《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_j.1151-2916.2003.tb03463.x[1]

High-Temperature Water Vapor Effects ournal ⊥ m Cera. Sot,868]1272-8102003) High-Temperature Stability of Sic-Based Composites i High-Water-Vapor-Pressure Environments Karren L. More, Peter F. Tortorelli,*and Larry R.Walker Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6064 Naren Miriyala, Jeffrey R Price, and Mark van Solar Turbines, Inc., San Diego, California Microstructural characterization of boron-containing si fibers)and fibers having compositional improvements(such as reinforced Sic composites exposed at high temperature in Hi-NicalonTM) have also been characterized ,y In general, the high-water-vapor-pressure environments was used to deter- newer SiC fibers tend to be closer to stoichiometric SiC compo mine surface recession rates and to understand the controlling sition and have fewer impurities. Chemical-vapor-deposited degradation processes under these conditions. Results showed(CVD) boron nitride is currently the preferred fiber-matrix inter- that composite degradation was controlled by a series of facial component used by CFCC manufacturers since BN has eactions involving the formation of silica, boria, borosilicate improved oxidation resistance compared with carbon. At room glass, and gaseous products. Comparison of results (from temperature, BN fiber coatings allow for the fiber-matrix debond characterization of composites exposed at 1200.C and 1.5 atm ing characteristic of tough composites, but at elevated tempera- of H,o in a laboratory furnace and in the combustion zone of tures in water-containing environments, BN undergoes rapid a gas turbine) showed that these reactions were common to egradation that can negatively impact the long-term stability and both exposure conditions and, consequently, there was little usefulness of the composite. Although the thin bn fiber coating effect of gas velocity on degradation rates of boron-containing (0.4 Hm)represent a low percentage constituent of the total SiC/SiC composite materials. material of the bulk composite(<5 vol%), these coatings can have a significant effect on composite oxidation behavior he deleterious effect of water vapor on the oxidation rate and SiO, growth of Sic ceramics has been well documented. 4-17 CENT advances in the development and manufacture of More recently, work has been conducted at higher water-vapor iC-based continuous-fiber-reinforced ceramic-matrix com- posites(CFCCs) have led to the use of these materials in several and matrix)comprises >90 vol% of the composites, accelerated high-temperature applications. Improvements in the ceramic oxidation of these materials in high-H-O environments was ex fibers, interfacial coatings, and matrix processing have resulted in ted to be similar to that observed for monolithic sic hanced resistance to oxidation at temperatures up to at least liners show that this is not the case 22 To further the use of these testing of these materials in gas turbine engine hot-section com- materials in high-temperature applications, the issue of under onents, there is relatively little understanding of the oxidation tanding composite degradation has become an extremely impor mechanisms for SiC-based CFCCs in high-pressure steam envi tant one. Several different composite processing routes have been nts typically experience used to produce SiC/BN/SiC CFCC materials for hot-section To develop a complete understanding of the oxidation behavior components in engines, thereby resulting in different matrix of Sic-based CFCCs in water-containing environments, the con- compositions and microstructures. As a result, it has become tributions of different bulk composite microstructures and individ mperative to fully evaluate the degradation mechanisms associ- ual composite constituents(fibers, fiber-matrix interface coatings, ated with the use of various Sic/BN/SiC CFCC materials in nd matrix) to overall CFCC stability must be understood. Silicon high-water-vapor-containing environments and to develop an un- carbide fibers typically comprise 20-40 vol% of the Sic-based derstanding of the primary issues associated with the long-term CFCCs. The microstructural stability and oxidation behavior of stability of several different commercially available composites of ceramic-grade(CG)Nicalon TM (Si-C-O)fibers have been evalu- this type. To this end, this paper describes recent work on the ated.5,As a result of poor strength retention of CG NicalonTMat characterization of several commercially available SiC/BN/SIC temperatures >1000oC, SiC fibers having improved thermal sta- opposites exposed to high water-va ressures in a h emperature bility via cost-effective processing methods(such as Tyranno SiC pressure laboratory-scale exposure facility. In ular. oxidation mechanisms for the various fCCS were aluated based on the composite processing method used and the arting composition Of primary importance was as the contributio E.J. Opila-contributing editor of bn and B-containing compounds to the overall rate of com- osite degradation in these high-H2O oxidizing environments Results obtained from laboratory exposures are compared with similarly processed materials exposed in long-term engine tests at the Symposi Materials at the131 Annual meet吗号 of The Minerals, Il. Experimental Procedure Secretary rsr ehe gy efciency da d rehe wable bnergy onf ct of ndus ras iech. Two different types of Sic/BN/SiC composites were evaluated Resources progra contract D0OR22725 with UT-Battel in this study. The first was a composite with continuous Hi- Nicalon M SiC fiber reinforcement, CVD BN interfacial coati 1272

