J.Am. Ceran.Soe.,90214050-4054(2007) Dol:10.ll1551-29162007.02060x c 2007 The American Ceramic Society journal Low-Temperature Oxidation Embrittlement of SiC ( Nicalon )/CAS Ceramic Matrix Composites Kevin p. plucknett*, f Materials Engineering Program, Department of Process Engineering and Applied Science, Dalhousie University Halifax, Canada, B3J 1Z1 Hua-Tay Ceramic Science and Technology Group, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6068 The influence of extended duration(up to 1000 h), low ter ture (i.e,>800C)exposure on materials with a carbon-based ture oxidation heat-treatments(375-600.C)has been assess fiber/matrix interlayers in an oxidizing environment, typically tibe g a model ceramic matrix composite system with a graphitic under fast fracture conditions. Subsequent to that work,the r matrix interphase. For this study a Nicalon"fiber effects of extended duration oxidation exposure were exam- reinforced CaO-Al2OxSiOz matrix composite was selected ined, and particular emphasis was placed upon composite ( Nicalon"/CAS), which possesses a thin (20-40 nm) stability in the intermediate temperature range(i.e, 5000-800oC) for both unloaded and static fatigue-loaded conditions. It is now ed under both unloaded and static fatigue-loaded conditions, fe well established that, in the absence of a protective coating. pipe-line "oxidation of the carbon- or boron nitride-based in- 1000 h terlayer occurs, resulting in considerable property degrada exposure, resulting in a transition to a nominally brittle tion. While high temperature pretreatments can inhibit this node (i.e, negligible fiber pull-out). The degree of effect, by"plugging"the exposed fiber ends with SiO2, they mechanical property degradation increases with increasing tem- are only useful for applications at stresses below the onset of perature, such that strength degradation, and a transition to matrix microcracking (i.e, the composite proportional limit) minally brittle failure, is apparent after just 10 h at 600C More recently, these studies of CMc behavior have been ex- Static fatigue loading between 450 and 600C demonstrated tended to assess cyclic fatigue loading at intermediate tempera- generally similar trends, with reduced lifetimes being observed tures. It was noted that a significant decrease in the fatigue with increasing temperature Based upon the unloaded oxidation limit occurs at 800oC, relative to room temperature, with a re- experiments, combined with previously obtained intermediate duction down to the microcracking stress, ome. In a manner nd high-temperature oxidation stability studies, a simple envi- similar to unstressed intermediate temperature aging, this degra- onmental embrittlement failure mechanism map is presen dation behavior was attributed to oxidation of the carbon inter for Nicalon"/CAS. The implications of this study for advanced phase and the subsequent formation of a silica bridge between omposite designs with multiple thin carbon-based interphase the fiber and matrix, which results in a strong fiber-matrix bond vers are also discussed While these previous studies have concentrated upon the effects of high and intermediate temperature oxidation expo- L. Introduction tability of such materials at lower temperatures (i.e, below 700C), where carbon oxidation can still occur. For example, it TINUoUS fiber-reinforced ceramic matrix composites known that carbon/carbon composites exhibit oxidation at (CMCs) are promising candidates for a number of ad- mperatures as low as 400 C, and anced structural applications, including use in gas turbine en- ticipated that CMCs with carbon-based interlayers could show gines, automotive brakes, and heat exchangers. While there ong-term property degradation at similar temperatures. Sim are now several classes of these materials, developed with both larly, there have been several recent studies on the cyclic thermal dense and porous matrices, the majority of prior work has ex- shock degradation of Nicalon"/CAs within comparable tem- plored the development and mechanical assessment of material erature ranges 24 25 In the present work, particular emphasis with compliant interphases between the fiber and matrix. For has been placed upon assessing extended duration environmen- nonoxide composites, these interphase layers are required to tal stability at these lower temperatures, for both unstressed and lower the interfacial fracture energy and allow fiber-matrix stressed conditions debonding, and are typically based on the use of materials ch as carbon or A number of studies have environmental stability of Sic- nforced Cl Il. Experimental Procedures from both experimental and theoretical p es.- Sever. al of these have primarily assessed the effects of high ter Oxidation heat treatments have been performed on a continuous Nicalon fiber-reinforced glass-ceramic composite, with a devitrified Cao-AlO3-SiO2 matrix(Nicalon"/CAS Type Il. R. Naslain--contnbuting editor 10, 90]3s, Corning, NY). The primary crystalline phase in this naterial is anorthite( CaAl,Si,O%), with a small amount of fine (<I um)zircon(ZrO4) precipitates present at the fiber /matrix interface. 6. I4A continuous, in situ formed carbon layer is preser between the fiber and matrix, with a thickness typically between addressed. e-mail: kevin plucknett(a dal. 20 and 40 nm. A summary of the as-received Nicalon"/CAS
