E驅≈3S Journal of the European Ceramic Society 20(2000)607-618 Processing and properties of an all-oxide composite with a porous matrix .. Haslam, K.e. berroth*ff la Materials Department, University of California at Santa Barbara, Santa Barbara, CA 93106, US.A 18 August 199 Abstract Processing and mechanical properties of an all-oxide fiber composite with a porous matrix are presented here. The processing approach for an all-oxide composite was developed to be simple and involve one sintering process. The composite uses a porous matrix instead of riser coatings to deflect cracks from the fibers. a processing method involving recently developed methods for reshaping and forming saturated high-volume fraction(> 50 vol%) particle bodies was used to form the composite Good infiltra- tion of the woven fiber tows was obtained. Sintering in a pure HCl gas atmosphere was used to produce a porous matrix without shrinkage during processing. The sintering process also produced coarsening which makes the microstructure stable against densi- fication during use and thereby prevents forming cracklike voids and retains sufficient porosity for crack deflection. Measurements of interlaminar shear strength and strength of the composite show that composite produced by this processing method is compar- able to previous all-oxide materials produced using the oxide fibers used here. The mechanical properties are rationalized in terms of the features on the fracture surfaces. Disintegration of the matrix to allow energy dissipation during fracture was apparent and correlates with the measurements of the fracture toughness of the material. Moderate notch insensitivity was demonstrated with a net section strength in the presence of a notch being 700% of the unnotched strength. c 2000 Elsevier Science Ltd. All rights reserved Keywords: Aluminosilicate fibres; Composites; Mechanical properties; Porosity: Sintering: ZrO2 matrix 1. Introduction matrix. A fiber within a good CMC is only expected to break when the applied load exceeds its strength An important property of any ceramic matrix com- In the late 1960s Phillips recognized that brittle, but posite(CMC) is that its strength should be relatively strong fibers could be isolated from one another within insensitive to the presence of notches. If the riders a brittle matrix by providing a path for cracks prova within a CMC are effective, the strength of a body with gating through the matrix to bypass the fibers. a"weak a notch(or hole)of any size or shape will be the same as interface between the matrix and fiber provides the path the unnotched strength of a body with same net (or for crack deflection, thus allowing the crack to propa- reduced) cross-section. That is, for an ideal CMC, one gate along the fiber /matrix interface instead of through drill the fiber. As described by He and Huchinson, the con reducing the failure load other than the effect of redu- dition for crack deflection depends on the ratio of the cing its cross sectional critical strain energy release rate for the interface and Since the failure strain of a strong fiber is generally fiber, and the elastic properties of the two materials. For much larger than a dense matrix, cracks generally first a few fiber/matrix combinations, the crack defecting extend within the matrix. In terms of crack extension, interface needs no special processing conditions. For notch insensitivity requires that the fibers must be iso- example. the carbon fibers in the CMCs produced by lated from the very high stress field of a crack within the Phillips et al. did not bond to the glass matrix. For most other CMCs. the fibers must be coated with either carbon or boron nitride films to achieve a crack deflecting inter- Researcher. High Performance face. Not only do fiber coatings introduce cost and proces- Ceramic Section, Swiss Federal Laboratories for Materials Testing sing complexity, but they are not stable in oxidizing and Research. EMPA. Dube environments and they can cause composite embrittlement 0955-2219/00/S- see front matter o 2000 Elsevier Science Ltd. All rights reserved PII:S0955-2219(99)00259-9
Processing and properties of an all-oxide composite with a porous matrix J.J. Haslam, K.E. Berroth*, F.F. Lange Materials Department, University of California at Santa Barbara, Santa Barbara, CA 93106, USA Accepted 18 August 1999 Abstract Processing and mechanical properties of an all-oxide ®ber composite with a porous matrix are presented here. The processing approach for an all-oxide composite was developed to be simple and involve one sintering process. The composite uses a porous matrix instead of riser coatings to de¯ect cracks from the ®bers. A processing method involving recently developed methods for reshaping and forming saturated high-volume fraction (>50 vol%) particle bodies was used to form the composite. Good in®ltration of the woven ®ber tows was obtained. Sintering in a pure HCl gas atmosphere was used to produce a porous matrix without shrinkage during processing. The sintering process also produced coarsening which makes the microstructure stable against densi- ®cation during use and thereby prevents forming cracklike voids and retains sucient porosity for crack de¯ection. Measurements of interlaminar shear strength and strength of the composite show that composite produced by this processing method is comparable to previous all-oxide materials produced using the oxide ®bers used here. The mechanical properties are rationalized in terms of the features on the fracture surfaces. Disintegration of the matrix to allow energy dissipation during fracture was apparent and correlates with the measurements of the fracture toughness of the material. Moderate notch insensitivity was demonstrated with a net section strength in the presence of a notch being 700% of the unnotched strength. # 2000 Elsevier Science Ltd. All rights reserved. Keywords: Aluminosilicate ®bres; Composites; Mechanical properties; Porosity; Sintering; ZrO2 matrix 1. Introduction An important property of any ceramic matrix composite (CMC) is that its strength should