High-Temperature Stability of SiC-Based Composites in High-Water-Vapor-Pressure Environments Karren L. More,* Peter F. Tortorelli,* and Larry R. Walker Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6064 Naren Miriyala, Jeffrey R. Price,* and Mark van Roode Solar Turbines, Inc., San Diego, California 92186 Microstructural characterization of boron-containing SiC￾reinforced SiC composites exposed at high temperature in high-water-vapor-pressure environments was used to deter￾mine surface recession rates and to understand the controlling degradation processes under these conditions. Results showed that composite degradation was controlled by a series of reactions involving the formation of silica, boria, borosilicate glass, and gaseous products. Comparison of results (from characterization of composites exposed at 1200°C and 1.5 atm of H2O in a laboratory furnace and in the combustion zone of a gas turbine) showed that these reactions were common to both exposure conditions and, consequently, there was little effect of gas velocity on degradation rates of boron-containing SiC/SiC composite materials. I. Introduction RECENT advances in the development and manufacture of SiC-based continuous-fiber-reinforced ceramic-matrix com￾posites (CFCCs) have led to the use of these materials in several high-temperature applications.1–4 Improvements in the ceramic fibers, interfacial coatings, and matrix processing have resulted in composites with relatively good mechanical properties and en￾hanced resistance to oxidation at temperatures up to at least 1200°C. However, even with the recent implementation and testing of these materials in gas turbine engine hot-section com￾ponents,1–4 there is relatively little understanding of the oxidation mechanisms for SiC-based CFCCs in high-pressure steam envi￾ronments typically experienced in a turbine combustor. To develop a complete understanding of the oxidation behavior of SiC-based CFCCs in water-containing environments, the con￾tributions of different bulk composite microstructures and individ￾ual composite constituents (fibers, fiber-matrix interface coatings, and matrix) to overall CFCC stability must be understood. Silicon carbide fibers typically comprise 20–40 vol% of the SiC-based CFCCs. The microstructural stability and oxidation behavior of ceramic-grade (CG) Nicalon™ (Si-C-O) fibers have been evalu￾ated.5,6 As a result of poor strength retention of CG Nicalon™ at temperatures
1000°C, SiC fibers having improved thermal sta￾bility via cost-effective processing methods (such as Tyranno SiC fibers)7 and fibers having compositional improvements (such as Hi-Nicalon™) have also been characterized.8,9 In general, the newer SiC fibers tend to be closer to stoichiometric SiC compo￾sition and have fewer impurities. Chemical-vapor-deposited (CVD) boron nitride is currently the preferred fiber-matrix inter￾facial component used by CFCC manufacturers since BN has improved oxidation resistance compared with carbon.10,11 At room temperature, BN fiber coatings allow for the fiber-matrix debond￾ing characteristic of tough composites, but at elevated tempera￾tures in water-containing environments, BN undergoes rapid degradation that can negatively impact the long-term stability and usefulness of the composite.12 Although the thin BN fiber coatings (0.4 m) represent a low percentage constituent of the total material of the bulk composite (5 vol%), these coatings can have a significant effect on composite oxidation behavior.13 The deleterious effect of water vapor on the oxidation rate and SiO2 growth of SiC ceramics has been well documented.14–17 More recently, work has been conducted at higher water-vapor pressures (
1 atm) more typical of those in stationary gas turbines.18–21 Based on these results and the fact that SiC (fibers and matrix) comprises
90 vol% of the composites, accelerated oxidation of these materials in high-H2O environments was ex￾pected to be similar to that observed for monolithic SiC.20,21 However, observations after field-testing of CFCC combustor liners show that this is not the case.22 To further the use of these materials in high-temperature applications, the issue of under￾standing composite degradation has become an extremely impor￾tant one. Several different composite processing routes have been used to produce SiC/BN/SiC CFCC materials for hot-section components in engines, thereby resulting in different matrix compositions and microstructures. As a result, it has become imperative to fully evaluate the degradation mechanisms associ￾ated with the use of various SiC/BN/SiC CFCC materials in high-water-vapor-containing environments and to develop an un￾derstanding of the primary issues associated with the long-term stability of several different commercially available composites of this type. To this end, this paper describes recent work on the characterization of several commercially available SiC/BN/SiC composites exposed to high water-vapor pressures in a high￾temperature, high-pressure laboratory-scale exposure facility. In particular, oxidation mechanisms for the various CFCCs were evaluated based on the composite processing method used and the starting composition. Of primary importance was the contribution of BN and B-containing compounds to the overall rate of com￾posite degradation in these high-H2O oxidizing environments. Results obtained from laboratory exposures are compared with similarly processed materials exposed in long-term engine tests. II. Experimental Procedure Two different types of SiC/BN/SiC composites were evaluated in this study. The first was a composite with continuous Hi￾Nicalon™ SiC fiber reinforcement, CVD BN interfacial coatings, E. J. Opila—contributing editor Manuscript No. 186662. Received September 13, 2002; approved April 2, 2003. Presented at the Symposium on Water Vapor Effects on Oxidation of High￾Temperature Materials at the 131st Annual Meeting of The Minerals, Metals & Materials Society (TMS), Seattle, WA, February 18–20, 2002. This research was sponsored by the U.S. Department of Energy, Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Industrial Tech￾nologies, as part of the Distributed Energy Resources Program under contract DE-AC05-00OR22725 with UT-Battelle LLC. *Member, American Ceramic Society. High-Temperature Water Vapor Effects J. Am. Ceram. Soc., 86 [8] 1272–81 (2003) 1272 journal

