CATERTALIA Pergamon Acta mater.48(2000)46194628 www.elsevier.com/locate/actamat INTERFACIAL CHARACTERIZATION OF A SLURRY-CAST MELT-INFILTRATED SiC/SIC CERAMIC-MATRIX COMPOSITE .. J. breNnaN United Technologies Research Center, East Hartford, CT 06108, USA being developed for combustor applications under the High Speed Civil Transport(HSCT) Enabling Propul on Material(EPM) Program. A major part of this effort has dealt with the characterization and optimizatio cussed in this paper include an overview of the differences in composite microstructure between the EPM SiC/SiC material and a more conventional CVI SiC/SiC composite material, the microstructure/property relationships for the EPM SiC/SiC composite with two different types of Sic fiber(High- Nicalon and Sylramic ), and the effect of moist, high-temperature environments on the stability of the BN interface. o 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Ceramic composites; Fibers; Interface; Microstructure; Mechanical properties 1 INTRODUCTION SiC-fiber-reinforced SiC matrix offers the best comb The Enabling Propulsion Material(EPM) Program is nation of high-temperature capability and high ther aimed at developing a gas turbine engine combustor mal conductivity. Chemical vapor deposited(CVD) liner so that an environmentally acceptable and econ- SiC has been reported to possess a thermal conduc- omically viable High Speed Civil Transport(HSCT) tivity of up to 325 Wim K [I], which is higher than for some metal superalloys. The most common aircraft can be achieved. One of the prime objectives SiC/SiC CMC system consists of a chemically vapor of this program has been to develop and demonstrat infiltrated(CVI) SiC matrix around a woven SiC-fiber a material system, design concept and manufacturing preform. The resultant microstructure of this CVI process that can meet the HSCT combustor s environ- Sic/Sic composite system, however, is not dense mental, thermal, structural, economic and durability requirements. Among these requirements are an No, instead it contains rather large regions of matrix emissions index <5 g/kg, a material that can with- porosity. This porosity lowers the thermal conduc- stand temperatures to 1200 C under a tensile stress tivity of the composite to an unacceptable level fo of up to 100 MPa, and be able to meet an 18,000 h the HSCT combustor liner. Therefore, it was decided life requirement. The combustor concepts under con- to concentrate SiC/SiC composite developmental sideration do not permit the use of film cooling. efforts on the system that consists of a woven SiC- which traditionally has been used to reduce combus- fiber preform( five-harness satin) that has a BN fiber tor liner temperatures to manageable levels for met- interface coating applied to it by CVI, followed by a allic combustor liners. Therefore. high-thermal-con- relatively thin(1-4 um)CVI coating of SiC. This ductivity materials with a significantly higher relatively porous rigidized preform is then infiltrated temperature capability than those of current metallic with a slurry of SiC particles, dried, and then infil- combustor liners are required for this application. trated with molten silicon metal. The result is a full Ceramic-matrix composites(CMCs)have been ident- dense matrix(except within the fiber tows)of Si/Sic. ified as having the highest potential to satisfy the as shown for the melt infiltration(MI) SiC/SiC com- design requirements and service conditions posite in Fig. 1, which is compared with a typical Of the CMC systems available, the system of an CVI SiC/SiC composite with its inherent matrix porosity. Both composites have a fiber volume frac tion of-0.35. Besides the higher thermal conductivity nnanjj@aol.com(J.J.Brennan)oftheMISiC/SiCcomposite,thelackofmatrix 1359-6454100/520.00@ 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved PI:S1359-6454(00)00248-2
