ournal Am. Ceram.Soc.85m2599-632(2002) Interface Design for Oxidation-Resistant Ceramic Composites Ronald J. Kerans, *f Randall S Hay, f Triplicane A. Parthasarathy, and Michael K. Cinibulk*f Materials and Manufacturing Directorate, AFRL/MLLN, Air Force Research Laboratory Wright-Patterson Air Force Base, Ohio 45433 fibers and matrices also suffer environmental through distributed damage mechanisms. These mechanisms r, changes in the mechanical properties of carbon- or lent on matrix cracks deflecting into fiber/matrix rolled interfaces after oxidation or enhanced oxidation of interfacial debonding cracks. Oxidation resistance of the fiber fibers or matrices after interface oxidation usually dominates CMC coatings often used to enable crack deflection is an importal behavior(see, for example, Refs. 5 and 18-20). This has mot limitation for long-term use in many applications. Research on vated research on more oxidation-resistant fiber coatings, viscous alternative, mostly oxide, coatings for oxide and non-oxide alant phases, and porous-matrix systems that do not require composites is reviewed. Processing issues, such as fiber coat ific interface control constituents (for concise reviews, see ngs and fiber strength degradation, are discussed. Mechanics Refs. 21 and 22). From a mechanistic standpoint, the substitution work related to design of crack deflecting coatings is also of Bn for carbon has been relatively straightforward; they have reviewed, and implications on the design of coatings and of very similar structures and elastic and fracture properties. BN and composite systems using alternative coatings are discussed. carbon are used as solid lubricants and can be expected to provide Potential topics for further research are identified low sliding friction. Substitution of oxides is a very different matter, and, unfortunately, lack of well-defined interface property L. Introduction equirements complicates the design and evaluation of alternative viable approaches for use in composites IE discovery that brittle ceramics can be made highly having non-oxide constituents can be further complicated by the tolerant by combining them in fiber/matrix composi sIte form need for stability and compatibility in strongly reducing processin (ceramic-matrix composite or CMC, continuous-fiber environments. In fact, most oxide-coating work to date has been on e or CFCC, and ceramic-fiber matrix composite or oxide fibers to be used in oxide matrices. Research on fiber CFMC) has spawned research spanning approximately three de- coating processes is also required. For example, coated fibers often ades. Early work revealed that deflection of matrix cracks to the dis splay severely degraded tensile strength, 3, 24 which has moti- fiber/matrix interface, leaving intact fibers behind the matrix crack vated research on mechanisms of degradation. tip, was essential for tough behavior. Crack deflection in mos Although development of oxidation-resistant interface control is CMCs has been effected by a relatively weak and compliant complex, there has been progress carbon coating applied to the fibers before matrix processing o (1) There are many interface design parameters, and they are formed in situ by fiber decomposition during matrix processing better understood However, long-term use of CMCs has been limited by several (2) Several more oxidation-resistant alternatives to carbon and forms of environmental degradation, the most pervasive of which bn have the correct crack deflection behavior. and some show as been oxidation of the fiber coatings promise for the correct fiber pullout behavior. 2-30 To improve oxidation resistance, BN has been substituted for (3) There has been progress toward viable fiber-coating carbon(see, for example, Refs. 7-17). Progress has been made on process ystems using BN, and the best Bn coatings demonstrate very (4) Definitive evidence of oxide coatings effecting character good properties. Nevertheless, although BN is a much better istic composite fracture and properties in true yarm-reinforced coating than carbon, it has much poorer oxidation resistance than composites has been observed for two different oxide coatings most candidate fiber and matrix constituents (Fig. 1). In this review, progress is summarized in a manner intended to ssist in developing guidelines for the design and evaluation of B. Marshall-contributing editor fiber coatings and to highlight the most interesting areas for further esearch. Strategies for oxidation-resistant coatings and relevant interface mechanics are critically reviewed. Progress and problems in coating of fibers are summarized. Section Il provides back- Manuscript No 188122 Received November 29, 2000, approved June 13, 2002. ground in the form of a brief review of historical aspects of interface oxidation. a discussion of the mechanics of crack Also affiliated with UES, Inc, Dayton, OH, under U.S. Air Force Contract No. deflection and sliding, the effects of coating properties F33615-96-C5258 posite behavior, and target values for interface parameters. Section Feature
Interface Design for Oxidation-Resistant Ceramic Composites Ronald J. Kerans,* ,† Randall S. Hay,* ,† Triplicane A. Parthasarathy,* ,‡ and Michael K. Cinibulk* ,† Materials and Manufacturing Directorate, AFRL/MLLN, Air Force Research Laboratory, Wright-Patterson Air Force Base, Ohio 45433 Fiber-reinforced ceramic composites achieve high toughness through distributed damage mechanisms. These mechanisms are dependent on matrix cracks deflecting into fiber/matrix interfacial debonding cracks. Oxidation resistance of the fiber coatings often used to enable crack deflection is an important limitation for long-term use in many applications. Research on alternative, mostly oxide, coatings for oxide and non-oxide composites is reviewed. Processing issues, such as fiber coatings and fiber strength degradation, are discussed. Mechanics work related to design of crack deflecting coatings is also reviewed, and implications on the design of coatings and of composite systems using alternative coatings are discussed. Potential topics for further research are identified. I. Introduction THE discovery that brittle ceramics can be made highly damage tolerant by combining them in fiber/matrix composite form (ceramic-matrix composite or CMC, continuous-fiber ceramic composite or CFCC, and ceramic-fiber matrix composite or CFMC) has spawned research spanning approximately three decades. Early work revealed that deflection of matrix cracks to the fiber/matrix interface, leaving intact fibers behind the matrix crack tip, was essential for tough behavior.1–6 Crack deflection in most CMCs has been effected by a relatively weak and compliant carbon coating applied to the fibers before matrix processing or formed in situ by fiber decomposition during matrix processing. However, long-term use of CMCs has been limited by several forms of environmental degradation, the most pervasive of which has been oxidation of the fiber coatings. To improve oxidation resistance, BN has been substituted for carbon (see, for example, Refs. 7–17). Progress has been made on systems using BN, and the best BN coatings demonstrate very good properties. Nevertheless, although BN is a much better coating than carbon, it has much poorer oxidation resistance than most candidate fiber and matrix constituents. CMC fibers and matrices also suffer environmental degradation. However, changes in the mechanical properties of carbon- or BN-controlled interfaces after oxidation or enhanced oxidation of fibers or matrices after interface oxidation usually dominates CMC behavior (see, for example, Refs. 5 and 18–20). This has motivated research on more oxidation-resistant fiber coatings, viscous sealant phases, and porous-matrix systems that do not require specific interface control constituents (for concise reviews, see Refs. 21 and 22). From a mechanistic standpoint, the substitution of BN for carbon has been relatively straightforward; they have very similar structures and elastic and fracture properties. BN and carbon are used as solid lubricants and can be expected to provide low sliding friction. Substitution of oxides is a very different matter, and, unfortunately, lack of well-defined interface property requirements complicates the design and evaluation of alternative interfaces. Identifying viable approaches for use in composites having non-oxide constituents can be further complicated by the need for stability and compatibility in strongly reducing processing environments. In fact, most oxide-coating work to date has been on oxide fibers to be used in oxide matrices. Research on fibercoating processes is also required. For example, coated fibers often display severely degraded tensile strength,23,24 which has motivated research on mechanisms of degradation. Although development of oxidation-resistant interface control is complex, there has been progress. (1) There are many interface design parameters, and they are better understood.25,26 (2) Several more oxidation-resistant alternatives to carbon and BN have the correct crack deflection behavior, and some show promise for the correct fiber pullout behavior.27–30 (3) There has been progress toward viable fiber-coating processes.23,31–37 (4) Definitive evidence of oxide coatings effecting characteristic composite fracture and properties in true yarn-reinforced composites has been observed for two different oxide coatings (Fig. 1). In this review, progress is summarized in a manner intended to assist in developing guidelines for the design and evaluation of fiber coatings and to highlight the most interesting areas for further research. Strategies for oxidation-resistant coatings and relevant interface mechanics are critically reviewed. Progress and problems in coating of fibers are summarized. Section II provides background in the form of a brief review of historical aspects of interface oxidation, a discussion of the mechanics of crack deflection and sliding, the effects of coating properties on composite behavior, and target values for interface parameters. Section D. B. Marshall—contributing editor Manuscript No. 188122. Received November 29, 2000; approved June 13, 2002. *Member, American Ceramic Society. † Air Force Research Laboratory. ‡ Also affiliated with UES, Inc., Dayton, OH, under U.S. Air Force Contract No. F33615-96-C-5258. journal J. Am. Ceram. Soc., 85 [11] 2599–632 (2002) Feature
2600 Journal of the American Ceramic Socien-Kerans et aL. Vol. 85. No. 11 fiber surfaces(a thin carbon layer over a thin SiO, layer) that is weak enough to deflect matrix cracks and protect the fibers from matrix crack stress concentrations. ,s Sliding between fiber and matrix, before and after the fibers fracture, further dissipates energy via friction. These mechanisms give CMCs the tolerance to local overload that makes them useful as structural materials Composites with no carbon layer fail catastrophically with low strength in the manner of poor-quality monolithics cally, had the Nicalon fiber actually been stoichiometric crystalline SiC, carbon layers would not have formed in situ, and attaining mechanically viable ceramic composites would have been more problematic, but perhaps hastened more-detailed understanding of the mechanics governing composite desi Early CMC studies measured strength and load-deflection behavior at room temperature, CMCs with carbon layers on the fibers demonstrated high strength, high strain-to-failure, and non- linear load-deflection behavior. However. when tested at (b temperatures, there was a substantial loss in strength above 900C (Fig. 2).,,+Initially, this was attributed to replacement of the carbon layer by Sio, that strongly bonded fibers to the matrix and allowed matrix cracks to propagate directly through fibers. 7 48 Recent work suggests that oxidative degradation of Nicalon fiber may contribute to composite strength loss to a degree comparable to direct effects of interface property changes. 49-52Nevertheles in either case, carbon interface oxidation allowing oxygen acces to the entire fiber surface area in a Cmc is the first degradation step. Above 1000oC, a self-sealing SiO, layer can prevent acces of oxygen to the interface. 48,53 However, at intermediate temper occurs from uninterrupted oxidation(Fig. 3). 