Am, Cern.So.852703-1002002) ournal Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composites Michael K. Cinibulk, *Triplicane A. Parthasarathy,,f Kristin A. Keller, f and Tai-lI Mah*f Air Force Research Laboratory, Materials and Manufacturing Directorate, Wright-Patterson Air Force Base, Ohio 45433-7817 A porous oxide fiber coating was investigated for Nextel 610 mullite, zircon, and rare-earth aluminates(garnets and magneto- fibers in an alumina matrix. Polymeric-solution-derived yt lumbites, are good candidate porous fiber coatings. Porous trium aluminum garnet (YAG, Y3Al5O12) with a fugitive coatings based on alumina, mixed zirconia and silica, mullite carbon phase was used to develop the porous fiber coating lanthanum hexaluminate have been applied to small-d Ultimate tensile strengths of tows and minicomposites follow alumina and alumina-mullite fibers in tows. Porous ing heat treatments in argon and/or air were used to evaluate and lanthanum hexaluminate coatings have been applied the effect of the porous fiber coating. The porous YAG fiber phire monofilament coatings did not reduce the strength of the tows when heated in If an energy-based crack-deflection criterion is considered argon,and they degraded tow strength by only -20% after within a porous coating,2 a minimum pore volume fraction of air at 1200C for 100 h. minico posites containing +0.3 is needed to reduce the fracture energy of polycrystalline porous YAG-coated fibers were nearly twice as strong as those alumina to 25% of that of the dense material. 5, 6 For crack containing uncoated fibers. However after heating at 1200oC deflection at a coating/fiber interface, the required pore fraction for 100 h, the porous YAG coatings densified to >90%, at may be lower because of the higher elastic modulus of the fiber which point they were ineffective at protecting the fibers, compared with that of the porous coating. This paper focuses on a resulting in identical strengths for minicomposites with and porous yttrium aluminum garnet (YAG, Y3 Al, O12) fiber coating without a fiber coating that does not degrade fiber strength in air. Fiber degradation has been a major concern with most oxide coatings. Aspects of precursor synthesis, fiber coating, coating-microstructure develop- L. Introduction T HAS been well-documented that increasing the porosity of a ceramic decreases its mechanical properties. -This concept has been used to weaken matrixes of oxide composites to the extent Il. Experimental Procedure that vs 10-14 coating is not needed to protect the fibers from matrix ()) Precursor and Coating Synthesis YAG is the most creep-resistant oxide known. 26-28YAG trength of the composite; its function is mainly to hold the fibers synthesis by conventional solid-state reaction of the element in place and prevent matrix cracks from developing enough energy oxides requires temperatures >1600oC for extended period to be able to penetrate the fibers. However, in many applications Commercially available alumina fibers, such as Nextel 610(3M where hermeticity, compressive and/or transverse strength, or wear Corp, Minneapolis, MN), cannot be exposed to temperatures 1200C for more than very short periods without degrading not adequate. For such applications, a dense matrix is preferred, strength via grain growth. This limits the processin ng window fo to provide crack deflection at the fiber/matrix interfacial region. the pplication of a coating and subsequent matrix processing of and, therefore, some type of fiber coating is likely to be required the ≤1200°C. or a dense matrix composite, the concept of porosity has been Recently, the synthesis of phase-pure YAG at temperatures applied sparingly to weaken the fiber/matrix interface. Porous 2800%C within I h has been reported. The polymeric precursor coatings have been applied to large-diameter monofilaments where crystallizes directly into the garnet structure from an amorphous discrete carbon or polymer particles are used as a fugitive phase powder starting at 600C when heated in an oxidizing atmosphere that is later burned out to create porosity. 5-8 The polymer In argon, peaks of primarily hexagonal YAlO, are present along particles are usually >100 nm in size and are of the appropriate with traces of YAG and amorphous alumina at temperatures of scale to provide a fugitive phase for fiber coatings that are >l um 700-900 C At 1000oC within I h, YAlO )3 reacts with the residual in thickness alumina to form YAG, which is then the only crystalline phase The requirement for a much thinner coating on the 10-um- present. The precursor is well-suited for this work because an diameter filaments in commercially available tows and cloths has intimate mixture of nanosized YAG and carbon is obtained when led to the development of porous coatings derived from intimate it is heated in an inert atmosphere at temperatures below which mixtures of oxide and carbon particles on the order of 10 nm in strength degradation of Nextel 610 fiber occurs. Details of the diameter. 9-2 With such a fine-scale microstructure, the ability of mixed-metal citric acid/ethylene glycol/ethanol solution precursor the porous coating to resist sintering and/or coarsening when synthesis have been discussed elsewhere. 3 In the present study onstrained between fiber and matrix during exposure to elevated two solutions are prepared to obtain coatings with carbon contents temperatures is paramount. Oxides with low self-diffusion, such as (derived from thermal decomposition of the organic components) of 30 and 50 vol% on a solids basis. The carbon acts as a fugitive phase to establish and maintain porosity during subsequent matrix infiltration and densification; only after matrix processing is the F W. Zok--contributing editor carbon removed by heating in air (2) Fiber Coating ly30,20 xtel 610 fiber tow was first pass 33615-96C-5258. furnace of a fiber coating apparatus, shown in Fig. 1, at 900oC in filiated with UES. In OH45432 air to remove the sizing. The hot zone of the in-line furnace was 2703
Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composites Michael K. Cinibulk,* Triplicane A. Parthasarathy,* ,† Kristin A. Keller,* ,† and Tai-Il Mah* ,† Air Force Research Laboratory, Materials and Manufacturing Directorate, Wright–Patterson Air Force Base, Ohio 45433-7817 A porous oxide fiber coating was investigated for Nextel™ 610 fibers in an alumina matrix. Polymeric-solution-derived yttrium aluminum garnet (YAG, Y3Al5O12) with a fugitive carbon phase was used to develop the porous fiber coating. Ultimate tensile strengths of tows and minicomposites following heat treatments in argon and/or air were used to evaluate the effect of the porous fiber coating. The porous YAG fiber coatings did not reduce the strength of the tows when heated in argon, and they degraded tow strength by only 20% after heating in air at 1200°C for 100 h. Minicomposites containing porous YAG-coated fibers were nearly twice as strong as those containing uncoated fibers. However, after heating at 1200°C for 100 h, the porous YAG coatings densified to >90%, at which point they were ineffective at protecting the fibers, resulting in identical strengths for minicomposites with and without a fiber coating. I. Introduction I T HAS been well-documented that increasing the porosity of a ceramic decreases its mechanical properties.1–9 This concept has been used to weaken matrixes of oxide composites to the extent that a fiber coating is not needed to protect the fibers from matrix cracks.10–14 In such composites, the matrix contributes little to the strength of the composite; its function is mainly to hold the fibers in place and prevent matrix cracks from developing enough energy to be able to penetrate the fibers. However, in many applications where hermeticity, compressive and/or transverse strength, or wear resistance is required, for example, porous materials are usually not adequate. For such applications, a dense matrix is preferred, and, therefore, some type of fiber coating is likely to be required to provide crack deflection at the fiber/matrix interfacial region. For a dense matrix composite, the concept of porosity has been applied sparingly to weaken the fiber/matrix interface. Porous coatings have been applied to large-diameter monofilaments where discrete carbon or polymer particles are used as a fugitive phase that is later burned out to create porosity.15–18 The polymer particles are usually 100 nm in size and are of the appropriate scale to provide a fugitive phase for fiber coatings that are 1 m in thickness. The requirement for a much thinner coating on the 10-mdiameter filaments in commercially available tows and cloths has led to the development of porous coatings derived from intimate mixtures of oxide and carbon particles on the