High-Temperature Stability of sic-Based Composites in High-Water-Vapor-Pressure Emvironmer 273 trix processed using the isothermal chemical vapo II. Results ation(CVi) technique at GE Power Systems Composites, (Newark, DE).The second type of SiC/BN/SIC composite The two SiC/SiC composites compared in this study had very material also contained Hi-Nicalon TM fibers and CVd BN inter- different starting microstructures. The CVI composites were char facial coatings, but had a SiC matrix fabricated using both a CVI acterized by a two-dimensional woven fiber structure with a(os processing step followed by matrix densification via a silicon fiber lay-up, as shown in the relatively low magnification BSI melt-infiltration(Mi) process similar to that described in Ref, 24 image in Fig. 1(a). In addition to the cvd bn interfacial coating The MI composites were manufactured by either Goodrich Cor sed for the laboratory-exposed coupons, a B-containing particu (Santa Fe Springs, CA) or GE Power Systems Composites, LLC filler was used within and between the fiber tows to minimize the siC matrix infiltration time The fiber-coating structure within each type of CFCC were exposed in a high-temperature, high- the fiber tows for the CVI SiC/BN/SiC composite is shown in Fig pressure exposure facility (referred to in text as the ORnl 1(b). The BN-coated fibrous preforms were densified with Sic by furnace), which has been described in detail elsewhere. 25 No the CVI process. The resulting CVI composite was%95% protective SiC seal coatings were applied to any surfaces of the omposite coupons exposed in the laboratory furnace To simulate A similar fiber lay-up was used for the MI SiC/BN/SIC the typical water-vapor pressure and maximum temperature of a composite, as shown in Fig. 2(a). The MI SiC/BN/SiC used a olar Turbines Centaur 50s engine combustor environment(use uniformly thin, continuous CVD BN coating, -0.4 um thick for long-term engine tests, as described below ), all laboratory round the fibers. To protect the BN-coated fibrous preform during exposures were conducted at 1200C in a slowly flowing(3 MI processing, a-1-2 um CVI SiC layer was applied around the cm/min) gas mixture of air+ 15%H,O at 10 atm. Each exposure fiber tows(see Fig. 2(b). This minimal CVI treatment left a fairly period was 500 h, after which the specimens were carefully pen structure for subsequent molten silicon MI processing while moved and selectively cut for microstructural analysis. A ill protecting the fibers and fiber coatings. After processing, the 20.3-0.4 cm thick cross section was cut from the end of each MI matrix was composed primarily of Si SiC and the composite coupon.The cut section was mounted (with cut face down)in was near full density Composite density differences between the epoxy and polished. After cutting, the same composite coupons two composites are evident when the bulk microstructure of the were then placed back in the furnace for additional exposures for times up to 3000 h in 500 h increments and another cross section was cut from each coupon after each additional 500 h e ure this way, a total of six cross sections (cut after 500, 1000, 1500 2000, 2500, and 3000 h) were cut from each exposed CVI and mI composite coupon for microstructural evaluation In addition to the laboratory furnace exposures, engine tests -eric Fiber tow vere conducted using Sic/BN/SiC CFCC combustor liners ma ufactured using both CVI-and Ml-processed composites in a Solar Turbines(San Diego, CA) SoLONOX Centaur 50S engine located at the Chevron engine test site in Bakersfield. CA. A full-scale set of cfCc combustor liners consisted of a 33 cm diameter inner liner(either a SiC/BN/SiC MI liner produced by Goodrich Corp Santa Fe Springs, CA, or a CVI liner produced by GE Power Systems Composites, LLC, Newark, DE)and a 76 cm diameter SiC/C/SiC CVI outer liner(manufactured by GE Power Systems Composites, LLC, Newark, DE). This paper includes results for the engine exposure of two inner liners from two separate engine tests the mi inner liner ran in an engine for 2758 h and the cvi inner combustor liner was exposed for 2266 h. The gas surface for each liner was coated with a protective CVD SiC seal coat(-200 and 215 um for the MI and CVI liners, respectively The engine-exposed CFCC combustor liners were exposed to a maximum temperature of "1200.C (in the hot spots associated with fuel impingement areas) and a gas velocity of 30 m/s Microstructural examinations were conducted on polished cross sections prepared after each of the six exposures (up to 3000 h in SiC(matrix) 500 h increments) conducted in the ornl furnace to measure BN fiber oxide product thickness and surface recession, and to asses coating ubsurface oxidation-induced damage to the different CCCs (specimen mass and dimensional changes were monitored but were not reliable quantitative indicators of composite degradation due to friability of oxide products and the nature of some ,6,c constituent interactions. ) Composite surface recession was deter lined microstructurally by measuring the thickness of composite Sic (fiber) material below the surface that was not affected by oxidation good material) and subtracting this amount from the as-processed composite coupon thickness. Results from these examinations were com 20 um t sed combustor liners. Aft-to-fore liner cross sections were en from both heavily damaged areas and areas that appeared damaged from the inner and outer liners. Electron probe microanalysis(EPMA) and backscatte were the analysis techni sections before and after exposure in the d to examin ine CFCC cross Fig.1.micr ure of CVI SiC/BN/SiC: (a) Bse image of bulk mace or Sol microstructure showing fiber lay-up and CVI SiC matrix and(b) Turbines Centaur 50S engine. These examinations were conducte magnification image showing the Bn fiber coating and boron-containing using a JEOL 733 Superprobe (JEOL USA, Inc, Peabody, MA) filler within fiber tows

and a SiC matrix processed using the isothermal chemical vapor infiltration (CVI) technique at GE Power Systems Composites, LLC (Newark, DE).23 The second type of SiC/BN/SiC composite material also contained Hi-Nicalon™ fibers and CVD BN inter￾facial coatings, but had a SiC matrix fabricated using both a CVI processing step followed by matrix densification via a silicon melt-infiltration (MI) process similar to that described in Ref. 24. The MI composites were manufactured by either Goodrich Corp. (Santa Fe Springs, CA) or GE Power Systems Composites, LLC. As-processed coupons (typically 2.5 cm  5.0 cm  0.3 cm) of each type of CFCC were exposed in a high-temperature, high￾pressure exposure facility (referred to in text as the ORNL furnace), which has been described in detail elsewhere.25 No protective SiC seal coatings were applied to any surfaces of the composite coupons exposed in the laboratory furnace. To simulate the typical water-vapor pressure and maximum temperature of a Solar Turbines Centaur 50S engine combustor environment (used for long-term engine tests, as described below), all laboratory exposures were conducted at 1200°C in a