Acta mater. 48 (2000) 4619–4628 www.elsevier.com/locate/actamat INTERFACIAL CHARACTERIZATION OF A SLURRY-CAST MELT-INFILTRATED SiC/SiC CERAMIC-MATRIX COMPOSITE J. J. BRENNAN* United Technologies Research Center, East Hartford, CT 06108, USA Abstract—An SiC-particulate, silicon-metal melt-infiltration-matrix composite reinforced with SiC fibers is being developed for combustor applications under the High Speed Civil Transport (HSCT) Enabling Propulsion Material (EPM) Program. A major part of this effort has dealt with the characterization and optimization of the boron nitride (BN) based fiber/matrix interface. BN was chosen as the primary interfacial material due to its inherently weak structure and thus good crack-deflecting ability, ease of deposition by chemical vapor infiltration (CVI) into woven fiber preforms, and relatively good environmental stability. Topics discussed in this paper include an overview of the differences in composite microstructure between the EPM SiC/SiC material and a more conventional CVI SiC/SiC composite material, the microstructure/property relationships for the EPM SiC/SiC composite with two different types of SiC fiber (High-Nicalon and Sylramic), and the effect of moist, high-temperature environments on the stability of the BN interface. 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Ceramic composites; Fibers; Interface; Microstructure; Mechanical properties 1. INTRODUCTION The Enabling Propulsion Material (EPM) Program is aimed at developing a gas turbine engine combustor liner so that an environmentally acceptable and economically viable High Speed Civil Transport (HSCT) aircraft can be achieved. One of the prime objectives of this program has been to develop and demonstrate a material system, design concept and manufacturing process that can meet the HSCT combustor’s environmental, thermal, structural, economic and durability requirements. Among these requirements are an NOx emissions index ,5 g/kg, a material that can withstand temperatures to 1200°C under a tensile stress of up to 100 MPa, and be able to meet an 18,000 h life requirement. The combustor concepts under consideration do not permit the use of film cooling, which traditionally has been used to reduce combustor liner temperatures to manageable levels for metallic combustor liners. Therefore, high-thermal-conductivity materials with a significantly higher temperature capability than those of current metallic combustor liners are required for this application. Ceramic-matrix composites (CMCs) have been identified as having the highest potential to satisfy the design requirements and service conditions. Of the CMC systems available, the system of an * E-mail address: Brennanjj@aol.com (J.J. Brennan) 1359-6454/00/$20.00 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S13 59-6454(00)00248-2 SiC-fiber-reinforced SiC matrix offers the best combination of high-temperature capability and high thermal conductivity. Chemical vapor deposited (CVD) SiC has been reported to possess a thermal conductivity of up to 325 W/m K [1], which is higher than for some metal superalloys. The most common SiC/SiC CMC system consists of a chemically vapor infiltrated (CVI) SiC matrix around a woven SiC-fiber preform. The resultant microstructure of this CVI SiC/SiC composite system, however, is not dense; instead it contains rather large regions of matrix porosity. This porosity lowers the thermal conductivity of the composite to an unacceptable level for the HSCT combustor liner. Therefore, it was decided to concentrate SiC/SiC composite developmental efforts on the system that consists of a woven SiC- fiber preform (five-harness satin) that has a BN fiberinterface coating applied to it by CVI, followed by a relatively thin (1–4 µm) CVI coating of SiC. This relatively porous rigidized preform is then infiltrated with a slurry of SiC particles, dried, and then infiltrated with molten silicon metal. The result is a fully dense matrix (except within the fiber tows) of Si/SiC, as shown for the melt infiltration (MI) SiC/SiC composite in Fig. 1, which is compared with a typical CVI SiC/SiC composite with its inherent matrix porosity. Both composites have a fiber volume fraction of |0.35. Besides the higher thermal conductivity of the MI SiC/SiC composite, the lack of matrix
BRENNAN: INTERFACIAL CHARACTERIZATION Environmeatal degradation dimcult 2 Fig. I Microstructural differences between CVI and MI SiC/SiC composites porosity also leads to better environmental stability SiC fiber from Dow Corning Corp, Midland, MI and a higher proportional limit (matrix Both of these fibers have very low oxygen content, microcracking)stress. Figure 2 shows a thin-foil with the Hi-Nicalon fiber being carbon-rich while the transmission electron micrograph of an MI composite Sylramic fiber is very close to stoichiometic SiC ith Sylramic SiC fibers. The crystalline nature of Since the Sylramic fiber consists of 100-500 nm crys the fibers and the feathery CvI SiC layer can be seen tals of B-SiC, while the Hi-Nicalon fiber is a mixture clearly, in contrast to the more amorphous to turbos- of microcrystalline(2-20 nm) B-Sic plus