order of 10 nm in diameter.19–22 With such a fine-scale microstructure, the ability of the porous coating to resist sintering and/or coarsening when constrained between fiber and matrix during exposure to elevated temperatures is paramount. Oxides with low self-diffusion, such as mullite, zircon, and rare-earth aluminates (garnets and magnetoplumbites), are good candidate porous fiber coatings. Porous coatings based on alumina, mixed zirconia and silica, mullite, and lanthanum hexaluminate have been applied to small-diameter alumina and alumina–mullite fibers in tows.19–23 Porous zirconia and lanthanum hexaluminate coatings have been applied to sapphire monofilaments.15,18,24 If an energy-based crack-deflection criterion is considered within a porous coating,25 a minimum pore volume fraction of 0.3 is needed to reduce the fracture energy of polycrystalline alumina to 25% of that of the dense material.5,6 For crack deflection at a coating/fiber interface, the required pore fraction may be lower because of the higher elastic modulus of the fiber compared with that of the porous coating. This paper focuses on a porous yttrium aluminum garnet (YAG, Y3Al5O12) fiber coating that does not degrade fiber strength in air. Fiber degradation has been a major concern with most oxide coatings. Aspects of precursor synthesis, fiber coating, coating-microstructure development, and composite processing and evaluation are discussed. II. Experimental Procedure (1) Precursor and Coating Synthesis YAG is the most creep-resistant oxide known.26–28 YAG synthesis by conventional solid-state reaction of the elemental oxides requires temperatures 1600°C for extended periods.29 Commercially available alumina fibers, such as Nextel™ 610 (3M Corp., Minneapolis, MN), cannot be exposed to temperatures 1200°C for more than very short periods without degrading strength via grain growth.30 This limits the processing window for the application of a coating and subsequent matrix processing of the composite to temperatures 1200°C. Recently, the synthesis of phase-pure YAG at temperatures 800°C within 1 h has been reported.31 The polymeric precursor crystallizes directly into the garnet structure from an amorphous powder starting at 600°C when heated in an oxidizing atmosphere. In argon, peaks of primarily hexagonal YAlO3 are present along with traces of YAG and amorphous alumina at temperatures of 700°–900°C. At 1000°C within 1 h, YAlO3 reacts with the residual alumina to form YAG, which is then the only crystalline phase present.31 The precursor is well-suited for this work because an intimate mixture of nanosized YAG and carbon is obtained when it is heated in an inert atmosphere at temperatures below which strength degradation of Nextel 610 fiber occurs. Details of the mixed-metal citric acid/ethylene glycol/ethanol solution precursor synthesis have been discussed elsewhere.31 In the present study, two solutions are prepared to obtain coatings with carbon contents (derived from thermal decomposition of the organic components) of 30 and 50 vol% on a solids basis. The carbon acts as a fugitive phase to establish and maintain porosity during subsequent matrix infiltration and densification; only after matrix processing is the carbon removed by heating in air. (2) Fiber Coating Nextel 610 fiber tow was first passed continuously through the furnace of a fiber coating apparatus,32 shown in Fig. 1, at 900°C in air to remove the sizing. The hot zone of the in-line furnace was F. W. Zok—contributing editor Manuscript No. 187773. Received April 16, 2001; approved July 30, 2002. Supported by AFRL, under Contract No. F33615-96-C-5258. *Member, American Ceramic Society. † Also affiliated with UES, Inc., Dayton, OH 45432. J. Am. Ceram. Soc., 85 [11] 2703–10 (2002) 2703 journal
2704 Journal of the American Ceramic Sociery-Cinibulk et al. Vol. 85. No. 11 solutions and a furnace temperature of 1000C. After problems with severe fiber bridging by excess entrained solution, subsequent coating runs were conducted at lower concentrations and lower temperatures. The 5, 8, and 10 g/L sols were applied three times with the coated tow passing through a furnace at 800C(10 s residence time)and then over a 2.5-cm-diameter guide wheel to bend the tow by "60.(Fig. 1), as summarized in Table I. The guide wheel allowed for any weak coating bridges between filaments to be broken. The lower temperature ensured that a partially pyrolyzed amorphous coating was deposited during each pass, which minimized stratification of the coating due to multiple passes through the coater, and was more likely to produce a Furnace thicker, single homogeneous layer. Postcoating heat treatment were conducted under various conditions to evaluate their effect the coatings and on fiber tow strength; these are given in Table I All postcoating heat treatments included an initial heating for I h in argon at 1000C to convert the coating to a homogeneously dispersed mixture of nanometer-sized YAG and carbon Fiber Tow ( Minicommposite Processing In the absence of being able to fabricate an oxide composite t Argor with a dense matrix that also had a modulus close to that of the fibers, an alumina matrix with 40 vol% porosity was chosen to Immiscible Displacing liquid evaluate the porous fiber coating concept. Control composites using the same fiber volume fraction and matrix were fabricated Coating Precursor Liquid and tested for comparison. Tows(tows 2, 6, and 8 in Table D) containing the homogeneously dispersed YAG and carbon fiber ol% alumina(AKP-53, Sumitomo Chemicals, Tokyo, Japan) along with a small amount of gel-casting agents. Minicomposites x75 mm in length were prepared by inserting four alumina "1.6 mm. On heating the tubing with a heat gun, the inner diameter decreased to -l mm, expelling excess matrix slurry from the composites, which increased the fiber volume fraction. The alumina matrix was allowed to gel and then was dried under 95% elative humidity. The resulting composites were sintered at Fig. 1. Schematic of continuous fiber coater Inset shows magnified view 1200.C in either argon or air to form unidirectional minicompos of wheels to break weak fiber bridges ites containing four tows, each with a fiber volume fraction of "30%.(The processing and mini Ite fabrication procedures are discussed in detail elsewhere. , )A number of conditions were used to sinter and heat-treat the minicomposites to evaluate 8 cm in length, which resulted in a residence time of <10 s. the porous coatings, as summarized in Table Il. The final heat Desized tows were the with YAG solutions having ar treatment was always in air, which oxidized the fugitive carbon i oxide concentration of 5 15.or2 and a final carbo the coatings to yield porous YAG fiber coatings ontent of either 30 or 50 ased on total solids content. All Datings were applied using a continuous coating apparatus with exadecane as an immiscible liquid to displace excess YAG/C 4 Tensile Testin solution from between the filaments to minimize bridging(Fig. Over the past five years, a tensile test procedure has been 1). Initial coating trials were conducted with 15 and 25 g/L developed in our laboratory to evaluate novel fiber coatings. This Table L. Fiber-Coating Precursor Solutions, Heat-Treatment Conditions, and Strengths Precursor solution/coating Postcoating heat treatment Carbon content YAG content Atmosphere l000/1200 gon 1000/1200 Argon/air 0.10 33333333311 l000/1200 1/100 Argon/ai 1000/1200 l000/1200 Argon/ai 1000/12001/100rgon/air 74 13 Air 1.1 0.20 14 1200 1.0 10 0.12 dEsized at 900%C, no coating