20,48 so ength los atures, typically between 700 and 900C, significant st experiments, analytical modeling s and experiments on Nica- lon/C/SiC composites have contributed to the current under standing of this intermediate-temperature degradation. It has been argued that fibers(and coatings)do not oxidize in a crack-free CMC used at design stresses less then the matrix- cracker Fig. 1. (a)Fracture surface of Nextel 720 fiber/monazite fiber coating/ stress. Such an approach might be acceptable for preservation of the aluminosilicate matrix indicating that crack deflection occurred at or near interface when overloads are infrequent and design stresses are low ber/coating interfaces. Energy dispersive spectroscopy(EDS)analysis enough that cracks are not held open, or if there are mechanisms to indicates that the light phase is monazite and that it is essentially al ways al lightly loaded cracks. A sensible design using this approach left in the trough.( Fiber coating by AFRL/ML; composite by Composi trives to have the regions most likely to crack, the more highly Optics, Inc )(b) Fracture surface of Nextel 610/scheelite fiber coating/ stressed regions, at temperatures that are relatively benign alumina CerablakM matrix indicating that crack deflection occurred at or Although this approach has merit if the cra near fiber/coating interfaces( Coating and composite by Mc Dermott, In and Applied Thin Films, Inc. made sufficiently high and the application environment is well- known. it seems far from an ideal solution. All design stress calculations are approximations based on an idealized situation including mating of perfectly matching surfaces, absence of Ill discusses the design and evaluation of coatings and composites lefects and foreign matter, and predictable environments. These Section IV discusses specific approaches to interface control. For approximations work for metals, because ductile materials blunt completeness, BN coatings and porous-matrix composites also are flaws by local plastic deformation that otherwise cause local stress briefly reviewed in Section IV. Section V discusses coating concentrations. For CMCs, the equivalent local deformation is process technology and fiber degradation. Section VI summarizes local matrix cracking and a few broken fibers, which allows acces and speculates on future options. This review is intended to be a of the atmosphere to the composite interior. Furthermore, there is comprehensive critical review and to provide some thought- evidence that matrix cracking occurs in some CMCs well below provoking speculation on composite design and useful future the proportional limit. The fact that introduction of monolithic ceramics into structural applications has been slow and limited despite very high strength and thorough proof testing, provides circumstantial evidence for this point of view. At least occasional Il. Interface Properties and Mechanics local stress concentrations greater than the matrix-cracking stress almost 's exist in practice. Hence, the ideal composite Initial interest in CFCCs was generated by marketing of equires all constituents to be oxidation resistant, including the Nicalon fiber(Nippon Carbon Co Japan) and the fiber/matrix interface erceived availability of a fiber that had the nsity, creep, and oxidation resistance of sic and the high and fabrication ease of small-diameter filaments in a fiber tow However Nicalon (2) Initiation of Interfacial Cracks and Deflection of is not crystalline SiC, but instead is carbon-and oxygen-rich and atrix cracks Although in most respects Ni Crack deflection is the most important event for achievin excellent fiber, when exposed to high temperatures, it crystallizes tough composites; however, the complexities of the problem and to SiC, rejects carbon and oxygen, and shrinks slightly. ,4During of real materials require simplification for analysis, and confirma matrix processing, this decomposition can form a coating on the tion by experiment is problematic. The details of crack deflection
III discusses the design and evaluation of coatings and composites. Section IV discusses specific approaches to interface control. For completeness, BN coatings and porous-matrix composites also are briefly reviewed in Section IV. Section V discusses coating process technology and fiber degradation. Section VI summarizes and speculates on future options. This review is intended to be a comprehensive critical review and to provide some thoughtprovoking speculation on composite design and useful future work. II. Interface Properties and Mechanics Initial interest in CFCCs was generated by marketing of NicalonTM fiber (Nippon Carbon Co., Tokyo, Japan) and the perceived availability of a fiber that had the low density, creep, and oxidation resistance of SiC and the high strength and fabrication ease of small-diameter filaments in a fiber tow. However, Nicalon is not crystalline SiC, but instead is carbon- and oxygen-rich and nearly amorphous.38–41 Although in most respects Nicalon is an excellent fiber, when exposed to high temperatures, it crystallizes to SiC, rejects carbon and oxygen, and shrinks slightly.40,42 During matrix processing, this decomposition can form a coating on the fiber surfaces (a thin carbon layer over a thin SiO2 layer) that is weak enough to deflect matrix cracks and protect the fibers from matrix crack stress concentrations.3,38 Sliding between fiber and matrix, before and after the fibers fracture, further dissipates energy via friction. These mechanisms give CMCs the tolerance to local overload that makes them useful as structural materials. Composites with no carbon layer fail catastrophically with low strength in the manner of poor-quality monolithics.38,43–45 Ironically, had the Nicalon fiber actually been stoichiometric crystalline SiC, carbon layers would not have formed in situ, and attaining mechanically viable ceramic composites would have been more problematic, but perhaps hastened more-detailed understanding of the mechanics governing composite design. (1) Oxidation History Early CMC studies measured strength and load–deflection behavior at room temperature.3,4 CMCs with carbon layers on the fibers demonstrated high strength, high strain-to-failure, and nonlinear load–deflection behavior. However, when tested at high temperatures, there was a substantial loss in strength above 900°C (Fig. 2).18,46,47 Initially, this was attributed to replacement of the carbon layer by SiO2 that strongly bonded fibers to the matrix and allowed matrix cracks to propagate directly through fibers.47,48 Recent work suggests that oxidative degradation of Nicalon fiber may contribute to composite strength loss to a degree comparable to direct effects of interface property changes.49–52 Nevertheless, in either case, carbon interface oxidation allowing oxygen access to the entire fiber surface area in a CMC is the first degradation step. Above 1000°C, a self-sealing SiO2 layer can prevent access of oxygen to the interface.48,53 However, at intermediate temperatures, typically between 700° and 900°C, significant strength loss occurs from uninterrupted oxidation (Fig. 3).20,48,50,53 Model experiments,54 analytical modeling,55 and experiments on Nicalon/C/SiC composites20 have contributed to the current understanding of this intermediate-temperature degradation. It has been argued that fibers (and coatings) do not oxidize in a crack-free CMC used at design stresses less then the matrix-cracking stress.56 Such an approach might be acceptable for preservation of the interface when overloads are infrequent and design stresses are low enough that cracks are not held open, or if there are mechanisms to seal lightly loaded cracks.57 A sensible design using this approach strives to have the regions most likely to crack, the more highly stressed regions, at temperatures that are relatively benign. Although this approach has merit if the cracking stress can be made sufficiently high and the application environment is wellknown, it seems far from an ideal solution. All design stress calculations are approximations based on an idealized situation, including mating of perfectly matching surfaces, absence of defects and foreign matter, and predictable environments. These approximations work for metals, because ductile materials blunt flaws by local plastic deformation that otherwise cause local stress concentrations. For CMCs, the equivalent local deformation is local matrix cracking and a few broken fibers, which allows access of the atmosphere to the composite interior. Furthermore, there is evidence that matrix cracking occurs in some CMCs well below the proportional limit.58 The fact that introduction of monolithic ceramics into structural applications has been slow and limited, despite very high strength and thorough proof testing, provides circumstantial evidence for this point of view. At least occasional local stress concentrations greater than the matrix-cracking stress almost always exist in practice. Hence, the ideal composite requires all constituents to be oxidation resistant, including the fiber/matrix interface. (2) Initiation of Interfacial Cracks and Deflection of Matrix Cracks Crack deflection is the most important event for achieving tough composites; however, the complexities of the problem and of real materials require simplification for analysis, and confirmation by experiment is problematic. The details of crack deflection Fig. 1. (a) Fracture surface of Nextel 720 fiber/monazite fiber coating/ aluminosilicate matrix indicating that crack deflection occurred at or near fiber/coating interfaces. Energy dispersive spectroscopy (EDS) analysis indicates that the light phase is monazite and that it is essentially always left in the trough. (Fiber coating by AFRL/ML; composite by Composite Optics, Inc.) (b) Fracture surface of Nextel 610/scheelite fiber coating/ alumina CerablakTM matrix indicating that crack deflection occurred at or near fiber/coating interfaces. (Coating and composite by McDermott, Inc., and Applied Thin Films, Inc.) 2600 Journal of the American Ceramic Society—Kerans et al. Vol. 85, No. 11
Interface Design for Oxidation-Resistant Ceramic Composites 2 (a) 30l -0.0 25 800C Time(10 s Sio cOcO 250 CROSS-HEAD DISPLACEMENT ( IN L 0254 (b) PYROCARBON Fig. 3. (a) Measured oxidative weight change of a Nicalon/carbon/SiC omposite. Oxidation starts with a weight loss from carbon oxidation bove 900C, the oxidation of SiC results in plugs that limit further carbon oxidation; however, at 700 and 800C, the sealing does not happen, as indicated in the(b) schematic and associated statistics of the nterfaces are not well-known. The second problem is that there are no proven local failure criteria. That is, even if the stresses can be U 2。0 calculated, there is no appropriate failure criterion for a very small volume of material where bulk c flaw distribution is unknown. Energy-based analyses assume particular virtual crack extensions and may not be appropriate for 20 ng behavior on this scale. Continued progress can be expected with increasing comparison of analyses to experiments It is also not yet understood what level of detail a model must 10 apture to properly predict actual behavior. Treating coating layers as being properly described by properties of an infinitesimally thin nterface surely must be misleading in many situations. For example, the He and Hutchinson analysis considered the cri- terion for deflection of a mode i matrix crack to an interfacial crack in an ideal planar interface perpendicular to the matrix crack plane. They found that deflection should be expected only for ratios of interface toughness to "fiber" toughness less than a 005 certain value dependant on elastic properties but genera than about 1/4 This often has beel died without CROSS-HEAD DISPLACEMENT (IN) the details of crack deflection However if cracks deflect inside the coating, propagation of the crack in either sense is determined Fig. 2. Effect of oxidation on the mechanical behavior at 900%C in air of by coating fracture properties, i.e., the ratio of coating debond Nicalon-reinforced lithium aluminosilicate(LAS) matrix composite is mode to coating transmission mode fracture energies (T/I shown from the load-stress-displacement behavior at(a)room temperature where c indicates coating, r and z indicate crack surface normals in and at(b)900°cin lindrical coordinates with z along the fiber axis, and applied tractions are along +z )25 Fiber toughnesses are typically a few determine the interface property that must be engineered MPam, therefore, coating toughness can be higher than fiber mechanics analyses have contributed much to the level of toughness. A coating can fail the test of (debond fracture energy M standing of composite behavior, but there are two fiber fracture energy)< /4, but can deflect cracks, because the problems that limit the utility of the analyses in guiding co ratio of coating toughnesses for the two types of cracks do satisfy design. The first of these problems is that the properties, ge the criterion; that is, the coating is sufficiently anisotropic in