turbostratic tratic nature of the bn interface coating domains of carbon, the Sylramic fiber has properties Once the fabrication method of the composite was much like those of sintered SiC; i.e., an elastic modu hosen, developmental efforts concentrated on the lus of -385 GPa and a thermal expansion coefficient type of SiC fiber to be utilized, the thickness and of -54x10-o/C [2]. The elastic modulus, therma chemistry of the BN fiber/matrix interface, and the expansion coefficient and thermal conductivity of the thermo-mechanical properties of the composite. Two Hi-Nicalon fiber are much lower than those of the different SiC fibers that were commercially available Sylramic fiber [2]. The following discussion covers were selected as the fiber candidates: Hi-Nicalon Sic the experimental results obtained for composites with fiber from Nippon Carbon Co., Japan, and Sylramic these two fibers, the rationale for the selection of th Sylramic fiber as the primary candidate, and the results of environmentally sensitive tests relating to the bn fiber/matrix interf 2. EXPERIMENTAl 2. 1. MI SiC/SiC composite fabrication The MI SiC/SiC composites utilized in the EPM program were initially fabricated by Carborundum nc, Niagara Falls, NY, and later in the program by REaver Allied Signal Composites, Inc, now called Ho Iss:ic Matris i. well Advanced Composites, Inc (HACD), Newark DE. The fabrication of these composites consisted of oating a two-dimensional(2D)woven SiC-fiber pre- form, usually five- or eight-harness satin, with an nterface coating of CVI BN, followed by an over- coating of CVI SiC. The rigidized preform was then subjected to infiltration by an SiC particulate slurry in order to fill the large residual porosity with SiC, followed by a melt infiltration of silicon metal which Fig. 2. TEM thin-foil analysis of the microstructure of a Syl. would fill the fine porosity left between the SiC grain particles. The fabrication and properties of similar MI
4620 BRENNAN: INTERFACIAL CHARACTERIZATION Fig. 1. Microstructural differences between CVI and MI SiC/SiC composites. porosity also leads to better environmental stability and a higher proportional limit (matrix microcracking) stress. Figure 2 shows a thin-foil transmission electron micrograph of an MI composite with Sylramic SiC fibers. The crystalline nature of the fibers and the feathery CVI SiC layer can be seen clearly, in contrast to the more amorphous to turbostratic nature of the BN interface coating. Once the fabrication method of the composite was chosen, developmental efforts concentrated on the type of SiC fiber to be utilized, the thickness and chemistry of the BN fiber/matrix interface, and the thermo-mechanical properties of the composite. Two different SiC fibers that were commercially available were selected as the fiber candidates: Hi-Nicalon SiC fiber from Nippon Carbon Co., Japan, and Sylramic Fig. 2. TEM thin-foil analysis of the microstructure of a Sylramic SiC fiber MI SiC/SiC composite. SiC fiber from Dow Corning Corp., Midland, MI. Both of these fibers have very low oxygen content, with the Hi-Nicalon fiber being carbon-rich while the Sylramic fiber is very close to stoichiometic SiC. Since the Sylramic fiber consists of 100–500 nm crystals of β-SiC, while the Hi-Nicalon fiber is a mixture of microcrystalline (2-20 nm) β-SiC plus turbostratic domains of carbon, the Sylramic fiber has properties much like those of sintered SiC; i.e., an elastic modulus of |385 GPa and a thermal expansion coefficient of |5.4×1026 /°C [2]. The elastic modulus, thermal expansion coefficient and thermal conductivity of the Hi-Nicalon fiber are much lower than those of the Sylramic fiber [2]. The following discussion covers the experimental results obtained for composites with these two fibers, the rationale for the selection of the Sylramic fiber as the primary candidate, and the results of environmentally sensitive tests relating to the BN fiber/matrix interface. 2. EXPERIMENTAL 2.1. MI SiC/SiC composite fabrication The MI SiC/SiC composites utilized in the EPM program were initially fabricated by Carborundum, Inc., Niagara Falls, NY, and later in the program by Allied Signal Composites, Inc., now called Honeywell Advanced Composites, Inc. (HACI), Newark, DE. The fabrication of these composites consisted of coating a two-dimensional (2D) woven SiC-fiber preform, usually five- or eight-harness satin, with an interface coating of CVI BN, followed by an overcoating of CVI SiC. The rigidized preform was then subjected to infiltration by an SiC particulate slurry in order to fill the large residual porosity with SiC, followed by a melt infiltration of silicon metal which would fill the fine porosity left between the SiC grain particles. The fabrication and properties of similar MI
