urna J Am Ceram Soc, 80 [3]609-14(1997) Intermediate-Temperature environmental effects on Boron Nitride-Coated silicon carbide-Fiber-Reinforced Glass-Ceramic Composites Ellen Y. Sun and Hua-Tay Lin Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6068 John brennan United Technologies Research Center, East Hartford, Connecticut 06108 The environmental effects on the mechanical properties of such as load-transfer, crack-deflection, and fiber-pullout pro- fiber-reinforced composites at intermediate temperatures cesses. However, when exposed to an oxidizing environment, were investigated by conducting flexural static-fatigue the carbon interfacial layer can oxidize, resulting in strong experiments in air at 600 and 950%C. The material that fiber-matrix bonding via silica formation at the fiber surface was studied was a silicon carbide/boron nitride(SiC/BN) and/or in fiber-strength degradation because of oxidation or dual-coated Nicalon-fiber-reinforced barium magnesium recrystallization of the fiber. In recent years, research has been aluminosilicate glass-ceramic Comparable time-dependent conducted to control the fiber/matrix interface in the composite failure responses were found at 600 and 950C when the by applying coatings on the fiber surfaces prior to composi maximum tensile stress applied in the bend bar was. 60% processing. The goal is to achieve a fiber coating of relatively of the room-temperature ultimate Flexural strength of low shear strength and good oxidation resistance such that as-received materials. At both temperatures, the materials composites with excellent mechanical properties and useful survived 500 h fatigue tests at lower stress levels. Among engineering lifetimes in oxygen-rich environments can be the samples that survived the 500 h fatigue tests, a 20% obtained -7 Following this approach, barium mag alumi- degradation in the room-temperature flexural strength nosilicate(BMAS)glass-ceramic composites reinforced with was measured in samples tested at 600 C, whereas no deg silicon carbide/boron nitride(SiC/BN) dual-coated Nicalon dation was observed for the samples tested at 950C. Microstructure and chemistry studies revealed interfacial fiber(Nippon Carbon Co., Tokyo, Japan) have been fabricated Previous studies have revealed that this material exhibits better oxidation in the samples that were fatigued at 600C. The mechanical properties and thermal stability at high tempera- growth rate of the Si-C-O fiber oxidation product at 600 C tures(21100 C), compared to composites reinforced with uncoated fibers or carbon-coated fibers, and, hence, is a promis- the interior of the material was oxidized and resulted in a ing candidate for high-temperature structural applications trength degradation and less fibrous fracture. In contrast. the interior of the material remained intact at 950%c More recently, oxidation effects on fiber-reinforced glass- because of crack sealing by rapid silicate formation, and ceramic composites with in-situl-formed carbon interlayers have strength/toughness of the composite was maintained. Also, been found to be more severe at intermediate temperatures at 600C, BN oxidized via volatilization, because no borosil (4000-800oC)than at high temperatures(=1000C). -4These cate was formed studies were conducted without applied stress. At high tempera tures, the carbon interfacial layer can be protected by the oxide scale that formed on the fiber surface at the exposed fiber end L. Introduction The interfacial opening that formed because of carbon removal O XIDATION embrittlement of fiber-reinforced glass and glass an be quickly sealed before oxidation extends into the interior ceramic matrix composites at high temperatures(21100.C of the material. However, at intermediate temperatures, the is well documented. With polymer-derived silicon carbide opening may not be sealed, because of lower rates of silicate (SiC-type fibers, the formation of a thin carbon layer(20- scale formation, resulting in property degradation. On the other 50 nm thick) at the fiber/matrix interface can be obtained during reinforced bMas glass-ceramic composites have indicated that processing at elevated temperatures during composi tion. This weak interfacial layer results in composites with high his composite system shows no strength degradation after strength and toughne ss vIa n multiple toughening mechanisms, annealing at 550C in oxygen for 100 h. Therefore, to examine the environmental effects on this composite system at interme- diate temperatures, the material has been subjected to stresses R.J. Kerans--contributing editor above that which produces microcracking in the matrix and, thus, allows the interior of the composite to be exposed to the environment. In the present study, static-fatigue experiments Manuscript No 191916 Received April 1, 1996: approved September 27, 1996 have been used. Composite materials have been exposed to various applied flexural stresses at 600 and 950C in air. The Research Associates Program admiment of resistance of the composite to stress-induced oxidation has been 960R22464 with Lockheed Martin Energy Research Corp and by an evaluated by observing the static-fatigue behavior at different temperatures and correlating the mechanical properties with Presented in part at the 199 Conference on Composites, Advanced Ceramics, interfacial microstructure and chemistry studies. A parallel or Materials and Structures(Coce Advanced Ceramics and Composites Symposium(Paper No. C-78-96F) come of the study is a better understanding of the oxidation Member, American Ceramic Society. behavior of bn and the sic fibers at these
J. Am. Ceram. Soc., 80 [3] 609–14 (1997) Intermediate-Temperature Environmental Effects on Boron Nitride-Coated Silicon Carbide-Fiber-Reinforced Glass-Ceramic Composites Ellen Y. Sun* and Hua-Tay Lin* Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831–6068 John J. Brennan* United Technologies Research Center, East Hartford, Connecticut 06108 The environmental effects on the mechanical properties of such as load-transfer, crack-deflection, and fiber-pullout pro- fiber-reinforced composites at intermediate temperatures cesses. However, when exposed to an oxidizing environment, were investigated by conducting flexural static-fatigue the carbon interfacial layer can oxidize, resulting in strong experiments in air at 600 and 950C. The material that fiber–matrix bonding via silica formation at the fiber surface was studied was a silicon carbide/boron nitride (SiC/BN) and/or in fiber-strength degradation because of oxidation or dual-coated Nicalon-fiber-reinforced barium magnesium recrystallization of the fiber. In recent years, research has been aluminosilicate glass-ceramic. Comparable time-dependent conducted to control the fiber/matrix interface in the composite failure responses were found at 600 and 950C when the by applying coatings on the fiber surfaces prior to composite maximum tensile stress applied in the bend bar was .60% processing. The goal is to achieve a fiber coating of relatively of the room-temperature ultimate flexural strength of low shear strength and good oxidation resistance such that as-received materials. At both temperatures, the materials composites with excellent mechanical properties and useful survived 500 h fatigue tests at lower