determine the interface property that must be engineered. Micromechanics analyses have contributed much to the level of understanding of composite behavior, but there are two pervasive problems that limit the utility of the analyses in guiding composite design. The first of these problems is that the properties, geometry, and associated statistics of the thin coatings and associated interfaces are not well-known. The second problem is that there are no proven local failure criteria. That is, even if the stresses can be calculated, there is no appropriate failure criterion for a very small volume of material where bulk properties do not apply and the flaw distribution is unknown. Energy-based analyses assume particular virtual crack extensions and may not be appropriate for predicting behavior on this scale. Continued progress can be expected with increasing comparison of analyses to experiments. It is also not yet understood what level of detail a model must capture to properly predict actual behavior. Treating coating layers as being properly described by properties of an infinitesimally thin interface surely must be misleading in many situations. For example, the He and Hutchinson analysis59 considered the criterion for deflection of a Mode I matrix crack to an interfacial crack in an ideal planar interface perpendicular to the matrix crack plane. They found that deflection should be expected only for ratios of interface toughness to “fiber” toughness less than a certain value dependant on elastic properties but generally less than about 1/4.59–61 This often has been applied without regard to the details of crack deflection. However, if cracks deflect inside the coating, propagation of the crack in either sense is determined by coating fracture properties, i.e., the ratio of coating debond mode to coating transmission mode fracture energies (c r/ c z, where c indicates coating, r and z indicate crack surface normals in cylindrical coordinates with z along the fiber axis, and applied tractions are along z.).25 Fiber toughnesses are typically a few MPam1/2; therefore, coating toughness can be higher than fiber toughness. A coating can fail the test of (debond fracture energy)/ (fiber fracture energy) 1/4, but can deflect cracks, because the ratio of coating toughnesses for the two types of cracks do satisfy the criterion; that is, the coating is sufficiently anisotropic in Fig. 2. Effect of oxidation on the mechanical behavior at 900°C in air of a Nicalon-reinforced lithium aluminosilicate (LAS) matrix composite is shown from the load-stress–displacement behavior at (a) room temperature and at (b) 900°C in air.18 Fig. 3. (a) Measured oxidative weight change of a Nicalon/carbon/SiC composite. Oxidation starts with a weight loss from carbon oxidation. Above 900°C, the oxidation of SiC results in plugs that limit further carbon oxidation; however, at 700° and 800°C, the sealing does not happen, as indicated in the (b) schematic.48 November 2002 Interface Design for Oxidation-Resistant Ceramic Composites 2601
Journal of the American Ceramic Socien-Kerans et aL. Vol. 85. No. 1I fracture energy to deflect cracks itself. This discussion also failure criterion leaves the matter open to speculation. This suggests that coatings that are intended to deflect cracks by failure sequence has been observed in laminates and in model fracture within the coating can be evaluated independently of the composites Nicalon/C/SiC composites made with fibers treated to promote rack direction from perpendicular to parallel to a fiber surface dence of interfacial failure preceding matrix crack impinge- Fig. 4(a), but there are other possibilities. Mode I interface cracks can form in the tensile stress field normal to the fiber surface ahead connect to debonded coating/fiber interfaces, with no deflection in f a matrix crack z(Fig 4(b). Modeling of an annular matrix the coating itself, i.e., the interface, not the coating, fails. In crack has predicted that, for most reasonable choices of properties, otherwise identical composites made with treated fibers, matrix cracks connect to diffuse cracks in the coating without debonding other failure event(e. g, fiber fracture)intervenes before the coating/fiber interface; ie, the coating itself fails.66, 68The the matrix crack can be driven to the interface. unless the inter interfacial toughness, friction, and composite strength are higher face is debonded ahead of it. 63 Interface stresses can be high for the treated-fiber CMC If the matrix crack runs through the enough to make interfacial debond ahead of the matrix crack a plausible mechanism, but lack of a completely understood local coatings on untreated fibers before the interface debonds, it deflects in the coating, as it does in the identical coating on the treated fiber; hence, the coating/fiber interface must fail before the matrix crack enters the coating. When the crack does pass through the coating, the elastic constraint of the fiber is mostly removed by the preexisting debond, and the crack runs directly to the debonded interface. If there are truly no material difference besides interface strength, a definitive sequence of events consistent with the model is implied. These composites have other interesting properties that are discussed in later sections Another deflection mechanism preceding matrix crack impinge ment has been suggested for a composite comprising SiC mono- filaments with successive coatings of carbon and TiB, in a glass matrix.9 The coating was calculated to be in triaxial compressio d model ested that matrix cracks would run to th coating only after coating or interfacial failure. Experiment re- LATRIX vealed that debond cracks ran very near the fiber surface except for ular-section C rings around the fiber with their peaks at the matrix crack planes. Shear stresses on planes in the approximate orientation of the sides of the C rings(about 45. from the matrix crack plane) were calculated to be the highest coating stresses that could lead to fracture. The suggested failure sequence was(i)the matrix crack approached the coating, (ii) the coating failed in shear on planes +45 from the crack plane, ahead of the crack, and formed C rings, (iii) the coating cracks turned parallel to the fiber urface at or near the fiber surface, and(iv) the matrix crack advanced until it joined the coating shear cracks at their intersec- tion(see Fig. 4(c)) A further possibility is growth of periodic echelon cracks In an analysis of thin laminates with a Mode I crack normal to the nterface. the maximum tensile stresses in the coatings were 45o allel to a "half-turn of the nto eries of parallel"periodic echelon"microcracks on these hi stress planes at the center of the coating. As the microcrack approached the coating/plate interfaces, they turned parallel to the interface and joined to form a debond. Evidence for a similar- appearing but different sequence of events has been observed in monazite interlayers in Al, O /Al,O3 laminates(Fig. 5).27In that case, the echelon cracks appeared after initial deflection of the main crack into the coating/ laminate interface. [二A Detailed fracture observations are difficult: therefore. the se- quence of events for fiber/matrix debonding in CMCs speculative Debonding mechanisms may vary with the particular composite and global stress state. The coating most ortant to crack deflection depends on subtle differences in onstituent elastic properties and residual stresses. In the ideal case, coatings are engineered material "components'" of a compo ite system selected for phase, microstructure, and geometry to promote a specific failure mechanism. Enhanced understanding of crack deflection is necessary to allow such a priori design Simple Fig. 4. Three possible sequences leading to crack deflection:(a)matrix crack grows into the coating and then bifurcates and turn models are often useful, but they can sometimes be misleading running parallel to the fiber surface in each direction(as well hence, detailed analysis and comparison of microstructure, crack on in the matrix),(b) coating or interface fails in the tensile deflection behavior, and analytical models may be necessary. The he matrix crack before arrival of the matrix crack at the interface region same considerations apply to the interpretation of micromechani- and (c)in the matrix crack at the coating, the crack bifurcates and turns as cal tests. For example, fiber pushout/pullout tests may not direct the coating fails in shear at an intermediate angle, then turns parallel to the measure the parameters that actually determine debonding during fiber surface at or near the fiber surface composite failure. It is even possible that debonding in single
fracture energy to deflect cracks itself. This discussion also suggests that coatings that are intended to deflect cracks by fracture within the coating can be evaluated independently of the fiber. Crack deflection is usually assumed to be a local change in crack direction from perpendicular to parallel to a fiber surface (Fig. 4(a)), but there are other possibilities. Mode I interface cracks can form in the tensile stress field normal to the fiber surface ahead of a matrix crack62 (Fig. 4(b)). Modeling of an annular matrix crack has predicted that, for most reasonable choices of properties, some other failure event (e.g., fiber fracture) intervenes before the matrix crack can be driven to the interface, unless the interface is debonded ahead of it.63 Interface stresses can be high enough to make interfacial debond ahead of the matrix crack a plausible mechanism, but lack of a completely understood local failure criterion leaves the matter open to speculation. This failure sequence has been observed in laminates64 and in model composites.65 Nicalon/C/SiC composites made with fibers treated to promote higher coating/fiber interface strengths also provide indirect evidence of interfacial failure preceding matrix crack impingement.66,67 In composites made with untreated fibers, matrix cracks connect to debonded coating/fiber interfaces, with no deflection in the coating itself; i.e., the interface, not the coating, fails. In otherwise identical composites made with treated fibers, matrix cracks connect to diffuse cracks in the coating without debonding the coating/fiber interface; i.e., the coating itself fails.66,68 The interfacial toughness, friction, and composite strength are higher for the treated-fiber CMC. If the matrix crack runs through the coatings on untreated fibers before the interface debonds, it deflects in the coating, as it does in the identical coating on the treated fiber; hence, the coating/fiber interface must fail before the matrix crack enters the coating. When the crack does pass through the coating, the elastic constraint of the fiber is mostly removed by the preexisting debond, and the crack runs directly to the debonded interface. If there are truly no material differences besides interface strength, a definitive sequence of events consistent with the model is implied.63 These composites have other interesting properties that are discussed in later sections. Another deflection mechanism preceding matrix crack impingement has been suggested for a composite comprising SiC monofilaments with successive coatings of carbon and TiB2 in a glass matrix.69 The coating was calculated to be in triaxial compression, and modeling suggested that matrix cracks would run to the coating only after coating or interfacial failure. Experiment revealed that debond cracks ran very near the fiber surface except for triangular-section C rings around the fiber with their peaks at the matrix crack planes. Shear stresses on planes in