8 cm in length, which resulted in a residence time of 10 s. Desized tows were then coated with YAG solutions having an oxide concentration of 5, 8, 10, 15, or 25 g/L and a final carbon content of either 30 or 50 vol%, based on total solids content. All coatings were applied using a continuous coating apparatus with hexadecane as an immiscible liquid to displace excess YAG/C solution from between the filaments to minimize bridging (Fig. 1).32 Initial coating trials were conducted with 15 and 25 g/L solutions and a furnace temperature of 1000°C. After problems with severe fiber bridging by excess entrained solution, subsequent coating runs were conducted at lower concentrations and lower temperatures. The 5, 8, and 10 g/L sols were applied three times with the coated tow passing through a furnace at 800°C (10 s residence time) and then over a 2.5-cm-diameter guide wheel to bend the tow by 60° (Fig. 1), as summarized in Table I. The guide wheel allowed for any weak coating bridges between filaments to be broken. The lower temperature ensured that a partially pyrolyzed amorphous coating was deposited during each pass, which minimized stratification of the coating due to multiple passes through the coater, and was more likely to produce a thicker, single homogeneous layer. Postcoating heat treatments were conducted under various conditions to evaluate their effect on the coatings and on fiber tow strength; these are given in Table I. All postcoating heat treatments included an initial heating for 1 h in argon at 1000°C to convert the coating to a homogeneously dispersed mixture of nanometer-sized YAG and carbon. (3) Minicomposite Processing In the absence of being able to fabricate an oxide composite with a dense matrix that also had a modulus close to that of the fibers, an alumina matrix with 40 vol% porosity was chosen to evaluate the porous fiber coating concept. Control composites using the same fiber volume fraction and matrix were fabricated and tested for comparison. Tows (tows 2, 6, and 8 in Table I) containing the homogeneously dispersed YAG and carbon fiber coatings were infiltrated with an aqueous slurry containing 45 vol% alumina (AKP-53, Sumitomo Chemicals, Tokyo, Japan) along with a small amount of gel-casting agents.33 Minicomposites 75 mm in length were prepared by inserting four aluminainfiltrated tows into heat shrink tubing with an inner diameter of 1.6 mm. On heating the tubing with a heat gun, the inner diameter decreased to 1 mm, expelling excess matrix slurry from the composites, which increased the fiber volume fraction. The alumina matrix was allowed to gel and then was dried under 95% relative humidity. The resulting composites were sintered at 1200°C in either argon or air to form unidirectional minicomposites containing four tows, each with a fiber volume fraction of 30%. (The processing and minicomposite fabrication procedures are discussed in detail elsewhere.33,34) A number of conditions were used to sinter and heat-treat the minicomposites to evaluate the porous coatings, as summarized in Table II. The final heat treatment was always in air, which oxidized the fugitive carbon in the coatings to yield porous YAG fiber coatings. (4) Tensile Testing Over the past five years, a tensile test procedure has been developed in our laboratory to evaluate novel fiber coatings. This Fig. 1. Schematic of continuous fiber coater. Inset shows magnified view of wheels to break weak fiber bridges. Table I. Fiber-Coating Precursor Solutions, Heat-Treatment Conditions, and Strengths Tow Precursor solution/coating Postcoating heat treatment Strength Carbon content (vol%) YAG content (g/L) Number of passes Temperature (°C) Time (h) Atmosphere Tensile strength (GPa) Weibull modulus Coefficient of variation 1 50 5 3 2.1 19 0.06 2 50 5 3 1000 1 Argon 2.2 22 0.05 3 50 5 3 1000/1200 1/2 Argon 1.6 9 0.13 4 50 5 3 1000/1200 1/2 Argon/air 1.3 12 0.10 5 50 5 3 1000/1200 1/100 Argon/air 0.81 5 0.22 6 50 8 3 1000 1 Argon 7 50 8 3 1000/1200 1/2 Argon/air 1.2 14 0.09 8 30 10 3 1000 1 Argon 2.0 18 0.07 9 30 10 3 1000/1200 1/2 Argon 1.7 29 0.04 10 30 10 3 1000/1200 1/2 Argon/air 0.78 13 0.09 11 30 10 3 1000/1200 1/100 Argon/air 0.74 8 0.15 12† 1 1.6 10 0.11 13† 1 1200 2 Air 1.1 6 0.20 14† 1 1200 100 Air 1.0 10 0.12 † Desized at 900°C, no coating. 2704 Journal of the American Ceramic Society—Cinibulk et al. Vol. 85, No. 11
November 2002 Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composite 2705 Table Il. Minicomposite Processing and Heat-Treatment Conditions Fiber tow perature (C Time(h) Atmosphere Strength(GPa) Weibull modulus Coefficient of variation ABCDE 200/1200 2/100 Argon/air 0.28 0.41 1200/1100 Argon/air 1200/1200 2/100 Argon/air 636468 0.21 1200 100 0.28 fFrom Table 1. *Control composites with no fiber coating. procedure has been described in detail previously.23, 34,35 Tows and vol% carbon and those with 50 vol%. Based on the absence of any inicomposites were mounted with epoxy on card stock that was gross segregation of phases, we assumed that the volume fraction used as a fixture. Uncoated tows were first pulled through ethanol of porosity in the YAG fiber coatings after oxidation of the in an attempt to align individual filaments before mounting. The fugitive carbon equaled the initial carbon volume fraction, as gauge length was 25.4 mm. At least 20 sections of fiber tow and determined by TGA. at least 10 minicomposites were tested for each sample where Heat-treating the coated tows in air resulted in a loss of carbon strengths were reported. Tow and minicomposite tensile strengths and an increase in grain size. As the temperature was increased te were determined from maximum loads and total fiber areas( based 1200C, the coatings sintered, and the grains reached a size on an average fiber diameter of 12 um and 420 filaments/tow), comparable to the thickness of some of the coatings, as shown in there was usually no porosity left. Thicker coatings still had (5) Phase and Microstructure Characterization residual porosity, but it was reduced from -30 and 50 vol%to Powders of the YAG pre were prepared by heating the 10 vol%. These results are similar to those found for fiber solutions on a hot plate to "200C to rate the solvent and coatings that were heat-treated within a porous matrix, as dis decompose the majority of the organics before heating in argon to cussed later temperatures of =1000C for I h Powder XRD was used to verify Table I contains the strengths of coated and uncoated (control) phase development in the precursor. Final carbon content was tows after undergoing various heat treatments to determine the effects of coating composition on fiber strength, As-received and determined by thermogravimetric analysis (tga) in air on pow- desized tows of Nextel 610(without a coating)have a strength of ders previously heated to 1000C Surfaces of coated and heat-treated fibers and urfaces 1.6 GPa. Heating the same tow at 1200C for 2 and 100 h of minicomposites were characterized by SEM 360FE, reduces its strength to 1. I and 1.0 GPa, respectively.The reduction LEO, Cambridge, U. K ). Coated fiber tows er/matrⅸx interfaces in the composites were examined by TEM(Model CM200FEG, Philips, Eindhoven, The Netherlands). Energy- dispersive X-ray spectroscopy (EDS) was used for elemental analysis. TEM specimens were prepared by impregnating coated tows with a high-temperature epoxy and thinning to electron transparency with diamond lapping films, followed by low-angle ion-beam milling, as described in detail elsewhere. 3, II. Results and discussion (1 Coated Fiber Tows Characterization of the coated tows by SEM and TEM indicated at minimal fiber bridging remained, and the filaments appeared well-coated(Fig. 2). Evidence of prior bridging during the coating process was present, but the wheels were effective at breaking them. Multiple passes of the tow through the coating liquid were used to increase coverage of fiber surfaces and to increase coating thickness. The coated tows were not stiffened by the coating and felt the same as the uncoated tows on handling. Coating thickness anged from <10 to 100 nm. A few fibers appeared to not contain any coatings at all (via TEM); however, yttrium was always detectable at fiber surfaces by EDs. Whether this was due to the limited amount of precursor that wet the fiber surface durin coating or whether it was due to residual yttrium present after the had spalled off was difficult to determin .e. It was not ssible to differentiate a lack of fiber coating from loss of fiber coating during handling. It was also difficult to correlate coating thickness with precursor concentration of such a narrow range because the number of coatings measured by tEM was smal e The coatings were amorphous and homogeneous as-coated(Fig D). After heating for I h at 1000.C in a a two-phase coatin of intimately mixed YAG and carbon was obtained with equiaxed 10 nm in size(Figs. 3(b)and(c). Over this dual-phase coating, a dense shell of YAG-10 nm thick was usually observed Similar features have also been observed in other oxide/carbon Fig. 2. SEM images of tow 1, indicating good covithin the tow due to the f the filaments ber coatings using the same coating procedure. 1, 22 39 There was by Y AG/C coating and minimal bridging of fibers little noticeable difference between the coatings that contained 30 use of lower-concentrated solution and wheels to break bridges