slowly flowing (3 cm/min) gas mixture of air  15% H2O at 10 atm. Each exposure period was 500 h, after which the specimens were carefully removed and selectively cut for microstructural analysis. A 0.30.4 cm thick cross section was cut from the end of each coupon. The cut section was mounted (with cut face down) in epoxy and polished. After cutting, the same composite coupons were then placed back in the furnace for additional exposures for times up to 3000 h in 500 h increments and another cross section was cut from each coupon after each additional 500 h exposure. In this way, a total of six cross sections (cut after 500, 1000, 1500, 2000, 2500, and 3000 h) were cut from each exposed CVI and MI composite coupon for microstructural evaluation. In addition to the laboratory furnace exposures, engine tests were conducted using SiC/BN/SiC CFCC combustor liners man￾ufactured using both CVI- and MI-processed composites in a Solar Turbines (San Diego, CA) SoLoNOx Centaur 50S engine located at the Chevron engine test site in Bakersfield, CA. A full-scale set of CFCC combustor liners consisted of a 33 cm diameter inner liner (either a SiC/BN/SiC MI liner produced by Goodrich Corp., Santa Fe Springs, CA, or a CVI liner produced by GE Power Systems Composites, LLC, Newark, DE) and a 76 cm diameter SiC/C/SiC CVI outer liner (manufactured by GE Power Systems Composites, LLC, Newark, DE). This paper includes results for the engine exposure of two inner liners from two separate engine tests; the MI inner liner ran in an engine for 2758 h and the CVI inner combustor liner was exposed for 2266 h.1 The gas-path surface for each liner was coated with a protective CVD SiC seal coat (200 and 215 m for the MI and CVI liners, respectively). The engine-exposed CFCC combustor liners were exposed to a maximum temperature of 1200°C (in the hot spots associated with fuel impingement areas) and a gas velocity of 30 m/s. Microstructural examinations were conducted on polished cross sections prepared after each of the six exposures (up to 3000 h in 500 h increments) conducted in the ORNL furnace to measure oxide product thickness and surface recession, and to assess subsurface oxidation-induced damage to the different CFCCs (specimen mass and dimensional changes were monitored but were not reliable quantitative indicators of composite degradation due to friability of oxide products and the nature of some constituent interactions.) Composite surface recession was deter￾mined microstructurally by measuring the thickness of composite material below the surface that was not affected by oxidation (good material) and subtracting this amount from the as-processed composite coupon thickness. Results from these examinations were compared with similar analyses performed on the engine￾exposed combustor liners. Aft-to-fore liner cross sections were taken from both heavily damaged areas and areas that appeared undamaged from the inner and outer liners. Electron probe microanalysis (EPMA) and backscatter electron (BSE) imaging were the primary analysis techniques used to examine CFCC cross sections before and after exposure in the ORNL furnace or Solar Turbines Centaur 50S engine. These examinations were conducted using a JEOL 733 Superprobe (JEOL USA, Inc., Peabody, MA). III. Results The two SiC/SiC composites compared in this study had very different starting microstructures. The CVI composites were char￾acterized by a two-dimensional woven fiber structure with a 0°/45° fiber lay-up, as shown in the relatively low magnification BSE image in Fig. 1(a). In addition to the CVD BN interfacial coating used for the laboratory-exposed coupons, a B-containing particu￾late filler was used within and between the fiber tows to minimize the SiC matrix infiltration time. The fiber-coating structure within the fiber tows for the CVI SiC/BN/SiC composite is shown in Fig. 1(b). The BN-coated fibrous preforms were densified with SiC by the CVI process. The resulting CVI composite was 85%–95% dense. A similar fiber lay-up was used for the MI SiC/BN/SiC composite, as shown in Fig. 2(a). The MI SiC/BN/SiC used a uniformly thin, continuous CVD BN coating, 0.4 m thick, around the fibers. To protect the BN-coated fibrous preform during MI processing, a 1–2 m CVI SiC layer was applied around the fiber tows (see Fig. 2(b)). This minimal CVI treatment left a fairly open structure for subsequent molten silicon MI processing while still protecting the fibers and fiber coatings. After processing, the MI matrix was composed primarily of Si  SiC and the composite was near full density. Composite density differences between the two composites are evident when the bulk microstructure of the Fig. 1. Microstructure of CVI SiC/BN/SiC: (a) BSE image of bulk microstructure showing fiber lay-up and CVI SiC matrix and (b) higher￾magnification image showing the BN fiber coating and boron-containing filler within fiber tows. August 2003 High-Temperature Stability of SiC-Based Composites in High-Water-Vapor-Pressure Environments 1273

1274 Journal of the American Ceramic Society-More et al Vol 86. No. 8 these constituents oxidizes at a different rate and different reaction products having very different product morphologies are formed An example of such is shown in Fig. 5. Once the CVI SiC layer around the fibers is breached. oxidation of the fiber tows is faster than that of the other constituents and occurs primarily via the Fiber tow dation along/through the bn interfaces. Figure 6 shows that the BN acts as a conduit for rapid transport of oxidant through the fiber tow and clearly oxidizes faster than the surrounding CVI SiC and Sic fibers. Within the tows, the fibers are completely consumed by the oxidized bn(reactions between boria, SiC, and silica.)The preferred BN oxidation path around individual fibers within a tow is shown clearly in Fig. 6. A fully consumed fiber tow serves as the path for the progression of the oxidation front further into the compos as shown in Fig. 7. Areas that are primarily Si SiC matrix will oxidize fairly rapidly but not as rapidly as th fiber tows since the tows contain the thin, but continuous, BN 100m A photograph of the SiC/BN/SiC MI inner liner after the 2758 h engine test is shown in Fig. 8. The "white damaged areas observed on the gas-path surface of the inner liner are regions of greater damage accumulation that corresponded directly