber surface 0.8 material(Fig. I(A). The edax results for this region also are 0.6 plotted in Fig. 4. Those results show that carbon disappeared and an SiO2 layer formed during high-temperature exposure in an inert environment. However, the layer was significantly thinner than that formed after I h of exposure at 1200C in Fig. 2(B)) The differences in the BN/fiber interlayer after I h at 1200oC in air and I h in argon manifest the critical contribution of oxygen from the environment to the modification of the inter- face microstructure. This relationship poses a question about Energy(kev) the diffusion pathway to the fiber/BN interface. Sun et al 12 suggested two possible diffusion pathways. Oxygen can diffuse TEM photogra to the interface through matrix microcracks that penetrate the here, energy spe SiC outer coating, or via pipeline diffusion along the fiber/BN with arrows interface, which starts from the cut ends of fiber exposed to the composite surface. Oxygen diffusion also is favored by th show that the Al O3 matrix in the as-received condition is and contained significant amounts of carbon and oxygen. The fractured, because of the thermal stresses generated during most significant difference from the previous composite was ooling, a result of the mismatch in the thermal expansion the absence of an interlayer region at the BN/fiber interface coefficients of the constituents. Those analyses also show that (Fig. 5(A). EDAX of the interfacial region revealed no carbon- the matrix cracks fail to penetrate the thick outer SiC coating, rich layer. which prevents oxygen access to the interface. Thus, pipeline The samples treated at 800 and 1200oC (Fig. 5(B))for I h diffusion along the BN/fiber interface is the most likely diffu- presented characteristics similar to those found in the Al2O3 ion pathway for the oxygen, in agreement with the results of matrix composite: Namely, the BN coating was unaffected by heat treatment, and the BN/SiC interface remained sharp and BN/Nicalon interfaces. I5 The carbon-rich interlayer may facili- free from any reaction products. The only noticeable difference tate this process, because the carbon is burned out rapidly at high temperature and oxygen access to the sample becomes after I h at 800oC and an interlayer developed only after Ihat easier. Changes in the interface microstructure as a function of 1200C(Fig. 5(B). However, the interlayer that developed the distance to the surface should be observable if pipeline was significantly thinner than that in the Al2O3-matrix com- diffusion is dominant but, unfortunately, all of the present TEM posite (20 nm versus 100 nm). The energy spectrum of this samples were taken from the center of the composites region(Fig. 5(B) showed an increase in the oxygen content, as compared to the same region in the as-received material. Nev- (2) Si-C-N-Matrir Composite ertheless, the interface oxidation was significantly lower in this The morphology of the Si-C-N te was siml- composite than in the Al2O3-matrix material to that of its Al2O3 counterpart. The Si-C-N matrix intro- Although the experimental evidence is limited, pipeline dif. duced by polymer infiltration and pyrolysis filled the gaps fusion along the BN/fiber interface again seems to be the most between the CVd SiC coatings, which surrounded the nica likely pathway for the access of oxygen to the interface in this fibers. The BN layer between the fiber and the external Sic composite, because the SiC outer coating prevents the penetra coating was thicker(300 nm) than that of the Al2O3-matrix tion of oxygen from pores and cracks in the matrix. According