stress levels. Among engineering lifetimes in oxygen-rich environments can be the samples that survived the 500 h fatigue tests, a 20% obtained.5–7 Following this approach, barium magnesium alumi- degradation in the room-temperature flexural strength nosilicate (BMAS) glass-ceramic composites reinforced with was measured in samples tested at 600C, whereas no degra- silicon carbide/boron nitride (SiC/BN) dual-coated Nicalon dation was observed for the samples tested at 950C. fiber (Nippon Carbon Co., Tokyo, Japan) have been fabricated. Microstructure and chemistry studies revealed interfacial Previous studies have revealed that this material exhibits better oxidation in the samples that were fatigued at 600C. The mechanical properties and thermal stability at high tempera- growth rate of the Si-C-O fiber oxidation product at 600C tures (1100C), compared to composites reinforced with was not sufficient to seal the stress-induced cracks, so that uncoated fibers or carbon-coated fibers, and, hence, is a promis- the interior of the material was oxidized and resulted in a ing candidate for high-temperature structural applications.8–10 strength degradation and less fibrous fracture. In contrast, More recently, oxidation effects on fiber-reinforced glass- the interior of the material remained intact at 950C ceramic composites with in-situ-formed carbon interlayers have because of crack sealing by rapid silicate formation, and been found to be more severe at intermediate temperatures strength/toughness of the composite was maintained. Also, (400–800C) than at high temperatures (1000C).11–14 These at 600C, BN oxidized via volatilization, because no borosili- studies were conducted without applied stress. At high tempera- cate was formed. tures, the carbon interfacial layer can be protected by the oxide scale that formed on the fiber surface at the exposed fiber ends. I. Introduction The interfacial opening that formed because of carbon removal can be quickly sealed before oxidation extends into the interior OXIDATION embrittlement of fiber-reinforced glass and glass- of the material. However, at intermediate temperatures, the ceramic matrix composites at high temperatures (1100C) opening may not be sealed, because of lower rates of silicate is well documented.1–4 With polymer-derived silicon carbide scale formation, resulting in property degradation. On the other (SiC)-type fibers, the formation of a thin carbon layer (20– hand, preliminary studies on the SiC/BN dual-coated fiber- 50 nm thick) at the fiber/matrix interface can be obtained during reinforced BMAS glass-ceramic composites have indicated that processing at elevated temperatures during composite fabrica- this composite system shows no strength degradation after tion. This weak interfacial layer results in composites with high annealing at 550C in oxygen for 100 h.15 Therefore, to examine strength and toughness via multiple toughening mechanisms, the environmental effects on this composite system at intermediate temperatures, the material has been subjected to stresses above that which produces microcracking in the matrix and, R. J. Kerans—contributing editor thus, allows the interior of the composite to be exposed to the environment. In the present study, static-fatigue experiments Manuscript No. 191916. Received April 1, 1996; approved September 27, 1996. have been used. Composite materials have been exposed to Research at ORNL sponsored by the U.S. Department of Energy, Division of various applied flexural stresses at 600 and 950C in air. The Materials Sciences, Office of Basic Energy Sciences, under Contract No. DE-AC05- resistance of the composite to stress-induced oxidation has been 96OR22464 with Lockheed Martin Energy Research Corp. and by an appointment of author EYS to the ORNL Postdoctoral Research Associates Program administered evaluated by observing the static-fatigue behavior at different jointly by ORISE and ORNL. Author JJB was supported at UTRC by the Air Force temperatures and correlating the mechanical properties with Office of Scientific Research. Presented in part at the 1996 Conference on Composites, Advanced Ceramics, interfacial microstructure and chemistry studies. A parallel out- Materials and Structures (Cocoa Beach, FL, Jan. 1996), Environmental Effects on come of the study is a better understanding of the oxidation Advanced Ceramics and Composites Symposium (Paper No. C-78-96F). * Member, American Ceramic Society. behavior of BN and the SiC fibers at these temperatures. 609
10 Vol. 80. No. 3 Il. Experimental Procedure load was applied using a constant crosshead speed of 0.5 The materials studied were Si-C-O Nicalon-fiber-reinforced mm/min. To elucidate the effect of applied stresses on the xidation behavior, an annealing experiment also was con- BMAS glass ceramics. The fibers were dual coated with Sic ducted at 600%C in air for 500 h with the flexural strength of the coating was applied to the fibers by Cvd using a proprietary annealed sample measured at room temperature precursor(3M Co., St. Paul, MN) chosen to give an approxi The fracture surfaces of composites that failed during the mate composition of 40 at. boron, 40 at.% atigue tests or fractured at room temperature after the fatigue 20 at. carbon. The bn coating exhibited a turbostratic struc tests were examined using a high-resolution scanning electron ture and was comprised of nanoscale crystallites. The oxygen croscopy(SEM)microscope(Model $4100, Hitachi, To content in the Sic and coating layers was measured apan)that had a field-emission gun and was equipped fo by Auger spectroscopy to be <3 at. % Composite panels energy-dispersive spectroscopy(EDS )capable of light-element (100 mm X 100 mm) were fabricated by hot pressing a layup detection Compositional analysis of the fracture surfaces was of0°/90° plies at1450° for 5 min under a pressure of69MPa conducted using scanning auger microscopy(SAM)(Model 0, PHI Electronics, Eden Prairie, MN). The areas that were ture of 1200C for 24 h to crystallize the BMAS matrix to the reduce carbon contamination. The beam current was in the barium osumilite phase(BaMg2Al3 (Si._ O3o). The proce- range of 27-47 nA. During the experiments, an accelerating dures have been described in detail in Prew et al. 16 The final voltage of 5-10 kV was used to locate the microstructural composite contained-50 vol%fibers features, then a lower accelerating voltage of 2 kv was used for Static-fatigue experiments were conducted at 600 and 950 quantification analysis to reduce specimen charging. A pure BN using four-poI lexural tests, with the outer ply of compound was used as a standard for quantitative analysi fibers parallel to the stress axis(0 plies). The maximum tensile The interfacial microstructures of composites before and after fatigue tests were examined using conventional transmis- MPa, above the matrix-cracking stress of the composite, which sion electron microscopy(TEM). For fatigued samples, TEM was-160 MPa at room temperature and 140 MPa at 1100C. S thin foils were prepared from areas near the tensile surface, The dimensions of the bend bars were 6 mm x 55 mm x with disks 3 mm in diameter cut from the tensile surface 2.5 mm. All the flexural bars contained twelve plies through the ound and polished 20 um off the tensile surface side, and then