be relatively insensitive to the presence of notches. If the riders within a CMC are eective, the strength of a body with a notch (or hole) of any size or shape will be the same as the unnotched strength of a body with same net (or reduced) cross-section. That is, for an ideal CMC, one should be able to drill a hole without signi®cantly reducing the failure load other than the eect of reducing its cross sectional area. Since the failure strain of a strong ®ber is generally much larger than a dense matrix, cracks generally ®rst extend within the matrix. In terms of crack extension, notch insensitivity requires that the ®bers must be isolated from the very high stress ®eld of a crack within the matrix. A ®ber within a good CMC is only expected to break when the applied load exceeds its strength. In the late 1960s Phillips1 recognized that brittle, but strong ®bers could be isolated from one another within a brittle matrix by providing a path for cracks propagating through the matrix to bypass the ®bers. A `weak' interface between the matrix and ®ber provides the path for crack de¯ection, thus allowing the crack to propagate along the ®ber/matrix interface instead of through the ®ber. As described by He and Huchinson,2 the condition for crack de¯ection depends on the ratio of the critical strain energy release rate for the interface and ®ber, and the elastic properties of the two materials. For a few ®ber/matrix combinations, the crack defecting interface needs no special processing conditions. For example. the carbon ®bers in the CMCs produced by Phillips et al. did not bond to the glass matrix. For most other CMCs, the ®bers must be coated with either carbon or boron nitride ®lms to achieve a crack de¯ecting interface. Not only do ®ber coatings introduce cost and processing complexity, but they are not stable in oxidizing environments and they can cause composite embrittlement. 0955-2219/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved. PII: S0955-2219(99)00259-9 Journal of the European Ceramic Society 20 (2000) 607±618 * Corresponding author. Visiting Researcher, High Performance Ceramic Section, Swiss Federal Laboratories for Materials Testing and Research, EMPA, DuÈbendorf, Switzerland
J. Haslam et al. Journal of the European Ceramic Society 20(2000)607-618 Cracking phenomena for the tensile loading of a uni within the fibers and matrix is identical. It is possible directional CMC containing crack deflecting interfaces that the failure strain(Em) of the porous matrix can be can be related to the composite's stress/strain behavior. equal or even larger than the failure strain of the fibers During the initial loading, the behavior is linear and (Ee). Applying Hook's Law, one can show that characterized by the combined elastic modulus of the fiber and matrix weighted by the appropriate volume Er= E (1) raction of each. As loading proceeds, matrix cracking initiates without fiber failure. Matrix cracking is char- where Emf and Em f are the failure stress and elastic acterized by a decreased slope of the stress-strain curve. modulus of the matrix(m) and fibers(f). Using proper Multiple matrix cracking generally occurs prior to the ties of a low density Al2O3 matrix material and Al2O3 initiation of fiber failure. Prior to and during fiber fail- fibers(e.g. 0m 200 MPa, o 2000 MPa, Emx40 GPa ure, the cracked matrix is held together by the fibers, and Er= 400 GPa) we can see that it is reasonable to be which now supports nearly all of the applied load. Fiber able to fabricate a porous matrix with a failure strain failure, and thus CMC failure occurs at a high strain that approaches that of a strong fiber. Therefore, in (0.5 to 1.0%), indicative of a strong fiber with a low tension, a large fraction of the strength and strain to elastic modulus. For many commercial and experi- failure of strong fibers can be achieved in a ceramic mental CMCs the stress for matrix cracking lies between composite that contains a porous matrix 40 and 100 MPa whereas composite failure(fiber fail- The second role of the porous matrix is to allow fibers ure)does not occur until the stresses exceed 150 to 300 to be isolated from cracks within the matrix In porous MPa. Thus, CMCs that have been developed over the materials the crack front can be non-continuous and last 25 years to contain crack deflecting interfaces can crack extension must occur by the continued breaking not only be relatively notch insensitive, but they can of the solid phase units, i.e. fracture has to be reinitiated also exhibit higher strains to failure relative to mono- in the solid phase within the high stress field of the pro- lithic ceramics(e.g Si3N4, with a mean tensile strength pagating crack. A comparable example of this fracture of 1000 MPa, has a strain to failure of 0.3%) mode is the extension of a crack within cloth. where the Approximately 5 years ago another type of CMC was fracture of each fiber is independent of the last to fail inadvertently discovered. 3.4 Unlike the CMCs with This mode of crack extension occurs in powder com- weak ' fiber matrix interfaces, the matrix and fibers are pacts that have been heated to produce necks between bonded together in these ' new'CMCs. The second touching particles. Observing the fr acture change is that the matrix in the 'new'CMCs is purpo- these very porous materials one can see that fracture(or sely made to be porous. Despite these two major chan- crack extension,)occurred by the breaking of grain ges, both of which are not taught by mechanics of pairs at grain boundaries. A continuous crack front conventional CMCs, the new CMC with well-bonded does not exist in these porous materials fiber/matrix interfaces and porous matrixes are notch The lack of a crack front in a porous matrix means insensitive. In addition, although not as high as the that embedded fibers never see an extending crack front conventional CMCs, their failure strain is larger than as the matrix fails. Fiber fracture within a very porous conventional monolithic ceramics. 