procedure has been described in detail previously.23,34,35 Tows and minicomposites were mounted with epoxy on card stock that was used as a fixture. Uncoated tows were first pulled through ethanol in an attempt to align individual filaments before mounting. The gauge length was 25.4 mm. At least 20 sections of fiber tow and at least 10 minicomposites were tested for each sample where strengths were reported. Tow and minicomposite tensile strengths were determined from maximum loads and total fiber areas (based on an average fiber diameter of 12 m and 420 filaments/tow), neglecting any contributions from either the coating or matrix.36,37 (5) Phase and Microstructure Characterization Powders of the YAG precursors were prepared by heating the solutions on a hot plate to 200°C to evaporate the solvent and decompose the majority of the organics before heating in argon to temperatures of 1000°C for 1 h. Powder XRD was used to verify phase development in the precursor. Final carbon content was determined by thermogravimetric analysis (TGA) in air on powders previously heated to 1000°C in argon. Surfaces of coated and heat-treated fibers and fracture surfaces of minicomposites were characterized by SEM (Model 360FE, LEO, Cambridge, U.K.). Coated fiber tows and fiber/matrix interfaces in the composites were examined by TEM (Model CM200FEG, Philips, Eindhoven, The Netherlands). Energydispersive X-ray spectroscopy (EDS) was used for elemental analysis. TEM specimens were prepared by impregnating coated tows with a high-temperature epoxy and thinning to electron transparency with diamond lapping films, followed by low-angle ion-beam milling, as described in detail elsewhere.38,39 III. Results and Discussion (1) Coated Fiber Tows Characterization of the coated tows by SEM and TEM indicated that minimal fiber bridging remained, and the filaments appeared well-coated (Fig. 2). Evidence of prior bridging during the coating process was present, but the wheels were effective at breaking them. Multiple passes of the tow through the coating liquid were used to increase coverage of fiber surfaces and to increase coating thickness. The coated tows were not stiffened by the coating and felt the same as the uncoated tows on handling. Coating thickness ranged from 10 to 100 nm. A few fibers appeared to not contain any coatings at all (via TEM); however, yttrium was always detectable at fiber surfaces by EDS. Whether this was due to the limited amount of precursor that wet the fiber surface during coating or whether it was due to residual yttrium present after the coating had spalled off was difficult to determine, i.e., it was not possible to differentiate a lack of fiber coating from loss of fiber coating during handling. It was also difficult to correlate coating thickness with precursor concentration of such a narrow range because the number of coatings measured by TEM was small. The coatings were amorphous and homogeneous as-coated (Fig. 3(a)). After heating for 1 h at 1000°C in argon, a two-phase coating of intimately mixed YAG and carbon was obtained with equiaxed particles 10 nm in size (Figs. 3(b) and (c)). Over this dual-phase coating, a dense shell of YAG 10 nm thick was usually observed. Similar features have also been observed in other oxide/carbon fiber coatings using the same coating procedure.21,22,39 There was little noticeable difference between the coatings that contained 30 vol% carbon and those with 50 vol%. Based on the absence of any gross segregation of phases, we assumed that the volume fraction of porosity in the YAG fiber coatings after oxidation of the fugitive carbon equaled the initial carbon volume fraction, as determined by TGA. Heat-treating the coated tows in air resulted in a loss of carbon and an increase in grain size. As the temperature was increased to 1200°C, the coatings sintered, and the grains reached a size comparable to the thickness of some of the coatings, as shown in Fig. 4. When YAG grains spanned the thickness of the coating, there was usually no porosity left. Thicker coatings still had residual porosity, but it was reduced from 30 and 50 vol% to 10 vol%. These results are similar to those found for fiber coatings that were heat-treated within a porous matrix, as discussed later. Table I contains the strengths of coated and uncoated (control) tows after undergoing various heat treatments to determine the effects of coating composition on fiber strength. As-received and desized tows of Nextel 610 (without a coating) have a strength of 1.6 GPa. Heating the same tow at 1200°C for 2 and 100 h in air reduces its strength to 1.1 and 1.0 GPa, respectively. The reduction Fig. 2. SEM images of tow 1, indicating good coverage of the filaments by YAG/C coating and minimal bridging of fibers within the tow due to the use of lower-concentrated solution and wheels to break bridges. Table II. Minicomposite Processing and Heat-Treatment Conditions Minicomposite Fiber tows† Temperature (°C) Time (h) Atmosphere Strength (GPa) Weibull modulus Coefficient of variation A 2 1200 2 Air 1.1 6 0.22 B 6 1200/1200 2/100 Argon/air 0.28 3 0.41 C 8 1200/1100 2/2 Argon/air 0.80 6 0.18 D 8 1200/1200 2/100 Argon/air 0.28 4 0.28 E‡ 12 1200 2 Air 0.66 6 0.21 F‡ 12 1200 100 Air 0.28 8 0.16 † From Table I. ‡ Control composites with no fiber coating. November 2002 Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composites 2705
Journal of the American Ceramic Sociery-Cinibulk et al. Vol. 85. No. 1I Fibe Fiber 200mn n 100nm Epoxy YAG YAG+C 100 nm Fig. 4. TEM images of tow 4 showing grain growth and sintering of YAG coating after heating at 1200C for 2 h in air. Where the uncon strained porous coating is thin, grains grow to film thickness and eliminate 100mm porosity. Epoxy with heat treatment, either of which may be attributable to enviro All tows initially increased in strength after coating with YAG/C and heat-treating in argon. Figure 5 contains a Weibul plot to illustrate the strength increase with 5 g/L YAG containing 50 vol% carbon then heat-treated under various conditions. After an initial strength of 2. 1 GPa for the as-coated fiber. the strength increased further to 2.2 GPa after heating at 1000C for I h in argon. Heating at 1200oC in argon for 2 h decreased the fiber strength to 1.6 GPa while. at 1200 C heating in air for 2 h reduced the strength to 1.3 GPa. a 100 h heat treatment at 1200% C reduced the strength of the coated fiber to 0.8 Pa, which was 80% of the strength of the uncoated fiber heated 20 nm under the same conditions. Similar results were obtained for tows coated with different solution concentrations and 30 vol% carbon The initial increase in strength that is often observed for fibers Fig 3. TEM images of (a)tow I, showing amorphous coating on fiber as with carbon-containing coatings could be attributed to flaw heal and 10 nm crystalline YAG after heating at 1000C for I h in argon, and straightening of filaments within the coated tow, which then allows (c)tow 8, showing intimate mixture of 30 vol% amorphous carbon and 10 nm crystalline YAG after heating at 1000 C for I h in argon. Note a greater fraction of the filaments to bear load during a tension test formation of a surface layer of dense Y AG following heat treatment. Inset however, the desized tows(without the coating) are pulled throu n Fig. 3(b)is a selected-area diffraction pattern of the coating that ethanol to help straighten the filaments before mounting for orresponds with randomly orient stalline YAG tension testing. As the carbon is removed by oxidation, strength decreases and approaches that of the uncoated fibers with a simila thermal history. This subsequent reduction in strength is most likely due to weakening of the fiber by either flaw-size increase or been attributed arily to grain growth; however, environmental effects. as is observed for uncoated fibers. rather the ure to than any detrimental effect due to the presence of the YAG coa correlation between size and fiber strength following heat tself. The presence of subsequent porosity also enhances the treatment. For example, in the present study, for a 1.6 GPa material to be reduced in strength to 1. 1 GPa, as measured for tows heated at 1200.C for 2 h, the critical flaw size would have to be increased lowever. a coated fiber heat-treated at 1200.C for 100 h does by over a factor of 2, given o /o,=a7al, where o and a are show strength reductions of 19%0-26% compared with a similarl strength and flaw size. res ely. Figure 3(a) shows a coate thermally processed uncoated tow. In this case, the coating ber after no heat treatment other than the brief time spent in the densify to the point of full density for thin coatings and to levels coating furmace, and Fig. 4 shows a fiber after heating for 2 h at 1200C. Clearly, there is not a twofold increase in grain size following heat treatment. Either critical flaws grow at different Weibull analysis does not necessarily imply that fiber-bundle prop rates than the grains or the toughness at the crack tip decreases follow weakest-link statistics
