1 tnt Cera So,861017340(2003 urna Processing and Properties of a Porous Oxide Matrix Composite Reinforced with Continuous Oxide Fibers lagnus G. Holmquist.and Fred F, Lange* s Materials Department, University of California, Santa Barbara, California 93106 Volvo Aero Corporation, 461 81 Trollhattan, Sweden a process to manufacture porous oxide matrix/polycrystalline environment, such as nitrogen oxides, carbon monoxide. and xide fiber composites was developed and evaluated. The burned hydrocarbons. method uses infiltration of fiber cloths with an aqueous slurry Most CFCCs that are commercially available are based on SiC of mullite/alumina powders to make prepregs. By careful fibers, with either oxide or non-oxide matrixes, and interphases manipulation of the interparticle pair potential in the slurry, a consisting of carbon, BN, SiC or combinations thereof. The consolidated slurry with a high particle density is produce interphases are designed to provide a crack-deflecting layer be- with a sufficiently low viscosity to allow efficient infiltration of tween the matrix and fibers that prevents matrix cracks from the fiber tows. Vibration-assisted infiltration of stacked cloth extending through the fibers, thus allowing crack bridging to occur prepregs in combination with a simple vacuum bag technique on matrix cracking enabling damage tolerance via notch insensi produced composites with homogeneous microstructures. The tivity. .5 SiC fiber based composites have attractive high method has the additional advantage of allowing complex temperature prope y sensiti, and high thermal conductivity shapes to be made. Subsequent infiltration of the powder es such as creep resistance, microstructura tability, high ten mixture with an alumina precursor was made to strengthen However, the oxidation sensitivity of the crack-deflecting inter phase will cause embrittlement of the composite after service at thermal stability and showed linear shrinkage of 0.9% on heat high temperature for long times. Embrittlement is most severe with treatment at 1200 C Mechanical properties were evaluated in lexural testing in a manner that precluded interlaminar shear cyclic loading beyond the proportional limit because oxygen that failure before failure via the tensile stresses. It was shown that penetrates via the matrix cracks will react with the interphase and the composite produced by this method was comparable to the fibers. This effect is most pronounced for carbon coatings, porous oxide matrix composites manufactured by other pr but the introduction of BN coatings and boron additives has cesses using the same fibers(N610 and N720). The ratio of oxidation products (liquid boron oxide) help in healing matrix improved the situation in oxidizing environments, where BN notch strength to unnotch strength for a crack to width ratio of cracks. However, in wet environments the problem persists since 0.5 was 0. 7-0.9, indicating moderate notch sensitivity. Inter- the boron oxidation products volatilize as boron hydroxides. . To laminar shear strength, which is dominated by matrix avoid degradation in oxidizing (especially wet) environments. strength, changed from 7 to 12 MPa for matrix porosity structural design strategies therefore usually require that the structure did not change after aging at 1200 C for 100 h Heat stresses remain below the matrix cracking stress. End-user treatment at 1300.C for 100 h reduced the strength for the experience indicates that stress excursions above the matrix N610 and N720 composites by 35% and 20%, respectively, and cracking stress is very difficult to avoid, and thus local increased their brittle nature embrittlement will be one of the dominant life-limiting phe nomena of non-oxide composites. These shortcomings have romoted the development of environmentally stable all-oxide I. Introduction composites, i.e., materials where all constituents(fiber, inter hase, and matrix)are oxides. 3.12, ONTINUoUS fiber ceramic composites( CFCCs) have attracted Two approaches have been used to develop damage-tolerant interest for a variety of high-temperature thermostructural all-oxide composites. The more traditional approach requires a applications in gas turbine engines, rocket engines, heat ex- crack-deflecting interface between the matrix and fibers. This can be changers, and hot filters. The reason is that they offer achieved by adding an interphase which either forms a crack combined with an inelastic deformation behavior rendering them toughness (e. g, "cleavable"oxides 6 or a porous layer7-19),or forms a gap between fiber and matrix(fugitive coating) walls. By taking advantage of the CFCCs ability to operate The use of a porous matrix to isolate fibers from matrix cracks at high temperatures with reduced need for cooling air, it is is a second, more recent approach for developing damage-tolerant the efficiency an nd also control the coml tion process to minimize formation of species harmful to the composites. In this approach, the crack does not have a contin- uous front, but, instead, the matrix is held together by grain pairs. ,23 Matrix failure occurs by the sequential failure of grain bers are isolated from the stress singula E, Lara-Curzio-comntributing editor crack because the matrix is not sufficiently continuous to support a crack. There are several examples of CFCCs which rely on a porous matrix for damage tolerance. 