BRENNAN: INTERFACIAL CHARACTERIZATION SiC/SiC composites have been presented previously frictional sliding and pull-out of the fibers-which 3-9 contribute to the composite toughness [ 121being more extensive for the hi-Nicalon fibers than for the 2.2. Composite testing Sylramic fibers. The reason for this may be related to Tensile testing was performed in accordance with the higher elastic modulus of the Sylramic fiber, or, High-Speed Research/Enabling Propulsion Materials more likely, the much higher surface roughness of the (HSR/EPM) consensus standard. The testing method Sylramic fiber compared with the Hi-Nicalon fiber,as was similar to ASTM standard C-1275-95 [10]. Ten- shown in Fig. 4. The surface roughness of the two sile specimen design was a contoured, face-loaded fibers is a reflection of the relative Sic grain size of specimen geometry with an overall length of 152 mm the fibers Fibers with higher surface roughness have and a gage-section width of 10.16 mm. The tensile been found to have onounced influence on the specimen design used a gradual radius from the tab fiber sliding behavior in ceramic-matrix composites to the gage section to reduce stress concentrations to [131 avoid grip and tab failures[11]. All of the mechanical While the Hi-Nicalon MI SiC/SiC composite sys property definitions can be found in the ASTM stan- tem appears to have better toughness than the Syl- dard. One of the most crucial and subjective measure ments is the proportional limit. The proportional limit well in high-temperature fatigue testing. As shown in was determined using the offset method. A line, run- Fig. 5, a Hi-Nicalon fiber composite failed after 38h ning parallel to the elastic modulus slope, was gener- and 19 cycles in a 2 h hold-time tensile fatigue test ated at a strain axis offset of 0.00005 mm/mm. The at 1200oC and a stress of 160 MPa, whereas a Syl proportional limit was defined as the stress level at ramic fiber composite ran out to 1000 h, 500 cycles, which point the offset line intersected the load versus under the same test conditions. The stress level during elongation curve. A servo-hydraulic machine was this test was chosen to be somewhat higher than the used for the vast majority of tensile testing. Testing proportional limit, or matrix microcrack stress of the from room temperature to elevated temperature was composites, which at this temperature is-140 MPa done in the same rig and differed only by the presence for the Sylramic fiber composite and -125 MPa for of a small furnace that heated the sample gage plus the Hi-Nicalon fiber composite. The residual 1200C radius region to within 1% of the desired temperature. tensile properties of the fatigued Sylramic fiber com Fatigue testing was conducted at elevated tempera- posite are not significantly different from those meas- ture in load control. Testing was done in a programm as-fabricated Sylramic fiber composite able servo-hydraulic machine. The testing consisted From the fracture surfaces shown in Fig. >1.fatigued an of a 2 h dwell fatigue test with the load being relieved seen that the amount of fiber pull-out for to 5% of the hold load at the end of each cycle. This Sylramic fiber composite is similar to that shown in simulated service conditions for the proposed appli- Fig 3 for the 1200C tensile sample, but the fatigued cation. Most tests were run to failure or 500 h of Hi-Nicalon composite shows a region around the fatigue exposure. All samples that made 500 h of edges of the sample that appears very brittle exposure were then tested for residual tensile proper- This embrittlement of the Hi-Nicalon fiber com- es at room temperature to note if degradation had posite is even more apparent after tensile fatigue test occurred. A selected number of fatigue tests was run ing at a lower temperature of 650%C, as shown in Fig to many thousands of hours before failure occurred, 6. While both composites ran out to the test limit of as will be described later 546h, 273 cycles, the residual 650C strength and strain-to-failure of the Hi-Nicalon fiber composite were significantly degraded, while those for the Syl- 3. RESULTS AND DISCUSSION ramic fiber composite were not 3. 1. Fracture characteristics of Hi-Nicalon versus 3. 