with fuel impingement areas on the liner surface. These areas also corre- sponded with higher-temperature regions where the liner surface was subjected to temperatures 21200C Surface recession of the MI liner, measured microstructurally from several aft-to-fore cross sections, is summarized in Fig. 9 for three different areas visually BN fiber showing different degrees of surface damage. The three areas used coating to generate data for this plot are circled in Fig. 8. The line at "3000 um on the graph in Fig 9 represents the CVD SiC seal coat/MI composite interface. While the majority of the inner liner urface did not exhibit the extensive damage associated with the white regions(see Fig 8), the white areas on the liner surface were SiC fiber clearly characteristic of regions showing break-through in the thick CVD SIC seal coat due to accelerated sic recession in the combustor, albeit at a somewhat lower rate than that given by siC/S Robinson and Smialek for recession of cVd Sic (silica volatil- matrix 10 um ization by Si(OH) formation) in a high-pressure burner rig at 1200.C.The CVD SiC seal coat recession rate for combustor liners engine tested in a Solar Turbines Centaur 50S engine as 0 11 um/h at 1200C, which was determined directly from microstructural measurements of the present as well as previous Fig.2. Microstructure of MI SiC/BN/SiC:(a)BSE image of bulk engine-tested C VD SiC seal-coated liners. In the areas surround magnification SEM image of 0.4 Hm BN coatings around Hi-Nicalon M Sic seal coat was still intact but did show varying de fibers and thin Cvd SiC laver surrounding fiber tow urface recession. The underlying MI composite material in these less damaged regions did not exhibit any microstructural damage Within the white areas on the inner liner surface. damage to the VI and MI SiC/BN/SiC composites are compared; large, inter- underlying MI composite was evident. In less damaged areas connected matrix porosity is characteristic of the CVI composites, occasional break-through of the sic seal coat resulted in localized whereas MI composites are much denser ubsurface composite damage. This is shown in Fig. 10. Greater MI composite damage depths were observed when the CVD SiC seal coat was completely removed by volatilization to fully expose (I MI SiC/BNAiC large regions of the CFCC substrate. However, the MI composite The Mi composite coupons were exposed in the ornl urface damage was limited to the first or second fiber layer in the to 1.5 atm of H,o at 1200.C for a total of 3000 h during th white areas, even in the highly damaged regions. The typical MI icrostructural damage was extensive. Figure 3 shows a composite damage observed in large areas of Sic seal coat sion of images of the bulk MI composite recession and concurrent break-through is shown in Fig. ll. From the data presented in Fig silica scale formation as a function of exposure time. The lines 9, the maximum depth of microstructural damage observed on the drawn approximately parallel to the original coupon faces repre- inner liner was -300 um below the seal coat, which corresponded sent the total amount of surface recession(composite degradation with the depth of the second layer of fibers from the working that is, the depth to which the structure has been altered due to urface. These areas were associated with oxidation of many of the oxidation and associated constituent and product reactions(see on te constituents, including the si Sic matrix, the Cvd Sic around the fiber tows. and the bn coating and Sic fibers The surface of the mi sic/BN/sic before oxidation wa within the fiber tows. A SiO,+ borosilicate glass scale developed composed primarily of the Si SiC MI matrix, whether thinly on the surface(Fig. I1) covering a fiber tow(CVI SiC) or as large-scale matrix areas During the steady-state recession of pure CVD SiC in the between tows, as shown in Fig 4. The composite can be consid- high-gas-velocity, high-H2O-pressure environment typical of a ered as a complex layered structure of the different constituents turbine engine combustor 1O, volatilizes at the same rat Ite see Fig 4) which are oxidized as the oxidation front progresses SiC oxidizes, and thus, thick surface scale does not form these through the composi te structure the mi matrix(Si SiC results are consistent with observations made on the remaining sic normally oxidizes first, followed by the CvI SiC layer around the eal coat on the gas-path surface of the inner liner. A thick, porous fiber tows, which is subsequently followed by oxidation within the scale was not observed; however, a thin vitreous SiO, layer was ber tows(composed of Bn coatings around SiC fibers ). Each of present on the oxidized SiC surface 20, 21 In the case of the MI

CVI and MI SiC/BN/SiC composites are compared; large, inter￾connected matrix porosity is characteristic of the CVI composites, whereas MI composites are much denser. (1) MI SiC/BN/SiC The MI composite coupons were exposed in the ORNL furnace to 1.5 atm of H2O at 1200°C for a total of 3000 h. During this time, microstructural damage was extensive. Figure 3 shows a succes￾sion of images of the bulk MI composite recession and concurrent silica scale formation as a function of exposure time. The lines drawn approximately parallel to the original coupon faces repre￾sent the total amount of surface recession (composite degradation), that is, the depth to which the structure has been altered due to oxidation and associated constituent and product reactions (see below). The surface of the MI SiC/BN/SiC before oxidation was composed primarily of the Si  SiC MI matrix, whether thinly covering a fiber tow (CVI SiC) or as large-scale matrix areas between tows, as shown in Fig. 4. The composite can be consid￾ered as a complex layered structure of the different constituents (see Fig. 4) which are oxidized as the oxidation front progresses through the composite structure. Thus, the MI matrix (Si  SiC) normally oxidizes first, followed by the CVI SiC layer around the fiber tows, which is subsequently followed by oxidation within the fiber tows (composed of BN coatings around SiC fibers). Each of these constituents oxidizes at a different rate, and different reaction products having very different product morphologies are formed. An example of such is shown in Fig. 5. Once the CVI SiC layer around the fibers is breached, oxidation