3.21 24-27 The failure mecl anisms have been examined in some detail. 22-24.28-3 On loading Manuscript No. 1875004 Received August 22, 2001: approved November 8. 2002. ice, DAAG55-98-1-0 the matrix experiences continuous microcracking during loading M. G.H. thanks the Hans Werthen Foundation for financial support and appears to have completely disintegrated at the onset of fiber failure. Contrary to the conventional weak interface CFCCs where irent affiliation: Advanced Engineering. SAAB, 461 80 Trollhattan, Sweden. the fibers slide out of the matrix. leaving distinct holes, when niversity of California. fibers fail in a porous matrix. they release a large volume of 1733
Journal of the American Ceramic Society-Holmquist and lange disintegrated matrix in the form of powder. Thus, the ability of the mullite grains and have higher creep resistance and high rous matrix to isolate the fibers from matrix cracks will allow temperature stability, but lower room temperature strength, com the fibers to fail in a manner similar to what is seen for di pared with N610 fibers, which are high purity (99%o) polycry bundles.The high failure strain of the fiber bundle will also talline a-alumina become the failure strain of the composite. The mullite powder used was MU-107 (Showa Denko KK. Two methods have been developed at University of California Tokyo, Japan), which has a mean particle size of-I um, a particl at Santa Barbara(UCSB) to produce porous matrix composites. size distribution of 0.5-2.5 um, and a BET surface area of 7 The first uses pressure filtration to pack particles around fibers 2/g. It has a chemistry of 75.5% AL,O, and 24% SiO,(by within a preform. The powder surrounding the fibers is then weight) with only trace amounts of TiO2. Fe=O3, and Na2O strengthened by the cyclic infiltration and pyrolysis of a precur sor.21..25.2 For this method, the slurry is formulated so that the pan)was selected as the alumina powder and has a mean particle particles are repulsive to themselves and also to the fibers. Mullite size of -0.2 Hm, a more narrow particle size distribution (0. 1-0.3 has been chosen as the matrix material because of its lack of um), and a BET surface area of 10.6 m /g(manufacturers data densification at temperatures below -1300oC. Levi et al. Its chemistry is essentially pure a-Al, O,(99.995%). An AL,O reported that alumina powder with a much smaller particle size precursor, aluminum( Ill) sec-butoxide, C12H27O, Al( Gelest, Inc (-200 nm)could be added to the mullite powder to aid in Tullytown, PA), was used to strengthen the matrix. The precursor strengthening the powder matrix. At processing temperatures is 95% pure and has a yield of 4% of Al]O, by volume. Before around 1200 C(chosen to avoid degradation of the fibers)the infiltration the precursor was diluted with 25%(by volume) of alumina sinters to form bridges between the larger mullite particles sec-butyl alcohol(Sigma-Aldrich, Milwaukee, wI) and between the mullite particles and the fibers. In this case, the 1200C, is sufficient (20.70)to prevent shrinkage of the mixed ratio was selected to achieve high packing densities of the powder powder matrix. body and low shrinkage in the following sintering step. Lower The second method was first introduced by Haslam et al alumina content would give fewer sintering necks between mullite Instead of packing the particles around the fibers via pressure rticles, whereas mixtures with higher alumina content will filtration, the powder is first consolidated to a very high volume densify relatively fast above 1200C. Tetraethylammonium hy raction and then infiltrated into the fiber preform via vibration- droxide (TEA-OH)was used to maintain the pH above I assisted flow. As detailed elsewhere. the initial slurry must be allowing electrostatic repulsive interactions to develop between formulated so that the particles are weakly attracted to one another. the oxide particles, A 2 wt% amount(relative to the solids) of The formulation of the the development of a short-range