2. Fiber/matrix interfacial characteristics of Hi- Sylramic SiC fiber MI composites Nicalon versus Sylramic SiC fiber MI composites Figure 3 shows the fracture surfaces of 1200C ten- From transmission electron microscopy (TEM) sile samples of MI SiC/SiC composites with either thin-foil analysis of the interfacial region in an as- Sylramic or Hi-Nicalon fibers. From this figure, one processed Hi-Nicalon fiber MI composite, as show can see that the fracture surface of the hi-nicalon Fig. 7, it was found that the probable reason fo fiber composite is much more fibrous in nature than the easy debonding and long fiber pull-out in these that of the Sylramic fiber composite. The measured composites, as was shown in Fig. 3, is that a thin tensile strengths of the two composites are similar, at (40 nm) carbon-rich layer has formed between th 286 MPa; however, the strain-to-failure of the Hi- BN and the Hi-Nicalon fiber during the MI composite Nicalon fiber composite(0.48%) is over twice that processing. This carbon-rich layer is probably a result of the Sylramic fiber composite(0.21%). Thus, the of interaction of the oxygen in the bn layer "toughness", or inherent matrix crack-stopping (-6 at%)with the excess carbon in the Hi-Nicalon ability, of the Hi-Nicalon fiber composite is greater fiber at the MI composite processing temperature of an that of the Sylramic fiber composite, due to the >1400C. Thicker, but similar, carbon layers have
BRENNAN: INTERFACIAL CHARACTERIZATION 4621 SiC/SiC composites have been presented previously [3–9]. 2.2. Composite testing Tensile testing was performed in accordance with High-Speed Research/Enabling Propulsion Materials (HSR/EPM) consensus standard. The testing method was similar to ASTM standard C-1275-95 [10]. Tensile specimen design was a contoured, face-loaded specimen geometry with an overall length of 152 mm and a gage-section width of 10.16 mm. The tensile specimen design used a gradual radius from the tab to the gage section to reduce stress concentrations to avoid grip and tab failures [11]. All of the mechanical property definitions can be found in the ASTM standard. One of the most crucial and subjective measurements is the proportional limit. The proportional limit was determined using the offset method. A line, running parallel to the elastic modulus slope, was generated at a strain axis offset of 0.00005 mm/mm. The proportional limit was defined as the stress level at which point the offset line intersected the load versus elongation curve. A servo-hydraulic machine was used for the vast majority of tensile testing. Testing from room temperature to elevated temperature was done in the same rig and differed only by the presence of a small furnace that heated the sample gage plus radius region to within 1% of the desired temperature. Fatigue testing was conducted at elevated temperature in load control. Testing was done in a programmable servo-hydraulic machine. The testing consisted of a 2 h dwell fatigue test with the load being relieved to 5% of the hold load at the end of each cycle. This simulated service conditions for the proposed application. Most tests were run to failure or 500 h of fatigue exposure. All samples that made 500 h of exposure were then tested for residual tensile properties at room temperature to note if degradation had occurred. A selected number of fatigue tests was run to many thousands of hours before failure occurred, as will be described later. 3. RESULTS AND DISCUSSION 3.1. Fracture characteristics of Hi-Nicalon versus Sylramic SiC fiber MI composites Figure 3 shows the fracture surfaces of 1200°C tensile samples of MI SiC/SiC composites with either Sylramic or Hi-Nicalon fibers. From this figure, one can see that the fracture surface of the Hi-Nicalon fiber composite is much more fibrous in nature than that of the Sylramic fiber composite. The measured tensile strengths of the two composites are similar, at |286 MPa; however, the strain-to-failure of the HiNicalon fiber composite (0.48%) is over twice that of the Sylramic fiber composite (0.21%). Thus, the “toughness”, or inherent matrix crack-stopping ability, of the Hi-Nicalon fiber composite is greater than that of the Sylramic fiber composite, due to the frictional sliding and pull-out of the fibers—which contribute to the composite toughness [12]—being more extensive for the Hi-Nicalon fibers than for the Sylramic fibers. The reason for this may be related to the higher elastic modulus of the Sylramic fiber, or, more likely, the much higher