the approximate orientation of the sides of the C rings (about 45° from the matrix crack plane) were calculated to be the highest coating stresses that could lead to fracture. The suggested failure sequence was (i) the matrix crack approached the coating, (ii) the coating failed in shear on planes 45° from the crack plane, ahead of the crack, and formed C rings, (iii) the coating cracks turned parallel to the fiber surface at or near the fiber surface, and (iv) the matrix crack advanced until it joined the coating shear cracks at their intersection (see Fig. 4(c)). A further possibility is growth of periodic echelon cracks. In an analysis of thin laminates with a Mode I crack normal to the interface, the maximum tensile stresses in the coatings were 45° to the interface plane, that is, parallel to a “half-turn” of the impinging crack into the interface plane.70 Failure initiated as a series of parallel “periodic echelon” microcracks on these highstress planes at the center of the coating. As the microcracks approached the coating/plate interfaces, they turned parallel to the interface and joined to form a debond. Evidence for a similarappearing but different sequence of events has been observed in monazite interlayers in Al2O3/Al2O3 laminates (Fig. 5).27 In that case, the echelon cracks appeared after initial deflection of the main crack into the coating/laminate interface. Detailed fracture observations are difficult; therefore, the sequence of events for fiber/matrix debonding in CMCs remains speculative. Debonding mechanisms may vary with the particular composite and global stress state. The coating property most important to crack deflection depends on subtle differences in constituent elastic properties and residual stresses. In the ideal case, coatings are engineered material “components” of a composite system selected for phase, microstructure, and geometry to promote a specific failure mechanism. Enhanced understanding of crack deflection is necessary to allow such a priori design. Simple models are often useful, but they can sometimes be misleading; hence, detailed analysis and comparison of microstructure, crack deflection behavior, and analytical models may be necessary. The same considerations apply to the interpretation of micromechanical tests. For example, fiber pushout/pullout tests may not directly measure the parameters that actually determine debonding during composite failure.71 It is even possible that debonding in singleFig. 4. Three possible sequences leading to crack deflection: (a) matrix crack grows into the coating and then bifurcates and turns with fronts running parallel to the fiber surface in each direction (as well as continuing on in the matrix); (b) coating or interface fails in the tensile field ahead of the matrix crack before arrival of the matrix crack at the interface region; and (c) in the matrix crack at the coating, the crack bifurcates and turns as the coating fails in shear at an intermediate angle, then turns parallel to the fiber surface at or near the fiber surface. 2602 Journal of the American Ceramic Society—Kerans et al. Vol. 85, No. 11
Interface Design for Oxidation-Resistant Ceramic Composites 2603 Alumina Alumina ng, or deflection, of a matrix crack can occur through the formation of"echelon"cracks, as shown in the optical micrograph of the phenomenon taking place within a monazite interlayer separating two Al,O, regions.(b)and(c)are schematics of the mechanisms filament microcomposites can be different from that in full debond length because of insufficient axial strength On the other composites because of the different constraints hand, the local stress state changes and this short coating crack nay not provide sufficient stress concentration to greatly influence ( Interfacial Crack Propagation fiber fracture. In the latter case. the "last" debond crack continues If debonding is along a fiber/coating or coating/matrix interface, to grow while the last coating layer develops multiple Mode I then debond propagation is determined by the interfacial energ cracks that are benign in the short term but presumably cause some d the friction generated by shear traction 72,73 If matrix crack decrease in apparent fiber strength. In the former case, this are deflected in the coating, debonding criteria and crack propa- scenario implies that(i)even if the coating deflects cracks, debond gation can be expected to be more complex. Deflection within the lengths may be short because of a nondeflecting coating/fit coating is attractive, because a layer of coating remains on the nterface, (ii) long debond lengths may require coatings with high fiber, slowing environmental degradation of the fiber. However, it xial strain to failure, as does fiber oxidation protection by a seems that the remaining coating is unlikely to remain intact coating,(iii) failure characterization may find coating/fiber inter- beyond some critical level of strain. This limits the protective face cracks even though crack deflection occurs in the coating, and function and may limit debond I Athought experiment "can (iv)residual coating layers should not be expected to"seal" fibers be illustrative(Fig. 6). We imagine that a matrix crack impinges or throughout their entire strain range. Although this discussion is through the coating away from the matrix crack plane. Eventually, the later sectore, it is consistent with the behavior observed in larg on easy-cleaving oxides, and it comprises the matrix crack bypasses ypothesis for comparison of fracture evidence coating; therefore, the matrix crack is bridged by a fiber with a thinner coating. (This remaining thickness continues to function to slow oxidation and other environmental degradation. )As the mposite is loaded further, the coated fiber is strained until the coating fails in Mode I via a surface-initiated crack. However,a coating that deflects cracks can be expected to again deflect a Mode I crack to Mode Il, leaving the fiber with a yet thinner intact coating. The strain-to-failure of thin coatings often increases with decreasing thickness: 4 therefore, the now thinner coating segment can tolerate higher strain before the deflection process repeats chaps many times. Even if the strain-to-failure does not increase as the layers become thinner. successive mode i cracks can be expected to initiate in a noncoplanar fashion, either because of random flaw distribution or biased strain fields at the tips of the debonding cracks. In either case, eventually, this Mode I coating crack impinges the fiber, where deflection is governed by a different criterion. T_/T where i refers to the coating/fiber Fig. 6. Schematic of a matrix crack impinging on a coated fiber in a interface and f to the fiber, r and z to the normals of crack planes mposite under increasing tension along the axis of the fiber(vertical):(a) initial crack deflection within a coating,(b) subsequent Mode I failure of in cylindrical coordinates with z along the fiber axis. Hence, a the coating, followed by a second deflection; and(c) additional Mode I coating can successfully deflect cracks but not provide sufficient failures and deflections. until the fiber/matrix interface is reached
filament microcomposites can be different from that in full composites because of the different constraints.72 (3) Interfacial Crack Propagation If debonding is along a fiber/coating or coating/matrix interface, then debond propagation is determined by the interfacial energy and the friction generated by shear traction.72,73 If matrix cracks are deflected in the coating, debonding criteria and crack propagation can be expected to be more complex. Deflection within the coating is attractive, because a layer of coating remains on the fiber, slowing environmental degradation of the fiber. However, it seems that the remaining coating is unlikely to remain intact beyond some critical level of strain. This limits the protective function and may limit debond length. A “thought experiment” can be illustrative (Fig. 6). We imagine that a matrix crack impinges on a coated fiber, is deflected in the coating (a debond), and advances through the coating away from the matrix crack plane. Eventually, the matrix crack bypasses the fiber, and the debond advances in the coating; therefore, the matrix crack is bridged by a fiber with a thinner coating. (This remaining thickness continues to function to slow oxidation and other environmental degradation.) As the composite is loaded further, the coated fiber is strained until the coating fails in Mode I via a surface-initiated crack. However, a coating that deflects cracks can be expected to again deflect a Mode I crack to Mode II, leaving the fiber with a yet thinner intact coating. The strain-to-failure of thin coatings often increases with decreasing thickness;74 therefore, the now thinner coating segment can tolerate higher strain before the deflection process repeats, perhaps many times. Even if the strain-to-failure does not increase as the layers become thinner, successive Mode I cracks can be expected to initiate in a noncoplanar fashion, either because of random flaw distribution or biased strain fields at the tips of the debonding cracks. In either case, eventually, this Mode I coating crack impinges the fiber, where deflection is governed by a different criterion, i r/ f z, where i refers to the coating/fiber interface and f to the fiber, r and z to the normals of crack planes in cylindrical coordinates with z along the fiber axis. Hence, a coating can successfully deflect cracks but not provide sufficient debond length because of insufficient axial strength. On the other hand, the local stress state changes and this short coating crack may not provide sufficient stress concentration to greatly influence fiber fracture. In the latter case, the “last” debond crack continues to grow while the last coating layer develops multiple Mode I cracks that are benign in the short term but presumably cause some decrease in apparent fiber strength. In the former case, this scenario implies that (i) even if the coating deflects cracks, debond lengths may be short because of a nondeflecting coating/fiber interface, (ii) long debond lengths may require coatings with high axial strain to failure, as does fiber oxidation protection by a coating, (iii) failure characterization may find coating/fiber interface cracks even though crack deflection occurs in the coating, and (iv) residual coating layers should not be expected to “seal” fibers throughout their entire strain range. Although this discussion is largely speculative, it is consistent with the behavior observed in the later section on easy-cleaving oxides, and it comprises a hypothesis for comparison of fracture evidence. Fig. 5. (a) Blunting, or deflection, of a matrix crack can occur through the formation of “echelon” cracks, as shown in the optical micrograph of the phenomenon taking place within a monazite interlayer separating two Al2O3 regions. (b) and (c) are schematics of the mechanisms.27 Fig. 6. Schematic of a matrix crack impinging on a coated fiber in a composite under increasing tension along the axis of the fiber (vertical): (a) initial crack deflection within a coating; (b) subsequent Mode I failure of the coating, followed by a second deflection; and (c) additional Mode I failures and deflections, until the fiber/matrix interface is reached. November 2002 Interface Design for Oxidation-Resistant Ceramic Composites 2603