repulsive particle potentialthat was added to induce a steric dispersing effect. The PEG-silane pressure filtration, thus allowing the network to retain its interpar--M-OH (M- metal atom)surface sites, y reacting with the prevents the particles from being pushed into contact during molecules chem-absorbed to the particles ticle potential in the consolidated state agglomerates was promoted by ultrasonic agitation for 5 min. The In the present investigation we used a powder slurry with a slurry was placed on a mechanical roller for 12 h, and then special interparticle pair potential (weakly attractive produced by a tetramethylammonium nitrate(TMA-N) salt was added (25M)to short-range repulsive)that allows a powder compact, which has form weakly attractive pair potentials between the particles. As described elsewhere, TMA counter ions aid in producing a The preconsolidated slurry with a relative density of 0.54 was weakly attractive particle network. The network can be packed on a fiber cloth, and since it shows shear rate thinning, to a high density via pressure filtration, and the so-formed ibration reduces the viscosity and allows rapid and effi consolidated body can subsequently be fluidized again via vibra- cient intrusion into the fiber tows. The technique to make c tion. The slurry was consolidated by pressure filtration at 4 MPa to posites, called VibroIntrusion, has been described elsewhere. 2 form disk-shaped bodies that were stored in sealed plastic bags The vibration of the preconsolidated slurry was conducted between The consolidation pressure of 4 MPa was lower than the critical could be frozen and stored. Once thawed they were flexible and which would obviate fluidization after consolidation, a>o contact astic sheets to avoid ev uld be bent, cut, and formed much like an epoxy/fiber prepreg Prepregs could be stacked and formed into cox?) Drying only x geometries the consolidated bodies was determined to 54+ 1%0 like T-joints, doubly curved shapes. and tubes The consolidated powder compact was fluidized by subjecting it causes minimal shrinkage since a high dry content could be to mechanical vibration Fiber cloths (cut to -60 X 60 mm")were reached in the matrix slurry. Subsequent precursor impregnation/ put in separate plastic bags, and an excess of the preconsolidated pyrolysis cycles were done to strengthen the matrix. slurry was dispensed to both sides of each fiber cloth. Assisted by This new method to manufacture CCCs produced material vibration, the slurry was manually rolled across the surface of th with similar matrixes as a previous route fiber cloth with a piece of aluminum rod until the cloth was fully between this new method and the older method to produce porous infiltrated. Since the preconsolidated slurry exhibits shear-rate thinning, the vibration reduces the viscosity difference is that the mullite/alumina volume ratio, 70/30 used intrusion of the particles into the fiber cloth a and allows rapid matrix CFCCs will be made throughout the text. The main here, is lower than that(80/20) used earlier. The purpose was to Prepregs were frozen to aid removal from the plastic bag, or show that the process produces composite material with similar they could be stored for later use. To produce the composite, properties but with the added advantage of allowing complex frozen prepregs were removed from the plastic bags and stacked pes to be made on top of each other. The pile of prepregs was packed in a plastic bag, evacuated, and placed between two flat steel plates using two spacer bars to fix the thickness, The number of prepregs was IL. Experimental Procedure chosen to give the desired fiber volume fraction of the composite (in this case 13 layers of prepregs and a nominal thickness of 3. 18 (I) Materials and Composite Processing mm). After thawing, the assembled layers were put on a vibrating Reinforcement fibers used in this work were Nextel 610 M and table and pressed lightly to cause the preconsolidated slurry to Nextel 720M(N610 and N720, 3M Corp. St. Paul MN)woven flow and complete the infiltration. To remove trapped air as much into eig harness satin fabrics. The tows in the fabric contain as possible(which later on could give rise to large-scale porosity ) -400 filaments with diameters between 10 and 12 um. N720 the vacuum was kept on a level of -10 torr during the vibration fibers are composed of a mixture of sub-micrometer alumina and step(-5 min). It should be noted that the so-formed green ceramic