Low-Temperature Oxidation Embrittlement of SiC (Nicalont)/CAS Ceramic Matrix Composites Kevin P. Plucknett* ,w Materials Engineering Program, Department of Process Engineering and Applied Science, Dalhousie University, Halifax, Canada, B3J 1Z1 Hua-Tay Lin* Ceramic Science and Technology Group, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6068 The influence of extended duration (up to 1000 h), low temperature oxidation heat-treatments (3751–6001C) has been assessed using a model ceramic matrix composite system with a graphitic fiber/matrix interphase. For this study a Nicalont fiber reinforced CaO–Al2O3–SiO2 matrix composite was selected (Nicalont/CAS), which possesses a thin (B20–40 nm) carbon-based interphase. Oxidation exposure has been conducted under both unloaded and static fatigue-loaded conditions. For unstressed oxidation exposure, degradation of the carbon-based interphase is apparent at temperatures as low as 3751C, after 1000 h exposure, resulting in a transition to a nominally brittle failure mode (i.e., negligible fiber pull-out). The degree of mechanical property degradation increases with increasing temperature, such that strength degradation, and a transition to nominally brittle failure, is apparent after just 10 h at 6001C. Static fatigue loading between 4501 and 6001C demonstrated generally similar trends, with reduced lifetimes being observed with increasing temperature. Based upon the unloaded oxidation experiments, combined with previously obtained intermediate and high-temperature oxidation stability studies, a simple environmental embrittlement failure mechanism map is presented for Nicalont/CAS. The implications of this study for advanced composite designs with multiple thin carbon-based interphase layers are also discussed. I. Introduction CONTINUOUS fiber-reinforced ceramic matrix composites (CMCs) are promising candidates for a number of advanced structural applications, including use in gas turbine engines, automotive brakes, and heat exchangers.1–6 While there are now several classes of these materials, developed with both dense and porous matrices, the majority of prior work has explored the development and mechanical assessment of materials with compliant interphases between the fiber and matrix. For nonoxide composites, these interphase layers are required to lower the interfacial fracture energy and allow fiber–matrix debonding, and are typically based on the use of materials such as carbon or boron nitride. A number of studies have addressed issues related to the environmental stability of SiC-based fiber-reinforced CMCs, from both experimental and theoretical perspectives.7–20 Several of these have primarily assessed the effects of high temperature (i.e., 48001C) exposure on materials with a carbon-based fiber/matrix interlayers in an oxidizing environment, typically under fast fracture conditions.7–10 Subsequent to that work, the effects of extended duration oxidation exposure were examined,11–17 and particular emphasis was placed upon composite stability in the intermediate temperature range (i.e., 5001–8001C) for both unloaded and static fatigue-loaded conditions. It is now well established that, in the absence of a protective coating, ‘‘pipe-line’’ oxidation of the carbon- or boron nitride-based interlayer occurs,10–20 resulting in considerable property degradation. While high temperature pretreatments can inhibit this effect,21 by ‘‘plugging’’ the exposed fiber ends with SiO2, they are only useful for applications at stresses below the onset of matrix microcracking (i.e., the composite proportional limit). More recently, these studies of CMC behavior have been extended to assess cyclic fatigue loading at intermediate temperatures.22 It was noted that a significant decrease in the fatigue limit occurs at 8001C, relative to room temperature, with a reduction down to the microcracking stress, smc. In a manner similar to unstressed intermediate temperature aging, this degradation behavior was attributed to oxidation of the carbon interphase and the subsequent formation of a silica bridge between the fiber and matrix, which results in a strong fiber–matrix bond. While these previous studies have concentrated upon the effects of high and intermediate temperature oxidation exposure, there is only minimal information regarding the long-term stability of such materials at lower temperatures (i.e., below 7001C), where carbon oxidation can still occur. For example, it is known that carbon/carbon composites exhibit oxidation at temperatures as low as 4001C,23 and consequently it can be anticipated that CMCs with carbon-based interlayers could show long-term property degradation at similar temperatures. Similarly, there have been several recent studies on the cyclic thermal shock degradation of Nicalont/CAS within comparable temperature ranges.24,25 In the present work, particular emphasis has been placed upon assessing extended duration environmental stability at these lower temperatures, for both unstressed and stressed conditions. II. Experimental Procedures Oxidation heat treatments have been performed on a continuous Nicalont fiber-reinforced glass–ceramic composite, with a devitrified CaO–Al2O3–SiO2 matrix (Nicalont/CAS Type II, [01, 901]3S, Corning, NY). The primary crystalline phase in this material is anorthite (CaAl2Si2O8), with a small amount of fine (o1 mm) zircon (ZrO4) precipitates present at the fiber/matrix interface.6,14 A continuous, in situ formed carbon layer is present between the fiber and matrix, with a thickness typically between 20 and 40 nm. A summary of the as-received Nicalont/CAS R. Naslain—contributing editor *Member, the American Ceramic Society w Author to whom correspondence should be addressed. e-mail: kevin.plucknett@dal.ca Manuscript No. 23158. Received May 2, 2007; approved July 26, 2007. Journal J. Am. Ceram. Soc., 90 [12] 4050–4054 (2007) DOI: 10.1111/j.1551-2916.2007.02060.x r 2007 The American Ceramic Society 4050
December 2007 Communications of the American Ceramic Society Table I. Basic Physical Properties of the As-Received after oxidation at 375.C(Fig. 1). However, a transition from Nicalon"/CAS Composite, Including Three-Point Flexure typical"pseudo-ductile"composite failure(with a failure"tail Mechanical Data observed on the load-deflection curve indicative of fiber pull- out)to a brittle failure mode (with no failure"tail")occurs be- Measured value tween 100 and 1000 h(Fig. 2). Similar behavior is apparent at Fiber fraction(vol%) 38.1+4.5 450oC. however the transition to a brittle failure mode is accel 2.67+0.0 erated, occurring between 10 and 100 h of oxidation exposure cn At 525 and 600C, the maximum flexure strength decreases Flexure modulus(GPa) 92.0+5.0 Flexure strength, or(MPa) 0l0+14. continuously with increasing heat-treatment duration, although a slight strength increase is initially noted after oxidation at Flexure proportional limit, omc(MPa) 1160+4.0 525C for 10 h; an essentially brittle fracture mode is observed for all these examples. The fracture transition is accompanied by hange from a fibrous failure surface (i.e, considerable fiber composite properties is presented in Table I. Bars for flexure pull-out)to one where there is minimal fiber