2604 Journal of the American Ceramic Socien-Kerans et aL. Vol. 85. No. 1I Consideration of protection of fibers by residual coating lay on interfacial stresses and sliding friction. Realization that rough- es the issue of the degree of protection that might be expecte ness misfit effects can be substantial in oxide coatings has led to nra has discussed the issue of Sic-fiber protection from reexamination of conventional composites for coating thicknes oxidation in some detail. It is evident that very thin coatings can compliance effects. Modeling has shown that roughness in slow oxidation only to a limited degree. Small-diameter fibers- creases the compressive radial stress in a hypothetical uncoated Sic filaments are typically 8-12 um in diameter--are desirable Nicalon fiber/SiC composite from "150 MPa before sliding to 450 for easy handling, weaving, and shape-making, but the surface/ MPa after sliding. These stresses are decreased by 1/3 by including volume ratio is very high. Consequently, oxidation depths that are a 0.5 um thick carbon coating, therefore, changes in coating thickness can be expected to affect debond length and composite properties. In general, oxides are less compliant than carbon and BN; therefore, thicker coatings are required to similarly accom- n modate misfit stresses. Assuming a Nicalon/SiC composite and a a. MC behavior also depends strongly on the fiber/matrix sliding practical lower limit of 70 GPa for the elastic modulus of a porous iction. The ultimate strength, strain-to-failure, matrix crack oxide, the compliance provided by 500 nm of carbon requires 2 spacing, and toughness are affecte Coulomb friction is um of oxide. If coatings of such thickness are not practical proportional to the radial clamping stress on the fiber, which can suitable friction levels may need to be engineered in other ways be caused by residual stress from differential thermal expansion or e.g., by controlling roughness, matrix compliance, and residual els and experiments focus on residual stresses, ,/- but, re- thickness are also a large volume fraction of the composite and can cently, more attention has been given to roughness-induced stress- affect other composite properties, such as modulus, thermal A large roughness effect on sliding friction has been conductivity, and thermal expansion. Astute design allows for the shown by fiber push back or"seating drop"measurements. 3 effects on composite properties. 6 nitial modeling of the roughness effectis based on an approx- mation that debond roughness of amplitude h causes a mismatch strain of h/R, where R, is the fiber radius, that adds to the thermal (6) Effects of Coating Properties on Composite Analysis mismatch strain. Experiments show that this captures major Many calculations of radial clamping stress during fiber/matrix debonding and sliding consider only the thermoelastic properties spects of the behavior for many interfacial crack roughness of the fiber and matrix. The discussion above implies that serious geometries and. for most systems. during sliding of long fiber lengths. However, modeling has shown that the effect of roug errors may result. A rigorous treatment of the coating elastic ness in the early stages of debond crack propagation(Fig. 7)can effects exists 7 but the results are not easily incorporated into be much more pronounced and can have a significant influence on existing models of behavior. An approach that utilizes an approx omposite properties. This effect is due to the initial unseating of imation of this work in a method that represents the coated fiber by the matching rough surfaces just behind the crack tip. In this an"effective "( transversely isotropic) fiber in simple fiber/matrix region, the work required to further compress the fiber and matrix composites allows simple inclusion of coating elasticity in exi to accommodate the misfit is done. Furthermore the sliding analyses. This work also indicates that many conventional urfaces are not parallel to the fiber axis therefore, there is a analyses that have neglected carbon and Bn coatings in a Nicalon/ Sic system are significantly in error, Plots of normalized elastic beehponent of applied force that increases the friction. Perhap the modulus and coefficient of thermal expansion(CTE)for isotro treated-fiber Sic composite system discussed earlier, a rough interface model is necessary to decrease pushout data, and rough ometries for which th work well for compliant(carbon, BN) coating thickness up to 6 ness appears to be the primary source of the high friction that of the fiber radius, and they give reasonable approximations for dictates the very good fracture properties. ,0 Models of such thickness up to 10%. The thickness constraints relax somewhat A-Tocesses are now available and can be used to study debonding with incr, sg S coating stiffness. Other limitations are discussed ughness contributions to composite behavior. 7, 4 Effects pre- dicted for oxide fiber coatings are discussed later elsewhere This approach is applicable to many models that assume transversely isotropic fibers. For example, effective fiber pr (5) Interfacial Layer Compliance ties can be directly used in the shear-lag models of fiber pullout Although the coating is not often explicitly considered in pushout, ,8 as well as the Budiansky-Hutchinson-Evans nalysis, the compliance of the coating can have significant effects (BHE)model for matrix-cracking stress (7 Necessary Values of Interfacial Toughness and Friction Many CMCs fit in one of two categories: those with negligible interfacial strength, moderate to low interfacial friction, and toug (bond Crack-tip behavior, and those with high interfacial strength and elastic behavior. From these categories, it often has been inferred that toughness. s, u, When combined with the ease of using one parameter to describe the interface, this practice has led to the ssumption of zero interfacial strength and constant low interfacial friction(T)in most fracture models. 75.,90 Nicalon/C/SiC composites made with fibers treated to enhand Matrix oating/fiber bond strength",o evidence interface properties that defy common assumptions regarding what is required for good te behavior. Composites made with treated fibers have 30% higher tensile strength(from 250 to 350 MPa) at the same strain-to-failure. much finer matrix crack Fig. 7. Illustration of the effect of interfacial cantly different stress-strain behavior(Fig. 9). The change is e debonding progressing away from a matrix e attributed to interfacial friction(T)that increases from -5 to -150 sion. Three different reg labeled I.II MPa. Strong and tough composites with high strain-to-failure Roughness amplitude, h, period, 2d, and R. are the (0.5%)are observed even when T 370 MPa. The high mportant parameters that influence interfacial friction. has been attributed to the decrease in effective
Consideration of protection of fibers by residual coating layers raises the issue of the degree of protection that might be expected. Luthra57 has discussed the issue of SiC-fiber protection from oxidation in some detail. It is evident that very thin coatings can slow oxidation only to a limited degree. Small-diameter fibers— SiC filaments are typically 8–12 m in diameter—are desirable for easy handling, weaving, and shape-making, but the surface/ volume ratio is very high. Consequently, oxidation depths that are insignificant in monolithics damage fibers. (4) Interfacial Friction CMC behavior also depends strongly on the fiber/matrix sliding friction. The ultimate strength, strain-to-failure, matrix crack spacing, and toughness are affected.75,76 Coulomb friction is proportional to the radial clamping stress on the fiber, which can be caused by residual stress from differential thermal expansion or misfit from roughness at the debonding interface.73,77 Most models and experiments focus on residual stresses,73,78–80 but, recently, more attention has been given to roughness-induced stresses.71,81–83 A large roughness effect on sliding friction has been shown by fiber push back or “seating drop” measurements.82,83 Initial modeling of the roughness effect73 is based on an approximation that debond roughness of amplitude h causes a mismatch strain of h/Rf , where Rf is the fiber radius, that adds to the thermal mismatch strain. Experiments show that this captures major aspects of the behavior for many interfacial crack roughness geometries and, for most systems, during sliding of long fiber lengths.77 However, modeling has shown that the effect of roughness in the early stages of debond crack propagation (Fig. 7) can be much more pronounced and can have a significant influence on composite properties. This effect is due to the initial unseating of the matching rough surfaces just behind the crack tip. In this region, the work required to further compress the fiber and matrix to accommodate the misfit is done. Furthermore, the sliding surfaces are not parallel to the fiber axis; therefore, there is a component of applied force that increases the friction. Perhaps the best example of a system where this effect is important is the treated-fiber SiC composite system discussed earlier; a roughinterface model is necessary to decrease pushout data, and roughness appears to be the primary source of the high friction that dictates the very good fracture properties.25,68 Models of such processes are now available and can be used to study debonding roughness contributions to composite behavior.71,84 Effects predicted for oxide fiber coatings are discussed later. (5) Interfacial Layer Compliance Although the coating is not often explicitly considered in analysis, the compliance of the coating can have significant effects on interfacial stresses and sliding friction. Realization that roughness misfit effects can be substantial in oxide coatings has led to reexamination of conventional composites for coating thickness/ compliance effects.85 Modeling has shown that roughness increases the compressive radial stress in a hypothetical uncoated Nicalon fiber/SiC composite from 150 MPa before sliding to 450 MPa after sliding. These stresses are decreased by 1/3 by including a 0.5 m thick carbon coating; therefore, changes in coating thickness can be expected to affect debond length and composite properties. In general, oxides are less compliant than carbon and BN; therefore, thicker coatings are required to similarly accommodate misfit stresses. Assuming a Nicalon/SiC composite and a practical lower limit of 70 GPa for the elastic modulus of a porous oxide, the compliance provided by 500 nm of carbon requires 2 m of oxide.86 If coatings of such thickness are not practical, suitable friction levels may need to be engineered in other ways, e.g., by controlling roughness, matrix compliance, and residual stress state, or by other deformation mechanisms. Coatings of such thickness are also a large volume fraction of the composite and can affect other composite properties, such as modulus, thermal conductivity, and thermal expansion. Astute design allows for the effects on composite properties.86 (6) Effects of Coating Properties on Composite Analysis Many calculations of radial clamping stress during fiber/matrix debonding and sliding consider only the thermoelastic properties of the fiber and matrix. The discussion above implies that serious errors may result. A rigorous treatment of the coating elastic effects exists,87 but the results are not easily incorporated into existing models of behavior. An approach that utilizes an approximation of this work in a method that represents the coated fiber by an “effective” (transversely isotropic) fiber in simple fiber/matrix composites allows simple inclusion of coating elasticity in existing analyses.88 This work also indicates that many conventional analyses that have neglected carbon and BN coatings in a Nicalon/ SiC system are significantly in error. Plots of normalized elastic modulus and coefficient of thermal expansion (CTE) for isotropic “effective” fibers are given in Fig. 8. There are limits to the geometries for which this approach yields good results. The plots work well for compliant (carbon, BN) coating thickness up to 6% of the fiber radius, and they give reasonable approximations for thickness up to 10%. The thickness constraints relax somewhat with increasing coating stiffness. Other limitations are discussed elsewhere.88 This approach is applicable to many models that assume transversely isotropic fibers. For example, effective fiber properties can be directly used in the shear–lag models of fiber pullout and pushout,72,81 as well as the Budiansky–Hutchinson–Evans (BHE) model89 for matrix-cracking stress. (7) Necessary Values of Interfacial Toughness and Friction Many CMCs fit in one of two categories: those with negligible interfacial strength, moderate to low interfacial friction, and tough behavior; and those with high interfacial strength and elastic behavior. From these categories, it often has been inferred that negligible interfacial strength and low friction are necessary for toughness.75,90,91 When combined with the ease of using one parameter to describe the interface, this practice has led to the assumption of zero interfacial strength and constant low interfacial friction () in most fracture models.75,90 Nicalon/C/SiC composites made with fibers treated to enhance coating/fiber bond strength25,68 evidence interface properties that defy common assumptions regarding what is required for good composite behavior.25 Composites made with treated fibers have 30% higher tensile strength (from 250 to 350 MPa) at the same strain-to-failure, much finer matrix crack spacing, and significantly different stress–strain behavior (Fig. 9). The change is attributed to interfacial friction () that increases from 5 to 150 MPa.67 Strong and tough composites with high strain-to-failure (0.5%) are observed even when 370 MPa. The high composite strength has been attributed to the decrease in effective Fig. 7. Illustration of the effect of interfacial roughness during progressive debonding progressing away from a matrix crack in a composite under tension. Three different regions, labeled I, II, and III, can be envisioned. Roughness amplitude, h, period, 2d, and fiber radius, R, are the most important parameters that influence interfacial friction.71 2604 Journal of the American Ceramic Society—Kerans et al. Vol. 85, No. 11