in strength has been attributed primarily to grain growth; however, there is little evidence in the literature to support a direct correlation between grain size and fiber strength following heat treatment. For example, in the present study, for a 1.6 GPa material to be reduced in strength to 1.1 GPa, as measured for tows heated at 1200°C for 2 h, the critical flaw size would have to be increased by over a factor of 2, given 1/2 a2 1/2/a1 1/2, where and a are strength and flaw size, respectively. Figure 3(a) shows a coated fiber after no heat treatment other than the brief time spent in the coating furnace, and Fig. 4 shows a fiber after heating for 2 h at 1200°C. Clearly, there is not a twofold increase in grain size following heat treatment. Either critical flaws grow at different rates than the grains or the toughness at the crack tip decreases with heat treatment, either of which may be attributable to environmental effects. All tows initially increased in strength after coating with YAG/C and heat-treating in argon. Figure 5 contains a Weibull plot to illustrate the strength increase of tows coated three times with 5 g/L YAG containing 50 vol% carbon then heat-treated under various conditions.‡ After an initial strength of 2.1 GPa for the as-coated fiber, the strength increased further to 2.2 GPa after heating at 1000°C for 1 h in argon. Heating at 1200°C in argon for 2 h decreased the fiber strength to 1.6 GPa while, at 1200°C, heating in air for 2 h reduced the strength to 1.3 GPa. A 100 h heat treatment at 1200°C reduced the strength of the coated fiber to 0.8 GPa, which was 80% of the strength of the uncoated fiber heated under the same conditions. Similar results were obtained for tows coated with different solution concentrations and 30 vol% carbon. The initial increase in strength that is often observed for fibers with carbon-containing coatings could be attributed to flaw healing. The strength increase could also be partially attributed to the straightening of filaments within the coated tow, which then allows a greater fraction of the filaments to bear load during a tension test; however, the desized tows (without the coating) are pulled through ethanol to help straighten the filaments before mounting for tension testing. As the carbon is removed by oxidation, strength decreases and approaches that of the uncoated fibers with a similar thermal history. This subsequent reduction in strength is most likely due to weakening of the fiber by either flaw-size increase or environmental effects, as is observed for uncoated fibers, rather than any detrimental effect due to the presence of the YAG coating itself. The presence of subsequent porosity also enhances the permeability of any trapped gases out of the coating that may otherwise cause stress corrosion at fiber grain boundaries.40 However, a coated fiber heat-treated at 1200°C for 100 h does show strength reductions of 19%–26% compared with a similarly thermally processed uncoated tow. In this case, the coatings densify to the point of full density for thin coatings and to levels ‡ Use of Weibull analysis does not necessarily imply that fiber-bundle properties follow weakest-link statistics. Fig. 3. TEM images of (a) tow 1, showing amorphous coating on fiber as coated; (b) tow 2, showing intimate mixture of 50 vol% amorphous carbon and 10 nm crystalline YAG after heating at 1000°C for 1 h in argon; and (c) tow 8, showing intimate mixture of 30 vol% amorphous carbon and 10 nm crystalline YAG after heating at 1000°C for 1 h in argon. Note formation of a surface layer of dense YAG following heat treatment. Inset in Fig. 3(b) is a selected-area diffraction pattern of the coating that corresponds with randomly oriented polycrystalline YAG. Fig. 4. TEM images of tow 4 showing grain growth and sintering of YAG coating after heating at 1200°C for 2 h in air. Where the unconstrained porous coating is thin, grains grow to film thickness and eliminate porosity. 2706 Journal of the American Ceramic Society—Cinibulk et al. Vol. 85, No. 11
Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composite 2707 N610-Desized 1.6, 10 日N610-1200C2h11,6 2.1,19 1.6,9 b-1200c2h(A)1.3.12 1200c2h(Air)0.8,5 12000,100h(Air) 与4 44.555.56657758 Ln stress, MPa I 5. Weibull plot of tows 1-5 that were coated with 50 vol% YAG/C precursor and tows 12 and 13(desized, without coating) heat-treated under various here residual porosity is <10% in much thicker coatings; any flaws in the coating(or matrix) now penetrate the fiber readily in the absence of any crack-deflecting mech Incomp Table ll summarizes the st ns of minicomposites Minicom- osite A had an average strength that was nearly twice as high as the control minicomposite F. Minicomposites B and D, heat- treated for 100 h in air at 1200oC, had strengths that were equal to that of the control minicomposite F heated under the same conditions but without a fiber coating. Figure 6 contains sem images of fracture surfaces of minicomposites A and B, showing the difference in fracture behavior of the two composites. Long ber pullout lengths and holes left behind by fibers were clearly visible in the fracture surfaces of minicomposite A, as opposed to nly fibers exposed by matrix disintegration, as is often seen in (a) orous matrix composites. Minicomposite B displayed a brittle fracture surface with little, if any, fiber pullout A TEM image of the fiber/matrix interfacial region from minicomposite A is shown in Fig. 7. T and pore sizes in the YAG fiber coating increase to 40-50 nm, and particles begi to sinter following matrix processing at 1200%C for 2 h(F b), (c)). Because the matrix is porous as well and could pected to deflect cracks, the presence of the porous YAG coati er tensile strengths. Either rosity distribution in the matrix alone is not adequate to deflect acks, whereas that in the fiber coating is, or the coating protects the fiber from degradation during matrix processing. Sintering of matrix alumina particles to the alumina fiber is observed in the absence of a coating, which could lead to fiber degradation in the control composites by enhanced fiber/matrix bonding and stress concentration Figures 8 and 9 show TEM images of the fiber/ in minicomposites C and D, which contain an I matrix ite aes of Yag that
where residual porosity is 10% in much thicker coatings; any flaws in the coating (or matrix) now penetrate the fiber readily in the absence of any crack-deflecting mechanism.35 (2) Minicomposites Table II summarizes the strengths of minicomposites. Minicomposite A had an average strength that was nearly twice as high as the control minicomposite F. Minicomposites B and D, heattreated for 100 h in air at 1200°C, had strengths that were equal to that of the control minicomposite F heated under the same conditions but without a fiber coating. Figure 6 contains SEM images of fracture surfaces of minicomposites A and B, showing the difference in fracture behavior of the two composites. Long fiber pullout lengths and holes left behind by fibers were clearly visible in the fracture surfaces of minicomposite A, as opposed to only fibers exposed by matrix disintegration, as is often seen in porous matrix composites.41 Minicomposite B displayed a brittle fracture surface with little, if any, fiber pullout. A TEM image of the fiber/matrix interfacial region from minicomposite A is shown in Fig. 7. The grain and pore sizes in the YAG fiber coating increase to 40–50 nm, and particles begin to sinter following matrix processing at 1200°C for 2 h (Fig. 3(b),(c)). Because the matrix is porous as well and could be expected to deflect cracks, the presence of the porous YAG coating results in composites with higher tensile strengths. Either the porosity distribution in the matrix alone is not adequate to deflect cracks, whereas that in the fiber coating is, or the coating protects the fiber from degradation during matrix processing. Sintering of matrix alumina particles to the alumina fiber is observed in the absence of a coating, which could lead to fiber degradation in the control composites by enhanced fiber/matrix bonding and stress concentration. Figures 8 and 9 show TEM images of the fiber/matrix interfaces in minicomposites C and D, which contain an initial 30 vol% porosity coating and are heat-treated at 1100°C for 2 h in air and 1200°C for 100 h in air, respectively. Clearly evident is the grain growth, pore coarsening, and sintering of YAG that is occurring. In Fig. 5. Weibull plot of tows 1–5 that were coated with 50 vol% YAG/C precursor and tows 12 and 13 (desized, without coating) heat-treated under various conditions. Fig. 6. SEM images of (a) minicomposite A and (b) minicomposite B following heating at 1200°C for 2 and 100 h in air, respectively. Note fiber pullout and troughs left by fiber pulled out in Fig. 6(a). November 2002 Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composites 2707