Dctober 2003 Porous Oxide Matrix Composite Reinforced with Oxide Fibers 1735 matrix composite(CMC) body is very flexible and could be shaped much like an epoxy/fiber prepreg After shaping, the water was removed from the powder matrix by drying at 70C. An initial sintering treatment was done at 900C for 2 h to promote the development of alumina bridges between the mullite network. The composites were subsequently npregnated with the alumina precursor solution under vacuum The impregnation step was performed in a dry nitrogen atmo sphere to prevent premature gelation of the precursor due to atmospheric water vapor. The composites were left in the precur sor solution for 2 h at atmospheric pressure and transferred into ammoniated water(pH 10) to gel the precursor throughout the body and prevent it from redistributing to the surface during the evaporation of the solvent. After 4 h the composites were removed, dried, and heated to 900C to pyrolyze the precursor This was repeated eight times and following the last cycle, the composites were given a final sintering treatment at 1200C for (b)W 2 h, which served to crystallize the precursor to the corundum(a) structure. In this way, the strength of the connection between mullite particles could be increased without any shrinkage of the mullite network taking place Knowing the volume fraction of fibers per unit area of cloth (data obtained from the manufacturer). the volume fraction of fibers in the composite was calculated by measuring the volume of the composite and counting the number of fiber layers in each specimen. The porosity of the composites was measured using the Archimedes technique Studies of the shrinkage of the matrix slurry were done rately by casting thin rods (-2 2X 10 mm )of slurry on (c) h different heat treatments. For each step, the specimens were treate similar to the composites (heated at 5C/min up to the mperature, held for 2 h, and then cooled down at Fig. I. Test bar geometries used for(a) in-plane bend testing for flexural strength and elastic modulus, (b) in-plane bend testing for notch sensitivity and work of fracture, and (c) interlaminar shear strength. (2) Mechanical Testing Fiber dominated composite properties were evaluated using an composites. A solution which also was used in this work)was mens(58 mm iong. 3 mm wide, 3.5 mm high) were processed as to place a rubber tween the loading pins and the specimen described above and were tested using a loading span of 53 mm urface to minimize stress concentrations 23.29 Interlaminar shear This configuration and loading mode precluded interlaminar shear strength, T, was calculated from the maximum load, Pmax, and test failure before failure via the tensile stresses on one surface. A bar dimensions using the equation for shear stress at the midplane servoelectric testing machine (Instron, Inc. Model 8562) with a of a flexural bar specimen described by beam theory high stiffness loading frame was used. A crosshead speed of 0.1 mm/min was used. Nylon rods were used as loading pins to reduce contact stress. Strain was calculated based on crosshead displace- ment and by correcting for the compliance in the load train. Since The tensile stress in the outer fibers in three-point bending is given the mode of testing is not pure tension, and because the matrix of b the composites are known to continually fracture during loading" which changes the modulus of the material during testing, the results of these tests should be considered as qualitative rather than quantitative. Although qualitative, the results will serve to com pare the different specimens in this stud sessed using the edge. Thus, the midplane shear stress to maximum tensile stress(T/o)is The in-plane notch sensitivity was as given by notched specimens shown in Fig. I(b), 58 mm long and 3 mm wide and 7 mm high. A notch with a length ao =3.5 mm(nominally half of the test bar height, a/w=0.50 + 0.002 mm)and a width section notch strength was compared with the unnotched strength To ensure failure by delamination rather than a tensile to assess the degree of notch sensitivity. Calculations of the energy failure originating from the surface, t (s)to thickness ratio, equired to break the specimens were also conducted to further s/t, is kept small. Failure modes of the bars were examined in characterize the work required for fracture. This was done by difterent microscope Fracture surfaces and the microstructure of the composites were tudied by optical microscopy and scanning electron microscopy I. Results and Discussion Interlaminar shear strength(a matrix dominated property fo 1) Composite Matrix and Composite Characteristics composites) was determined using a short beam shear tes The results from the sintering of pure matrix rods are n in Fig. I(c). The specimens were 30 mm long, 5 mm wide, ummarized in Fig. 2. During dry mean linear shrinkage of 3 mm high. and the loading span was 15 mm. It has been -1. 2% was observed. The ad change on sintering at erved that local stress concentrations due to the loading pins temperatures up to 1200C is-0.9%. Since the processing can give premature failure at low loads in porous oxide matrix temperature is of this order and the observed shrinkage is small