November 2002 Interface Design for Oxidation-Resistant Ceramic Composites 2605 1.3 0.9 10(c) t/R<6% 0.8 20(BN) s.12/Q。=12575100C 0.7 0.75 0.6 0 0.5 t<0.5μm:E≤E 04 00.2040.60.8 -0.1-0.0500.050.10.150.20.25 R}{(Ect4}1+0.5h}tR}{(axa)-1}{1+0.5(E,/E} Fig. 8. Universal plots can be used to obtain the properties of an"effective" transverse fiber that can be substituted for a fiber pl odels that use transversely isotropic moduli and CTE: plots of (a)effective modulus and CTE in the transverse direction for fibers lots are a good approximation for up to 0.5 um thick(n)coating on an 8 um fiber radiL bols e and a refer to the modulus and Cte, subscripts t, c, and f stand for the transverse, coating, and fiber, respectively. Asterisk(* ffective properties.(Plot(b) was corrected for an er in the oniginal reference, where the matrix on the right-hand side in Eq. (8)should be gauge length of bridging fibers resulting from short debond lengths oating-it can be as high as 0.7 for an elastic anisotropy that are, in turn, a consequence of high T. As discussed earlier. om the He and Hutchinson%,60 analysis is not satisfied matrix cracks in high-strength material deflect into multiple the coating. A similar discrepancy has been noted for interfacial cracks, rather than a single debond. Therefore, crack on criteria using a laminate geometry. Although this deflection for this CMC is decided primarily by fracture anisotropy result is not well understood, it is encouraging with regard to the within the coating, rather than at the coating/fiber or coating/ development of alternative coatings in that the fracture energies matrix interface. Unusual fiber pushout load-deflection curves and the sliding friction may not be required to be as low as uggest substantial effects of rough interfaces, and subsequent previously thought. In any event, many of the coating approaches analysis implies that the critical strain energy to propagate cracks discussed later are likely to exhibit sufficiently high fracture in this interfacial region may be as high as 25 J/m". This is more energy and friction to greatly restrict debond lengths. It is helpful than half the fracture energy across the strongest graphite planes to know that, although the composites discussed above exhibit The criterion of fracture energy anisotropy of 1/4 or less( for an matrix crack spacings of from one to three fiber diameters, J d prc fiber 0.6 LONGITUDINAL TENSILE STRAIN(%) Fig 9. Tensile stress-strain behaviors in tension measured on the two-dimensional SiC/SiC composites fabricated from ()untreated or()treated Nicalon (ceramic grade)fibers. Complex crack deflection within the coating on treated fibers(schematic upper left)leads to higher friction than smooth interfacial ailure with untreated fibers (lower right)
gauge length of bridging fibers resulting from short debond lengths that are, in turn, a consequence of high . As discussed earlier, matrix cracks in high-strength material deflect into multiple interfacial cracks, rather than a single debond.67 Therefore, crack deflection for this CMC is decided primarily by fracture anisotropy within the coating, rather than at the coating/fiber or coating/ matrix interface. Unusual fiber pushout load–deflection curves suggest substantial effects of rough interfaces, and subsequent analysis implies that the critical strain energy to propagate cracks in this interfacial region may be as high as 25 J/m2 . 66 This is more than half the fracture energy across the strongest graphite planes. The criterion of fracture energy anisotropy of 1/4 or less (for an isotropic coating—it can be as high as 0.7 for an elastic anisotropy of 6) from the He and Hutchinson59,60 analysis is not satisfied, even in the coating. A similar discrepancy has been noted for deflection criteria using a laminate geometry.64 Although this result is not well understood, it is encouraging with regard to the development of alternative coatings in that the fracture energies and the sliding friction may not be required to be as low as previously thought. In any event, many of the coating approaches discussed later are likely to exhibit sufficiently high fracture energy and friction to greatly restrict debond lengths. It is helpful to know that, although the composites discussed above exhibit matrix crack spacings of from one to three fiber diameters, Fig. 8. Universal plots can be used to obtain the properties of an “effective” transversely isotropic fiber that can be substituted for a fiber plus coating in models that use transversely isotropic moduli and CTE: plots of (a) effective modulus and (b) effective CTE in the transverse direction for fibers with coatings. Plots are a good approximation for up to 0.5 m thick (t) coating on an 8 m fiber radius (R). Symbols E and refer to the modulus and CTE, respectively; subscripts t, c, and f stand for the transverse, coating, and fiber, respectively. Asterisk () denotes effective properties.88 (Plot (b) was corrected for an error in the original reference, where the matrix on the right-hand side in Eq. (8) should be inverted.) Fig. 9. Tensile stress–strain behaviors in tension measured on the two-dimensional SiC/SiC composites fabricated from (I) untreated or (J) treated Nicalon (ceramic grade) fibers. Complex crack deflection within the coating on treated fibers (schematic upper left) leads to higher friction than smooth interfacial failure with untreated fibers (lower right).350 November 2002 Interface Design for Oxidation-Resistant Ceramic Composites 2605
2606 Journal of the American Ceramic Socien-Kerans et aL. Vol. 85. No. 1I implying very short debond lengths, they also demonstrate high and coating surface roughnesses. o Therefore, if debonding is strength and toughness. Nevertheless, there is such a thing as within a coating and the crack meanders in the coating, a thinner debond lengths that are too short, even though that value is coating may decrease the fracture surface roughness and, there- considerably less than has been widely assumed before analysis of fore, increase toughness. If debonding initiates and remains at the these composites. coating/fiber interface, fracture surface roughness can be varied If the coating cracks ultimately reach the coating/fiber interface, only by modifying the fiber surface roughness as discussed in Section I(), the result is apparently benign. That However, if the debonding crack tends to approach th is, either(i)the interface, although stronger than the coating itself, surface via Mode I steps as it propagates and the interfa is weak enough to fail before the fiber, (ii) the changed local stress debond criterion is not satisfied( the situation discussed in state and short crack do not pose substantial stress concentration Il(), then greater coating thickness leads to longer debond on the fiber, or (iii) the resulting failure event is sufficiently late to ngths and higher toughness. yet allow excellent composite properties There are conflicts between some coating design parameters For example, a thicker coating can provide a route to lower friction by decreasing the compressive residual stresses, but it counters that Ill. Coating System Design and Evaluation effect by allowing higher fracture surface roughness, conversely, a thin coating may contribute to decreasing friction by minimizing (I General Interface Considerations roughness, but it may fail to relieve compressive residual stress In Ideally, the choice of composite constituents and geometry such cases, two coating layers might be considered. The weak should lead to the best balance of properties throughout the crack-deflecting layer would be thin, and the compliant layer for omponent service lifetime. In fact, many possibilities must await relief of residual compressive stress would be thick. The added the development of more constituent options, and optimizing complexity and expense is not desirable, but it may not be properties requires more highly sophisticated models. Eventually, prohibitive there may be more fibers, coatings, and matrices to choose from, but, presently, composite design is constrained by constituents for which there are no viable alternatives. Likewise. mechanistic (2) CMC Design Steps understanding is incomplete and often speculative. Nevertheless, it The first step in a logical CMC design sequence might be the is useful to take a logical approach that develops a framework into choices of fiber, coating, and matrix that are thermochemical which new tools can be fitted as they become available and that stable individually and in combination in the temperature range an provide insight for the refinement of approaches and environment of interest. In practice, that condition is often The first function of the coating or interface is that it must fail elaxed to include materials that react at acceptably slow rates. In before the fiber fails, thereby removing matrix -imposed stress fact, almost no structural materials are at thermodynamic equili concentrations on the fiber. The second function is that the coating rium in their use environments. a common example of acceptable must allow some sliding along the fiber/matrix debond after environmental instability is SiC 20,- Sio,+ co,, whe deflection. As discussed earlier. results from carbon- and bN- oxidation of SiC is defined by the low diffusion rates of oxygen in interface CMCs and models for their behavior suggest that the the SiO, scale. -s The second step that must be considered in debond may be at either the fiber/coating interface or within the design is processing. Processing should not excessively degrade coating. Coating design strategies can be based on either possibil the fiber or coating, therefore, matrix choice can be, and often is For debonding at the fiber/coating interface, allowable T-/ limited by the processing requirements values based on the He and Hutchinson criterion 9,60 vary with Excessive thermal stress in the coating may cause it to sp fiber/coating elastic modulus mismatch from -0.25 for matrix processing. This is particularly important for CMC mismatch to almost 0.7 when the fiber is 6 times stiffer than the atings, because they are designed to be weak, or weaki coating or matrix, as in SiC-reinforced glass-matrix CMCs.A to the fiber. Many excellent review articles discuss similar criterion based on interface strengths also can be used. 9 debonding of coatings from thermal stress(see, for example, Ref. For debonding within the coating, fracture anisotropy of the 96). If possible, choice of a fiber-coating combination with coating is the most important parameter. Although the He and minimal thermal stress should be considered. Debonding of Hutchinson criterion is a very useful guide as discussed earlier, it coatings during handling or weaving of coated fibers might be may not always be relevant because of effects such as debonding decreased by eliminating steps that bend fibers excessively ahead of the matrix crack. Excessive handling can be avoided by applying fiber coatings to Once debonding starts. it continue to propagate as a woven cloth or, better yet, the final fiber preform, as is often done cylindrical Mode Il crack between the fiber and matrix. The length in chemical vapor infiltration(CVD) processing, rather