of the fiber tows is faster than that of the other constituents and occurs primarily via the oxidation along/through the BN interfaces. Figure 6 shows that the BN acts as a conduit for rapid transport of oxidant through the fiber tow and clearly oxidizes faster than the surrounding CVI SiC and SiC fibers. Within the tows, the fibers are completely consumed by the oxidized BN (reactions between boria, SiC, and silica.) The preferred BN oxidation path around individual fibers within a tow is shown clearly in Fig. 6. A fully consumed fiber tow serves as the path for the progression of the oxidation front further into the composite, as shown in Fig. 7. Areas that are primarily Si  SiC matrix will oxidize fairly rapidly,20 but not as rapidly as the fiber tows since the tows contain the thin, but continuous, BN.12 A photograph of the SiC/BN/SiC MI inner liner after the 2758 h engine test is shown in Fig. 8. The “white” damaged areas observed on the gas-path surface of the inner liner are regions of greater damage accumulation that corresponded directly with fuel impingement areas on the liner surface. These areas also corre￾sponded with higher-temperature regions where the liner surface was subjected to temperatures
1200°C. Surface recession of the MI liner, measured microstructurally from several aft-to-fore cross sections, is summarized in Fig. 9 for three different areas visually showing different degrees of surface damage. The three areas used to generate data for this plot are circled in Fig. 8. The line at 3000 m on the graph in Fig. 9 represents the CVD SiC seal coat/MI composite interface. While the majority of the inner liner surface did not exhibit the extensive damage associated with the white regions (see Fig. 8), the white areas on the liner surface were clearly characteristic of regions showing break-through in the thick CVD SiC seal coat due to accelerated SiC recession in the combustor, albeit at a somewhat lower rate than that given by Robinson and Smialek for recession of CVD SiC (silica volatil￾ization by Si(OH)4 formation) in a high-pressure burner rig at 1200°C.18 The CVD SiC seal coat recession rate for combustor liners engine tested in a Solar Turbines Centaur 50S engine was 0.11 m/h at 1200°C, which was determined directly from microstructural measurements of the present as well as previously engine-tested CVD SiC seal-coated liners.26 In the areas surround￾ing the white areas on the inner liner gas-path surface, the CVD SiC seal coat was still intact but did show varying degrees of surface recession. The underlying MI composite material in these less damaged regions did not exhibit any microstructural damage. Within the white areas on the inner liner surface, damage to the underlying MI composite was evident. In less damaged areas, occasional break-through of the SiC seal coat resulted in localized subsurface composite damage. This is shown in Fig. 10. Greater MI composite damage depths were observed when the CVD SiC seal coat was completely removed by volatilization to fully expose large regions of the CFCC substrate. However, the MI composite surface damage was limited to the first or second fiber layer in the white areas, even in the highly damaged regions. The typical MI composite damage observed in large areas of SiC seal coat break-through is shown in Fig. 11. From the data presented in Fig. 9, the maximum depth of microstructural damage observed on the inner liner was 300 m below the seal coat, which corresponded with the depth of the second layer of fibers from the working surface. These areas were associated with oxidation of many of the composite constituents, including the Si  SiC matrix, the CVD SiC around the fiber tows, and the BN coating and SiC fibers within the fiber tows. A SiO2  borosilicate glass scale developed on the surface (Fig. 11). During the steady-state recession of pure CVD SiC in the high-gas-velocity, high-H2O-pressure environment typical of a turbine engine combustor,16,18 SiO2 volatilizes at the same rate as SiC oxidizes, and thus, a thick surface scale does not form. These results are consistent with observations made on the remaining SiC seal coat on the gas-path surface of the inner liner. A thick, porous scale was not observed; however, a thin vitreous SiO2 layer was present on the oxidized SiC surface.20,21 In the case of the MI Fig. 2. Microstructure of MI SiC/BN/SiC: (a) BSE image of bulk microstructure showing fiber lay-up and dense MI matrix and (b) higher￾magnification SEM image of 0.4 m BN coatings around Hi-Nicalon™ fibers and thin CVD SiC layer surrounding fiber tows. 1274 Journal of the American Ceramic Society—More et al. Vol. 86, No. 8

August 2003 High-Temperature Stability of Sic-Based Composites in High-Water-Vapor-Pressure Emironments 275 As processed 1500h 2500h lica scale mm Fig 3. Cross-section BSE images of MI SiC/BN/SiC after exposure in a high-pressure furnace at 1200C and 1.5 atm of H2O for 0, 1500, and 2500 h showing damage to composite and Sio, scale development mposite material with bn as a constituent within the fiber tows, (2) CV SiC/BNAiC the borosilicate surface product (primarily a low-B-containing The CVI SiC/BNSiC composite was exposed in the ORNL orosilicate glass) was relatively stable in the combustor environ- furnace for a total of 1500 h at 1200 C and 1.5 atm of H,O before t volatility of boria based or uilibrium thermochemical considerations. 2 These observations extensive composite damage prevented reintroduction into the expo- sure facility. Figure 13 is an image of the coupon cross section after indicate that there is a low boria activity in the glass and/or the rate exposure for 1500 h at 1200C. Note the very thick, porous scale on of formation of the borosilicate glass was significantly greater than the volatilization rate the surface of the CVI SiC/BN/SiC. The as-oxidized morphology was quite different from that observed on the MI composite(compare Figure 12 is a higher-magnification image of the oxidation- Figs. 5 and 13). Fiber tows near the exposed surface were oxidized in associated with the various reactions are very simllar to phologies a manner similar to that observed for the MI SiC/BN/SiC, albeit more nduced microstructura bserved for the mi composite exposed at 1200.C to 1.5 atm of severely and rapidly, as shown in Fig. 14 (compare with Fig. 7)as a water vapor in the laboratory