dimpled from the back side. The foil was argon-ion milled In the following text, the ply on the tensile surface(0% ply)is to perforation and carbon coated to prevent electrostatic charg ing effects. Specimens were examined using a microscope 90 ply. A pneumatic-type loading system was used to apply Model EM400T, Phillips Electronic Instruments, Mahwah, the load to the test bars through an alumina pushrod. The test N) that was operated at 100 kv. The microscope had a field- outer spans of 6.35 and 40 mm, respectively. The test bars were spectrometry(EDAX)Model 9100, EDAX, Mahwah, NJ)and eld in the fixture with a small load (outer fiber tensile stress of allel electron-energy-loss spectrometry(PEELS)(Model <15 MPa) and heated to the desired test temperature. The 7. Gatan, Pleasanton, CA). Interfacial chemistry was ana- sample was then allowed to equilibrate for at least 30 min lyzed using EDS and PEELS, and the microscope was operated before increasing the applied load to the selected level. The in the scanning TEM (STEM) mode with the smallest probe applied load was held constant until the test bar failed. At that size(<2 nm) oint, sensors interrupted the furnace power supply circuit so that the bend bars were able to cool to room temperature within 20 min(under applied load), minimizing damage and oxidation II. Results after 500 h, the experiment was terminated and the retained ( Static-Fatigue Behavior room-temperature strength of the sample was evaluated using a The results of static-fatigue experiments indicated that no four-point flexural test. During the room-temperature tests, the distinguishable difference existed in the stress dependence of Four-Point Flexure in air 600 500 之400 300 200 101001000 Timc to Failure(hour) Fig. 1. Stress dependence of the composite lifetime at(o)950 and(B)600C. ( Arrows indicate that the specimens did not fail after 500 h at the
610 Journal of the American Ceramic Society— Sun et al. Vol. 80, No. 3 II. Experimental Procedure load was applied using a constant crosshead speed of 0.5 mm/min. To elucidate the effect of applied stresses on the The materials studied were Si-C-O Nicalon-fiber-reinforced oxidation behavior, an annealing experiment also was con- BMAS glass ceramics. The fibers were dual coated with SiC ducted at 600C in air for 500 h with the flexural strength of the over BN using chemical vapor deposition (CVD). The BN annealed sample measured at room temperature. coating was applied to the fibers by CVD using a proprietary The fracture surfaces of composites that failed during the precursor (3M Co., St. Paul, MN) chosen to give an approxi- fatigue tests or fractured at room temperature after the fatigue mate composition of 40 at.% boron, 40 at.% nitrogen, and tests were examined using a high-resolution scanning electron 20 at.% carbon. The BN coating exhibited a turbostratic struc- microscopy (SEM) microscope (Model S4100, Hitachi, Tokyo, ture and was comprised of nanoscale crystallites.8 The oxygen Japan) that had a field-emission gun and was equipped for content in the SiC and BN coating layers was measured energy-dispersive spectroscopy (EDS) capable of light-element by Auger spectroscopy to be 3 at.%. Composite panels detection. Compositional analysis of the fracture surfaces was (100 mm 100 mm) were fabricated by hot pressing a layup conducted using scanning Auger microscopy (SAM) (Model of 0/90 plies at 1450C for 5 min under a pressure of 6.9 MPa 660, PHI Electronics, Eden Prairie, MN). The areas that were in argon with a graphite die. After hot pressing, the composite studied were sputtered by the electron beam for 1–15 min to panels were cut into bars and heat treated in argon at a tempera- reduce carbon contamination. The beam current was in the ture of 1200C for 24 h to crystallize the BMAS matrix to the range of 27–47 nA. During the experiments, an accelerating barium osumilite phase (BaMg2Al3(Si9Al3O30)). The proce- voltage of 5–10 kV was used to locate the microstructural dures have been described in detail in Prewo et al.16 The final features, then a lower accelerating voltage of 2 kV was used for composite contained 50 vol% fibers. Static-fatigue experiments were conducted at 600 and 950C quantification analysis to reduce specimen charging. A pure BN compound was used as a standard for quantitative analysis. in air using four-point flexural tests, with the outer ply of fibers parallel to the stress axis (0 plies). The maximum tensile The interfacial microstructures of composites before and stresses applied in the bend bars were in the range of 250–450 after fatigue tests were examined using conventional transmisMPa, above the matrix-cracking stress of the composite, which sion electron microscopy (TEM). For fatigued samples, TEM was 160 MPa at room temperature and 140 MPa at 1100C. thin foils were prepared from areas near the tensile surface, 15 The dimensions of the bend bars were 6 mm 55 mm with disks 3 mm in diameter cut from the tensile surface, 2.5 mm. All the flexural bars contained twelve plies through the ground and polished 20 m off the tensile surface side, and thickness, with 0 plies on the tensile and compressive surfaces. then dimpled from the back side. The foil was argon-ion milled In the following text, the ply on the tensile surface (0 ply) is to perforation and carbon coated to prevent electrostatic chargcalled the “top ply” and the 90 ply beneath it is called the “first ing effects. Specimens were examined using a microscope 90 ply.” A pneumatic-type loading system was used to apply (Model EM400T, Phillips Electronic Instruments, Mahwah, the load to the test bars through an alumina pushrod. The test NJ) that was operated at 100 kV. The microscope had a field- fixtures were fabricated from sintered -SiC with inner and emission gun and was equipped for energy-dispersive X-ray outer spans of 6.35 and 40 mm, respectively. The test bars were spectrometry (EDAX) (Model 9100, EDAX, Mahwah, NJ) and held in the fixture with a small load (outer fiber tensile stress of parallel electron-energy-loss spectrometry (PEELS) (Model 15 MPa) and heated to the desired test temperature. The 607, Gatan, Pleasanton, CA). Interfacial chemistry was anasample was then allowed to equilibrate for at least 30 min lyzed using EDS and PEELS, and the microscope was operated before increasing the applied load to the selected level. The in the scanning TEM (STEM) mode with the smallest probe applied load was held constant until the test bar failed. At that size (2 nm). point, sensors interrupted the furnace power supply circuit so that the bend bars were able to cool to room temperature within III. Results 20 min (under applied load), minimizing damage and oxidation of the fracture surface. However, if the specimen did not fail (1) Static-Fatigue Behavior after 500 h, the experiment was terminated and the retained room-temperature strength of the sample was evaluated using a The results of static-fatigue experiments indicated that no four-point flexural test. During the room-temperature tests, the distinguishable difference existed in the stress dependence of Fig. 1. Stress dependence of the composite lifetime at () 950 and () 600C. (Arrows indicate that the specimens did not fail after 500 h at the applied stress level and, therefore, the tests were terminated.)