Journal of the American Ceramic Sociery-Cinibulk et al. Vol. 85. No. 1I are thicker than 150 nm (YAG grain size), some porosity is still present, but the amount of porosity has decreased to "10 vol% This work shows that a porous fiber coating increase the strength of composites produced with small-diameter fibers in a polycrystalline oxide matrix. However, the improve- ment in strength diminishes when the composites are exposed to air at high temperatures for extended (100 h) periods. After 100 h at 1200C in air, the composites containing the porous YAG coating have the same strengths as those composites produced without a fiber coating From previous work on microcomposites with porous matrixes, we find that a pore volume fraction of at least "15% is needed to give tough, composite-like behavior 200 nm likely penetrate the fibers without being deflected or trapped The densified yag coatings now can introduce flaws that most within the porous coating itself. The dramatic loss in strength of the porous Y AG-containing composites after 100 h at 1200C in Matrix air when the coating has densified suggests that the higher strength after only short heat treatments is due to the presence of a porous coating. The fact that both the control minicomposite and the minicomposite with the porous YAG fiber coating yield the same strength indicates that once the fiber coating has densified the twe composites behave the same, and there is no apparent benefit to simply having a dense YAG fiber coating Pores trapped within a dense solid are constrained, and it is generally very difficult to remove them. Similarly a porous fiber coating, trapped between a dense fiber and dense matrix, would be equally constrained, and little densification of the porous coating would be expected to occur. Because the porous coating in Fig. 7. TEM images of minicomposite A heated at 1200C for 2 h in air. a dense matrix is constrained from densifying, only pore coarser Where coating is thin, porosity is reduced to zero as grains grow to equal ing is expected to occur. However, the minicomposites in the film thickness and sinter to full density. In thicker coatings, coarsening of present study contain a porous alumina matrix that cannot be rains and porosity occul expected to constrain the porous fiber coating from sintering. A Figs. 7-9 show, in addition to grain growth and pore coarsening, there is a significant amount of densification that occurs. In fact, the rate of densification approaches that of the fiber coatings in the nicomposite D(Fig. 9), the grain size of Y AG(100-150 nm) absence of a matrix(Fig. 4). We estimate that the final porosity in has approached the coating thickness in some cases, leaving either a dense or discontinuous coating behind. When the fiber coatings where grain size is comparable to coating thickness to -10 volo for the thickest coatings. If a critical volume fraction of pores is required to encourage crack deflection, a reduction in porosity Matrix atrix p-YAG YAG Fiber Matrix fibe Fig. 8. TEM images of fiber coatings in m osite C heated at I 100oC for 2 h in air. porosity retained after short exposures to high Fig. 9. TEM images of fiber coatings in minicomposite D heated at 1200C for 100 h in air. Porosity is reduced to <10 vol% after 100 h
minicomposite D (Fig. 9), the grain size of YAG (100–150 nm) has approached the coating thickness in some cases, leaving either a dense or discontinuous coating behind. When the fiber coatings are thicker than 150 nm (YAG grain size), some porosity is still present, but the amount of porosity has decreased to 10 vol%. This work shows that a porous fiber coating can be used to increase the strength of composites produced with small-diameter fibers in a polycrystalline oxide matrix. However, the improvement in strength diminishes when the composites are exposed to air at high temperatures for extended (100 h) periods. After 100 h at 1200°C in air, the composites containing the porous YAG coating have the same strengths as those composites produced without a fiber coating. From previous work on microcomposites with porous matrixes, we find that a pore volume fraction of at least 15% is needed to give tough, composite-like behavior.35 The densified YAG coatings now can introduce flaws that most likely penetrate the fibers without being deflected or trapped within the porous coating itself. The dramatic loss in strength of the porous YAG-containing composites after 100 h at 1200°C in air when the coating has densified suggests that the higher strength after only short heat treatments is due to the presence of a porous coating. The fact that both the control minicomposite and the minicomposite with the porous YAG fiber coating yield the same strength indicates that once the fiber coating has densified, the two composites behave the same, and there is no apparent benefit to simply having a dense YAG fiber coating. Pores trapped within a dense solid are constrained, and it is generally very difficult to remove them. Similarly a porous fiber coating, trapped between a dense fiber and dense matrix, would be equally constrained, and very little densification of the porous coating would be expected to occur. Because the porous coating in a dense matrix is constrained from densifying, only pore coarsening is expected to occur. However, the minicomposites in the present study contain a porous alumina matrix that cannot be expected to constrain the porous fiber coating from sintering. As Figs. 7–9 show, in addition to grain growth and pore coarsening, there is a significant amount of densification that occurs. In fact, the rate of densification approaches that of the fiber coatings in the absence of a matrix (Fig. 4). We estimate that the final porosity in the YAG coatings ranges from zero, for the thinnest coatings where grain size is comparable to coating thickness, to 10 vol%, for the thickest coatings. If a critical volume fraction of pores is required to encourage crack deflection, a reduction in porosity Fig. 7. TEM images of minicomposite A heated at 1200°C for 2 h in air. Where coating is thin, porosity is reduced to zero as grains grow to equal film thickness and sinter to full density. In thicker coatings, coarsening of grains and porosity occur. Fig. 8. TEM images of fiber coatings in minicomposite C heated at 1100°C for 2 h in air. Porosity is retained after short exposures to high temperatures. Fig. 9. TEM images of fiber coatings in minicomposite D heated at 1200°C for 100 h in air. Porosity is reduced to 10 vol% after 100 h. 2708 Journal of the American Ceramic Society—Cinibulk et al. Vol. 85, No. 11
November 2002 Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composite 2709 below that critical level may have occurred and could be respon- References sible for the reduction in composite strength compared to that of composites without a fiber coating. Clearly, in minicomposite D E. Ryshkewitch, "Compression Strength of Porous Sintered Alumina and Zirco- the volume fraction of porosity in the coating has been reduced to nia”J.Am. Ceran.Soc,36|2]65-68(1953) less than the porosity present in the alumina matrix(Fig. 9) J. Am Ceram Soc., 39[11]377-85(1956). that th BL. A Simpson, "Effect of Microstructure on Measurements of Fracture Energy of orous Y AG coating gives a higher-strength composite, even if the Ai pna d Am Ceram soc 56mi-il( 1975) nce of Elastic Moduli on Porosity matrix is porous. Porosity alone in the matrix is not sufficient to JAm. Ceram.Soc,607-8]345-49(1977) give a high-strength composite at 1200oC. Whether the fine ngh, and R. B Poeppel, "Dependence of Ceramic Fracture porosity of the porous YAG fiber coating is needed to provide a Popertescon J. Mater.Sci,28,3589-93(1993) weaker interface than a porous matrix alone can provide, or C. Lam, F. F. Lange, and A. G. Evans, "Mechanical Properties of Partiall na Produced from Powder Compacts,J.Am. Cera. Soc., 7718 whether the fibers are damaged by the alumina matrix(including 2113-17(1994) its processing)in the absence of a porous Y AG fiber coating, is not S C Nanjangud, R. Brezny, and D J, Green,"Strength and Young's Modulu clear. The level of porosity present in the matrix (-40 vol%)(1996). Behavior of a Partially Sintered Porous Alumina,"JAm. Ceram Soc., 78[1 266-68 should be adequate to give composite-like behavior and high of Physical Property-Porosity Models strength In commercial porous oxide matrix composites, nearly Based on minimum solid fully retained fiber strengths are obtained after processing and 9A. P. Roberts and E. J Properties of Model Porous Ceramics, long-term heat treatment at 1200C. However, the commercial Soc,83[12 TOM. G. Harrison. M composites are produced with Nextel 720 diphasic alumina- Matrix Composite Member and Method of Making, "U.S. Pat. No. 5 306 554, Ap mullite fibers, not single-phase alumina, so a direct compariso cannot be made because of the differences in fiber moduli thermal Iw.