even at 1400°C. The stability of the BN coating and the presence of SiO2 (and not silicates) at the BN/fiber interface suggest that the observed changes resulted from oxidation of the Nicalon fiber surface. Ample experimental evidence showed the formation of an amorphous SiO2 layer on the fiber surface when samples were exposed to high temperature in air.23–25 Finally, the appearance of the interface after 1 h at 1200°C under an argon atmosphere is shown in Fig. 4. The contrast of the BN/fiber interlayer is more marked than in the as-received material (Fig. 1(A)). The EDAX results for this region also are plotted in Fig. 4. Those results show that carbon disappeared and an SiO2 layer formed during high-temperature exposure in an inert environment. However, the layer was significantly thinner than that formed after 1 h of exposure at 1200°C in air (Fig. 2(B)). The differences in the BN/fiber interlayer after 1 h at 1200°C in air and 1 h in argon manifest the critical contribution of oxygen from the environment to the modification of the interface microstructure. This relationship poses a question about the diffusion pathway to the fiber/BN interface. Sun et al.12 suggested two possible diffusion pathways. Oxygen can diffuse to the interface through matrix microcracks that penetrate the SiC outer coating, or via pipeline diffusion along the fiber/BN interface, which starts from the cut ends of fiber exposed to the composite surface. Oxygen diffusion also is favored by the submicrometer-scale porosity of the BN coatings deposited by CVD. Previous microstructural analyses of this composite13,26 show that the Al2O3 matrix in the as-received condition is fractured, because of the thermal stresses generated during cooling, a result of the mismatch in the thermal expansion coefficients of the constituents. Those analyses also show that the matrix cracks fail to penetrate the thick outer SiC coating, which prevents oxygen access to the interface. Thus, pipeline diffusion along the BN/fiber interface is the most likely diffusion pathway for the oxygen, in agreement with the results of other investigations on the oxidation of carbon/Nicalon9 and BN/Nicalon interfaces.15 The carbon-rich interlayer may facilitate this process, because the carbon is burned out rapidly at high temperature and oxygen access to the sample becomes easier. Changes in the interface microstructure as a function of the distance to the surface should be observable if pipeline diffusion is dominant but, unfortunately, all of the present TEM samples were taken from the center of the composites. (2) Si–C–N-Matrix Composite The morphology of the Si–C–N-matrix composite was similar to that of its Al2O3 counterpart. The Si–C–N matrix introduced by polymer infiltration and pyrolysis filled the gaps between the CVD SiC coatings, which surrounded the Nicalon fibers. The BN layer between the fiber and the external SiC coating was thicker (∼300 nm) than that of the Al2O3-matrix composite, and its energy spectrum was analogous to that plotted in Fig. 1(B). The BN also exhibited a turbostratic structure and contained significant amounts of carbon and oxygen. The most significant difference from the previous composite was the absence of an interlayer region at the BN/fiber interface (Fig. 5(A)). EDAX of the interfacial region revealed no carbonrich layer. The samples treated at 800° and 1200°C (Fig. 5(B)) for 1 h presented characteristics similar to those found in the Al2O3- matrix composite: Namely, the BN coating was unaffected by heat treatment, and the BN/SiC interface remained sharp and free from any reaction products. The only noticeable difference was at the BN/fiber interface. No distinct interlayer was seen after 1 h at 800°C and an interlayer developed only after 1 h at 1200°C (Fig. 5(B)). However, the interlayer that developed was significantly thinner than that in the Al2O3-matrix composite (20 nm versus 100 nm). The energy spectrum of this region (Fig. 5(B)) showed an increase in the oxygen content, as compared to the same region in the as-received material. Nevertheless, the interface oxidation was significantly lower in this composite than in the Al2O3-matrix material. Although the experimental evidence is limited, pipeline diffusion along the BN/fiber interface again seems to be the most likely pathway for the access of oxygen to the interface in this composite, because the SiC outer coating prevents the penetration of oxygen from pores and cracks in the matrix. According Fig. 3. Energy spectrum of BN coating after 1 h at 1400°C. Fig. 4. TEM photograph of BN coating after 1 h at 1200°C in argon atmosphere; energy spectrum corresponding to the BN/fiber interlayer (marked with arrows in the micrograph) also is plotted. December 1999 Oxidation of BN/Nicalon Fiber Interfaces in Ceramic-Matrix Composites 3497