furnace(compare Figs result of substantial borosilicate glass formation in these areas The both cases. d was significantly accelerated because of the fiber tows oxidized more rapidly in the C vi composite compared with esence of the BN fiber coatings, which acted as preferred path the MI composite because of the presence of significant amounts of for the transport of oxidants well into the fiber tows B-containing phases(filler) within the fiber tows as well as greater amounts of BN. The amount of BN and B-containing phases present within the CVi composite fiber tows was >10 times that in the 0.4 um BN fiber coatings in the MI composite. In addition to the extensive surface oxidation, subsurface reaction within the large, interconnected porosity was observed in the CVI SiC/BN/SiC. During the first 500 h exposure, silica scales formed on the Cvi SiC matrix surrounding the subsurface poros- ity, as shown in Fig. 15(a). With additional exposure, the CVI SiC matrix around the subsurface porosity was breached. As shown Fig. 15(b), rapid degradation of the B-containing filler material and the cvd bn fiber coatings within the tows then occurred a manner similar to that described previously for the mi compos ite. However, because of the greater amount of boron-containing phases in the CVI composite, subsurface fiber tows were con sumed during this process and pools of glass were formed within these areas well below the exposed surface. Thus, for the CVI SIC/BN/SIC, oxidation was not simply limited to the exposed surface of the composite. Rapid CVI composite degradation was compounded by subsurface accumulation of amage due to greater amounts of boron(B-co filler and Fig4. Cross-section BSE image of as-processed MI SiC/BN/SiC show- bn fiber coatings) as well as a network connected ing complex layered structure of composite constituents

composite material with BN as a constituent within the fiber tows, the borosilicate surface product (primarily a low-B-containing borosilicate glass) was relatively stable in the combustor environ￾ment despite the expected significant volatility of boria based on equilibrium thermochemical considerations.12 These observations indicate that there is a low boria activity in the glass and/or the rate of formation of the borosilicate glass was significantly greater than the volatilization rate. Figure 12 is a higher-magnification image of the oxidation￾induced microstructural damage in the MI liner. The morphologies associated with the various reactions are very similar to what was observed for the MI composite exposed at 1200°C to 1.5 atm of water vapor in the laboratory furnace (compare Figs. 6 and 12). In both cases, damage was significantly accelerated because of the presence of the BN fiber coatings, which acted as preferred paths for the transport of oxidants well into the fiber tows. (2) CVI SiC/BN/SiC The CVI SiC/BN/SiC composite was exposed in the ORNL furnace for a total of 1500 h at 1200°C and 1.5 atm of H2O before extensive composite damage prevented reintroduction into the expo￾sure facility. Figure 13 is an image of the coupon cross section after exposure for 1500 h at 1200°C. Note the very thick, porous scale on the surface of the CVI SiC/BN/SiC. The as-oxidized morphology was quite different from that observed on the MI composite (compare Figs. 5 and 13). Fiber tows near the exposed surface were oxidized in a manner similar to that observed for the MI SiC/BN/SiC, albeit more severely and rapidly, as shown in Fig. 14 (compare with Fig. 7) as a result of substantial borosilicate glass formation in these areas. The fiber tows oxidized more rapidly in the CVI composite compared with the MI composite because of the presence of significant amounts of B-containing phases (filler) within the fiber tows as well as greater amounts of BN. The amount of BN and B-containing phases present within the CVI composite fiber tows was
10 times that in the 0.4 m BN fiber coatings in the MI composite. In addition to the extensive surface oxidation, subsurface reaction within the large, interconnected porosity was observed in the CVI SiC/BN/SiC. During the first 500 h exposure, silica scales formed on the CVI SiC matrix surrounding the subsurface poros￾ity, as shown in Fig. 15(a). With additional exposure, the CVI SiC matrix around the subsurface porosity was breached. As shown in Fig. 15(b), rapid degradation of the B-containing filler material and the CVD BN fiber coatings within the tows then occurred in a manner similar to that described previously for the MI compos￾ite. However, because of the greater amount of boron-containing phases in the CVI composite, subsurface fiber tows were con￾sumed during this process and pools of glass were formed within these areas well below the exposed surface. Thus, for the CVI SiC/BN/SiC, oxidation was not simply limited to the exposed surface of the composite. Rapid CVI composite degradation was compounded by subsurface accumulation of oxidation-induced damage due to greater amounts of boron (B-containing filler and BN fiber coatings) as well as a network of interconnected, subsurface porosity. Fig. 3. Cross-section BSE images of MI SiC/BN/SiC after exposure in a high-pressure furnace at 1200°C and 1.5 atm of H2O for 0, 1500, and 2500 h showing damage to composite and SiO2 scale development. Fig. 4. Cross-section BSE image of as-processed MI SiC/BN/SiC show￾ing complex layered structure of composite constituents. August 2003 High-Temperature Stability of SiC-Based Composites in High-Water-Vapor-Pressure Environments 1275