-C. Tu, F. F. Lange, and A. G. Evans, "Concept for a Damage-Tolerant Ceramic Composite with Strong Interfaces,J.Am. Ceram. Soc., 79[2]417-24 expansion, and bonding/interaction with the matrix. Ideally, fibers would need to be extracted from the matrix after processing to test AC. G. Levi, J. Y. Yang, B J. Dalgleish, F. w. Zok, and A. G. Evans,"Processin for strength degradation; however, extraction of suitable lengths of and Performance of an All-Oxide Ceramic Composite, J. Am. Ceram Soc., 81 [81 alumina fibers from an alumina matrix for testing is difficult, if not 2077-86(1998). S. G. Steel, L. P. Zawada, and S. Mall, ""Fatigue Behavior of a Nextel impossible, without damaging the fibers In recent work, the presence of a YAG second phase in a porous Pro 2 i3 695-702 te o0 Room and Elevated lemperature, Ceram. Eng. alumina matrix composite without a separate fiber coating also M. K. Cinibulk, K. A. Keller, T. Mah, and T. A. Parthasarathy, "Nextel 610 and 650 Fiber Reinforced Porous Alumina-YAG Matrix Composites, Ceram. Eng. Sci. control composites were found to be of equal strength at 1100.C PasM 2 3 2000). M. H. Jaskowiak, S. I. Eldridge, J. B. Hurst, and J. A. Setlock, "Interfacial Coatings for Sapphire/AL,O3, p. 84 in HITEMP Review. NASA Conference composites displayed greater strengths at 1200C for times of up to U J. T. Ogbuji, "A Porous Oxidation-Resistant Fiber Coating for CMC 100h. YAG was distributed as a dense phase bonding the alumina Interphase, "Ceram. Eng. Sci. Proc., 16[4]497-505(199 particles of the matrix and also separating the fiber from the matrix PL. U. J. T. Ogbuji, "Evaluation of a Porous Fiber Coating in SiC-Si3N4 alumina. These results, along with those of the present stud uggest that Y ag inhibits densification of the matrix at 1200 C. In IO Sudre, A G Razzell, L. Molliex, and M. Holmquist, "Alumina Single-Crystal Fiber Reinforced Alumina Matrix for Combustor Tiles, Ceram. Eng. Sci. Proc., 19 the absence of YAG in the matrix, however, a porous YAG fiber [4]273-80(1998) ating is sufficient to deflect nly until a critical coating density and/or matrix density is Lanthanum Phosphate Fiber Cou Proc,1753-60(19 Carbon-Aluminum Oxide between 5 and 100 h at 1200C later. Res. Soc. Symp. Proc., 432 E. Boakye, R. S. Hay, M. D. Petry, and T. A. Parthasarathy, "Sol-Gel Zircon-Carbon Precursors and Coating of Nextel 720 Fiber Tows, "Cera. Eng IV. Conclusions Sc.Poc,20B3165-72(1999 4-M. K. Cinibulk, T. A, Parthasarathy, K. A. Keller, and T. Mah, "Porous Polymeric solutions were used to apply a porous yttrium er Coatings for Oxide-Oxide Composites, "Ceram. Eng aluminum garnet(YAG,Y,Al, O12) coating to fiber tows at Sci, Proc, 21F41219-28(2000). temperatures that did not degrade fiber strength. The use of the Porous Zirconia-Silica and Monazite Coatings us Keller, and R S Hay,"Evaluation o xtel 720 Fiber-Reinforce polymeric solution also allowed for a fugitive carbon phase that lackglas Minicomposites, J. Am. Ceran. Soc., 84[7] 1526-32(2001) was intimately mixed with the oxide to provide homogeneously 243. D. Sibold, R. L. Cook, K. Bader, and I. Reimanis, "Porous Hexaluminate dispersed porosity. The coatings were initially amorphous, but, Advances in Ceramic Matrir Composites 1. Edited by N. P. Bansal, J.P. Singh, and ille, OH, 2001 crystallized to an intimate mixture of nanometer-sized YAG and 2M. Y. He and JW "Crack Deflection at the Interface between residual amorphous carbon. Further heat treatment in air resulted Dissimilar Materials, "Int J Solids Struct, 25, 1053-67(1989). T.A. Parthasarathy, T Mah, and KKeller, "Cree chanism of Polycrystalline n a porous YAG coating Strengths of coated tows were initially Y trium Aluminum Garnet,"J. Am. Ceram Soc, 7517)1756-59(1992) in air, the strengths of the coated tows were reduced but compa- Sci Lert, 12, 379-82(192 Yttrium Aluminium Garnet Single Crystals,".Mater. higher than those of as-received tows with longer heat treatments rable to those of uncoated tows heated under similar conditions for 2>S. Karato, Z. Wang, and K. Fujino, "High-Temperature Creep of Yttrium- short times. For times of 100 h, the coated tows had a strength of on of Aluminum for silicon 20% less than the uncoated tows the System 3MnO- A2O, SiOx-3Y2O3 5Al2O3, "Am. Mineral, 7, 519(1951). As-processed minicomposites containing porous YAG fiber coatings had strengths that were nearly twice as strong as mini- atrix Composites. Edited by w. Krenkel, R Naslain, and Schneider. Wiley-VCH, Weinheim, Germany, 200 composites prepared without a fiber coating, despite the presence 3M. K. Cinibulk, "Synthesis of Yttrium Aluminum Garnet from a Mixed-Metal of high levels of matrix porosity. After heat-treating for 100 h in air, the strengths of minicomposites with and without fiber 32R. S. Hay, "Sol-Gel Coating of Fiber Tows, "Ceram. Eng. Sci. Proc., 12 [7-8] coatings were the same. This was attributed to densification of the 064-74(1991) 3K. A. Keller, T Mah, E. E. Boakye, and T. A Parthasarathy, "Gel-Casting and YAG coating, whic Reaction Bonding of Oxide-Oxide Minicomposites with Monazite Interphase vol% that no longe ned to deflect matrix cracks. hence Ceram. Eng. Sci Proc., 21 [3]525-34(2000) after long times at without 34K.A. Keller, T. A. Parthasarathy, T. Mah, M. K. Cinibulk, and E.E. boakye valuation of mona ber Coatings in a Dense Matrix Composite, Ceram. Eng coatings behaved similarly Sc.Proc,2013]451-61(1999
below that critical level may have occurred and could be responsible for the reduction in composite strength compared to that of composites without a fiber coating. Clearly, in minicomposite D, the volume fraction of porosity in the coating has been reduced to less than the porosity present in the alumina matrix (Fig. 9). One particularly interesting finding is that the presence of a porous YAG coating gives a higher-strength composite, even if the matrix is porous. Porosity alone in the matrix is not sufficient to give a high-strength composite at 1200°C. Whether the fine porosity of the porous YAG fiber coating is needed to provide a weaker interface than a porous matrix alone can provide, or whether the fibers are damaged by the alumina matrix (including its processing) in the absence of a porous YAG fiber coating, is not clear. The level of porosity present in the matrix (40 vol%) should be adequate to give composite-like behavior and high strength. In commercial porous oxide matrix composites, nearly fully retained fiber strengths are obtained after processing and long-term heat treatment at 1200°C.13 However, the commercial composites are produced with Nextel 720 diphasic alumina– mullite fibers, not single-phase alumina, so a direct comparison cannot be made because of the differences in fiber moduli, thermal expansion, and bonding/interaction with the matrix. Ideally, fibers would need to be extracted from the matrix after processing to test for strength degradation; however, extraction of suitable lengths of alumina fibers from an alumina matrix for testing is difficult, if not impossible, without damaging the fibers. In recent work, the presence of a YAG second phase in a porous alumina matrix composite without a separate fiber coating also gave increased strengths over control composites at 1200°C.42 The control composites were found to be of equal strength at 1100°C but decreased rapidly at 1200°C, whereas the YAG-containing composites displayed greater strengths at 1200°C for times of up to 100 h. YAG was distributed as a dense phase bonding the alumina particles of the matrix and also separating the fiber from the matrix alumina. These results, along with those of the present study, suggest that YAG inhibits densification of the matrix at 1200°C. In the absence of YAG in the matrix, however, a porous YAG fiber coating is sufficient to deflect cracks only until a critical coating density and/or matrix density is reached, which occurs sometime between 5 and 100 h at 1200°C. IV. Conclusions Polymeric solutions were used to apply a porous yttrium aluminum garnet (YAG, Y3Al5O12) coating to fiber tows at temperatures that did not degrade fiber strength. The use of the polymeric solution also allowed for a fugitive carbon phase that was intimately mixed with the oxide to provide homogeneously dispersed porosity. The coatings were initially amorphous, but, when heated at 1000°C for 1 h in inert atmospheres, they crystallized to an intimate mixture of nanometer-sized YAG and residual amorphous carbon. Further heat treatment in air resulted in a porous YAG coating. Strengths of coated tows were initially higher than those of as-received tows. With longer heat treatments in air, the strengths of the coated tows were reduced but comparable to those of uncoated tows heated under similar conditions for short times. For times of 100 h, the coated tows had a strength of 20% less than the uncoated tows. As-processed minicomposites containing porous YAG fiber coatings had strengths that were nearly twice as strong as minicomposites prepared without a fiber coating, despite the presence of high levels of matrix porosity. After heat-treating for 100 h in air, the strengths of minicomposites with and without fiber coatings were the same. This was attributed to densification of the YAG coating, which resulted in residual porosity levels of 10 vol% that no longer functioned to deflect matrix cracks; hence, after long times at 1200°C, the composites with and without coatings behaved similarly. References 1 E. Ryshkewitch, “Compression Strength of Porous Sintered Alumina and Zirconia,” J. Am. Ceram. Soc., 36 [2] 65–68 (1953). 