surface roughness of the Sylramic fiber compared with the Hi-Nicalon fiber, as shown in Fig. 4. The surface roughness of the two fibers is a reflection of the relative SiC grain size of the fibers. Fibers with higher surface roughness have been found to have a pronounced influence on the fiber sliding behavior in ceramic-matrix composites [13]. While the Hi-Nicalon MI SiC/SiC composite system appears to have better toughness than the Sylramic fiber composite system, it does not perform well in high-temperature fatigue testing. As shown in Fig. 5, a Hi-Nicalon fiber composite failed after 38 h and 19 cycles in a 2 h hold-time tensile fatigue test at 1200°C and a stress of 160 MPa, whereas a Sylramic fiber composite ran out to 1000 h, 500 cycles, under the same test conditions. The stress level during this test was chosen to be somewhat higher than the proportional limit, or matrix microcrack stress of the composites, which at this temperature is |140 MPa for the Sylramic fiber composite and |125 MPa for the Hi-Nicalon fiber composite. The residual 1200°C tensile properties of the fatigued Sylramic fiber composite are not significantly different from those measured for an as-fabricated Sylramic fiber composite. From the fracture surfaces shown in Fig. 5, it can be seen that the amount of fiber pull-out for the fatigued Sylramic fiber composite is similar to that shown in Fig. 3 for the 1200°C tensile sample, but the fatigued Hi-Nicalon composite shows a region around the edges of the sample that appears very brittle. This embrittlement of the Hi-Nicalon fiber composite is even more apparent after tensile fatigue testing at a lower temperature of 650°C, as shown in Fig. 6. While both composites ran out to the test limit of 546 h, 273 cycles, the residual 650°C strength and strain-to-failure of the Hi-Nicalon fiber composite were significantly degraded, while those for the Sylramic fiber composite were not. 3.2. Fiber/matrix interfacial characteristics of HiNicalon versus Sylramic SiC fiber MI composites From transmission electron microscopy (TEM) thin-foil analysis of the interfacial region in an asprocessed Hi-Nicalon fiber MI composite, as shown in Fig. 7, it was found that the probable reason for the easy debonding and long fiber pull-out in these composites, as was shown in Fig. 3, is that a thin (|40 nm) carbon-rich layer has formed between the BN and the Hi-Nicalon fiber during the MI composite processing. This carbon-rich layer is probably a result of interaction of the oxygen in the BN layer (|6 at%) with the excess carbon in the Hi-Nicalon fiber at the MI composite processing temperature of >1400°C. Thicker, but similar, carbon layers have
BRENNAN: INTERFACIAL CHARACTERIZATION SylramicTM 5HS 1200°CUTs=288MPa,Ef=0.48% 1200CUTS=286MPa,f=0.21 UTRC Fig 3. Fracture surface comparison between MI SiC/SiC composites with Hi-Nicalon and Sylramic SiC fibers Hi-Nicalon Fiber Sylramie M Fiber 1331510,8010u05 18.Bk I RMS 2.137 nM I 17,561 Fig 4. Fiber surface roughness differences between Sylramic and Hi-Nicalon SiC fibers
4622 BRENNAN: INTERFACIAL CHARACTERIZATION Fig. 3. Fracture surface comparison between MI SiC/SiC composites with Hi-Nicalon and Sylramic SiC fibers. Fig. 4. Fiber surface roughness differences between Sylramic and Hi-Nicalon SiC fibers
BRENNAN: INTERFACIAL CHARACTERIZATION Hi-Nicalon 5HS SyIramicTM 5HS 1200"C, 160 MPa, 19 cycles, 38 hrs to failure 1200.C, 160 MPa, 501 cycles, 1002 hrs to runout (Residual 1200C UTS=248 MPa, Ef=0.21%) Fig. 5. Fracture surface comparison between Hi-Nicalon and Sylramic SiC fiber MI SiC/SiC composites after 200° C tensile fatigue been found in many ceramic-matrix composites con- much less fibrous and with much shorter fiber pull taining the oxygen-rich Nicalon SiC fibers [14]. This out. However, the mode of fracture does not change carbon-rich layer is not seen in CVI SiC-matrix com- after elevated-temperature fatigue testing, nor does posites with Hi-Nicalon fibers and BN interfaces, the residual tensile strength change significantly which are processed at much lower temperatures TEM thin-foil analysis was performed on both as-fab- (1000oC) than the MI composites ricated and high-temperature-fatigued Sylramic fiber TEM thin-foil analysis of the Hi-Nicalon fiber MI composites, with the results indicating that no carbon composite that was subjected to tensile fatigue at rich layer forms during composite processing, nor 650C, as was shown in Fig. 6, and resulted in a quite does a glassy silica layer form as a result