Journal of the American Ceramic SociefyPeres-Rigueiro et al. Vol. 82. No. 12 (A) Fibe Fiber 100nn 10 0.8 T0.6 0.4 0.2 00 05 Energy(ke E Fig. 5. TEM photographs of SH-C-N-matrix composite, showing BN coating(A)in the as-received condition and(B)after I h at 1200C in air. Energy spectra of the BN/fiber interlayer in the as-received condition and after I h at 1200oC are shown below the corresponding micrographs to this assumption, the differences in oxidation between this surface. The formation of SiO on the fiber surface immedi- composite and its Al, O3-matrix counterpart should be in the ately led to the nucleation of defects, severely reducing the structure and/or composition of this interface. In particular, the fiber strength and, consequently the composite strength and high carbon content of the interfacial layer, which is beneficial toughness. To assess the importance of this phenomenon, the in to fiber/matrix debonding in the presence of a crack, may be a situ fiber strength in the as-received composite, at 25C, and negative factor from the point of view of oxidation resistance, after I h at 8000, 1000, and 1200C was determined from the because carbon is rapidly removed at high temperature, open- radius of the mirror region on the fiber fracture surface. An ing the path for the entry of oxygen. This hypothesis is sup empirical relationship between the fracture mirror radius, ams ported by the results reported by Sun et al.2 on a BMAs and the fiber strength, S, was shown by numerous authors to be glass-ceramic matrix reinforced with Nicalon fibers with a dual of the form, Sic/BN coating. The composite in that study was subjected to a heat treat- Sa A (1) ment("ceraming')at 1200@C in argon before high-temperature where Am usually is denoted the mirror constant, estimated at testing to crystallize the matrix. This treatment also affected the 2.51 MPam for Nicalon SiC fibers. 16, 7, The fracture su interface: Two thin sublayers(one carbon rich and another faces of -130 fibers were analyzed for each temperature. Most SiO2 rich) that were present at the BN/fiber interface in the of the fibers exhibited a distinct mirror-mist-hackle structure as-pressed composite disappeared in the ceramed material. 9(type II), in which the mirror radius could be determined easily As a result, the oxidation resistance of the bn and the bn/fiber However, this radius was too short to be measured accurately interface was excellent, as shown by the analyses of samples in a small fraction of fibers(type 1). In addition, no distinct subjected to creep tests in air at 1100.C over hundreds of fracture mirror boundary was seen in other fibers, and the whole fiber fracture surface was specular(type Ill). The pro- portions of fibers with type L, Il, and Ill fracture surfaces V. Fiber degradation shown in Table I for each temperature These results can be used to compute the fiber-failure prob- As indicated already, the BN/fiber interlayer in the Al2O3. ability, F, as a function of the fiber strength, for each tempera- natrix composite seemed to change by oxidation of the fiber ture. The strength data computed from the mirror radius
to this assumption, the differences in oxidation between this composite and its Al2O3-matrix counterpart should be in the structure and/or composition of this interface. In particular, the high carbon content of the interfacial layer, which is beneficial to fiber/matrix debonding in the presence of a crack, may be a negative factor from the point of view of oxidation resistance, because carbon is rapidly removed at high temperature, opening the path for the entry of oxygen. This hypothesis is supported by the results reported by Sun et al.12 on a BMAS glass-ceramic matrix reinforced with Nicalon fibers with a dual SiC/BN coating. The composite in that study was subjected to a heat treatment (“ceraming”) at 1200°C in argon before high-temperature testing to crystallize the matrix. This treatment also affected the interface: Two thin sublayers (one carbon rich and