than to fiber of the debond crack(distance from the matrix crack plane to the tows. Preform-coating processes using other than cvi or in situ debond crack tip) depends on the interfacial sliding friction. The ocesses using fiber constituents have not been demonstrated lower the friction, the longer the crack and the greater the distance opposites that perform poorly may require careful evaluation to from the matrix crack plane required to transfer the excess load on determine if an ineffective coating, a damaged coating, or a the fiber back to the matrix. Higher friction along this Mode l damaged fiber is responsible ack causes the fiber stress to decrease faster with distance from Thermal expansion mismatch, roughness, and coating com the matrix crack plane. That is, the highly stressed portion of the ance interplay determine the postslidin ber is shorter, and there is a higher probability that fibers fracture and they should be considered simultaneously. For example, if the at or near the matrix crack plane. Therefore, toughne fiber is known to have a comparatively rough surface, residual decrease with increasing friction. Friction is controlled by residual stresses should be low and coating compliance should be high and applied stress, the fracture surface and the coeffi- cient of friction Residual stress is d CTEs. the coating thickness the fiber volup tion and the use ( Coating Evaluation emperature. In many systems, the coating is the most compliant The properties a coating must possess to provide good compo component; therefore, coating thickness can provide some adjust ite properties are not well-known. Hence, coating evaluation is ment of residual stresses. Specifically, where the coating is more most convincingly done via behavior of a composite that is compliant and/or has higher thermal expansion than the other analogous to a practically usable material form: for example, in constituents, thicker coatings can be expected to provide higher sheet form with fiber volume fraction >25%. This process can be toughness development of new fiber-coating and matrix-processing metha c time consuming and expensive. Each new approach can requi Potential opposite effects of coating thickness on crack pat should be considered. The maximum fracture surface roughness is Replacement of the CMC matrix with a glass matrix that is easier bounded by the sum of the coating thickness as well as the fiber to process also can be considered for coating evaluation, although
implying very short debond lengths, they also demonstrate high strength and toughness. Nevertheless, there is such a thing as debond lengths that are too short, even though that value is considerably less than has been widely assumed before analysis of these composites. If the coating cracks ultimately reach the coating/fiber interface, as discussed in Section II(3), the result is apparently benign. That is, either (i) the interface, although stronger than the coating itself, is weak enough to fail before the fiber, (ii) the changed local stress state and short crack do not pose substantial stress concentration on the fiber, or (iii) the resulting failure event is sufficiently late to yet allow excellent composite properties. III. Coating System Design and Evaluation (1) General Interface Considerations Ideally, the choice of composite constituents and geometry should lead to the best balance of properties throughout the component service lifetime. In fact, many possibilities must await the development of more constituent options, and optimizing properties requires more highly sophisticated models. Eventually, there may be more fibers, coatings, and matrices to choose from, but, presently, composite design is constrained by constituents for which there are no viable alternatives. Likewise, mechanistic understanding is incomplete and often speculative. Nevertheless, it is useful to take a logical approach that develops a framework into which new tools can be fitted as they become available and that can provide insight for the refinement of approaches. The first function of the coating, or interface, is that it must fail before the fiber fails, thereby removing matrix-imposed stress concentrations on the fiber. The second function is that the coating must allow some sliding along the fiber/matrix debond after deflection. As discussed earlier, results from carbon- and BNinterface CMCs and models for their behavior suggest that the debond may be at either the fiber/coating interface or within the coating. Coating design strategies can be based on either possibility. For debonding at the fiber/coating interface, allowable i r/ f z values based on the He and Hutchinson criterion59,60 vary with fiber/coating elastic modulus mismatch from 0.25 for zero mismatch to almost 0.7 when the fiber is 6 times stiffer than the coating or matrix, as in SiC-reinforced glass-matrix CMCs. A similar criterion based on interface strengths also can be used.91 For debonding within the coating, fracture anisotropy of the coating is the most important parameter. Although the He and Hutchinson criterion is a very useful guide, as discussed earlier, it may not always be relevant because of effects such as debonding ahead of the matrix crack. Once debonding starts, it must continue to propagate as a cylindrical Mode II crack between the fiber and matrix. The length of the debond crack (distance from the matrix crack plane to the debond crack tip) depends on the interfacial sliding friction. The lower the friction, the longer the crack and the greater the distance from the matrix crack plane required to transfer the excess load on the fiber back to the matrix. Higher friction along this Mode II crack causes the fiber stress to decrease faster with distance from the matrix crack plane. That is, the highly stressed portion of the fiber is shorter, and there is a higher probability that fibers fracture at or near the matrix crack plane. Therefore, toughness may decrease with increasing friction. Friction is controlled by residual and applied stress, the fracture surface roughness, and the coefficient of friction.81 Residual stress is determined by constituent CTEs, the coating thickness, the fiber volume fraction, and the use temperature.88 In many systems, the coating is the most compliant component; therefore, coating thickness can provide some adjustment of residual stresses. Specifically, where the coating is more compliant and/or has higher thermal expansion than the other constituents, thicker coatings can be expected to provide higher toughness. Potential opposite effects of coating thickness on crack path should be considered. The maximum fracture surface roughness is bounded by the sum of the coating thickness as well as the fiber and coating surface roughnesses.86 Therefore, if debonding is within a coating and the crack meanders in the coating, a thinner coating may decrease the fracture surface roughness and, therefore, increase toughness. If debonding initiates and remains at the coating/fiber interface, fracture surface roughness can be varied only by modifying the fiber surface roughness. However, if the debonding crack tends to approach the fiber surface via Mode I steps as it propagates and the interface/fiber debond criterion is not satisfied (the situation discussed in Section II(3)), then greater coating thickness leads to longer debond lengths and higher toughness. There are conflicts between some coating design parameters. For example, a thicker coating can provide a route to lower friction by decreasing the compressive residual stresses, but it counters that effect by allowing higher fracture surface roughness; conversely, a thin coating may contribute to decreasing friction by minimizing roughness, but it may fail to relieve compressive residual stress. In such cases, two coating layers might be considered. The weak, crack-deflecting layer would be thin, and the compliant layer for relief of residual compressive stress would be thick. The added complexity and expense is not desirable, but it may not be prohibitive. (2) CMC Design Steps The first step in a logical CMC design sequence might be the choices of fiber, coating, and matrix that are thermochemically stable individually and in combination in the temperature range and environment of interest. In practice, that condition is often relaxed to include materials that react at acceptably slow rates. In fact, almost no structural materials are at thermodynamic equilibrium in their use environments. A common example of acceptable environmental instability is SiC 2O2 3 SiO2 CO2, where oxidation of SiC is defined by the low diffusion rates of oxygen in the SiO2 scale.93–95 The second step that must be considered in design is processing. Processing should not excessively degrade the fiber or coating; therefore, matrix choice can be, and often is, limited by the processing requirements. Excessive thermal stress in the coating may cause it to spall during matrix processing. This is particularly important for CMC fiber coatings, because they are designed to be weak, or weakly bonded, to the fiber. Many excellent review articles discuss debonding of coatings from thermal stress (see, for example, Ref. 96). If possible, choice of a fiber–coating combination with minimal thermal stress should be considered. Debonding of coatings during handling or weaving of coated fibers might be decreased by eliminating steps that bend fibers excessively. Excessive handling can be avoided by applying fiber coatings to woven cloth or, better yet, the final fiber preform, as is often done in chemical vapor infiltration (CVI) processing, rather than to fiber tows. Preform-coating processes using other than CVI or in situ processes using fiber constituents have not been demonstrated. Composites that perform poorly may require careful evaluation to determine if an ineffective coating, a damaged coating, or a damaged fiber is responsible. Thermal expansion mismatch, roughness, and coating compliance interplay to determine the postsliding stresses and friction, and they should be considered simultaneously. For example, if the fiber is known to have a comparatively rough surface, residual stresses should be low and coating compliance should be high. (3) Coating Evaluation The properties a coating must possess to provide good composite properties are not well-known. Hence, coating evaluation is most convincingly done via behavior of a composite that is analogous to a practically usable material form: for example, in sheet form with fiber volume fraction 25%. This process can be time consuming and expensive. Each new approach can require development of new fiber-coating and matrix-processing methods. Replacement of the CMC matrix with a glass matrix that is easier to process also can be considered for coating evaluation, although 2606 Journal of the American Ceramic Society—Kerans et al. Vol. 85, No. 11