March 1997 the lifetime at950°and60°℃C. As shown in Fig.l, the compos Under high applied stress(appl 2 400 MPa), samples failed ites survived 500 h at 950" and 600C at stress levels of 400 MPa, the lifetime range was changes in the composites should have occurred during the 1-100 h, regardless of the testing temperatures. However, dif- fatigue tests rather than because of exposure to high tempera ferent retained room-temperature strengths were obtained in the ture after failure, because(i) the typical lifetime under stresses samples surviving the 500 h tests. Compared to the as-received of 400 and 450 MPa is >5 h, and (i) the specimens w re cooled es, samples that had been static-fatigue tested at 950C to room temperature within 20 min after failure. In samples that for 500 h under an applied stress of 350 MPa failed at a similar were fatigued at 950 C under a stress of 450 MPa, fiber pullout a result of tests at 950C with an applied stress Appl 5 0.6oo, RT in the corner regions. No substantial interfacial reaction was of as-received materials ). On the other hand, an-20% degrada observed in the top 0 ply at depths of >50 um from the tensile samples that had been fatigued at 600C for 500 h under an side surface. In contrast, ie n within 150 um from the external tion in the retained room-temperature strength was measured in surface. In the first 90 ply near the tensile surface, reaction only occurred in the regio ed stress of 350 MPa amples that failed at 600C under a No degradation in the room-temperature flexural strength stress of 450 MPa, no fiber pullout was observed in the entire d in the o°and90° was observed in the sample that was simply annealed in air at plies, and the fiber surface had an appearance that was similar 600C for 500 h without applied stress. The stress-displacement curve and the appearance of the fracture surface of the annealed to that of the sample that had been fatigued under a stress of e were similar to that of the sample that was fatigued at 350 MPa 950%C. This is consistent with the previous results from aging The chemistry of the fiber surfaces in the brittle -failure zone experiments at 550 C, 5 showing that the mechanical properties (in a sample that had been fatigued at 600"C) was analyzed of this composite system are not affected by annealing at inter- mediate temperatures when no external stress is applied to surface were oxygen-bonded silicon, carbon-bonded silicon, the material boron, carbon, nitrogen, and oxygen. The boron nitrogen ratio was -1. and the concentration variance from one area to (2) Microstructural Observations another was in the range of 5-13 at. % Thus, the residual fiber (A Fracture Surface--SEM and SAM Studies: The frac coatings(BN and SiC)and silicon oxides are on the fiber ture surface of composites that were fatigued at 950 and 600C surface. No area was found that had boron, silicon, and oxygen differed in appearance. The sample that was fractured at room but no nitrogen, indicating that no borosilicate glass was formed temperature after exposure to Appl values of s350 MPa at and that BN oxidized via volatilization C exhibited extensive fiber pullout across the entire frac (B) Interfacial Microstructure and Chemistry--TEM Studies ure surface. Reaction at the fiber surface was only observed Reaction occurred in the bn coating layer and at the BN/fiber the first 90 ply on the tensile side and very close to the exposed fiber ends at the external side surfaces. Away from the external interface in samples that were static fatigued at 600C, as shown Figs. 4(A)and (B). The sample was under an applied stress fairly smooth and apparently no interfacial reaction occurred In Fig. 4(A), the reaction has occurred in the region between t room mperature after exposure to oal values of 350 in the light-contrast layer beneath the Sic layer and a dark- MPa at 600C, no fiber pullout was observed in the region on the tensile side(Fig. 2(B). This affected region contrast reaction product growing on the surface of the fiber extended from the tensile surface to the third ply in depth The compositions of these two reaction regions were analyzed and 1-2 mm inward from the side surfaces. There Figure 5(A) is the PEELS spectrum taken from area I in fore, the affected zone that developed during the fatigue test at Fig. 4(A), indicating that the light-contrast area is the BN 600C was at least one order of magnitude larger than that coating layer; Fig. 5(B)is from area 2 in Fig. 4(A), showing no occurred at the fiber/matrix interface and glassy reaction prod- peaks are from the carbon coating that has been depositedomp which developed at 950C. In the affected zone, reaction trace of either boron or nitrogen. (In both spectra, the carbe Ict(s)formed in the 0 and 90 plies, as shown in Figs. 3(A) the TEM foil. In Fig. 4(B), large voids and amorphous liga- and (B) ments have formed in the area between the Sic overlayer and Tensile Surface (A) (B) Fig. 2. Fracture surface of samples static fatigued under a stress of 350 MPa for 500 h and then fractured at room temperature(A)no interfacial eaction was observed in the interior of the composite tested at 950C; (B)a brittle-failure zone with no fiber pullout was observed in the sample
March 1997 Environmental Effects on BN-Coated SiC-Fiber-Reinforced Glass-Ceramic Composites 611 the lifetime at 950 and 600C. As shown in Fig. 1, the compos- Under high applied stress (appl 400 MPa), samples failed ites survived 500 h at 950 and 600C at stress levels of 350 during the static-fatigue tests. Any major microstructural MPa; whereas at stresses 400 MPa, the lifetime range was changes in the composites should have occurred during the 1–100 h, regardless of the testing temperatures. However, dif- fatigue tests rather than because of exposure to high temperaferent retained room-temperature strengths were obtained in the ture after failure, because (i) the typical lifetime under stresses samples surviving the 500 h tests. Compared to the as-received of 400 and 450 MPa is 5 h, and (ii) the specimens were cooled samples, samples that had been static-fatigue tested at 950C to room temperature within 20 min after failure. In samples that for 500 h under an applied stress of 350 MPa failed at a similar were fatigued at 950C under a stress of 450 MPa, fiber pullout stress level. Thus, little stress-induced degradation occurred as still occurred in most of the area on the tensile side except a result of tests at 950C with an applied stress appl 0.60,RT in the corner regions. No substantial interfacial reaction was (where 0,RT is the room-temperature ultimate flexural strength observed in the top 0 ply at depths of 50 m from the tensile of as-received materials). On the other hand, an 20% degrada- surface. In the first 90 ply near the tensile surface, reaction tion in the retained room-temperature strength was measured in only occurred in the region within 150 m from the external samples that had been