1736 Journal of the American Ceramic Sociery-Holmquisf and Lange ol Drying shrinkage yield: yp =0.03 1000 1500 Number of precursor impregnations, N Fig. 2. Linear shrinkage after heat treatments for 2 h at various temper. Fig 3. Matrix porosity for the composites measured initially, after three atures of cast mullite/alumina (70/30) powder slurry. The volume fraction and eight precursor impregnation and pyrolysis cycles, ((" indicates of solids in the slurry was 54- 1%e porosity as derived trom Eq- (4) in text. the results indicate that the chosen composition will produce a and 6 14 after eight impregnations, ared with 5.67 for the microstructurally stable matrix. However, at 1400 C the shrinkage N720 fiber. This suggests that the n720 fiber and the matrix increases more rapidly, suggesting reduced long-term stability of have similar coefficients of thermal expansion (CTEs) and the the material at temperatures above 1200C. These measurements residual stresses might be very small. The composites containing are in agreement with observations made by others made for the N610 fibers had similar crack patterns, strongly suggesting that similar materials. 2 the cracks were caused by constrained drying/densification shrink A summary of the composite panels and their fiber content and age and not differential CtE porosity is given in Table L. Fiber volume fraction did not vary a between the panels, with values Ve "40%6-42%. The porosity levels were less uniform; the first two manufactured panels(A and B)had an initial matrix porosity of -51.5%, whereas the last two panels(C and D) had a matrix porosity of -46.5%(Fig. 3).This was attributed to a processing improvement made in the vacuum bagging step: the vacuum level was reduced from -300 to-10 pyrolysis cycles is shot/ rosity with the following impregnation/ torr. The change in in Fig 3. Assuming that all the available voids in the composite are filled in each impregnation cycle, the remaining porosity Pm, after N cycles should be given by where p'm is the initial matrix porosity and yn is the volume yield of the precursor solution(measured to 3%), Equation (4) is plotted in Fig. 3 and agrees well with measurements of matrix porosity after three impregnations. The decrease in measured porosity after eight impregnations is lower than the calculations predict, suggest- ng the formation of closed pores that would prevent subsequent precursor infiltration Micrographs of the composite structure reveal that the tows were well-infiltrated and only a few large-scale pores produce from trapped air were evident, as shown in Fig 4. As expected, the large pores were more frequent in panels A and B. Cracklike flaws, (bE erpendicular to the fibers and with regular spacing, were also bserved. These were more than likely due to the constraint the densification 8.4 The matrix had an Al, O /SiO, weight ratio ot 5.4I before the multiple impregnations with the AL,O, precursor Table L. Summary of Composite Panels Panel 100pm N720 42.1 23.2 42.0 N720 NolO 38 V,= fiher volume fraction, p= composite porosity. pan Fig. 4. Microstructural views of cross sections of N720 composite HatnA po panels: (a) as processed and (b) after 1300C/( 100 h) heat treatment
October 2003 Porous Oxide Matrix Composite Reinforced with Oxide Fibers 1737 15 Matrix porosity 38. 2%, 1200"C/100 o"D◆ 的的布∽ DN101200°cPn rix porosity 38. 2%, 1200C/ h 9DN6101200c100h Matrix porosity 43. 2%, 1200 C/100h Matrix porosity 43.2%, 1200C/2h 060810 Matrⅸ porosity(% on in maximum shear stress with matrix porosity in lition and after aging: open symbols, as-processed spec- Fig, 5. shear behavior for all-oxide composites with two bols, after aging at 1200 C/( 100 h) different osity levels, An increase in shear strength is seen after that high matrix strength can be obtained by strengthening the network without decreasing porosity EM observations showed no signif in the ma structure after heat treatment at 1200C for 100 h. However, aging (3) In-Plane Flexure Testing noted an evolution of f ws caused by matrix densification (Fig. 4(b). Others reported matrix Table II reports the in-plane mechanical properties for un- densification with associated matrix damage at 1200 C for longer notched composite specimens as processed and after aging. The periods(1000 h) in similar materials. strength of the N720 specimens was >170 MPa, and the N61o specimens was >280 MPa in