pull out testing were cut from a hot-pressed plate(thickness 2.35 mm) to dimensions of 50-mm length x 5-mm width. Weight change The influence of oxidation exposure time and temperature, after oxidation was determined to +0.0001 g. Preliminary ex- from 375. to 600C, upon the mass loss of the Nicalon"/CAS amination of samples oxidized at 300C for 1000 h showed neg- exure bars is demonstrated in Fig. 3. At 375C, the mass loss ligible change in weight (i.e, within measuring error).A exhibits a nominally linear dependence upon time. Increasing the consequence, oxidation temperatures from 375 to 600C were oxidation temperature results in significant deviation from the behavior of the material observed at 375C. at both 450. and (model 1200, Barnstead/Thermolyne, Dubuque, IA). for dura- however the samples have subsequently increased in mass after tions of up to 1000 h. Test bars were sited on a high-purity 1000 h exposure. This effect is even more apparent at 600C Al2O3 support, and inserted into and removed from the furnace where weight loss is observed within the first 10 h, with subse- It temperature. There was no evidence of sample/support read quent weight gain after this time. It was noted that for oxidation tion with this arrangement, for any of the oxidation condition periods in excess of 100 h, samples treated at 600C showed tudied. Flexure testing was performed at a displacement rate of essentially no further weight gain within experimental error a three-point bend fixture, with an ou loading span of 40 mm, giving a span to thickness ratio of 17: 1. At least two tests were performed for each combination 40-nm thick, which has a turbostratic-type layered graph tructure.In previous work on Nicalon"/CAs, scanning treatments, the samples were preloaded to the desired stress in a Auger microscopy analysis demonstrated that the carbon layer Sic four-point bend fixture, with inner and outer spans of 6.35 d40 ely, held within a Mosi, elemen t furnace in ccordance with the observation of weight loss noted in the air. Samples were then heated to the desired test temperature, present work( Fig. 3). The weight loss data demonstrates similar and held at this temperature until failure. Microstructures of the oxidative removal of carbon at all examined temperatures received and oxidized specimens were examined by field emis between 375 and 600C. However, it is also apparent that, sion scanning electron microscopy(FE-SEM: Hitachi S-4500. with the exception of exposure at 375.C, mass gain occurs after Tokyo, Japan) n initial period of loss. This behavior is indicative of oxidation of the exposed SiO-C fiber surface, forming amorphous Sio with an associated increase in mass. Several previous studies II. Results and discussion have noted the formation of sio, ligaments that bridge across The effects of low-temperature oxidation heat treatments the fiber/matrix interface within this intermediate temperature the maximum flexure strength of Nicalon"/ CAS are presented ange 12,13 scanni er microscopy has also demonstrate of the load/deflection curves obtained for he presence of Sio, regions on the fiber surface in the case of Nicalon"/CAS after this series of heat treatments is show icalon"/ CAS oxidized at 600C for 100 h. schematically in Fig. 2. It is apparent that strength is retained Ultimately, the exposed fiber ends will be sealed via Sio and even slightly increased, for exposure times up to 1000 h, formation and plugging, preventing further oxidation of the will essential ease, as noted for samples held at 600C for times in excess of 100 h(Fig. 3). It is clear that care must be taken in imparting any significance upon the apparent transition from linear oxi dation kinetics at 375C to the behavior at 450. and 525C. as carbon weight loss through oxidation will be masked by weight gain through SiO formation, thus complicating analysis. It is now well established that in this situation carbon removal occurs via a"pipe-line"oxidation process. Huger et al.developed a simple geometrical model to estimate the critical time, Ie, taken for sealing of the exposed fiber ends, such that carbon oxidation ceases. This model was based on oxidation of SiC-SiC compos- ites, where both the fiber and matrix can oxidize. Their a 国 proach was subsequently adapted for the case of high temperature sealing of composites where only the fiber can oxidize. For a nonoxidizing matrix the critical sealing time Ic, Is given by Fig. 1. Three point flexure strength of Nicalon"/CAS after isothermal B(-(1/01) xidation at temperatures between 375 and 600C, for up to 1000 h The as-received flexure strength, before oxidation exposure, was 501+14 where Br is the rate constant for oxidation of fibers in air. H is on interphase thickness, and er is
composite properties is presented in Table I. Bars for flexure testing were cut from a hot-pressed plate (thickness B2.35 mm) to dimensions of 50-mm length 5-mm width. Weight change after oxidation was determined to 70.0001 g. Preliminary examination of samples oxidized at 3001C for 1000 h showed negligible change in weight (i.e., within measuring error). As a consequence, oxidation temperatures from 3751 to 6001C were subsequently examined. Unstressed oxidation heat treatments were performed in a conventional laboratory muffle furnace (model 1200, Barnstead/Thermolyne, Dubuque, IA), for durations of up to 1000 h. Test bars were sited on a high-purity Al2O3 support, and inserted into and removed from the furnace at temperature. There was no evidence of sample/support reaction with this arrangement, for any of the oxidation conditions studied. Flexure testing was performed at a displacement rate of 0.5 mm/min using a three-point bend fixture, with an outer loading span of 40 mm, giving a span to thickness ratio of B17:1. At least two tests were performed for each combination of oxidation temperature and time. For the stressed fatigue heat treatments, the samples were preloaded to the desired stress in a SiC four-point bend fixture, with inner and outer spans of 6.35 and 40 mm, respectively, held within a MoSi2 element furnace in air. Samples were then heated to the desired test temperature, and held at this temperature until failure. Microstructures of the as-received and oxidized specimens were examined by field emission scanning