5 matrix must initiate within the fiber itself. ie. from flaws The mode of failure of these new composites is dif- either on the surface or within the fiber, and not by the ferent from the older CMCs. The stress/strain behavior propagation of a crack within the matrix. Thus, fibers in of tensile specimens is nearly linear to failure, indicating a very porous matrix can fracture in the same manner as that both the matrix and fibers fail at about the same they do when they exist as a bundle, without a matrix. failure strain. Tensile failure can occur at 200 MPa: The high failure strain of the fibers becomes the failure CMngth is also relatively notch insensitive. 5 These strain of the composite because the matrix will have a oxide matrix and can be very stable in air to tempera- As detailed elsewhere, -5 one method to produce the tures where the fibers begin to degrade. It can be fiber composites described above is to pack particles expected that the processing of the new CMCs is much around the fibers within a fiber preform by pressure fil less complex and less costly. Eliminating fiber coatings is tration and then strengthen the porous matrix. In this a significant advantage in processing and reducing cost. method, a fiber preform(3-D weave, stacked layers of The porous matrix appears to play a critical role in cloth, etc. ) is mounted on a filter within a die cavity. A achieving a notch insensitive strength and a high failure pressure is exerted to a dispersed slurry to cause the strain. One role concerns the strain to failure. When a particles to stream though the preform to become trap- composite is loaded in tension, the fibers will support ped at the filter, and to build up a consolidated layer much of the load due to their much larger elastic mod- within the fiber preform. The slurry must be formulated ulus relative to the porous matrix. Although the fiber such that the particles are repulsive with respect to must carry the major portion of the stress, the strain themselves and the fibers. The particles must also be
Cracking phenomena for the tensile loading of a unidirectional CMC containing crack de¯ecting interfaces can be related to the composite's stress/strain behavior. During the initial loading, the behavior is linear and characterized by the combined elastic modulus of the ®ber and matrix weighted by the appropriate volume fraction of each. As loading proceeds, matrix cracking initiates without ®ber failure. Matrix cracking is characterized by a decreased slope of the stress±strain curve. Multiple matrix cracking generally occurs prior to the initiation of ®ber failure. Prior to and during ®ber failure, the cracked matrix is held together by the ®bers, which now supports nearly all of the applied load. Fiber failure, and thus CMC failure occurs at a high strain (0.5 to 1.0%), indicative of a strong ®ber with a low elastic modulus. For many commercial and experimental CMCs the stress for matrix cracking lies between 40 and 100 MPa whereas composite failure (®ber failure) does not occur until the stresses exceed 150 to 300 MPa. Thus, CMCs that have been developed over the last 25 years to contain crack de¯ecting interfaces can not only be relatively notch insensitive, but they can also exhibit higher strains to failure relative to monolithic ceramics (e.g. Si3N4, with a mean tensile strength of 1000 MPa, has a strain to failure of 0.3%). Approximately 5 years ago another type of CMC was inadvertently discovered.3,4 Unlike the CMCs with `weak' ®ber matrix interfaces, the matrix and ®bers are bonded together in these `new' CMCs. The second change is that the matrix in the `new' CMCs is purposely made to be porous. Despite these two major changes, both of which are not taught by mechanics of conventional CMCs, the new CMC with well-bonded ®ber/matrix interfaces and porous matrixes are notch insensitive. In addition, although not as high as the conventional CMCs, their failure strain is larger than conventional monolithic ceramics.5 The mode of failure of these new composites is different from the older CMCs.4 The stress/strain behavior of tensile specimens is nearly linear to failure, indicating that both the matrix and ®bers fail at about the same failure strain.5 Tensile failure can occur at 200 MPa; the strength is also relatively notch insensitive.5 These new CMCs can be processed with oxide ®bers in an oxide matrix and can be very stable in air to temperatures where the ®bers begin to degrade. It can be expected that the processing of the new CMCs is much less complex and less costly. Eliminating ®ber coatings is a signi®cant advantage in processing and reducing cost. The porous matrix appears to play a critical role in achieving a notch insensitive strength and a high failure strain. One role concerns the strain to failure. When a composite is loaded in tension, the ®bers will support much of the load due to their much larger elastic modulus relative to the porous matrix. Although the ®ber must carry the major portion of the stress, the strain within the ®bers and matrix is identical. It is possible that the failure strain ("m) of the porous matrix can be equal or even larger than the failure strain of the ®bers ("f). Applying Hook's Law, one can show that "f "m; or f Ef m "m ; 1 where "m;f and Em,f are the failure stress and elastic modulus of the matrix (m) and ®bers (f). Using properties of a low density Al2O3 matrix material and Al2O3 ®bers (e.g. m200 MPa, f2000 MPa, Em40 GPa and Ef 400 GPa) we can see that it is reasonable to be able to fabricate a porous matrix with a failure strain that approaches that of a strong ®ber. Therefore, in tension, a large fraction of the strength and strain to failure of strong ®bers can be achieved in a ceramic composite that contains a porous matrix. The second role of the porous matrix is to allow ®bers to be isolated from cracks within the matrix. In porous materials the crack front can be non-continuous and crack extension must occur by the continued breaking of the solid phase units, i.e. fracture has to be reinitiated in the solid phase within the high stress ®eld of the propagating crack. A comparable example of this