as-processed condition, which is consistent with data presented by Levi, Zok, and co-workers" (2) Interlaminar Shear Strength for the previously described(see Section Ill( I) method of pro- Typical stress/displacement responses obtained for the inter- cessing porous mullite/alumina matrix composites. Figure 7 shows laminar shear tests are shown in Fig. 5. After an initial linear representative stress-strain curves for the composites ortion, a slight reduction in stiffness was observed after the maximum load was obtained, followed by a number of load drops In general the observed behavior is similar to the phenomenon of locations producing fiber brushes. The locations of fracture within sequential delamination failure. The failure mode was delani- individual tows also show a distribution(Fig. 8(a). These obser- nation in all tests. and the maximum interlaminar shear stresses vations show that the porous matrix is an efficient crack deflector calculated according to Eq. (3)are a measure of the matrix both within and between fiber tows. In more conventional CMCs rength. Interlaminar shear strength as a function of matrix with crack-deflecting matrix/fiber interfaces, one can observe porosity is shown in Fig. 6. These data suggest that the shea holes that contained fibers that fracture within the matrix. 7In rength increases with decreasing porosity. Heat treatment at composites with porous matrixes, no such holes can be observed strength. Although no matrix densification could be observed after which still are bonded to the fibers. 2. 24.2 The amount of fiber this heat treatment (see Section Ill(1), the results imply that the pull-out appears to be relatively uniform over the fracture surface, mullite/alumina matrix network has strengthened. The delamina- whereas others have observed more coplanar fracture near the tion stress measured for the least porous material (38 vol%) was edges of the test bar. This behavior has been explained by redistribution of the precursor(used to strengthen the matrix) in a composite with a more dense matrix(34 vol% and-12 MPa). the absence of a gelling step in the manufacturing process mullite/alumina composition of 80/20. This observation suggests specimens in air for 100 h at 1200 and 1300 C and then tested ag The composite used by Mattoni et al. had a matrix with a Table IL. Properties of In-Plane Unnotched Composite Three-Point Bend Tests Flexure strengt Elastic modulus Fiber Heat treatment Failure mode' N720 1200°C(2h) T N720 200°C/(2h) N720 1200°C/(100h) 1200°C/(100h) N720 1300C/(100h N720 300C/(100h 200°C/2h 85 N610 1200°C2h N 1200°c(100h) N6101200°C(100h N610 1300°C(100h) N60 300°C/100h 18 Failure mode; T, tensile: M, mixed mode (e. g, buckling, delamination)
1738 Journal of the American Ceramic Sociery-Holmqguist and Lange Vol. 86. No. 10 (a) N720.1200°c/00h N720.1200°c2h N720.1300c/10oh 300pm Nominal Flexure Strain (%) 0.1200c00h 10.1300c/100h (b 300pm Nominal Flexure Strain(%) Fig. 7. Flexural stress versus flexural strain plots for unnotched compos- Fig 8. SEM micrograph of fracture surfaces of N720 composites: (a)as ites reinforced with (a) N720 fibers and(b) N610 fibers, in as-processed processed, and (b) heat-treated at 1300C for 100 h. condition and after aging In-plane mechanical properties of the notched specimens are room temperature. The results are also summarized in Table lI and reported in Table Ill. Net-section strengths were evaluated by Fig. 7. There is no significant effect on either stiffness or strength assuming that the bar dimensions used to calculate the maximum after thermal exposure at 1200 C for 100 h. The fracture surfaces tensile stress at failure(Eq (2)do not include the notched portion are similar to those of the as-processed specimens, showing a of the bar; i.e., the length of the notch is subtracted from the height fibrous nature. However, the pull-out lengths of the broken fibers of the beam, and this value is used as the height used to determine appear to be somewhat shorter. Such a trend was also observed by Mattoni et al. 2 when the matrix strength was increased by obtained by comparing the net-section strengths of the notched repeated precursor impregnations to reduce the porosity. The specimens to the strength of the unnotched specimens(strength proposed explanation was that the reduced matrix porosity induced ratio, Table IID). The strength ratio was -0.75-0.90 for the N720 higher stress concentrations around fiber breaks and thereby the composites and-0.70-0.8.5 for the N610 composites, indicating failure probability of adjacent fibers was increased. A similar moderate notch sensitivity for the material. These values are in as-processed specimens tended to fail in a mixed