electron microscopy (FE-SEM; Hitachi S-4500, Tokyo, Japan). III. Results and Discussion The effects of low-temperature oxidation heat treatments upon the maximum flexure strength of Nicalont/CAS are presented in Fig. 1. A summary of the load/deflection curves obtained for Nicalont/CAS after this series of heat treatments is shown schematically in Fig. 2. It is apparent that strength is retained, and even slightly increased, for exposure times up to 1000 h, after oxidation at 3751C (Fig. 1). However, a transition from typical ‘‘pseudo-ductile’’ composite failure (with a failure ‘‘tail’’ observed on the load–deflection curve indicative of fiber pullout) to a brittle failure mode (with no failure ‘‘tail’’) occurs between 100 and 1000 h (Fig. 2). Similar behavior is apparent at 4501C, however the transition to a brittle failure mode is accelerated, occurring between 10 and 100 h of oxidation exposure. At 5251 and 6001C, the maximum flexure strength decreases continuously with increasing heat-treatment duration, although a slight strength increase is initially noted after oxidation at 5251C for 10 h; an essentially brittle fracture mode is observed for all these examples. The fracture transition is accompanied by a change from a fibrous failure surface (i.e., considerable fiber pull-out) to one where there is minimal fiber pull out. The influence of oxidation exposure time and temperature, from 3751 to 6001C, upon the mass loss of the Nicalont/CAS flexure bars is demonstrated in Fig. 3. At 3751C, the mass loss exhibits a nominally linear dependence upon time. Increasing the oxidation temperature results in significant deviation from the behavior of the material observed at 3751C. At both 4501 and 5251C, mass loss increases to a maximum after 100 h exposure, however the samples have subsequently increased in mass after 1000 h exposure. This effect is even more apparent at 6001C, where weight loss is observed within the first 10 h, with subsequent weight gain after this time. It was noted that for oxidation periods in excess of 100 h, samples treated at 6001C showed essentially no further weight gain within experimental error. Nicalont/CAS composites possess an in situ formed carbon interlayer between the fiber and matrix, typically between 20 and 40-nm thick, which has a turbostratic-type layered graphite structure.21,26 In previous work on Nicalont/CAS, scanning Auger microscopy analysis demonstrated that the carbon layer is completely removed after oxidation at 4501C for 100 h,27 in accordance with the observation of weight loss noted in the present work (Fig. 3). The weight loss data demonstrates similar oxidative removal of carbon at all examined temperatures between 3751 and 6001C. However, it is also apparent that, with the exception of exposure at 3751C, mass gain occurs after an initial period of loss. This behavior is indicative of oxidation of the exposed Si–O–C fiber surface, forming amorphous SiO2 with an associated increase in mass. Several previous studies have noted the formation of SiO2 ligaments that bridge across the fiber/matrix interface within this intermediate temperature range.12,13 Scanning Auger microscopy has also demonstrated the presence of SiO2 regions on the fiber surface in the case of Nicalont/CAS oxidized at 6001C for 100 h.27 Ultimately, the exposed fiber ends will be sealed via SiO2 formation and plugging, preventing further oxidation of the interlayer.21 At this point, further mass gain will essentially cease, as noted for samples held at 6001C for times in excess of 100 h (Fig. 3). It is clear that care must be taken in imparting any significance upon the apparent transition from linear oxidation kinetics at 3751C to the behavior at 4501 and 5251C, as carbon weight loss through oxidation will be masked by weight gain through SiO2 formation, thus complicating analysis. It is now well established that in this situation carbon removal occurs via a ‘‘pipe-line’’ oxidation process.18–20 Huger et al.28 developed a simple geometrical model to estimate the critical time, tc, taken for sealing of the exposed fiber ends, such that carbon oxidation ceases. This model was based on oxidation of SiC–SiC composites, where both the fiber and matrix can oxidize. Their approach was subsequently adapted for the case of high temperature sealing of composites where only the fiber can oxidize.27 For a nonoxidizing matrix the critical sealing time, tc, is given by: tc ¼ 1 Bf H ð Þ 1 ð1=yfÞ 2 where Bf is the parabolic rate constant for oxidation of the fibers in air, H is the carbon interphase thickness, and yf is the Table I. Basic Physical Properties of the As-Received Nicalont/CAS Composite, Including Three-Point Flexure Mechanical Data Property Measured value Fiber fraction (vol%) 38.174.5 Density (g/cm3 ) 2.6770.03 Flexure modulus (GPa) 92.075.0 Flexure strength, sf (MPa) 501.0714.0 Flexure proportional limit, smc (MPa) 116.074.0 0 100 200 300 400 500 600 10 100 1000 375 450 525 600 Oxidation time (hrs) Flexure strength (MPa) Fig. 1. Three point flexure strength of Nicalont/CAS after isothermal oxidation at temperatures between 3751 and 6001C, for up to 1000 h. The as-received flexure strength, before oxidation exposure, was 501714 MPa. December 2007 Communications of the American Ceramic Society 4051
Commmunications of the American Ceramic Society VoL. 90. No. 12 Sreceived 525°C igd boo ch for up o eoo h.T he behavi oad-deflection curves obtained after isothermal oxidation of Nicalon"/CAS at temperatures between 375 of the as-received material is also shown for reference expansion coefficient of the fiber due to oxidation.2. For the significant effect upon these interfacial properties. After heat case of Nicalon NLM-202, the fiber used in the present Cas treatment at 375C the debonding energy is significantly lower matrix composite, er, is determined to be 1.48.