of high weak and brittle composite, found that a glassy silica temperature fatigue testing. This is illustrated in Fig layer had replaced the carbon-rich layer between the 9 for the Sylramic fiber composite sample that was BN and Hi-Nicalon fiber surface. This glassy silica subjected to 650C tensile fatigue testineoon-nich layer, as shown in Fig. 8, appears to be a result of shown previously in Fig. 6. The lack of a car oxidation of the carbon-rich layer and then the Hi- layer next to fiber surface is undoubtedly due to the Nicalon fiber surface due to matrix cracks forming higher temperature stability of the stoichiometric, during the fatigue testing. This glassy layer appar- crystalline SiC Sylramic fiber, when compared with ently bonds the Bn strongly to the fiber and may the carbon-rich Hi-Nicalon SiC fiber, During fatigue weaken the Hi-Nicalon fiber itself. At higher tempera- at elevated temperatures, any matrix cracks that may tures, such as that experienced during the 1200oC form do not cause pipeline oxidation down the fatigue test, this glassy layer can actually start to con- fiber/BN interface without the presence of the carbon sume the BN layer, totally bonding up the fiber/matrix rich layer interface, as was seen around the periphery of the At high stresses during high-temperature tensile 1200 C fatigue fracture surface shown in Fig. 5 fatigue, cracks do form in the Sylramic fiber MI com- As shown previously, in contrast to the fibrous fast posites, as shown for a sample in Fig. 10 that was fracture surface of Hi-Nicalon fiber MI composites, tensile fatigued at 815C at a stress of 186 MPa, the fracture surface of Sylramic fiber composites is which is well above the matrix microcracking stress
BRENNAN: INTERFACIAL CHARACTERIZATION 4623 Fig. 5. Fracture surface comparison between Hi-Nicalon and Sylramic SiC fiber MI SiC/SiC composites after 1200°C tensile fatigue. been found in many ceramic-matrix composites containing the oxygen-rich Nicalon SiC fibers [14]. This carbon-rich layer is not seen in CVI SiC-matrix composites with Hi-Nicalon fibers and BN interfaces, which are processed at much lower temperatures (|1000°C) than the MI composites. TEM thin-foil analysis of the Hi-Nicalon fiber MI composite that was subjected to tensile fatigue at 650°C, as was shown in Fig. 6, and resulted in a quite weak and brittle composite, found that a glassy silica layer had replaced the carbon-rich layer between the BN and Hi-Nicalon fiber surface. This glassy silica layer, as shown in Fig. 8, appears to be a result of oxidation of the carbon-rich layer and then the HiNicalon fiber surface due to matrix cracks forming during the fatigue testing. This glassy layer apparently bonds the BN strongly to the fiber and may weaken the Hi-Nicalon fiber itself. At higher temperatures, such as that experienced during the 1200°C fatigue test, this glassy layer can actually start to consume the BN layer, totally bonding up the fiber/matrix interface, as was seen around the periphery of the 1200°C fatigue fracture surface shown in Fig. 5. As shown previously, in contrast to the fibrous fast fracture surface of Hi-Nicalon fiber MI composites, the fracture surface of Sylramic fiber composites is much less fibrous and with much shorter fiber pullout. However, the mode of fracture does not change after elevated-temperature fatigue testing, nor does the residual tensile strength change significantly. TEM thin-foil analysis was performed on both as-fabricated and high-temperature-fatigued Sylramic fiber composites, with the results indicating that no carbonrich layer forms during composite processing, nor does a glassy silica layer form as a result of hightemperature fatigue testing. This is illustrated in Fig. 9 for the Sylramic fiber composite sample that was subjected to 650°C tensile fatigue testing, as was shown previously in Fig. 6. The lack of a carbon-rich layer next to fiber surface is undoubtedly due to the higher temperature stability of the stoichiometric, crystalline SiC Sylramic fiber, when compared with the carbon-rich Hi-Nicalon SiC fiber. During fatigue at elevated temperatures, any matrix cracks that may form do not cause pipeline oxidation down the fiber/BN interface without the presence of the carbonrich layer. At high stresses during high-temperature tensile fatigue, cracks do form in the Sylramic fiber MI composites, as shown for a sample in Fig. 10 that was tensile fatigued at 815°C at a stress of 186 MPa, which is well above the matrix microcracking stress