mode rather than oth other processes. ,0 Load-displacement data in Fig. 9 shows that in a pure tensile mode the fracture occurs in a stable manor after the load maximum has Despite the higher thermal stability of the N720 fibers, these opposites exhibited a significantly lower strength at room tem- perature compared with those fabricated with N6l0 fibers: the lower strength of the N720 fibers is expected to be the cause of the Table Ill. Properties of In-Plane Notched Composite lower strength composite. Three- Point Bend Test Increasing the heat treatment to 1300 C/(100 h) had a more Net-section posites and -35% for N610 composites, The degree of fiber Heat treatment pull-out was significantly reduced. as shown in Fig. 8(b), A small AN7201200°C/2h) 1666 0.90 increase in stiffness could be observed for the N720 composite BN7201200°C12h) The N610 fibers are known to loose strength rapidly at high temperatures, which might explain the decrease in composite BN7201200°C/100h)18190.9 16 strength. Also the N720 composites are affected, but to less extent CN6101200°C/(2h) because of the higher thermal stability of N720 fibers compared CN6101200°C/100h)254 with N610 fibers. The heat treatment will also give a denser and DN6101200°C/100h) stronger matrix as described above aW=0.50±002
October 2003 Porous Oxide Matrix Composite Reinforced with Oxide Fibers 1739 ' National Materials Advisory Board, Committee on Advanced Fibers for High Temperature Ceramic Composites, Publication NMAB-494. National Academy Press, 110 T.J. McMahon, "Advanced Hot Gas Filter Development, "Ceram. Eng. Sei, Pre WOF 3601 J/m2 211347-57(2000 R, Warren(Ed). Comprehensive Composite Materials. Elsevier, Amsterdam, N610,1200°c100h Netherlands, 2000 WOF= 2542 J/m? Price.""Ceramic Stationary Gas Turbine Development Program-Fourth Annual 201200c/100h Summary, "AnN. Soc. Mech. Eng /Pap, 1. 97-GT-317(Is WOF= 1819J'm 'D. Filsinger, S. Munz, A. Schulz, S. Wittig, and G. Andrees, "Experimental Assessment of Fiber Reinforced Ceramics for Combustor Walls. Am. Soc. Mech N720.@C/2h Eng.Pap9GT·154(1997 K. Nishio, K.L. Igashira, K. Take, and T. Suemitsu, " Development of a 30 Eng,Pap198GT-104(1998 quist, L. Molliex, and O. Sudre, "Oxide/Oxide Ceramic 10 Matrix Composites in Gas Turbine Combustors, " Am. Soc. Mech. Eng, /Pap. 1, 98GT-301998) oW. H. Glime and J. D. Cawley."Stress Concentration Due to Fiber-Matni Fusion in Ceramic-Matrix Composites, " J, Am. Ceram. Soc., $1 [101 2597-604 Displacement(mm R. H. Jones, C. H. Henager Jr.C. A. Lewinsohn, and C. F, windisch. Jr, Fig. 9. Load versus displacement plots for notched composites reinforced ress.Corrosion Cracking of Silicon Carbide Fiber/Silicon Carbide Composite, with N720 or N610 fibers. in as-processed condition and after aging. JAm. Ceran.Soc,838l1999-2005(2000 R, E. Tressler. "Recent Developments in Fibers and D. B. Marshall, J. B. Davis, P. E. D. Morgan, and J Materials for Damage-Tolerant Oxide Composites, Key 127-131 d reached and, consequently, the area uI curve can be 27-36(1997) as a measure of the work of fractur WOF was measured to-1 6-2. 0 kJ/m for the N720 mEM1641631519 ing values for the N610 composite were.5-3.6 kJ/m". A slight Fiber Coatings for Oxide-Oxide CFCC, "Ceram Eng. Sci Prac. 18 131 279-86 decrease in WOF could be seen after heat treatment at 1200 C for 100 h. as shown in Table III M. K. Cinibulk, "Hexaluminate as a Cleavable Fiber- Matrix Interphase thesis, Texture Development, and Phase Compatibility. "J. Eur. Ceram. Soc., L1M Holmquist,RLundberg.OSudre,AGRazzell.L.Molliex,J,Benoit,and IV. Conclusions J. Adlerbom,"Alumina/Alumina Composite with a Porous Zirconia Interphase ssing. Properties and Component Testing. "J. Eur. Ceram. Soc. process to manufacture porous oxide mat oxide fiber composites was developed and evaluated. The method IO Sudre, A. G Razzell, L, Molliex, and M. Holmquist, "Alumina SingleCryst uses a preconsolidated slurry with a very high volume fraction of Fibre Reinforced Alumina Matrix for Combustor Tiles, Cera. Eng. Sci. Proc. powder to infiltrate fiber cloths. These infiltrated fiber cloths can M. J. O Brien and B. W. Sheldon,"Porous Alumina Coating with Tailored be frozen and used latter to fabricate engineering shapes. Mechan- Fracture Resistance for Alumina Composites. "J. Am. Ceram, Soc., 82| 3567- ical property measurements suggest that the processing method (1999) used here was comparable to porous oxide matrix composites K. A. Keller, T-1. Mah, T. A. Parthasarathy, and C. M. Cooke, "Fugitive Interfacial Carbon Coatings for Oxide/Oxide Composites. "J. Am Cerum. Soc., 8[2] manufactured by other processes using the same fibers 329-36(2000) In-plane flexural tests of composites based on alumina N610 W.