- Based on this, than for the as-received material, and essentially unchange an estimate for the sealing time can be determined from the ox- upon subsequent reloading, indicating that there is no rea idation kinetics of Nicalon NLM-202 fibers At the lowest interfacial chemical bond present after this heat treatment. temperature examined in that study, namely 700%C, sealing for a However, the frictional sliding stress is comparable with the in approximately 35 h for a 50-nm thick carbon layer, 5 h for from Fig 3 that only partial carbon removal occurs after 1.. glass-ceramic matrix composite can be predicted to occu as-received composite with a carbon interlayer. It can be seen a 20-nm thick layer, and I h 20 min for a 10-nm thick layer. The as weight loss continues essentially linearly up to 1000 h, which residual stress state of this composite is such that the matrix will infers that a thin carbon layer appears to be retained after 100 h. lamp down onto the fiber if the carbon layer is removed, thus The interfacial debonding energy and frictional sliding stress are aking the apparent fiber-matrix gap smaller and the sealing both increased by heat treatment at 450C (Table In). Given time shorter. Oxidation kinetics data are not readily available the absence of a lubricating carbon interlayer after this heat for lower temperatures, but extrapolation of the behavior ob- treatment, a dramatically increased frictional sliding stress can served at higher temperatures indicates that sealing will occur be anticipated if the matrix and fiber come into direct contact after approximately 20 h at 600C for a 20-nm thick fiber-m The increased debond energy is also believed to be due to the rix gap This estimation is in general accordance with the weigh arge frictional sliding stress that must be overcome before the gain response noted in Fig. 3, where weight gain ceases before fiber can be displaced, although the formation of isolated SiO2 0o h exposure. A more thorough determination of sealing be- bridges between the matrix and fiber cannot be completely dis- 5-600oC)would require counted (particularly given the observed weight gain after ex- continuous thermogravimetric analysis of both the fiber and tended heat-treatment at 450.C(Fig. 3). It was interesting to composite oxidation weight changes as a function of time, which note from this prior study that, upon reloading. the debond load is beyond the scope of the present study is essentially the same as the first loading cycle (i.e, 200 mN Prior work has demonstrated the effects of oxidation(for 10 which again indicates the lack of any interfacial chemical bond h)upon the interfacial micro-mechanical properties, namely the Examination of the surface of a polished sample, after sub- debonding energy (Gi) and frictional sliding stress([). Data sequent unstressed oxidation at 450.C for 100 h, demonstrates obtained in that prior study is presented in Table II for refer- that the Nicalon fibers protrude from the matrix by 1-2 um. ence. It is apparent that heat-treatment temperature has a This behavior arises from the mismatch in thermal expansion coefficient between the fibers, af. and the CAs matrix, am. Based upon published data, 31.32 a thermal expansion coefficient mis- 围 match,△x, can be determined, such that aa=am-ar≈1.9×10-6 K. Consequently, clamping of the fiber by the matrix can be nticipated when cooling from the initial processing temper ture. For example, for a 15-um diameter fiber the difference in shrinkage relative to the matrix sur 这0.04 17 nm, based on a AT of 1175.C (i.e, a processing temperature of 1200.C cooled to room temperature). At the same time. the 0.06 u>\able ll. Fiber Push Down Interfacial Properties of alon"/CAS in Both As-Received and Aged Conditions at treatment condition Debond energy, G,(/m) Sliding stress, t(MPa) 8.0±30 25+5 0.10 1.1+1.0 1200 40+14 Oxidation time(hrs) 450°C/100h 4.7+3.7 l44+63 525°C/100h 6.3+4.3 177+54 Fig 3. NIC loss after isothermal oxidation at 60 C for up to 1000 h 93±55 193+5
expansion coefficient of the fiber due to oxidation.27,28 For the case of Nicalont NLM-202, the fiber used in the present CAS matrix composite, yf, is determined to be 1.48.28 Based on this, an estimate for the sealing time can be determined from the oxidation kinetics of Nicalont NLM-202 fibers.29 At the lowest temperature examined in that study, namely 7001C, sealing for a glass–ceramic matrix composite can be predicted to occur in approximately 35 h for a 50-nm thick carbon layer, 5 h for a 20-nm thick layer, and 1 h 20 min for a 10-nm thick layer. The residual stress state of this composite is such that the matrix will clamp down onto the fiber if the carbon layer is removed, thus making the apparent fiber–matrix gap smaller and the sealing time shorter. Oxidation kinetics data are not readily available for lower temperatures, but extrapolation of the behavior observed at higher temperatures indicates that sealing will occur after approximately 20 h at 6001C for a 20-nm thick fiber–matrix gap. This estimation is in general accordance with the weight gain response noted in Fig. 3, where weight gain ceases before 100 h exposure. A more thorough determination of sealing behavior at these lower temperatures (3751–6001C) would require continuous thermogravimetric analysis of both the fiber and composite oxidation weight changes as a function of time, which is beyond the scope of the present study. Prior work has demonstrated the effects of oxidation (for 100 h) upon the interfacial micro-mechanical properties, namely the debonding energy (Gi) and frictional sliding stress (t).30 Data obtained in that prior study is presented in Table II for reference. It is apparent that heat-treatment temperature has a significant effect upon these interfacial properties. After heattreatment at 3751C the debonding energy is significantly lower than for the as-received material, and essentially unchanged upon subsequent reloading, indicating that there is no real interfacial chemical bond present after this heat treatment.30 However, the frictional sliding stress is comparable with the as-received composite with a carbon interlayer. It can be seen from Fig. 3 that only partial carbon removal occurs after 100 h, as weight loss continues essentially linearly up to 1000 h, which infers that a thin carbon layer appears to be retained after 100 h. The interfacial debonding energy and frictional sliding stress are both increased by heat treatment at 4501C (Table II). Given the absence of a lubricating carbon interlayer after this heat treatment,27 a dramatically increased frictional sliding stress can be anticipated if the matrix and fiber come into direct contact. The increased debond energy is also believed to be due to the