fracture mode is the extension of a crack within cloth, where the fracture of each ®ber is independent of the last to fail. This mode of crack extension occurs in powder compacts that have been heated to produce necks between touching particles. Observing the fracture surface of these very porous materials one can see that fracture (or `crack extension') occurred by the breaking of grain pairs at grain boundaries.6 A continuous crack front does not exist in these porous materials. The lack of a crack front in a porous matrix means that embedded ®bers never see an extending crack front as the matrix fails. Fiber fracture within a very porous matrix must initiate within the ®ber itself. i.e. from ¯aws either on the surface or within the ®ber, and not by the propagation of a crack within the matrix. Thus, ®bers in a very porous matrix can fracture in the same manner as they do when they exist as a bundle, without a matrix. The high failure strain of the ®bers becomes the failure strain of the composite because the matrix will have a comparable strain to failure. As detailed elsewhere,3±5 one method to produce the ®ber composites described above is to pack particles around the ®bers within a ®ber preform by pressure ®ltration and then strengthen the porous matrix. In this method, a ®ber preform (3-D weave, stacked layers of cloth, etc.) is mounted on a ®lter within a die cavity. A pressure is exerted to a dispersed slurry to cause the particles to stream though the preform to become trapped at the ®lter, and to build up a consolidated layer within the ®ber preform. The slurry must be formulated such that the particles are repulsive with respect to themselves and the ®bers. The particles must also be 608 J.J. Haslam et al. / Journal of the European Ceramic Society 20 (2000) 607±618
J. Haslam et al. Journal of the European Ceramic Society 20(2000)607-618 much smaller than the fiber diameter to ensure good tows produced by the 3M Corporation( St Paul, MN) particle packing. 7 To avoid large cracklike voids from Each tow nominally contains 420 fibers. The Nextel 720 developing within the matrix, the powder should not fiber is an experimental fiber composed of a mixture of densify during subsequent heat treatments and at applica- submicron alumina and mullite grains. The two inter tion temperatures. 8.9 For this reason, in our previous work penetrating phases ensure a small grain size during pro- we used a mullite powder that did not begin to shrink until cessing. The mullite in the fiber contributes to high A1300oC, the maximum fiber application temperature. creep resistance compared to a similar all-alumina fiber After removing the liquid via evaporation, the powder (Nextel 610). The strength of the Nextel 720 fiber matrix was strengthened by infiltrating the composite with about 30% less than the 610 for single filament proper a solution containing precursor molecules. After evapor- ties, but it was selected because of its greater creep ating the liquid, heating causes the precursor molecules to resistance. I decompose to form an inorganic material that bonds the as detailed below laminated ceramic cloth was infil particles together. The inorganic phase that bonds and trated with a previously consolidated mixture of 70 strengthens the powder matrix also bonds the particles vol% cubic zirconia(solid solution with 8 mol%Y203 (matrix)to the fibers. Cyclic solution precursor infiltration, TZ8YS, Toso Ceramics, average particle diameter of evaporation, and decomposition further strengthens the(0.4 um) and 30 vol% mullite(MU-107, Showa Denko) matrix phase. Care must be taken to avoid precursor As detailed below and elsewhere, zirconia was used as molecules from migrating to the surface during evapora- the matrix because it can be sintered, without shrinkage, tion. This condition produces surface cracking during when heat treated in HCI at temperatures as low as drying due to a thin layer of precursor molecules that 1 C. 3 Mullite was introduced because previous form on the surface. 0 An all-oxide. fiber reinforced cera- work has shown that mullite does not allow the sintered mic composite can be processed in this method. Extensive and coarsened zirconia to shrink(densify)after expo- mechanical testing by Levi et al. has shown that this sure to air at 1200 C for 100 h.6 The zirconia was com- type of composite can exhibit a significant notch insen- osed of agglomerated particles which contained sitive strength in tensile loading. It also has all the primary particles of 50-100 nm in diameter. Infrared attributes found for fiber reinforced ceramics fabricated spectroscopy indicated that the mullite contained an with dense matrixes and weak fiber/matrix interfaces organic contaminant that had to be removed from the Here we report a much simpler and less time consuming powder before it was formulated as a slurry. A 10 h heat fabrication method for processing these new CMCs with treatment in air at 800C was sufficient to remove the porous matrices. In the new method, the powder is treated contaminant. The particle size(average=0.7um) did not to produce a special interparticle pair potential which change during the heat treatment. llows the powder compact(previously consolidated by Dispersed, aqueous slurries containing 20 vol% of the pressure filtration)to be fluidized. It can then be formed two powders were formed by adding 1.3 vol% poly nto a thin sheet by vibrating between plastic sheets. The ethylene oxide urethane silane(PEg-silane, Gelest, Inc) plastic sheets help to avoid evaporation and the con- at pH 10.5. This was found to be sufficient to coat the sequent drying of the thin particle layers. The ceramic particle surfaces. As detailed elsewhere, the PEG-silane sheet is then frozen to enable removal from between the molecules chem-adsorb to the particles by reacting with plastic sheets. The frozen ceramic sheet of powder is the -M-OH(M=metal atom) surface sites. 