2 R. L. Coble and W. D. Kingery, “Effect of Porosity on Physical Properties of Sintered Alumina,” J. Am. Ceram. Soc., 39 [11] 377–85 (1956). 3 L. A. Simpson, “Effect of Microstructure on Measurements of Fracture Energy of Al2O3,” J. Am. Ceram. Soc., 56 [1] 7–11 (1973). 4 E. A. Dean and J. A. Lopez, “Empirical Dependence of Elastic Moduli on Porosity for Ceram Materials,” J. Am. Ceram. Soc., 60 [7–8] 345–49 (1977). 5 A. S. Wagh, J. P. Singh, and R. B. Poeppel, “Dependence of Ceramic Fracture Properties on Porosity,” J. Mater. Sci., 28, 3589–93 (1993). 6 D. C. C. Lam, F. F. Lange, and A. G. Evans, “Mechanical Properties of Partially Dense Alumina Produced from Powder Compacts,” J. Am. Ceram. Soc., 77 [8] 2113–17 (1994). 7 S. C. Nanjangud, R. Brezny, and D. J. Green, “Strength and Young’s Modulus Behavior of a Partially Sintered Porous Alumina,” J. Am. Ceram. Soc., 78 [1] 266–68 (1996). 8 R. W. Rice, “Evaluation and Extension of Physical Property–Porosity Models Based on Minimum Solid Area,” J. Mater. Sci., 31, 102–18 (1996). 9 A. P. Roberts and E. J. Garboczi, “Elastic Properties of Model Porous Ceramics,” J. Am. Ceram. Soc., 83 [12] 3041–48 (2000). 10M. G. Harrison, M. L. Millard, and A. Szweda, “Fiber Reinforced Ceramic Matrix Composite Member and Method of Making,” U.S. Pat. No. 5 306 554, Apr. 26, 1994. 11W.-C. Tu, F. F. Lange, and A. G. Evans, “Concept for a Damage-Tolerant Ceramic Composite with ’Strong’ Interfaces,” J. Am. Ceram. Soc., 79 [2] 417–24 (1996). 12C. G. Levi, J. Y. Yang, B. J. Dalgleish, F. W. Zok, and A. G. Evans, “Processing and Performance of an All-Oxide Ceramic Composite,” J. Am. Ceram. Soc., 81 [8] 2077–86 (1998). 13S. G. Steel, L. P. Zawada, and S. Mall, “Fatigue Behavior of a Nextel 720/Alumina Composite at Room and Elevated Temperature,” Ceram. Eng. Sci. Proc., 22 [3] 695–702 (2001). 14M. K. Cinibulk, K. A. Keller, T. Mah, and T. A. Parthasarathy, “Nextel 610 and 650 Fiber Reinforced Porous Alumina–YAG Matrix Composites,” Ceram. Eng. Sci. Proc., 22 [3] 677–86 (2001). 15M. H. Jaskowiak, S. I. Eldridge, J. B. Hurst, and J. A. Setlock, “Interfacial Coatings for Sapphire/Al2O3”; p. 84 in HITEMP Review. NASA Conference Publication 10082, 1991. 16L. U. J. T. Ogbuji, “A Porous Oxidation-Resistant Fiber Coating for CMC Interphase,” Ceram. Eng. Sci. Proc., 16 [4] 497–505 (1995). 17L. U. J. T. Ogbuji, “Evaluation of a Porous Fiber Coating in SiC–Si3N4 Minicomposites,” J. Mater. Res., 12, 1287–96 (1997). 18O. Sudre, A. G. Razzell, L. Molliex, and M. Holmquist, “Alumina Single-Crystal Fiber Reinforced Alumina Matrix for Combustor Tiles,” Ceram. Eng. Sci. Proc., 19 [4] 273–80 (1998). 19E. Boakye, M. D. Petry, and R. S. Hay, “Porous Aluminum Oxide and Lanthanum Phosphate Fiber Coatings,” Ceram. Eng. Sci. Proc., 17 [4] 53–60 (1996). 20E. Boakye, R. S. Hay, and M. D. Petry, “Mixed Carbon-Aluminum Oxide Coatings from Aqueous Sols and Solutions,” Mater. Res. Soc. Symp. Proc., 432, 363–68 (1997). 21E. Boakye, R. S. Hay, M. D. Petry, and T. A. Parthasarathy, “Sol–Gel Synthesis of Zircon–Carbon Precursors and Coating of Nextel 720 Fiber Tows,” Ceram. Eng. Sci. Proc., 20 [3] 165–72 (1999). 22M. K. Cinibulk, T. A. Parthasarathy, K. A. Keller, and T. Mah, “Porous Rare-Earth Aluminate Fiber Coatings for Oxide–Oxide Composites,” Ceram. Eng. Sci. Proc., 21 [4] 219–28 (2000). 23T. A. Parthasarathy, E. E. Boakye, K. A. Keller, and R. S. Hay, “Evaluation of Porous Zirconia–Silica and Monazite Coatings using Nextel 720 Fiber-Reinforced Blackglas Minicomposites,” J. Am. Ceram. Soc., 84 [7] 1526–32 (2001). 24J. D. Sibold, R. L. Cook, K. Bader, and I. Reimanis, “Porous Hexaluminate Coatings for Oxide/Oxide Composites”; pp. 3–13 in Ceramic Transactions, Vol. 103, Advances in Ceramic Matrix Composites V. Edited by N. P. Bansal, J. P. Singh, and E. Ustundag. American Ceramic Society, Westerville, OH, 2001. 25M. Y. He and J. W. Hutchinson, “Crack Deflection at the Interface between Dissimilar Materials,” Int. J. Solids Struct., 25, 1053–67 (1989). 26T. A. Parthasarathy, T. Mah, and K. Keller, “Creep Mechanism of Polycrystalline Yttrium Aluminum Garnet,” J. Am. Ceram. Soc., 75 [7] 1756–59 (1992). 27G. S. Corman, “Creep of Yttrium Aluminium Garnet Single Crystals,” J. Mater. Sci. Lett., 12, 379–82 (1993). 28S. Karato, Z. Wang, and K. Fujino, “High-Temperature Creep of YttriumAluminium Garnet Single Crystals,” J. Mater. Sci., 29, 6458–62 (1994). 29H. S. Yoder and M. L. Keith, “Complete Substitution of Aluminum for Silicon: the System 3MnOAl2O3SiO2–3Y2O35Al2O3,” Am. Mineral., 7, 519 (1951). 30D. M. Wilson, “New High Temperature Oxide Fibers”; pp. 3–12 in High Temperature Ceramic Matrix Composites. Edited by W. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, Weinheim, Germany, 2001. 31M. K. Cinibulk, “Synthesis of Yttrium Aluminum Garnet from a Mixed-Metal Citrate Precursor,” J. Am. Ceram. Soc., 83 [5] 1276–78 (2000). 32R. S. Hay, “Sol–Gel Coating of Fiber Tows,” Ceram. Eng. Sci. Proc., 12 [7–8] 1064–74 (1991). 33K. A. Keller, T. Mah, E. E. Boakye, and T. A. Parthasarathy, “Gel-Casting and Reaction Bonding of Oxide–Oxide Minicomposites with Monazite Interphase,” Ceram. Eng. Sci. Proc., 21 [3] 525–34 (2000). 34K. A. Keller, T. A. Parthasarathy, T. Mah, M. K. Cinibulk, and E. E. Boakye, “Evaluation of Monazite Fiber Coatings in a Dense Matrix Composite,” Ceram. Eng. Sci. Proc., 20 [3] 451–61 (1999). November 2002 Porous Yttrium Aluminum Garnet Fiber Coatings for Oxide Composites 2709
2710 Journal of the American Ceramic Sociery-Cinibulk et al. Vol. 85. No. 11 SST. A. Parthasarathy, E. Boakye, M. K. Cinibulk, and M. D. Petry, "Fabrication "M. K. Cinibulk, J. R. Welch, and R S. Hay, " Transmission Electron Microscopy sites with Monazite and Hibonite as Specimen Preparation of Ceramic Coatings on Ceramic Fibers, Mater. Res. Soc. Interlayers,J. Am. Ceram Soc., 82 [12]3575-83(1999). roc,480,3-17(1997) Weibull Modulus of Ceramic Filamen Temn.S. Hay, E. Boakaye, and M. D. Petry, "Effects of Coating Deposition J.Am. Ceram.Soe,80102741-44(1997) (2000). IT. A. Parthasarathy ."Extraction of Weibull Parameters of Fiber Strength from V. A. Kramb, R, John, and L. P. Zawada, "Notched Fracture Behavior of an rd Deviations of Failure Loads and Fiber Diameters. J. Am. Oxide/Oxide Ceramic-Matrix Composite, " J. Am. Ceram. Soc.,82 [111 3087-96 K. Cinibulk, J.R. Welch, and R. S. Hay, "Preparation of Thin Sections of 2M. K. Cinibulk, K. A Keller, T, Mah, and T. A. Parthasarathy, "Nextel6 Coated Fibers for Characterization by Transmission Electron Microscopy, J.Am ber-Reinforced Alumini-YAG Porous Matrix Composites, Ceram. Eng. So Ceran.Soc,79[92481-84(1996 Poc,23[3629-36(2002)
35T. A. Parthasarathy, E. Boakye, M. K. Cinibulk, and M. D. Petry, “Fabrication and Testing of Oxide/Oxide Microcomposites with Monazite and Hibonite as Interlayers,” J. Am. Ceram. Soc., 82 [12] 3575–83 (1999). 36M. D. Petry, T. Mah, and R. J. Kerans, “Validity of Using Average Diameter for Determination of Tensile Strength and Weibull Modulus of Ceramic Filaments,” J. Am. Ceram. Soc., 80 [10] 2741–44 (1997). 37T. A. Parthasarathy, “Extraction of Weibull Parameters of Fiber Strength from Means and Standard Deviations of Failure Loads and Fiber Diameters,” J. Am. Ceram. Soc., 84 [3] 588–92 (2001). 38M. K. Cinibulk, J. R. Welch, and R. S. Hay, “Preparation of Thin Sections of Coated Fibers for Characterization by Transmission Electron Microscopy,” J. Am. Ceram. Soc., 79 [9] 2481–84 (1996). 39M. K. Cinibulk, J. R. Welch, and R. S. Hay, “Transmission Electron Microscopy Specimen Preparation of Ceramic Coatings on Ceramic Fibers,” Mater. Res. Soc. Symp. Proc., 480, 3–17 (1997). 40R. S. Hay, E. Boakaye, and M. D. Petry, “Effects of Coating Deposition Temperature on Monazite Coated Fiber,” J. Eur. Ceram. Soc., 20 [5] 589 –97 (2000). 41V. A. Kramb, R. John, and L. P. Zawada, “Notched Fracture Behavior of an Oxide/Oxide Ceramic-Matrix Composite,” J. Am. Ceram. Soc., 82 [11] 3087–96 (1999). 42M. K. Cinibulk, K. A. Keller, T. Mah, and T. A. Parthasarathy, “Nextel™ 610 Fiber-Reinforced Alumina-YAG Porous Matrix Composites,” Ceram. Eng. Sci. Proc., 23 [3] 629–36 (2002). 2710 Journal of the American Ceramic Society—Cinibulk et al. Vol. 85, No. 11