large frictional sliding stress that must be overcome before the fiber can be displaced, although the formation of isolated SiO2 bridges between the matrix and fiber cannot be completely discounted (particularly given the observed weight gain after extended heat-treatment at 4501C (Fig. 3)). It was interesting to note from this prior study that, upon reloading, the debond load is essentially the same as the first loading cycle (i.e., B200 mN), which again indicates the lack of any interfacial chemical bond.30 Examination of the surface of a polished sample, after subsequent unstressed oxidation at 4501C for 100 h, demonstrates that the Nicalont fibers protrude from the matrix by 1–2 mm. This behavior arises from the mismatch in thermal expansion coefficient between the fibers, af, and the CAS matrix, am. Based upon published data,31,32 a thermal expansion coefficient mismatch, Da, can be determined, such that Da 5 am–af1.9 106 K1 . Consequently, clamping of the fiber by the matrix can be anticipated when cooling from the initial processing temperature. For example, for a 15-mm diameter fiber the difference in shrinkage relative to the matrix surrounding it is approximately 17 nm, based on a DT of 11751C (i.e., a processing temperature of 12001C cooled to room temperature). At the same time, the 375°C 450°C 525°C 600°C As-received 1 mm 100 N 1000 h 100 h 10 h Fig. 2. Schematic overview of the typical load-deflection curves obtained after isothermal oxidation of Nicalont/CAS at temperatures between 3751 and 6001C, for up to 1000 h. The behavior of the as-received material is also shown for reference. 0.00 0.02 0.04 0.06 0.08 0.10 0 200 400 600 1000 1200 375 450 525 600 Oxidation time (hrs.) 800 Mass loss (%) Fig. 3. Nicalont/CAS test bar weight loss after isothermal oxidation at temperatures between 3751 and 6001C for up to 1000 h. Table II. Fiber Push Down Interfacial Properties of Nicalont/CAS in Both As-Received and Aged Conditions30 Heat treatment condition Debond energy, Gi (J/m2 ) Sliding stress, t (MPa) As-received 8.073.0 2575 3751C/100 h 1.171.0 40714 4501C/100 h 4.773.7 144763 5251C/100 h 6.374.3 177754 6001C/100 h 9.375.5 193757 4052 Communications of the American Ceramic Society Vol. 90, No. 12
December 2007 Communications of the American Ceramic Society 4053 10000 (strong 1000 300 Composit i Composite Oxidation temperature (C Fig 4. Static fatigue lifetimes for Nicalon"/CAS tested in air at 4500 Fig. 5. Environmental embrittlement 525, and 600C. Arrows indicate samples that survived after 1000 h dation of Nicalon"/CAS, demonstrat Note that all testing is conducted above the unction of oxidation temperature and aterial, which was determined to be 116+4 MPa in three-point bend. failure: .. brittle failure. Note that the ided into regions of high, medium, and low strength. oughness of Nicalon fibers are typically on the order degradation noted in prior studies is readily apparent, and dem- Powell et al. presented a shear-lag model that de- onstrates the issues relating to interphase stability quite clearly he residual thermal stress state in Nicalon/ CAS after While Nicalon"/ CAS can now be viewed primarily as a model It was noted in that study that surface effects are sig- material for studies such as these, the present observations of such that the residual stress after fabrication and cool- low temperature degradation have significant implications for om temperature will be sufficient to result in a surface more advanced generations of non-oxide ceramic matrix com- egion equivalent to at least one-or two-fiber diam- posites. For example, the recent approach of developing com- eters in depth. Clearly, in the present case with pa emoval posites with multiple Sic/c interphase layers, where the carbon of the carbon layer this effect is exacerbated, and there will b yers are designed to be very thin to allow rapid sealing, may complete debonding for the relatively small sample size of a ot prevent degradation when flexure bar tion at low temperatures in the absence of a protective coating. It is apparent that the interfacial instabilities have a signifi Even when such a coating is present, the demands on it will cant effect upon the c ite macro-mechanical behavior. Re- gnificant as it must protect the material at wide extremes of moval of the carbon interface results in clamping of the fiber by temperature, in environments where stress is likely to be applied the surrounding matrix. In this instance both the interfacial At high temperatures some degree of coating compliance is like- debond energy and the frictional sliding stress are increased ly to be a typical feature however this may not be the case at substantially. The result of this interface modification is that nterfacial debonding criteria, where the interface fracture latively low temperatures, where the simple glass-sealing tech- iques described earlier will be ineffective. energy is significantly lower than the fiber fracture energy (i.e. Gi<0.25Gr, where Gr is the fiber fracture energy), are no longer atisfied. Consequently, debonding and fiber slie ng nt occur. and a transition to brittle failure is observed Figure 4 demonstrates the effect of combined applied stress Nicalon"/CAS composites have been exposed to oxidizing heat d oxidation temperature upon the static fatig treatments between 375 and 600 C. for both unstressed and CAS/Nicalon"composites. It is clear that, for a static fatigue-loaded conditions. While strength is essentially re- tress,there is a strong correlation between the composite life- tained after unstressed oxidation at the lowest temperatures (i. e, time and the oxidation temperature, in accordance with the un- 375 and 450C), a transition to brittle failure occurs with in- loaded oxidation response. In each of these cases the static time (i.e. between 100 and 1000 h at 375( applied stress exceeds the microcracking stress, omc, of the com- and between 10 and 100 h at 450C). At 525 and 600oC posite in flexure(-116.0+4.0 MPa); consequently the compos- strength decreases with increasing oxidation time, and the fail- ite matrix is microcracked for the entire test duration un ure mode is brittle for all examples. These changes in mechanical failure. In this instance, the presence of microcracking in the behavior arise from removal of the