fatigued at 600C for 500 h under an side surface. In contrast, in samples that failed at 600C under a applied stress of 350 MPa. stress of 450 MPa, no fiber pullout was observed in the entire No degradation in the room-temperature flexural strength tensile region. Interfacial reaction occurred in the 0 and 90 was observed in the sample that was simply annealed in air at plies, and the fiber surface had an appearance that was similar 600C for 500 h without applied stress. The stress–displacement to that of the sample that had been fatigued under a stress of curve and the appearance of the fracture surface of the annealed 350 MPa (Figs. 3(A) and (B)). sample were similar to that of the sample that was fatigued at The chemistry of the fiber surfaces in the brittle-failure zone 950C. This is consistent with the previous results from aging (in a sample that had been fatigued at 600C) was analyzed experiments at 550C,15 showing that the mechanical properties using SAM. The elements that were detected on the fiber of this composite system are not affected by annealing at intermediate temperatures when no external stress is applied to surface were oxygen-bonded silicon, carbon-bonded silicon, boron, carbon, nitrogen, and oxygen. The boron:nitrogen ratio the material. was 1, and the concentration variance from one area to (2) Microstructural Observations another was in the range of 5–13 at.%. Thus, the residual fiber (A) Fracture Surface—SEM and SAM Studies: The frac- coatings (BN and SiC) and silicon oxides are on the fiber ture surface of composites that were fatigued at 950 and 600C surface. No area was found that had boron, silicon, and oxygen differed in appearance. The sample that was fractured at room but no nitrogen, indicating that no borosilicate glass was formed temperature after exposure to appl values of 350 MPa at and that BN oxidized via volatilization. 950C exhibited extensive fiber pullout across the entire frac- (B) Interfacial Microstructure and Chemistry—TEM Studies: ture surface. Reaction at the fiber surface was only observed in Reaction occurred in the BN coating layer and at the BN/fiber the first 90 ply on the tensile side and very close to the exposed interface in samples that were static fatigued at 600C, as shown fiber ends at the external side surfaces. Away from the external in Figs. 4(A) and (B). The sample was under an applied stress side surfaces, typically beyond 100 m, the fiber surface was of 450 MPa and exposed to air at 600C for 32 h before failure. fairly smooth and apparently no interfacial reaction occurred In Fig. 4(A), the reaction has occurred in the region between (Fig. 2(A)). On the other hand, in samples that were fractured the Nicalon fiber and the SiC overcoat layer, with voids forming at room temperature after exposure to appl values of 350 in the light-contrast layer beneath the SiC layer and a dark- MPa at 600C, no fiber pullout was observed in the corner contrast reaction product growing on the surface of the fiber. region on the tensile side (Fig. 2(B)). This affected region The compositions of these two reaction regions were analyzed. extended from the tensile surface to the third ply in depth Figure 5(A) is the PEELS spectrum taken from area 1 in (500 m) and 1–2 mm inward from the side surfaces. Therefore, the affected zone that developed during the fatigue test at Fig. 4(A), indicating that the light-contrast area is the BN 600C was at least one order of magnitude larger than that coating layer; Fig. 5(B) is from area 2 in Fig. 4(A), showing no which developed at 950C. In the affected zone, reaction trace of either boron or nitrogen. (In both spectra, the carbon occurred at the fiber/matrix interface and glassy reaction prod- peaks are from the carbon coating that has been deposited on uct(s) formed in the 0 and 90 plies, as shown in Figs. 3(A) the TEM foil.) In Fig. 4(B), large voids and amorphous ligaand (B). ments have formed in the area between the SiC overlayer and (A) (B) Fig. 2. Fracture surface of samples static fatigued under a stress of 350 MPa for 500 h and then fractured at room temperature ((A) no interfacial reaction was observed in the interior of the composite tested at 950C; (B) a brittle-failure zone with no fiber pullout was observed in the sample tested at 600C).
612 Journal of the American Ceramic Sociery-Sun et al Vol. 80. No. 3 (A) (B matrix er Interfacial reaction occurred in the sample fatigued at 600C for 500 h under an applied stress of 350 MPa(A) in the top 0 ply, glassy product(s)formed at the fiber/matrix interface(arrow) and no fiber pullout was observed; (B)in the first 90 ply, viscous glass phase(s) on the fiber surfaces) the Nicalon fiber, where the bn coating layer has been oxi- initiated at the tensile surface and propagated inward, deflected dized. The areas shown in Figs. 4(A)and(B)are two different at the BN/fiber interface, and left the 0 fibers bridging the interfacial regions along the same fiber, with the area in rack surfaces in the crack wake; whereas in the first 90 ply, Fig. 4(B) closer to the external surface. Considering the fact cracks initiated at the external side surface and propagated that interfacial reactions proceed along the fibers from the along the BN/fiber interface At 950C, the BN coating layer in a SiC/BN system remained intact and the crack opening at the the composite, Fig 4(A) shows an early stage of the reaction, BN/fiber interface was quickly sealed by the oxidation products reaction was initiated at the BN/fiber interface and proceeded For the 0 bridging fibers in the top ply, fiber oxidation occurred into the BN layer. Again, neither boron nor nitrogen w only in a very short zone that had been exposed to the matrix detected in the amorphous ligaments. Electron-diffraction pa crack, and the fiber strength might not degrade significantl tern and EDS analyses indicated that this reaction product was Thus, cracks in the matrix in the top ply were unlikely to grow amorphous silica. These observations are consistent with the with the fibers bridging the crack surfaces in the crack wake SAM results described above This is evidenced by the fact that, in the samples that were tested at 950C (500 h and 350 MPa), the brittle zone IV. Discussion limited within a depth of 50 um from the tensile surface Similarly, in the first 90 ply, the crack opening at the bN/fibe The rimental results presented above indicate that, for interface was sealed and the affected zone did not extend the Sic/bn dual-coated Nicalon-fiber-reinforced BMAS glass- eyond 100 um. However, at lower temperatures(600%), the ceramic composites, stress-induced degradation in an oxidizing opening at the BN/fiber interface could not be fully sealed by environment occurred more readily at intermediate tempera- the reaction products, because of the lower oxidation rates of tures than at high temperatures. Because of the high stress level the SiC fibers than those at high temperatures. Oxidation stud used in the present study(appl value of -60%-70% of oorr) es of SiC fibers indicated that the thickness of the Sio, film matrix cracking was induced under the applied load, especially that formed on the fiber was proportional to the oxidation rate near the tensile surface, where the applied tensile stress was constant, which was a factor of -3 larger at 950C than that at maximum. Previous studies have revealed that interfacial 600.. In addition, the bn coating layer on the SiC fibe debonding in this composite system occurred mostly at the oxidized via volatilization at 600.. Therefore, the cracks at the BN/fiber interface. o Therefore, in the top ply (0), the crack BN/fiber interface were likely to remain open and function as a 150nm 200nm void-forming SiC SiC BN ea area 2 fiber reaction product Fig. 4. Interfacial microstructures in a sample static fatigued at 600C under a stress of 450 MPa for 32 h before failure(A)an earlier stage and