14, 15The then sandwiched between sheets of ceramic fibers(e.g. zirconia slurry was attrition milled for 15 min after the woven cloth). After thawing, the powder sheet is flui- powder was added. The mullite slurry was sonicated dized by vibrating. It then flows to surround all fibers in with an ultrasonic horn for 5 min prior to the final the adjacent fiber sheets After evaporation, the powder adjustment of the pH. Tetraethylammonium chloride urrounding the fibers can be made strong either by the TEACI)salt(0. 1 molar) was added to form weakly use of precursors described above or by an HCl eva- attractive pair potentials between the particles. As poration/condensation treatment described below. Pre- detailed elsewhere, TMA+ counter ions aid in produ- liminary mechanical measurements show that this new cing a weakly attractive particle network which can be route can result in similar properties as the previous packed to a high density via pressure filtration and route to manufacture CMCs with porous matrices allow the consolidated body to be fluidized via vibra- tion 16-18 TEA+ counter ions were used in this work these counter ions are slightly larger than TMA. Other 2. Experimental methods can be used to produce weakly attractive net- works such as surfactants or chemi-sorption of alco- 2.1. Composite processing hols. 9 The PEG-silane plus TEACI approach was appropriate here due to the two different powders used Composites were formed from layers of two dimen- to form a composite slurry The two slurries were mixed sional, 8 harness woven cloth of Nextel M 720 fiber in appropriate portions described above. The mixed
much smaller than the ®ber diameter to ensure good particle packing.7 To avoid large, cracklike voids from developing within the matrix, the powder should not densify during subsequent heat treatments and at application temperatures.8,9 For this reason, in our previous work we used a mullite powder that did not begin to shrink until 1300C, the maximum ®ber application temperature. After removing the liquid via evaporation, the powder matrix was strengthened by in®ltrating the composite with a solution containing precursor molecules. After evaporating the liquid, heating causes the precursor molecules to decompose to form an inorganic material that bonds the particles together. The inorganic phase that bonds and strengthens the powder matrix also bonds the particles (matrix) to the ®bers. Cyclic solution precursor in®ltration, evaporation, and decomposition further strengthens the matrix phase. Care must be taken to avoid precursor molecules from migrating to the surface during evaporation. This condition produces surface cracking during drying due to a thin layer of precursor molecules that form on the surface.10 An all-oxide, ®ber reinforced ceramic composite can be processed in this method. Extensive mechanical testing by Levi et al.5 has shown that this type of composite can exhibit a signi®cant notch insensitive strength in tensile loading. It also has all the attributes found for ®ber reinforced ceramics fabricated with dense matrixes and weak ®ber/matrix interfaces. Here we report a much simpler and less time consuming fabrication method for processing these new CMCs with porous matrices. In the new method, the powder is treated to produce a special interparticle pair potential which allows the powder compact (previously consolidated by pressure ®ltration) to be ¯uidized. It can then be formed into a thin sheet by vibrating between plastic sheets. The plastic sheets help to avoid evaporation and the consequent drying of the thin particle layers. The ceramic sheet is then frozen to enable removal from between the plastic sheets. The frozen ceramic sheet of powder is then sandwiched between sheets of ceramic ®bers (e.g. woven cloth). After thawing, the powder sheet is ¯uidized by vibrating. It then ¯ows to surround all ®bers in the adjacent ®ber sheets After evaporation, the powder surrounding the ®bers can be made strong either by the use of precursors described above or by an HCl evaporation/condensation treatment described below. Preliminary mechanical measurements show that this new route can result in similar properties as the previous route to manufacture CMCs with porous matrices. 2. Experimental 2.1. Composite processing Composites were formed from layers of two dimensional, 8 harness woven cloth of NextelTM 720 ®ber tows produced by the 3M Corporation (St. Paul, MN). Each tow nominally contains 420 ®bers. The Nextel 720 ®ber is an experimental ®ber composed of a mixture of submicron alumina and mullite grains. The two interpenetrating phases ensure a small grain size during processing. The mullite in the ®ber contributes to high creep resistance compared to a similar all-alumina ®ber (Nextel 610). The strength of the Nextel 720 ®ber is about 30% less than the 610 for single ®lament properties,11 but it was selected because of its greater creep resistance.12 As detailed below, laminated ceramic cloth was in®ltrated with a previously consolidated mixture of 70 vol% cubic zirconia (solid solution with 8 mol% Y2O3, TZ8YS, Toso Ceramics, average particle diameter of (0.4 mm) and 30 vol% mullite (MU-107, Showa Denko). As detailed below and elsewhere, zirconia was used as the matrix because it can be sintered, without shrinkage, when heat treated in HCl at temperatures as low as 1100C.13 Mullite was introduced because previous work has shown that mullite does not allow the sintered and coarsened zirconia to shrink (densify) after exposure to air at 1200C for 100 h.6 The zirconia was composed of agglomerated particles which contained primary particles of 50±100 nm in diameter. Infrared spectroscopy indicated that the mullite contained an organic contaminant that had to be removed from the powder before it was formulated as a slurry. A 10 h heat treatment in air at 800C was sucient to remove the contaminant. The particle size (average=0.7mm) did not change during the heat treatment. Dispersed, aqueous slurries containing 20 vol% of the two powders were formed by adding 1.3 vol% polyethylene oxide urethane silane (PEG-silane, Gelest, Inc.) at pH 10.5. This was found to be sucient to coat the particle surfaces. As detailed elsewhere, the PEG-silane molecules chem-adsorb