carbon-based fiber /matrix matrix significantly increases the number of paths for oxygen terphase, via"pipe-line oxidation ngress to the fiber-matrix interface. Generally similar observa creased friction present at the fiber/matrix interface(arisin tions to these were made by Yasmin and Bowen, for cyclic due to the residual stress state in this particular system, with fatigues loading at 800C. At room temperature, they noted that the higher CTE matrix effectively"clamping-down""onto the II composite was 200 MPa(defined as successful completion under static fatigue loading in oxidizing environments, with life- of 10 cycles). However, the fatigue limit was reduced signifi- times significantly reduced under increasing temperature when to 100 MPa at 800C, which corresponds with an ples are loaded above the materials proportional limit, ome stress just below the microcracking stress of this mate when matrix microcracking will occur. Based upon the current r applied cyclic stresses above omc, the fatigue lifetimes study and previously published data, an environmental embrit were dramatically reduced at the elevated test temperature. tlement failure mechanism map has been developed for Nica- Based upon the data outlined in Fig. 1, in combination with lon"/CAS, which highlights the region of intermediate that presented in previous work for the oxidation behavior of temperature embrittlement as a function of ex /CAs at higher temperature ture and time. The present work has demonstrated that oxida to develop a failure mechanism map for Nicalon"/CAS after tion degradation can occur in CMCs with a carbon-based oxidation exposure (Fig. 5). The intermediate temperature terphase, even at temperatures as low as 375%C, which may
surface roughness of Nicalont fibers are typically on the order of 5 nm.33 Powell et al. 34 presented a shear-lag model that describes the residual thermal stress state in Nicalont/CAS after synthesis. It was noted in that study that surface effects are significant, such that the residual stress after fabrication and cooling to room temperature will be sufficient to result in a surface debonded region equivalent to at least one- or two-fiber diameters in depth. Clearly, in the present case with partial removal of the carbon layer this effect is exacerbated, and there will be near complete debonding for the relatively small sample size of a flexure bar. It is apparent that the interfacial instabilities have a signifi- cant effect upon the composite macro-mechanical behavior. Removal of the carbon interface results in clamping of the fiber by the surrounding matrix. In this instance, both the interfacial debond energy and the frictional sliding stress are increased substantially. The result of this interface modification is that interfacial debonding criteria, where the interface fracture energy is significantly lower than the fiber fracture energy (i.e., Gio0.25Gf, where Gf is the fiber fracture energy),35 are no longer satisfied. Consequently, debonding and fiber sliding no longer occur, and a transition to brittle failure is observed. Figure 4 demonstrates the effect of combined applied stress and oxidation temperature upon the static fatigue lifetime of CAS/Nicalont composites. It is clear that, for a given applied stress, there is a strong correlation between the composite lifetime and the oxidation temperature, in accordance with the unloaded oxidation response. In each of these cases the static applied stress exceeds the microcracking stress, smc, of the composite in flexure (B116.074.0 MPa); consequently the composite matrix is microcracked for the entire test duration until failure. In this instance, the presence of microcracking in the matrix significantly increases the number of paths for oxygen ingress to the fiber–matrix interface. Generally similar observations to these were made by Yasmin and Bowen,22 for cyclic fatigues loading at 8001C. At room temperature, they noted that the fatigue limit for an identical cross-ply Nicalont/CAS TypeII composite was B200 MPa (defined as successful completion of 106 cycles). However, the fatigue limit was reduced signifi- cantly to B100 MPa at 8001C, which corresponds with an applied stress just below the microcracking stress of this material. For applied cyclic stresses above smc, the fatigue lifetimes were dramatically reduced at the elevated test temperature. Based upon the data outlined in Fig. 1, in combination with that presented in previous work for the oxidation behavior of Nicalont/CAS at higher temperatures,21,26,27 it is possible to develop a failure mechanism map for Nicalont/CAS after oxidation exposure (Fig. 5). The intermediate temperature degradation noted in prior studies is readily apparent, and demonstrates the issues relating to interphase stability quite clearly. While Nicalont/CAS can now be viewed primarily as a model material for studies such as these, the present observations of low temperature degradation have significant implications for more advanced generations of non-oxide ceramic matrix composites. For example, the recent approach of developing composites with multiple SiC/C interphase layers, where the carbon layers are designed to be very thin to allow rapid sealing,36 may not prevent degradation when the material is exposed to oxidation at low temperatures in the absence of a protective coating. Even when such a coating is present, the demands on it will be significant as it must protect the material at wide extremes of temperature, in environments where stress is likely to be applied. At high temperatures some degree of coating compliance is likely to be a typical feature, however this may not be the case at relatively low temperatures, where the simple glass-sealing techniques described earlier will be ineffective.21 IV. Conclusions Nicalont/CAS composites have been exposed to oxidizing heattreatments between 3751 and 6001C, for both unstressed and static fatigue-loaded conditions. While strength is essentially retained after unstressed oxidation at the lowest temperatures (i.e., 3751 and 4501C), a transition to brittle failure occurs with increasing exposure time (i.e., between 100 and 1000 h at 3751C, and between 10 and 100 h at 4501C). At 5251 and 6001C, strength decreases with increasing oxidation time, and the failure mode