another SiO2 rich) that were present at the BN/fiber interface in the as-pressed composite disappeared in the ceramed material.19 As a result, the oxidation resistance of the BN and the BN/fiber interface was excellent, as shown by the analyses of samples subjected to creep tests in air at 1100°C over hundreds of hours. V. Fiber Degradation As indicated already, the BN/fiber interlayer in the Al2O3- matrix composite seemed to change by oxidation of the fiber surface. The formation of SiO2 on the fiber surface immediately led to the nucleation of defects, severely reducing the fiber strength and, consequently, the composite strength and toughness. To assess the importance of this phenomenon, the in situ fiber strength in the as-received composite, at 25°C, and after 1 h at 800°, 1000°, and 1200°C was determined from the radius of the mirror region on the fiber fracture surface. An empirical relationship between the fracture mirror radius, am, and the fiber strength, S, was shown by numerous authors to be of the form, Sam 4 Am (1) where Am usually is denoted the mirror constant, estimated at 2.51 MPazm1/2 for Nicalon SiC fibers.16,17,27 The fracture surfaces of ∼130 fibers were analyzed for each temperature. Most of the fibers exhibited a distinct mirror-mist-hackle structure (type II), in which the mirror radius could be determined easily. However, this radius was too short to be measured accurately in a small fraction of fibers (type I). In addition, no distinct fracture mirror boundary was seen in other fibers, and the whole fiber fracture surface was specular (type III). The proportions of fibers with type I, II, and III fracture surfaces are shown in Table I for each temperature. These results can be used to compute the fiber-failure probability, F, as a function of the fiber strength, for each temperature. The strength data computed from the mirror radius are Fig. 5. TEM photographs of Si–C–N-matrix composite, showing BN coating (A) in the as-received condition and (B) after 1 h at 1200°C in air. Energy spectra of the BN/fiber interlayer in the as-received condition and after 1 h at 1200°C are shown below the corresponding micrographs. 3498 Journal of the American Ceramic Society—Pe´rez-Rigueiro et al. Vol. 82, No. 12
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-rich
arranged in ascending order, and a corresponding failure probability, 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 fibers 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 unaltered by the oxidation processes on the fiber surface at high temperature, and the Weibull modulus remained almost constant 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 region and did not nucleate defects on the fiber. The microstructural 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 temperature effect, seemed to be the main cause of fiber-strength degradation. This conclusion is supported by a previous investigation, 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 depended on the nature of the BN/fiber interface and the temperature. A carbon-rich interlayer ∼10 nm thick was found at the BN/fiber interface in the as-received Al2O3-matrix composite. 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 oxidation 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
Journal of the American Ceramic Sociery-Pere--Rigueiro et al. Vol. 82. No. 12 nterlayer at the BN/fiber interface of the as-received compos- Fabricated by Direct Metal Oxidation, "Ceram. Eng. Sci. Proc., 14 19-10 ite. No distinct oxidation interlayer was seen in the samples treated at 800oC, and such a layer developed only after I h at 12E. Y. Sun, S.R. Nutt, and JJ an,"High-Temperature Tensile Be- havior of a boron Nitride- Coated Silicon Carbide- Fiber Glass-Ceramic Com- 1200C. However, the layer that developed was significantly posite, J.Amt Ceram. Soc., 79[6] 1521-29(1996 thinner than that developed in the Al2O3-matrix composite at the same temperature(20 nm versus 100 nm). This difference ior at 20 and 1200.C of Nicalon-Silicon-Carbide- Fiber-Reinforced Alumina- suggests that the presence of a carbon-rich layer at the BN/fiber 14F E. Heredia, J C. McNulty, F. W. Zok, and A.G. Evans, "Oxidation nterface is deleterious to interface stability at high tempera- brittlement Probe for Ceramic-Matrix Composites, ". Am. Ceram. Soc., 78 ture, because carbon is rapidly burned out, facilitating the entry of oxygen into the sample ISE. Y. Sun, H.