Journal of the American Ceramic Society-More et al Vol 86. No. 8 Scale due to matrix (Si+SiC)oxidation ale due to Si oxidation (SiC+BN) oxidation 100 Fig. 5. Cross-section BSE image of MI SiC/BN/SiC after exposure in high-pressure furnace for 500 h at 1200C and 1.5 atm of H,O showing different scale morphologies that form on the different composite constituents Areas of white surface damage were also observed on the cvi section and was associated with subsurface CVI composite dam- SiC/BN/SIC inner liner that was exposed for 2266 h in the gas age to a depth of -250 um urbine. Again, these regions were indicative of varying tempera- The nature of the CVI composite damage observed in areas tures across the liner's gas-path surface and a localized increase in where the seal coat had completely recessed ( via silica volatiliza- liner/composite damage associated with the fuel impingement tion) was similar to that observed for the MI SiC/BN/SiC. The the engine test, the temperature associated with the white surface the oxidation of CVD SiC surrounding subsurface pores in the CVI damage was assumed to be -1200oC based on temperature composite and(2)more extensive damage to the underlying CVi measurements made during the early stages of a previous engine composite in areas of Sic seal coat break-through. Figure 18 test. However, the temperature in some of the areas showing shows the subsurface oxidation within pores well below the maximum damage may be slightly higher than 1200.. A photo gas-path surface. The subsurface porosity within the bulk CVI aph of the gas-path surface of the CvI inner liner after engine composite(below gas-path surface) may be considered a slow testing is shown in Fig. 16(fuel impingement areas are designate gas-flow area. For this reason, oxidation of the SiC surfaces of the by the 12 small arrows). Many different areas of this liner were pores below the surface resulted in the formation and retention of mpared in cross section, but the primary areas used for compar a crystalline Sio,(cristobalite, as determined by transmission ison were from regions showing the maximum amount of white electron microscopy) scale(designated area I in Fig. 18)which urface damage(area I in Fig. 16) and a region showing minimal was structurally similar to the sio, formed during exposure of surface damage(area 3 in Fig. 16) where the CVD SiC seal coat CVD SiC in the ORNL furnace 2,2 Composite damage just below appeared to still be intact. The amount of surface recession areas of CVD SiC seal coat break-through(designated area 2 in easured microstructurally from these two areas is summarized in Fig. 18)was more extensive than in similar areas observed for the 16. Note that the black line at liner thickness =0 in the plot MI composite(see Fig 10). Considering that the CVI SiC/BN/SIC in Fig 17 corresponds to the CVD SiC seal coat/CVI SiC/BN/SiC inner liner had a slightly thicker as-processed CVD SiC seal coat interface. Area 3(minimum damage area) showed c100 (215 um for CVI VS -20 forM) and that the cⅥ recession of the SiC seal coat with no break-through, whereas area composite liner was exposed in an engine for 300 h less, the I exhibited seal coat break-through in the middle of the aft-to-fore oxidation behavior and total amount of damage accumulation observed for the CVI composite material were significantly worse B-Si-O-C B-Si-O ○ 100um

Areas of white surface damage were also observed on the CVI SiC/BN/SiC inner liner that was exposed for 2266 h in the gas turbine. Again, these regions were indicative of varying tempera￾tures across the liner’s gas-path surface and a localized increase in liner/composite damage associated with the fuel impingement areas. While no accurate temperature data were recorded during the engine test, the temperature associated with the white surface damage was assumed to be 1200°C based on temperature measurements made during the early stages of a previous engine test. However, the temperature in some of the areas showing maximum damage may be slightly higher than 1200°C. A photo￾graph of the gas-path surface of the CVI inner liner after engine testing is shown in Fig. 16 (fuel impingement areas are designated by the 12 small arrows). Many different areas of this liner were compared in cross section, but the primary areas used for compar￾ison were from regions showing the maximum amount of white surface damage (area 1 in Fig. 16) and a region showing minimal surface damage (area 3 in Fig. 16) where the CVD SiC seal coat appeared to still be intact. The amount of surface recession measured microstructurally from these two areas is summarized in Fig. 16. Note that the black line at liner thickness  0 in the plot in Fig. 17 corresponds to the CVD SiC seal coat/CVI SiC/BN/SiC interface. Area 3 (minimum damage area) showed 100 m recession of the SiC seal coat with no break-through, whereas area 1 exhibited seal coat break-through in the middle of the aft-to-fore section and was associated with subsurface CVI composite dam￾age to a depth of 250 m. The nature of the CVI composite damage observed in areas where the seal coat had completely recessed (via silica volatiliza￾tion) was similar to that observed for the MI SiC/BN/SiC. The major differences between the CVI and MI composites were (1) the oxidation of CVD SiC surrounding subsurface pores in the CVI composite and (2) more extensive damage to the underlying CVI composite in areas of SiC seal coat break-through. Figure 18 shows the subsurface oxidation within pores well below the gas-path surface. The subsurface porosity within the bulk CVI composite (below gas-path surface) may be considered a slow gas-flow area. For this reason, oxidation of the SiC surfaces of the pores below the surface resulted in the formation and retention of a crystalline SiO2 (cristobalite, as determined by transmission electron microscopy) scale (designated area 1 in Fig. 18) which was structurally similar to the SiO2 formed during exposure of CVD SiC in the ORNL furnace.20,21 Composite damage just below areas of CVD SiC seal coat break-through (designated area 2 in Fig. 18) was more extensive than in similar areas observed for the MI composite (see Fig. 10). Considering that the CVI SiC/BN/SiC inner liner had a slightly thicker as-processed CVD SiC seal coat (215 m for CVI vs 200 m for MI) and that the CVI composite liner was exposed in an engine for 300 h less, the oxidation behavior and total amount of damage accumulation observed for the CVI composite material were significantly worse Fig. 5. Cross-section BSE image of MI SiC/BN/SiC after exposure in high-pressure furnace for 500 h at 1200°C and 1.5 atm of H2O showing different scale morphologies that form on the different composite constituents. Fig. 6. BSE image of MI SiC/BN/SiC showing the accelerated oxidation along/through BN interfacial coatings during exposure in high-pressure furnace for 2000 h at 1200°C and 1.5 atm of H2O. Fig. 7. BSE cross-section image of the surface of MI SiC/BN/SiC composite showing several consumed fiber tows after exposure for 2000 h at 1200°C and 1.5 atm of H2O that oxidized and formed borosilicate glass. 1276 Journal of the American Ceramic Society—More et al. Vol. 86, No. 8