to the particles by reacting with the -M±OH (M=metal atom) surface sites.14,15 The zirconia slurry was attrition milled for 15 min after the powder was added. The mullite slurry was sonnicated with an ultrasonic horn for 5 min prior to the ®nal adjustment of the pH. Tetraethylammonium chloride TEACl) salt (0.1 molar) was added to form weakly attractive pair potentials between the particles. As detailed elsewhere, TMA+ counter ions aid in producing a weakly attractive particle network which can be packed to a high density via pressure ®ltration and allow the consolidated body to be ¯uidized via vibration.16±18 TEA+ counter ions were used in this work; these counter ions are slightly larger than TMA+. Other methods can be used to produce weakly attractive networks such as surfactants or chemi-sorption of alcohols.19 The PEG-silane plus TEACI approach was appropriate here due to the two dierent powders used to form a composite slurry The two slurries were mixed in appropriate portions described above. The mixed J.J. Haslam et al. / Journal of the European Ceramic Society 20 (2000) 607±618 609
J. Haslam et al. Journal of the European Ceramic Society 20(2000)607-618 slurry was then consolidated by pressure filtration at 5 wiched between two layers of woven, fiber cloth. Sand- MPa to form disc shaped bodies that were fully satu- wiches with up to 27 layers, (14 weaves and 13 frozen rated with water. The saturated bodies were stored in ceramic tapes) were piled up, packed in plastic, evac- sealed plastic bags containing a small paper towel satu- uated, and sealed in plastic. After thawing, the assem- rated with water to help prevent drying. The volume bled layers were vibrated and pressed lightly in between fraction of powder within the saturated, consolidated two steel plates with appropriate spacers to cause the bodies was determined by weight difference method as fluidized powder to flow and intrude the finer layers 52%. 5 At a later time the consolidated powder com- Multi-layer composites in sizes of 40x 100x 3 mm could pact (or a portion cut with a razor blade) was placed be fabricated by this vibration supported single step between two plastic sheets(e. g. a bag) and fluidized with impregnation. The multi-layer cloth composite could an air-powered vibrator into uniform 300 um thick then be dried or frozen for later use. It should be noted esof consolidated particles that the layers of fiber cloth, impregnated with the flui- An illustration of the composite processing steps is dized powder ceramic compact as described were very shown in Fig. 1. Initially, tapes are formed by pressing flexible and could be shaped much like a sheet of un- the fluidized sheets in between two flat steel plates using cross-linked carbon fiber/epoxy prepreg wo spacer bars to fix the thickness. The pressed tape Further processing requires removing the water from were flexible due to the weakly attractive particle poten- the saturated powder matrix by drying in an oven at tial. The tapes, still between the plastic sheets, were fro- 70C, and then sintering the ZrO 2 in a dry HCl gas zen to facilitate composite processing and/or storage environment at temperatures between 1200 and To produce the composite, the frozen tapes were 1300 C20 As reported elsewhere the HCI gas heat removed from between the plastic sheets and sand- treatment did not affect the strength of fiber bundles With the knowledge of the volume of fibers per unit area of cloth the volume fraction of foors within the Vibro Impregnation Process composite was determined by measuring the volume of with Tape Freezing the composite and counting the number of fiber layers sll ach specimen. For composites fabricated for this ly, the average volume fraction of fibers was 0.37±0.02 2.2 Interlaminar shear tests queeze For some design considerations, a desirable property of a woven, layered composite is to have sufficient interlaminar shear strength to resist delamination. This type of failure might be encountered in a bending type pile + pack of loading through the thickness as encountered with a through-thickness temperature gradient. Interlaminar shear strength was determined with 0/90 bar speci mens(3.5x7x20 mm nominal dimensions) diamond cut from larger plates fabricated with 12 or more cloth lay ers. The specimen edges were diamond ground(400 grit) vibrate squeeze to remove a minimum 300 um of damage introduced by the diamond cutting. 3-Point flexural tests were the span was changed for reasons dis- cussed below. The fiber weave orientation wa horizontal with the loading in the vertical direction as layered composite shown in Fig. 2(a). Nylon rods(6.45 mm diameter)were Matrix material Fiberweave 2 Impregnated composite 口 Steel plates o Plastic baa (b) Fig. 2. Schematic of fiber orientation of composite for bending tests. Fig 1. This illustration shows the processing steps used to form the (a)Interlaminar Shear Strength tests.(b)In-plane bend testing composite. flexural strength and elastic modulus
slurry was then consolidated by pressure ®ltration at 5 MPa to form disc shaped bodies that were fully saturated with water. The saturated bodies were stored in sealed plastic bags containing a small paper towel saturated with water to help prevent drying. The volume fraction of powder within the saturated, consolidated bodies was determined by weight dierence method as 52%.15 At a later time, the consolidated powder compact (or a portion cut with a razor blade) was placed between two plastic sheets (e.g. a bag) and ¯uidized with an air-powered vibrator into uniform 300 mm thick `tapes' of consolidated particles. An illustration of the composite processing steps is shown in Fig. 1. Initially, tapes are formed by pressing the ¯uidized sheets in between two ¯at steel plates using two spacer bars to ®x the thickness. The pressed tapes were ¯exible due to the weakly attractive particle potential. The tapes, still between the plastic sheets, were frozen to facilitate composite processing and/or storage. To produce the composite, the frozen tapes were removed from between the plastic sheets