612 Journal of the American Ceramic Society— Sun et al. Vol. 80, No. 3 (A) (B) Fig. 3. Interfacial reaction occurred in the sample fatigued at 600C for 500 h under an applied stress of 350 MPa ((A) in the top 0 ply, glassy reaction product(s) formed at the fiber/matrix interface (arrow) and no fiber pullout was observed; (B) in the first 90 ply, viscous glass phase(s) formed on the fiber surfaces). the Nicalon fiber, where the BN coating layer has been oxi- initiated at the tensile surface and propagated inward, deflected dized. The areas shown in Figs. 4(A) and (B) are two different at the BN/fiber interface, and left the 0 fibers bridging the interfacial regions along the same fiber, with the area in crack surfaces in the crack wake; whereas in the first 90 ply, Fig. 4(B) closer to the external surface. Considering the fact cracks initiated at the external side surface and propagated that interfacial reactions proceed along the fibers from the along the BN/fiber interface. At 950C, the BN coating layer in exposed fiber ends at the external surface toward the interior of a SiC/BN system remained intact and the crack opening at the the composite, Fig. 4(A) shows an early stage of the reaction, BN/fiber interface was quickly sealed by the oxidation products whereas Fig. 4(B) shows a more-advanced stage. Therefore, the of the SiC fiber. The reactions did not extend along the fibers. reaction was initiated at the BN/fiber interface and proceeded For the 0 bridging fibers in the top ply, fiber oxidation occurred into the BN layer. Again, neither boron nor nitrogen was only in a very short zone that had been exposed to the matrix detected in the amorphous ligaments. Electron-diffraction pat- crack, and the fiber strength might not degrade significantly. tern and EDS analyses indicated that this reaction product was Thus, cracks in the matrix in the top ply were unlikely to grow amorphous silica. These observations are consistent with the with the fibers bridging the crack surfaces in the crack wake. SAM results described above. This is evidenced by the fact that, in the samples that were tested at 950C (500 h and 350 MPa), the brittle zone was limited within a depth of 50 m from the tensile surface. IV. Discussion Similarly, in the first 90 ply, the crack opening at the BN/fiber The experimental results presented above indicate that, for interface was sealed and the affected zone did not extend the SiC/BN dual-coated Nicalon-fiber-reinforced BMAS glass- beyond 100 m. However, at lower temperatures (600C), the ceramic composites, stress-induced degradation in an oxidizing opening at the BN/fiber interface could not be fully sealed by environment occurred more readily at intermediate tempera- the reaction products, because of the lower oxidation rates of tures than at high temperatures. Because of the high stress level the SiC fibers than those at high temperatures. Oxidation studused in the present study (appl value of 60%–70% of 0,RT), ies of SiC fibers indicated that the thickness of the SiO2 film matrix cracking was induced under the applied load, especially that formed on the fiber was proportional to the oxidation rate near the tensile surface, where the applied tensile stress was constant, which was a factor of 3 larger at 950C than that at 600C.17,18 maximum. Previous studies have revealed that interfacial In addition, the BN coating layer on the SiC fiber debonding in this composite system occurred mostly at the oxidized via volatilization at 600C. Therefore, the cracks at the BN/fiber interface. BN/fiber interface were likely to remain open and function as a 8,10 Therefore, in the top ply (0), the crack (A) (B) Fig. 4. Interfacial microstructures in a sample static fatigued at 600C under a stress of 450 MPa for 32 h before failure ((A) an earlier stage and (B) a more-advanced stage of the reaction); the two regions in these figures were along a same fiber, with the region in Fig. 4(B) closer to the external surface. Therefore, these figures indicate that oxidation initiated at the BN/fiber interface and developed into the BN coating layer.
March 1997 613 Area 1 410 3.510 2.510 210 g1510 c 点11 5000 200 300 400 500 600 Energy Loss (ev Area 2 410 3.510 310 2.510 210 1.510 5000 200 300 400 500600 Energy Loss ( ev) was deposited on the res of regions labeled(A)area I and (B)area 2 in Fig. 4(A)(in both spectra, the carbon peaks are from the carbon coating that Fig. 5. PEELS an diffusion path at600°C ation proceeded along the fibers the simultaneous oxidation of BN and SiC at high temperatures and affected the entire debonded length, resulting in strength (1100C and 950%C)and intermediate temperatures(600oC degradation of the fibers. In the top ply (0), the cracks in the The calculations were initiated with equal amounts of SiC and matrix were likely to propagate when the bridging fibers in the Bn and under a constant pressure of 10 Pa(I atm). The crack wake failed. In the first 90 ply, oxidation extended into results indicated that, in the temperature range of 600-1100C the interior of the composite. These are evidenced by the large the oxygen partial pressure that was necessary to oxidize SiC brittle-failure zone that is observed in the samples that have and form solid-state SiO2(Po. was lower than the oxygen been fatigued at600°C partial pressure that was necessary for solid-/liquid-state B2O, Different oxidation behaviors were observed for the BN formation (Po, ) In other words, solid SiO2 formed before con- oated SiC fibers in a glass-ceramic matrix at intermediate and densed B,O,. At high temperatures, the crack opening was igh temperatures. At 1100C, the BN/fiber interfaces remained quickly closed by the oxidation product of SiC. The availability his intact under static loading conditions and substantial interfacial of oxygen in the sealed crack was limited because solid-state reaction occurred more readily under cyclic loading conditions diffusion of oxygen through the silica sealing is much slower when the maximum applied stress was 103 MPa ,o The inter- than the gas-phase transport prior to crack closure. The oxygen facial reaction at 1 100C exhibited an oxidation sequence such partial pressure in the crack(po, may satisfy such a condition that the BN layer remained unaffected, whereas the SiC fiber <Po, <po, that oxidation of BN is prevented. However, oxidized. On the other hand, the interfacial reaction that was the po, level in the crack at 600"C can be much higher than that nduced at 600C exhibited spontaneous oxidation of BN and at 950C. This is due to two reasons. First, the crack opening is SiC. Thermodynamic calculations were performed to analyze less likely to be sealed by the silicate scale at the external