-C. Tu, F, F. Lange. and A. G. Evans, " Concept for a Damage Tolerant and mullite/alumina N720 fibers showed nonbrittle fracture behav Ceramic Composite with "Strong"Interfaces, J. Am. Ceran. Soc. 79121417-24 ior and strengths of >280 and >170 MPa, respectively, and moderate notch sensitivity. Interlaminar shear strength, which is F. F. Lange, C, G. Levi, and F, w. Zok, " Processing Fiber Reinforced Ceramics with Porous Matrices": Ch. 14 in Comprehensive Composite Materials. Edited by R, dominated by the porous matrix, ranged between 7 and 12 MPa for Warren. Elsevier, Amsterdam, Netherlands. 2000 matrix porosity ranging from 38% to 43%o, respectively J. Haslam, K E Berroth, and F F. Lange, "Processing and Properties of an The composites possessed good thermal stability: the micro- All-Oxide Composite with a Porous Matrix, J. Eur. Ceram. Soc., 20, 607-18 structure was stable after aging at 1200 C for 100 h, showing no 2C, G, Levi, F. W, Zok, J -Y. Yang, M. Mattoni, and J. P. A. Lofvander visible signs of densification. This heat treatment was found to le Porous Matrices for All-Oxide Ceramic Compo slightly increase the interlaminar shear strength, which was ites." Z Metalled,9012]1037-47(1999 attributed to a strengthening of the matrix network, and was F. F. Lange, W. C. Tu, and A. G. Evans, U.S. Pat. No, 5856 252, 1999. Several ccompanied by a reduction in composite toughness. However DL. P. Zawada. "Longitudinal and Transthickness Tensile Behavi the 1200C/(100 h)treatment did not significantly change the Oxide/Oxide Composites, "Ceram Eng. Sci. Proc., 19 B31327-39(1998). 27R.A Jurf and S C. Butner,"Advances in Oxide- Oxide CMC, "An. Soc. Mech mposite strength or strain to failure. Heat treatment at 1300 C for 100 h reduced the strength for the N610 and N720 G. Levi, J, -Y. Yang, B J Dalgleish, F. w, Zok, and A, G. Evans, "Processin composites by 35% and 20%, respectively, and increased their and Performance of an All-Oxide Ceramic Composite. " J. Am. Ceram Soc., 81 [81 brittle nature Derived Alumina on the Mechanical Properties of a PorousMatrix. All-Oxide Ceramic Composite. "J Am, Ceram. Soc., in revie Acknowledgments A. Heathcote, X.Y. Gong, I -Y, Yang, U. Ramamurty, and F. w. Zok, " In-plane Mechanical Properties of an All-Oxide Ceramic Composite. " J Am Ceran We thank R. Harrysson for SEM work and Professor F. W. Zok, Dr. J -Y, Yang. Soc,82012721-30 A. V. Carelli, H, Fujita, J. Y. Yang, and F. w. Zok, "Effects of Thermal Aging aG. W. ranks and F. E. Lange. "Plastic Clay-like Flow Stress of Saturated References SM. Holmquist. T C. Radsick, O. Sudre, FF. Lange, and F W. Zok. " Fabrication and Testing of All-Oxide CFCC Tubes, to be submitted C, P Beesley, " The Application of CMCs in High Integrity Gas Turbine F. F. Lange, T. C. Radsick, and M. Holm nes." Key eng.Maer,127-131.165-74(1997) Control of Microstructure and Properties": pp. 587-99 Beyer, H. Knabe. and F. Strobel. "Development and Testing of C/SiC ternational Conference on High Temperature Ceramic Matrix Components for Liquid Rocket Propulsion Applications. " AlAA Pap,,99-2896 Edited by w. Krenkel, R. Naslain, and H. Schneider. Wiley-VCH, New York, (1999
1740 Journal of the American Ceramic Society-Holmquist and La Vol. 86. No, 10 D. M. Wilson S. L. Lieder, and D. C. Lueneberg, "Microstructure and High w.C. Tu and F. F. Lange, "Liquid Precursor Infiltration Processing of Powde Temperature Properties of Nextel 720 Fibers. Ceram. Eng. Sci. Proc., 16 [51 II, Fracture Toughness and Strength. "J. Am. Cermet. Soc.. 78 1121 3283-89(1995 xD.M. wilson, "Statistical Tensile Strength of Nextel 610 and 720 Fibres, 4D. C. C. Lam and F. F. Lange. "Microstructural Observations on Constrained M. Colic. G. Franks, M. Fisher, and F, F. Lange, "Chemisorption of Organo- Am. Ceram. Soc. 77 1711976-78(1994). a Periodic Array of Sapphire Fibers funetional Silanes on Silicon Nitride for Improved Aqueous Processing, J.Am J. B. Davis. A. Kristoflersson, E. Carlstrom, and W. J. Clegg. "Fabrication and sS. Klein, M. Fisher, G. Franks, M. Colic, and F. Lange, " Effect of the Crack Deflection in Ceramic Laminates with Porous Interlayers. "J. Am. Ceram. Sone. 831012369-74(2000 terparticle Pair Potential on the Rheological Behavior of Zirconia Powders: IL. Th Influence of Chem-Adsorbed Silanes. " J. Am. Ceram. Soc., in pr w.-C. Tu and F. F. Lange. "Liquid Precursor Infiltration Processing of Powder Tests Using Optical Fluorescence, "Prox. R. Sox. London, Ser. A, 453. 1881-901(1997). Kinetic Studies and Microstructure Development Composites: Localized Load- Sharing and Associated Size Effects. "Int. J. Solids 781213277-82(1995 Sra,34[212649-68(1997
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