and sandwiched between two layers of woven, ®ber cloth. Sandwiches with up to 27 layers, (14 weaves and 13 frozen ceramic tapes) were piled up, packed in plastic, evacuated, and sealed in plastic. After thawing, the assembled layers were vibrated and pressed lightly in between two steel plates with appropriate spacers to cause the ¯uidized powder to ¯ow and intrude the ®ner layers. Multi-layer composites in sizes of 401003 mm could be fabricated by this vibration supported single step impregnation. The multi-layer cloth composite could then be dried or frozen for later use. It should be noted that the layers of ®ber cloth, impregnated with the ¯uidized powder ceramic compact as described were very ¯exible and could be shaped much like a sheet of uncross-linked carbon ®ber/epoxy prepreg. Further processing requires removing the water from the saturated powder matrix by drying in an oven at 70C, and then sintering the ZrO2 in a dry HCl gas environment at temperatures between 1200 and 1300C.20 As reported elsewhere21 the HCI gas heat treatment did not aect the strength of ®ber bundles. With the knowledge of the volume of ®bers per unit area of cloth, the volume fraction of ¯oors within the composite was determined by measuring the volume of the composite and counting the number of ®ber layers in each specimen. For composites fabricated for this study, the average volume fraction of ®bers was 0.370.02. 2.2. Interlaminar shear tests For some design considerations, a desirable property of a woven, layered composite is to have sucient interlaminar shear strength to resist delamination. This type of failure might be encountered in a bending type of loading through the thickness as encountered with a through-thickness temperature gradient. Interlaminar shear strength was determined with 0/90 bar specimens (3.5720 mm nominal dimensions) diamond cut from larger plates fabricated with 12 or more cloth layers. The specimen edges were diamond ground (400 grit) to remove a minimum 300 mm of damage introduced by the diamond cutting. 3-Point ¯exural tests were performed where the span was changed for reasons discussed below. The ®ber weave orientation was horizontal with the loading in the vertical direction as shown in Fig. 2(a). Nylon rods (6.45 mm diameter) were Fig. 1. This illustration shows the processing steps used to form the composite. Fig. 2. Schematic of ®ber orientation of composite for bending tests. (a) Interlaminar Shear Strength tests. (b) In-plane bend testing for ¯exural strength and elastic modulus. 610 J.J. Haslam et al. / Journal of the European Ceramic Society 20 (2000) 607±618
J. Haslam et al. Journal of the European Ceramic Society 20(2000)607-618 used as loading pins to accommodate the inherent composite strength in bending without interlaminar roughness of the woven fibers on the specimen surface shear type failures. Strain was calculated based on and to reduce contact loading stress. No permanent measurements of the bottom beam displacement. In deformation of the loading rods was observed after addition, based on a technique used by Heathcote et testing, which indicates that they remained elastic under al., 22 notched bending tests were performed in a similar the stresses encountered in testing. A servo-electric test- manner. Two diferent bar specimens(3.5x7x90 mm ing machine (Instron, Inc. model 8562) with a high and 3. 5x7x45 mm nominal dimensions)were tested in stifness frame was used to load the specimens at a cross 3-point flexural load with an outer span of 35 mm and head speed of 0. 1 mm/ min 89 mm for the shorter and longer specimens respec The shear stress at the mid-plane of a flexural bar tively. The cross head displacement rate was 0. 1 mm/ pecimen can be calculated from beam theory as min. two different fiber with0°/90° and one with+/-45° fiber directions rela r=3/4L/(W*1) (2) tive to the direction of the bar lengtH Notches were cut in the center of bars with a diamond where L=load. w=width, and t= thickness The mea wheel with a resulting notch thickness of 0.55 mm. The surement of delamination stress for CMCs usually pre- notch depth was nominally one half of the sample scribes a span(S)to thickness ratio(S/ n)of greater than height, a/W=0.485+0.015. The high stifness of the 10 to help insure that the specimen fails by delamination testing machine/load cell, and the non-catastrophic fail- (shear)rather than a tensile failure on the surface, given ure of the specimens allowed for careful measurement of the energy required to break the notched specimens The measured projected surface area of the sample was 3/2PS (3) used to calculate energy per unit area to produce frac- ture The stress-strain response of the un-notched compo- The midplane shear stress to maximum tensile stress sites was nearly linear elastic with some significant ratio(t/o)is given by deviations in the +/-45 fiber direction tests. Due to the low loads encountered in testing specimens of this r/a=1/2(/S) (4) length, no damage was observed at the loading po The specimens were tested in the same high-stifiness Therefore, for a given specimen thickness, the shorter universal testing machine used for the interlaminar the span, the greater the probability that failure will shear strength tests take place by a delamination of the cloth layers, rather than crack extension through the layers. Flexural testing with small values of S/t is called short beam bend test esults ing and is used to characterize the interlaminar shear trength 3. 1. Interlaminar shear strength 2.3. In-plane flexure testing of notched and un-notched Fig. 3 reports the apparent interlaminar shear strength as a function of span to thickness ratio(S/n) for individual CMC specimens heat treated in HCI for dif- Attempts to perform tensile tests on 100 mm long ferent time periods and temperatures. As shown, the specimens(same material as above) with a reduced delamination stress was 10+2 MPa for all heat treat gauge section(5. 