March 1997 Environmental Effects on BN-Coated SiC-Fiber-Reinforced Glass-Ceramic Composites 613 (A) (B) Fig. 5. PEELS analyses of regions labeled (A) area 1 and (B) area 2 in Fig. 4(A) (in both spectra, the carbon peaks are from the carbon coating that was deposited on the TEM foil). diffusion path at 600C. Oxidation proceeded along the fibers the simultaneous oxidation of BN and SiC at high temperatures (1100C19 and affected the entire debonded length, resulting in strength and 950C) and intermediate temperatures (600C). degradation of the fibers. In the top ply (0), the cracks in the The calculations were initiated with equal amounts of SiC and BN and under a constant pressure of 105 matrix were likely to propagate when the bridging fibers in the Pa (1 atm). The crack wake failed. In the first 90 ply, oxidation extended into results indicated that, in the temperature range of 600–1100C, the interior of the composite. These are evidenced by the large the oxygen partial pressure that was necessary to oxidize SiC brittle-failure zone that is observed in the samples that have and form solid-state SiO2 ( pSiC O2 ) was lower than the oxygen been fatigued at 600 partial pressure that was necessary for solid-/liquid-state B2O3 C. Different oxidation behaviors were observed for the BN- formation (pBN O2 ). In other words, solid SiO2 formed before condensed B2O3 coated SiC fibers in a glass-ceramic matrix at intermediate and . At high temperatures, the crack opening was high temperatures. At 1100C, the BN/fiber interfaces remained quickly closed by the oxidation product of SiC. The availability intact under static loading conditions and substantial interfacial of oxygen in the sealed crack was limited because solid-state reaction occurred more readily under cyclic loading conditions diffusion of oxygen through the silica sealing is much slower when the maximum applied stress was 103 MPa. than the gas-phase transport prior to crack closure. The oxygen 9,10 The interfacial reaction at 1100C exhibited an oxidation sequence such partial pressure in the crack (pO2 ) may satisfy such a condition pSiC O2 pO2 pBN O2 that oxidation of BN is prevented.19 that the BN layer remained unaffected, whereas the SiC fiber However, oxidized.10 On the other hand, the interfacial reaction that was the pO2 level in the crack at 600C can be much higher than that induced at 600C exhibited spontaneous oxidation of BN and at 950C. This is due to two reasons. First, the crack opening is SiC. Thermodynamic calculations were performed to analyze less likely to be sealed by the silicate scale at the external
Journal of the American Ceramic Society-Sun et Vol. 80. No. 3 surface at lower temperatures; thus, oxygen can easily diffuse References along the crack via gas-phase J. J. Brennan, " Interfacial Characterization of Glass and Glass-Ceramic consumed by the oxidation of SiC fiber, because of the lower Matrix/Nicalon SiC Fiber Composites": pp. 549-60 in Materials Science reaction rate at the lower temperature, i.e., more oxygen is Research, Vol 20. Plenum, New York, 19 R. F. Cooper and K. Chyung, "Structure and Chemistry of Fibre-Matrix accumulated in the crack The condition be satisfied at 600C, and oxidation may possibly occur both in Microscopy Study,J Mater. ScL, 22, 3148-60(1987 SiC and BN. Also. no solid-state boron oxides formed in the J. Brennan, Fiber Reinforced Ceramic Composites; Ch. 8(Glass and Glass-Ceramic Matrix Composites"). Edited by K S. Mazdiyasni. Noyes, Par oxidized interfacial region in the samples that were tested at Ridge, NJ, 1990 600C, although the main reaction product of bn is the con- C. Cao. E. Bischoff. O. Sbaizero M. Ruhle. A. G. Evans. D. B. Marshall densed B,O, phase, according to thermodynamic calculations and J. J. Brennan, " Effect of Interfaces on the Properties of Fiber-Reinforced This could be caused by convection of gas-phase species. Diffu- ce,"BN Coating of Ceramic Fibers for Ceramic Fiber Com sion of BO(g)and B,O,(g) along the crack into the external U.S. Pat. No 4642271,Feb 10, 1987 environment could reduce the corresponder essures eristics in a Fiber-Reinforced in the crack and promote volatilization of solid BN. Overall, the Ga Narain omposeteand amG cere m bOron N tide interbase in( different oxidation behaviors at high and intermediate tem ratures and volatilization of solid BN at intermediate temper- SE. Y. Sun, S.R. Nutt, and J. J. Brennan, " Interfacial Microstructure and atures contributed to the different static-fatigue behaviors that Chemistry of SiC/BN Dual-Coated Nicalon-F ave been observed 里sm.sRa.,m Fiber-Reinforced Glass-Ceramic Composites, "J. Am. Ceram. Soc., 78 [5]1233- V. Conclusions Y. Sun, SR Nutt, and J J. Brennan, "High-Temperature Tensile Behavior The oxidation behavior of bn-coated Nicalon -fiber-reinforced of a Coated SiC Fiber Glass-Ceramic Composite, " J. Am. Ceram. Soc., 79 [6 1521-29(199 BMAS glass-ceramic composites at intermediate temperatures IR. T. Bhatt, Oxidation Effects on the Mechanical Properties of a Sic-Fiber- was studied. For this composite system, stress-induced degra- Reinforced Reaction-Bonded Si, N, Matrix Composite, "J. Am. Ceram. Soc., 75 dation in an oxidizing environment occurred more readily at 2]405-12(1992) 2K. P Plucknett, S. Sutherland, A M. Daniel, R L. Cain, G. West, D. M.R. intermediate temperatures than at high temperatures. No degra- in, and M. H. Lewis, ""Environmental Ageing Effects in a Silicon Carbide dation of the retained room-temperature strength was observed einforced Glass-Ceramic Matrix Composite, " J. Microsc., 177 [3 251 in composites that were static fatigued at 950%C under an appl ess of 350 MPa for 500 h, whereas an% decrease in the Degradation of Calcium K.P. Plucknett and M. H. Lewis, "Inhibition of Intermediate Temperature Pretreat retained strength was measured after static fatigue at 600oC under similar stress conditions. a porous silicon oxide glassy F. E. Heredia, J. C. McNulty, F. w. Zok, and A, G. Evans, " Oxidation layer was formed at the BN/fiber interface in the sample tested mbrittlement Probe for Ceramic-Matrix Composites, J.Amm. Ceram. Soc., 78 at 600C. because of surface oxidation of the fiber and volatil- S. R Nutt, and E. Y Sun, " Interfacial Studies of Coated Fiber zation of the bn coating layer. The present study suggested Reinforced Glass-Ceramic Matrix Composites, "Annual Rept. R93-970150-2 on that consideration of the kinetic effects, such as oxygen diffu- AFOSR Contract F49620-92-C-000 sion through the condensed phases and convection of gas-phase an, and G. K. Layden, "Fiber Reinforced Glasses species along the crack, is necessary to predict the intermediate- 65(21305-13, 322 gh Performance Applications, Am. Ce temperature oxidation behavior of BN-coated fiber in a ceramic T. Shimon, H. Chen, and K. Okamura, "Mechanism of Oxidation of matrix accurately Si-C-O Fibers, "J. Ceram. Soc. Jpn.