is brittle for all examples. These changes in mechanical behavior arise from removal of the carbon-based fiber/matrix interphase, via ‘‘pipe-line’’ oxidation, and the subsequently increased friction present at the fiber/matrix interface (arising due to the residual stress state in this particular system, with the higher CTE matrix effectively ‘‘clamping-down’’ onto the Nicalont fiber). A similar embittlement mechanism is apparent under static fatigue loading in oxidizing environments, with lifetimes significantly reduced under increasing temperature when samples are loaded above the materials proportional limit, smc, when matrix microcracking will occur. Based upon the current study and previously published data, an environmental embrittlement failure mechanism map has been developed for Nicalont/CAS, which highlights the region of intermediate temperature embrittlement as a function of exposure temperature and time. The present work has demonstrated that oxidation degradation can occur in CMCs with a carbon-based interphase, even at temperatures as low as 3751C, which may 1 10 100 1000 0 100 200 300 400 500 600 450 525 600 0.001 0.01 0.1 Lifetime (hrs) Applied stress (MPa) Fig. 4. Static fatigue lifetimes for Nicalont/CAS tested in air at 4501, 5251, and 6001C. Arrows indicate samples that survived after 1000 h. Note that all testing is conducted above the proportional limit for this material, which was determined to be 11674 MPa in three-point bend. 0.1 1 10 100 1000 10000 0 200 400 600 800 1000 1200 Composite Brittle (strong) Brittle (medium strength) Brittle (weak) Composite Oxidation temperature (°C) Oxidation time (hrs) Fig. 5. Environmental embrittlement failure mechanism map for oxidation of Nicalont/CAS, demonstrating the failure mode observed as a function of oxidation temperature and exposure duration. m, composite failure; , brittle failure. Note that the brittle failure regions are subdivided into regions of high, medium, and low strength. December 2007 Communications of the American Ceramic Society 4053
4054 Commmunications of the American Ceramic Society VoL. 90. No. 12 have significant implications for the long term use of such ma- aslan, " Oxidation Mechanisms and Kinetics of ID-SiCl an extremely wide temperature range (i.e, 300-1200.C C/SiC Composite-Materials, 2: Modeling, J. An. Ceram. Soc., 77[2] 467-80 A. J. Eckel, J. D. Cawley, and T. A. Parthasarathy, "Oxidation-Kinetics of a Continuous Carbon Phase in a Non-Reactive Matrix. " J. Anm Ceram Soc., 78[ Acknowledgments 972-80(1995) Morscher and J. D. Cawley The authors would thank David Larsen. Corn 豇DT K. P ur. Ceran.Soc2214-15277-87(2002) (Corming. NY), for provision of the Nicalon"/CAS Type Il co ilicate/Nicalon by High-Temperature Pre- Treatment,J.Mater. Sci. Lett., 14 [171 1223-6(190 References Cross-Ply Nicalon/CAS-Il omposite at Room and Elevated Temperature, Com- M. van Roode, J. Price, J. Kimmel, N. Miriyala. D. Leroux, A. Fahme, and Immary of Field E. Westwood. J D H Evaluations. "J. Eng. Gas Turbines Power, 129 [1] 21-30(2007). "Oxidation Protection for Carbon Fibre Composites. "J Mater. Sci. 31(6]1389- P. Barnard, M. B. Henderson, and N. Rhodes. "" CMC Integration and Dem- J. Blissett. P. A. Smith. and J. A. Yeomans. "Flexural mechanical Properties of Non-Oxide CMCs for Ap- einforced Calcium Aluminosilicate Composites, "J. Mater. Sci., 33[16]4181 plication in Engines and Nuclear Reactors: An Overview. "Camp. Sci. Tech, 64[2] 90(1998) Kastntseas P.A. Smith, and J. A. Yeomans. "Damage Characterisation o Ch. Zuber and B. Heidenreich. ""Development of a Net Shape Manufacturing hermally Shocked Cross-Ply Ceramic Composite Laminates, "J. Mater. Sci, ngle piece stofftech, 37[]301-8(2006) K.P. Plucknett, R. L. Cain, and M. H. Lewis. ""Interface Degradat Friction Systems,"a.denreich, and R Renz. "C/C-SiC Composites for Advanced Nicalon During Elevated Temperature Ageing": pp. 421-6 in Ceramic Composites: Advanced High-Te Structural materials Materials I 6T. M. Besmann D. P Stinton. R. A. Lowden and K. J. Probst."Near-Net Society Symposium Proceedings Vol. 365, Materials Research Society, Pittsburgh, Shape Fabrication by Forced-Flow Thermal Gradient CVI. "Ind. Cera., 20 [2] PA 199 K.P. Plucknett. H -T. Lin. D. N. Braski, and P. F. Becher. ""Environmental andell. D. H. Grande, and J. Jacobs, " Tensile Behavior of osite-Materials at Elevated Temperatures. "J. Eng. Gas Turbines roceedings of the 10th international Conference on Composite Materials, VoL. II Power,109267-73(1987) Characterist ion and Ceramic Matrix Composites. Woodhead Publishing Limited, dation Mechanisms in Nicalon/C/SiC Composites. "J. Am. Ceran. Soc. 77[10] of a Sic Fibre-Reinforced Magnesiu silicate Glass -Ceramic matrix omposite. " J Mater. Sci. 27 [11]3075-81(1992). asmin and P. Bowen, ""Fracture Behaviour of Cross-Ply Nicalon/ CAS-ll Glass-Ceramic Matrix Composite Laminate at Room and Elevated Tempera- 30A. M. Daniel. A. Martin Meizoso. K. P. Plucknett and D N. Braski"In- Composites,A33间9 M. Boussuge, ""Oxidizing Aging Eflects on SiC-SiC mposite, Ceram. Eng. Sci. Proc., 17[4]280-7(1996 M. w. Pharaoh. A. M. Daniel and M. H. Lewis. Stability of Interfaces in mal-Expansion Calcium Aluminosilicate Matrix Nicalon SiC Fiber Composites, "J. Mater. Sci. Composites. " J. Mater. Sci., 25 [11]4836-46(1990). Lo J-G. Duh. a k ell, s. Sutherland, A.M. Daniel R L. Cain, G West, D: M R ring and Microwave Dielectric Properties of Anorthite-Based Glass-Ceramics, "J. Am. Ceran.Soe,8592230-502002) ced Glass-ceramic Matrix Composite,. Microscop), 1713)251-63 terizat ion o Nicaon" and H - anical on (Ceramic Fibers y tonic sorbe mi. E.Heredia,J.MeNt邮WAEm“mxm1Am入图Rmm Mechanical Behavior of a Si-C-o Fiber-Reinforced Magnesium Aluminosili- S.Teoh,47435967(199 35A.G. Evans, M.Y. He, and J W. Hutchinson, "Ir bonding and Fiber ”J.Am. Ceran.Soe.,792375-8(199 Cracking in Brittle Matrix Composites. "J.Am Soc,72Ⅱl22300-3 IeP. F. Becher, H.-T. Lin, and K. L. More, "Lifetime-Applied Stress Response Betrand. C. Droillard. R. Pailler, X. Bourrat, and R. Naslain, " TEM ial Layer: Effects of Temperature(300 to 1150.C). " J. Am. Ceram Soc., 81[71 tructure of (PyC/ SiCh Multilayered Interphases in SiC/SiC Composites, "J. Eur. 1919-25(1998) Ceran. Soc., 20 [11-13(2000)
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