1 mm wide, 40 mm long, produced with ments, and that specimens produced from one heat a 152 mm diameter diamond grinding wheel) were not treatment(1250C/5 h) failed in tension and did not successful with our limited amount of material. Despite delaminate. Fig. 3 also reports the delamination stress the use of fiberglass tabs that were epoxied to the ends reported by Levi et al. for a CMC with a porous of the specimen and double knife-edge universal joints matrix, but fabricated by the older method(pressure within the tensile train, most specimens failed either in filtration, multiple precursor infiltration and pyrolysis the non-reduced gauge section or adjacent to the cycles ). Its delamination strength of 8 MPa is a little clamping grip lower than most of the values reported for our newer Because of the limited amount of fiber cloth available method but Levi et al. used harder, steel loading pins, for fabricating specimens, the implementation of an which could have produced a stress concentration and a improved tensile test was not possible. The testing mode lower delamination stress. was changed to an in-plane flexural test, subjected to 3- Figs. 4 and 5 illustrate typical stress versus strain plot point flexural loading as shown in Fig. 2(b). This con- for specimens that delaminated prior to tensile failure figuration and loading mode allowed for testing of the In general, one or two load drops were observed similar
used as loading pins to accommodate the inherent roughness of the woven ®bers on the specimen surface and to reduce contact loading stress. No permanent deformation of the loading rods was observed after testing, which indicates that they remained elastic under the stresses encountered in testing. A servo-electric testing machine (Instron, Inc. model 8562) with a high stiness frame was used to load the specimens at a cross head speed of 0.1 mm/min. The shear stress at the mid-plane of a ¯exural bar specimen can be calculated from beam theory as: 3=4 L= Wt; 2 where L=load, w=width, and t=thickness. The measurement of delamination stress for CMCs usually prescribes a span (S) to thickness ratio (S/t) of greater than 10 to help insure that the specimen fails by delamination (shear) rather than a tensile failure on the surface, given by 3=2PS bt2 : 3 The midplane shear stress to maximum tensile stress ratio (t/s) is given by = 1=2 t=S: 4 Therefore, for a given specimen thickness, the shorter the span, the greater the probability that failure will take place by a delamination of the cloth layers, rather than crack extension through the layers. Flexural testing with small values of S/t is called short beam bend testing and is used to characterize the interlaminar shear strength. 2.3. In-plane ¯exure testing of notched and un-notched specimens Attempts to perform tensile tests on 100 mm long specimens (same material as above) with a reduced gauge section (5.1 mm wide, 40 mm long, produced with a 152 mm diameter diamond grinding wheel) were not successful with our limited amount of material. Despite the use of ®berglass tabs that were epoxied to the ends of the specimen and double knife-edge universal joints within the tensile train, most specimens failed either in the non-reduced gauge section or adjacent to the clamping grip. Because of the limited amount of ®ber cloth available for fabricating specimens, the implementation of an improved tensile test was not possible. The testing mode was changed to an in-plane ¯exural test, subjected to 3- point ¯exural loading as shown in Fig. 2(b). This con- ®guration and loading mode allowed for testing of the composite strength in bending without interlaminar shear type failures. Strain was calculated based on measurements of the bottom beam displacement. In addition, based on a technique used by Heathcote et al.,22 notched bending tests were performed in a similar manner. Two dierent bar specimens (3.5790 mm and 3.5745 mm nominal dimensions) were tested in 3-point ¯exural load with an outer span of 35 mm and 89 mm for the shorter and longer specimens respectively. The cross head displacement rate was 0.1 mm/ min. Two dierent ®ber alignments were tested, one with 0/90 and one with +/ÿ45 ®ber directions relative to the direction of the bar length. Notches were cut in the center of bars with a diamond wheel with a resulting notch thickness of 0.55 mm. The notch depth was nominally one half of the sample height, a/W=0.4850.015. The high stiness of the testing machine/load cell, and the non-catastrophic failure of the specimens allowed for careful measurement of the energy required to break the notched specimens. The measured projected surface area of the sample was used to calculate energy per unit area to produce fracture. The stress±strain response of the un-notched composites was nearly linear elastic with some signi®cant deviations in the +/ÿ45 ®ber direction tests. Due to the low loads encountered in testing specimens of this length, no damage was observed at the loading points. The specimens were tested in the same high-stiness universal testing machine used for the interlaminar shear strength tests. 3. Results 3.1. Interlaminar shear strength Fig. 3 reports the apparent interlaminar shear strength as a function of span to thickness ratio (S/t) for individual CMC specimens heat treated in HCI for different time periods and temperatures. As shown, the delamination stress was 102 MPa for all heat treatments, and that specimens produced from one heat treatment (1250C/5 h) failed in tension and did not delaminate. Fig. 3 also reports the delamination stress reported by Levi et al.5 for a CMC with a porous matrix, but fabricated by the older method (pressure ®ltration, multiple precursor in®ltration and pyrolysis cycles). Its delamination strength of 8 MPa is a little lower than most of the values reported for our newer method but Levi et al. used harder, steel loading pins, which could have produced a stress concentration and a lower delamination stress.5 Figs. 4 and 5 illustrate typical stress versus strain plot for specimens that delaminated prior to tensile failure. In general, one or two load drops were observed similar J.J. Haslam et al. / Journal of the European Ceramic Society 20 (2000) 607±618 611