-T. Lin, and JJ. Brennan, "Interme Envi- ronmental Effects on Boron Nitride-Coated Silicon iber-Reinforced Acknowledgment: The authors are indebted to Dr. J. Y Pastor for his 16A. J. Eckel and R C. Bradt, "Strength Distributio rcing Fibers in help with the high-temperature mechanical test Infiltrated Silicon Carbide Matrix Compos- ite, " J. Am. Ceram Soc., 72 3J455-58(1989 racture Mirrors, "J. Am. Ceram Soc., 77[4J1075-78(1994 References JJ. Brennan and K. M. Prewo, "Silicon Carbide Fibre-Reinforced Glas acterization of Si-C(O) Fibe (CVI)Matrix Composites with a BN Inter- -K M. Prewo, JJ. Brennan, and G K. Layden, "Fiber-Reinforced Glasses Chemistry of SiC/BN Dual-Coated Nicalon-Fiber-Reinforced Glass-Ceramic 2R. Naslain, O. Dugne, A. Guette, J. Sevely, C R. Brosse, J.-P. Rocher, and amico, G. A. Bernhart, M. M. Dauchier, and J H. Mace, "SIC/SIC J. Cotteret, "Boron Nitride Interphase in Ceramic-Matrix Composites, "J.Am. Composites with Multilayered Interfaces,"J. Am. Ceram. Soc., 79 14] 849-58 with Carbon and Boron Nitride Interphases at Elevated Temperatures in Air, J. K M. Prewo, JJ. Brennan, B. Johnson, and S. Starrett, "Silicon Carbide H -T Lin and P. F. Becher "Effect of Fiber Coating on Lifetime of Nicalon Fibre-Reinforced Glass-Ceramic Composite Tensile Behaviour at Elevated Fiber-Silicon Carbide Composites, " Mater Sci Eng, A, A231, 143-50(1997). osed to High-Temperature Gaseous Environments, "J. Am. Ceram. Soc., 74 rials", pp. 111-22 in Fracture Mechanics of Ceramics, VoL 9. Edited by R C. J Clark, M. Jaffe, J. Rabe, and N. R. Langley, "Thermal Stability Char- radt, D P H Hasselman, D Munz, M. Sakai, and v. Y Schevchenko. Plenum acterization of SiC Ceramic Fibers: I, Mechanical Property and Chemical Struc- Timms, and J. 2L. C. Sawyer, R. T. Chen, F. Haimbach, P. J. Harget, E P. Prack, and M. [),si-6astng the Tensile Properties of Ceramic-Matrix Composites,"J.Met raphy and Structure, "Ceram. Eng. Sci. Proc., 719-10J914-30(1986) N. Frety and M. Boussuge, "Relationship between High-Temperature De E. Heredia, A. G. Evans, and C.A. Anderson, "Tensile and Shear Prop- ed SiC/AL,O, Composites Processed by op p. psucknens, s. sutherland. 4. M. Daniel. R.L. Cain, G. West. D. M. R der and R.T Chen,Strength, Structure, and Fracture Properties of Ceramic Fibers Produced Fibre-Reinforced Glass-Ceramic Matrix Composite, "J. Microsc., 177[3 251 from Polymeric Precursors: I, Base-Line Studies, "J. Am. Ceram. Soc., 70 [111 IoJ. LLorca, M. Elices, and J.A n,Toughness and Microstructural E. Erauzkin, M. Gutierrez, and J. LLorca, "Mechanical Properties of Ce- opposites for Aeroengine Applications"; pp. 761-68 in Ceramic and Metal Matrix Composites, CMMC 96. Edited by M. Fuentes, J M. Mar- oky, and C.R. ent of BN/ tinez-Esnaola, and A M. Daniel. Trans Tech Publications, Zurich, Switzerland SiC Duplex Fiber Coatings in Fiber-Reinforced Alumina Matrix Composites 1996
interlayer at the BN/fiber interface of the as-received composite. No distinct oxidation interlayer was seen in the samples treated at 800°C, and such a layer developed only after 1 h at 1200°C. However, the layer that developed was significantly thinner than that developed in the Al2O3-matrix composite at the same temperature (20 nm versus 100 nm). This difference suggests that the presence of a carbon-rich layer at the BN/fiber interface is deleterious to interface stability at high temperature, because carbon is rapidly burned out, facilitating the entry of oxygen into the sample. Acknowledgment: The authors are indebted to Dr. J. Y. Pastor for his help with the high-temperature mechanical tests. References 1 J. J. Brennan and K. M. Prewo, “Silicon Carbide Fibre-Reinforced GlassCeramic Matrix Composites Exhibiting High Strength and Toughness,” J. Mater. Sci., 17, 2371–83 (1982). 2 K. M. Prewo, J. J. Brennan, and G. K. 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