(Nippon Seramikkusu Kyokai Gakuyut Ronbunsh),100,918-24(1992) Oxidation Kinetics Acknowledgmen The authors wish to thank Dr. D. N. Braski for CHaSICI3/H, under CVI Conditions, J. Mater. Sci., 27 uger analyses. Drs. P. F. Becher. E. Lara-Curzio and P. F. Tortorelli at ORNL heldon, E. Y. Sun, S. R Nutt, and J. J 当证 Oxidat nd Prof. B. w. Sheldon at Brown University(Providence, RI)are acknowledged BN-Coated SiC Fibers in Ceramic-Matrix Composites, "JAm Ceram Soc. 2]539-43(1996
614 Journal of the American Ceramic Society— Sun et al. Vol. 80, No. 3 surface at lower temperatures; thus, oxygen can easily diffuse References 1 J. J. Brennan, “Interfacial Characterization of Glass and Glass-Ceramic along the crack via gas-phase transport. Second, less oxygen is Matrix/Nicalon SiC Fiber Composites”; pp. 549–60 in Materials Science consumed by the oxidation of SiC fiber, because of the lower Research, Vol. 20. Plenum, New York, 1986. reaction rate at the lower temperature; i.e., more oxygen is 2 R. F. Cooper and K. Chyung, “Structure and Chemistry of Fibre–Matrix accumulated in the crack. The condition p Interfaces in SiC Fibre-Reinforced Glass-Ceramic Composites: An Electron SiC O2 pBN O2 pO2 may Microscopy Study,” J. Mater. Sci., 22, 3148–60 (1987). be satisfied at 600C, and oxidation may possibly occur both in 3 J. J. Brennan, Fiber Reinforced Ceramic Composites; Ch. 8 (“Glass and SiC and BN. Also, no solid-state boron oxides formed in the Glass-Ceramic Matrix Composites”). Edited by K. S. Mazdiyasni. Noyes, Park oxidized interfacial region in the samples that were tested at Ridge, NJ, 1990. 4 H. C. Cao, E. Bischoff, O. Sbaizero, M. Ru¨ 600 hle, A. G. Evans, D. B. Marshall, C, although the main reaction product of BN is the con- and J. J. Brennan, “Effect of Interfaces on the Properties of Fiber-Reinforced densed B2O3 phase, according to thermodynamic calculations. Ceramics,” J. Am. Ceram. Soc., 73 [6] 1691–99 (1989). This could be caused by convection of gas-phase species. Diffu- 5 R. W. Rice, “BN Coating of Ceramic Fibers for Ceramic Fiber Composites,” sion of BO(g) and B U.S. Pat. No. 4 642 271, Feb. 10, 1987. 2O3(g) along the crack into the external 6 R. N. Singh, “Fiber–Matrix Interfacial Characteristics in a Fiber-Reinforced environment could reduce the corresponding partial pressures Glass-Ceramic Composite,” J. Am. Ceram. Soc., 72 [9] 1764–67 (1989). in the crack and promote volatilization of solid BN. Overall, the 7 R. Naslain, O. Dugne, and A. Guette, “Boron Nitride Interphase in Ceramic- different oxidation behaviors at high and intermediate tem- Matrix Composite,” J. Am. Ceram. Soc., 74 [10] 2482–88 (1991). 8 peratures and volatilization of solid BN at intermediate temper- E. Y. Sun, S. R. Nutt, and J. J. Brennan, “Interfacial Microstructure and Chemistry of SiC/BN Dual-Coated Nicalon-Fiber-Reinforced Glass-Ceramic atures contributed to the different static-fatigue behaviors that Matrix Composites,” J. Am. Ceram. Soc., 77 [5] 1329–39 (1994). have been observed. 9 E. Y. Sun, S. R. Nutt, and J. J. Brennan, “Flexural Creep of Coated SiCFiber-Reinforced Glass-Ceramic Composites,” J. Am. Ceram. Soc., 78 [5] 1233– 39 (1995). V. Conclusions 10E. Y. Sun, S. R. Nutt, and J. J. Brennan, “High-Temperature Tensile Behavior of a Coated SiC Fiber Glass-Ceramic Composite,” J. Am. Ceram. Soc., 79 [6] The oxidation behavior of BN-coated Nicalon-fiber-reinforced 1521–29 (1996). 11 BMAS glass-ceramic composites at intermediate temperatures R. T. Bhatt, “Oxidation Effects on the Mechanical Properties of a SiC-FiberReinforced Reaction-Bonded Si was studied. For this composite system, stress-induced degra- 3N4 Matrix Composite,” J. Am. Ceram. Soc., 75 [2] 405–12 (1992). dation in an oxidizing environment occurred more readily at 12K. P. Plucknett, S. Sutherland, A. M. Daniel, R. L. Cain, G. West, D. M. R. intermediate temperatures than at high temperatures. No degra- Taplin, and M. H. Lewis, “Environmental Ageing Effects in a Silicon Carbide dation of the retained room-temperature strength was observed Fibre-Reinforced Glass-Ceramic Matrix Composite,” J. Microsc., 177 [3] 251– 63 (1995). in composites that were static fatigued at 950C under an applied 13K. P. Plucknett and M. H. Lewis, “Inhibition of Intermediate Temperature stress of 350 MPa for 500 h, whereas an 20% decrease in the Degradation of Calcium Aluminosilicate/Nicalon by High-Temperature Pretreatretained strength was measured after static fatigue at 600C ment,” J. Mater. Sci. Lett., 14, 1223–26 (1995). 14 under similar stress conditions. A porous silicon oxide glassy F. E. Heredia, J. C. McNulty, F. W. Zok, and A. G. Evans, “Oxidation Embrittlement Probe for Ceramic-Matrix Composites,” J. Am. Ceram. Soc., 78 layer was formed at the BN/fiber interface in the sample tested [8] 2097–100 (1995). at 600C, because of surface oxidation of the fiber and volatil- 15J. J. Brennan, S. R. Nutt, and E. Y. Sun, “Interfacial Studies of Coated Fiber ization of the BN coating layer. The present study suggested Reinforced Glass-Ceramic Matrix Composites,” Annual Rept. R93-970150-2 on AFOSR Contract F49620-92-C-0001, Nov. 30, 1993. that consideration of the kinetic effects, such as oxygen diffu- 16K. M. Prewo, J. J. Brennan, and G. K. Layden, “Fiber Reinforced Glasses sion through the condensed phases and convection of gas-phase and Glass-Ceramics for High Performance Applications,” Am. Ceram. Soc. Bull., species along the crack, is necessary to predict the intermediate- 65 [2] 305–13, 322 (1986). 17 temperature oxidation behavior of BN-coated fiber in a ceramic T. Shimoo, H. Chen, and K. Okamura, “Mechanism of Oxidation of matrix accurately. Si-C-O Fibers,” J. Ceram. Soc. Jpn. (Nippon Seramikkusu Kyokai Gakujutsu Ronbunshi), 100, 918–24 (1992). 18L. Filipuzzi and R. Naslain, “Oxidation Kinetics of SiC Deposited from Acknowledgments: The authors wish to thank Dr. D. N. Braski for CH3SiCl3 /H2 under CVI Conditions,” J. Mater. Sci., 27, 3330–34 (1992). Auger analyses. Drs. P. F. Becher, E. Lara-Curzio, and P. F. Tortorelli at ORNL 19B. W. Sheldon, E. Y. Sun, S. R. Nutt, and J. J. Brennan, “Oxidation of and Prof. B. W. Sheldon at Brown University (Providence, RI) are acknowledged BN-Coated SiC Fibers in Ceramic-Matrix Composites,” J. Am. Ceram. Soc., 79 for helpful discussions. [2] 539–43 (1996).