November 2002 Interface Design for Oxidation-Resistant Ceramic Co 2607 the change in chemistry, and probably elastic properties, may environmental resistance have been studied. Periodic matrix introduce some ambiguity in interpretation of results. cracks, nonlinear load displacement, and hysteresis during unload- Porous-matrix CMCs without fiber coatings can have attractive reload cycles have been observed, from which debond energies properties via distributed damage mechanisms, because cracks and the average friction(T)have been estimated. o However, full deflect around fibers without need for a coating (see Section confidence in validity of the results for property prediction in a full Tv(9). Matrix pore volume fractions at which significant tough- CMC has not been established. 100 ening is observed range from >30% to 15%.98,9%Hence, porous Oxide/oxide microcomposites have been fabricated and tested matrices complicate evaluation of fiber coatings, because the to evaluate the effectiveness of monazite (LapO,)and hibonite porous matrix and the coating can contribute to toughening. CaAl12Ojg)as interlayers in sapphire reinforced/Al,O,matrix matrix composites may be necessary for complete understanding as the control composites, the fractography and fracture strengths of damage mechanisms in coated-fiber composites with imper- were compared. For interlayer thicknesses of 0.3-0.5 um, bot fectly densified matrices-usually the case interlayers showed evidence of crack deflection; however debond lengths in hibonite-coated specimens were limited to just a smal (4) Micro- and Mini-CMCs fraction of the fiber diameter. Monazite-coated specimens showed Use of micro- or mini-CMCs for more-rapid evaluation has multiple matrix cracks and extensive debonding at the coating atrix interface. In both cases, the load-displacement curves were cylindrical matrix reinforced with one fiber, while a mini-CMC almost linear to failure, therefore, there was no unload-reload ses one or multiple fiber tows(200-3000 fibers/tow and up to hysteresis from which to measure interfacial friction. Failure four tows). The mechanical behavior of a mini-CMC is more strength was the only measurable mechanical parameter. The difficult to interpret, but it includes the statistical nature of fiber extent of nonlinearity in tension of specimens of any type with fracture and is more representative of a real composite. These high fiber modulus, straight fibers d low matrix volume fractio micro- and mini-CMCs are easier to fabricate than full cMcs. and must be small. The evaluation of the results was based on the relaxed sintering constraints on matrix densification can allow hypothesis that, even if the coating and matrix volume fraction is denser matrices to be more easily made. 4 03 Most such tests very low, there is severe degradation in apparent fiber strength if ave been limited to carbon and bn fiber/matrix interfaces and there is no mechanism to deflect cracks. The matrix and coating ostly CVI-SiC matrices. Effects of fiber surface treatments or crack at relatively low strain, and, unless the crack deflects, it ac coating procedures on interface properties and evaluation of as a large flaw in the fiber. In this experiment, composite strengths Sapph I mm Hibonite TM-DAR :bonded Su √ atrix dislodge AlumIt 2 2 CMC-Control 1. 18 GPa 0 1.18 G -1 Eiber-1450.C,2h 2 2.33 GPa m-lI. Fiber-Hibonite 3 Fiber CMCMonazite 2.25 GPa m=103 m=577 -10.500.51 1.5 1-0.500.511.5 Ln I Stress, GPa I Ln Stress, GPa Fig. 10. Single-filament sapphire fiber reinforced/Al,O, matrix microcomposites tested in tension (a) cracks deflect within the hibonite interface but by matrix regions that fell off during the test; (c)and(d) ites with coatings have almost the same mean strengths as the control composite tes with coatings, but the Weibull modulus is higher, about the same as the coated fibers. Results imply that the matrix is not sufficiently dense for evaluation of the coatings, because even the control samples have high microcomposite strength
the change in chemistry, and probably elastic properties, may introduce some ambiguity in interpretation of results. Porous-matrix CMCs without fiber coatings can have attractive properties via distributed damage mechanisms, because cracks deflect around fibers without need for a coating (see Section IV(9)). Matrix pore volume fractions at which significant toughening is observed range from 30% to 15%.98,99 Hence, porous matrices complicate evaluation of fiber coatings, because the porous matrix and the coating can contribute to toughening. Therefore, better understanding of damage mechanisms in porousmatrix composites may be necessary for complete understanding of damage mechanisms in coated-fiber composites with imperfectly densified matrices—usually the case. (4) Micro- and Mini-CMCs Use of micro- or mini-CMCs for more-rapid evaluation has received increasing attention.100,101 A micro-CMC is defined as a cylindrical matrix reinforced with one fiber, while a mini-CMC uses one or multiple fiber tows (200–3000 fibers/tow and up to four tows). The mechanical behavior of a mini-CMC is more difficult to interpret, but it includes the statistical nature of fiber fracture and is more representative of a real composite. These micro- and mini-CMCs are easier to fabricate than full CMCs, and relaxed sintering constraints on matrix densification can allow denser matrices to be more easily made.102,103 Most such tests have been limited to carbon and BN fiber/matrix interfaces and mostly CVI-SiC matrices. Effects of fiber surface treatments or coating procedures on interface properties100 and evaluation of environmental resistance have been studied.101 Periodic matrix cracks, nonlinear load displacement, and hysteresis during unload– reload cycles have been observed, from which debond energies and the average friction () have been estimated.100 However, full confidence in validity of the results for property prediction in a full CMC has not been established.100 Oxide/oxide microcomposites have been fabricated and tested to evaluate the effectiveness of monazite (LaPO4) and hibonite (CaAl12O19) as interlayers in sapphire reinforced/Al2O3 matrix composites.99 Using sapphire monofilaments in an Al2O3 matrix as the control composites, the fractography and fracture strengths were compared. For interlayer thicknesses of 0.3–0.5 m, both interlayers showed evidence of crack deflection; however debond lengths in hibonite-coated specimens were limited to just a small fraction of the fiber diameter. Monazite-coated specimens showed multiple matrix cracks and extensive debonding at the coating/ matrix interface. In both cases, the load–displacement curves were almost linear to failure; therefore, there was no unload–reload hysteresis from which to measure interfacial friction.99 Failure strength was the only measurable mechanical parameter. The extent of nonlinearity in tension of specimens of any type with high fiber modulus, straight fibers, and low matrix volume fraction must be small. The evaluation of the results was based on the hypothesis that, even if the coating and matrix volume fraction is very low, there is severe degradation in apparent fiber strength if there is no mechanism to deflect cracks. The matrix and coating crack at relatively low strain, and, unless the crack deflects, it acts as a large flaw in the fiber. In this experiment, composite strengths Fig. 10. Single-filament sapphire fiber reinforced/Al2O3 matrix microcomposites tested in tension: (a) cracks deflect within the hibonite interface but debond lengths are very short, much less than a fiber diameter, because of the roughness; (b) debonds present at the monazite/matrix interface are revealed by matrix regions that fell off during the test; (c) and (d) microcomposites with coatings have almost the same mean strengths as the control composites with no coatings, but the Weibull modulus is higher, about the same as the coated fibers. Results imply that the matrix is not sufficiently dense for evaluation of the coatings, because even the control samples have high microcomposite strength.99 November 2002 Interface Design for Oxidation-Resistant Ceramic Composites 2607
2608 Journal of the American Ceramic Socien-Kerans et aL. Vol. 85. No. 11 were relatively high for both coatings, considering the fiber the nature of the process. Although the desired Al,O, phase strengths were not significantly different from that of the fiber/ How ever, the stability of the interface coating during CVD matrix control specimens, although coated-fiber composites had processing is unknown and likely to be a major issue because of higher Weibull moduli. The lack of difference in strength is the use of gaseous hydrogen and Co in the process attributed to the porosity in the matrix; porous-matrix composites This has led to the use of glass matrices to test coating concepts are known to perform well without interface treatments(see next section). The results imply that the matrix density needs to be Allied Signal, Inc(now 85% to evaluate novel interface strategies reliably. 9 Honeywell), Morristown, N)polymer-derived glass as matrix shows some promise, although the matrices remain far from ideal The processing of even minicomposites having a dense oxide blackglas yields a matrix that is locally dense but filled with matrix can be challenging. Use of chemical vapor deposition (CVD) to deposit oxides remains in the developmental stage array of shrinkage microcracks. Oxide-fiber-reinforced minicom- CVD-deposited Al2O, matrices are amorphous and do not bond posites having a dense but microcracked glassy matrix of Black- readily to coated or uncoated fiber tows, which causes debonding, glas have been used in two studies to test oxidation-resistant even in control specimens. There is no known work on In one study, hnique was used t polycrystalline-oxide-matrix composites with high enough matrix 610(3M Corp, St Paul, MNyBlackglas composites with and densities to definitively suppress the mechanism of debonding vi matrix cracking(say 90%) CVI of dense stable polycrystallin minicomposites with the fiber coatings had significar oxides is made difficult by the formation of amorphous ultimate strengths than the uncoated control specime metastable oxides(which later crystallize or transform, introducin another study, porous oxide(zrO2-SiO2 mixture) and significant stresses and cracking)and by the inability to reach were evaluated in Nextel 720m-reinforced Blackglas porosity levels below the permeation threshold (-15%)because of coated and uncoated fibers were used as controls for comparison (a Control (uncoated b)BN CMC- Control (uncoated) 265MPa;m=88 0 742MP 383 MPa 55 5.5 5.5 65 Ln stress, MPa 1 Ln Stress, MPa] c) Porous ZrO2-SiO2 (d Monazite CMC. Control cMc· Contro coated) 265MPa;m=8.8 MPa; m=8.8 234 己5,5 CMC-porous Zro-sio cMc· Monazite 356MPa;m=6.0 353MPa;m=88 55 65 5.5 6.5 Ln Stress, MPa Ln stress, MPa] Fig. 11. Weibull plots of the strengths of minicomposites using dense Blackglas as the matrix show that porous(c)ZrO2-SiO2 and(d) monazite coatings on Nextel 720 fibers are as effective as the(b) BN-coated fibers (a) Control is significantly weaker than the fiber, showing that Blackglas might be a good model matrix to evaluate interface coatings
were relatively high for both coatings, considering the fiber strength degradation during processing; the strengths were greater than the matrix-cracking stresses (Fig. 10). However, the mean strengths were not significantly different from that of the fiber/ matrix control specimens, although coated-fiber composites had higher Weibull moduli. The lack of difference in strength is attributed to the porosity in the matrix; porous-matrix composites are known to perform well without interface treatments (see next section). The results imply that the matrix density needs to be 85% to evaluate novel interface strategies reliably.99 The processing of even minicomposites having a dense oxide matrix can be challenging. Use of chemical vapor deposition (CVD) to deposit oxides remains in the developmental stage. CVD-deposited Al2O3 matrices are amorphous and do not bond readily to coated or uncoated fiber tows, which causes debonding, even in control specimens.104 There is no known work on polycrystalline-oxide-matrix composites with high enough matrix densities to definitively suppress the mechanism of debonding via matrix cracking (say 90%).99 CVI of dense stable polycrystalline oxides is made difficult by the formation of amorphous or metastable oxides (which later crystallize or transform, introducing significant stresses and cracking) and by the inability to reach porosity levels below the permeation threshold (15%) because of the nature of the process. Although the desired Al2O3 phase remains difficult to process, work has been reported where almost 85%-dense ZrO2 has been deposited around woven preforms.105 However, the stability of the interface coating during CVD processing is unknown and likely to be a major issue because of the use of gaseous hydrogen and CO in the process. This has led to the use of glass matrices to test coating concepts. Preliminary work using BlackglasTM (Allied Signal, Inc. (now Honeywell), Morristown, NJ) polymer-derived glass as matrix shows some promise, although the matrices remain far from ideal. Blackglas yields a matrix that is locally dense but filled with an array of shrinkage microcracks. Oxide-fiber-reinforced minicomposites having a dense but microcracked glassy matrix of Blackglas have been used in two studies to test oxidation-resistant coatings. In one study, the technique was used to evaluate Nextel 610TM (3M Corp., St. Paul, MN)/Blackglas composites with and without porous lanthanum hexaluminate fiber coatings.106 The minicomposites with the fiber coatings had significantly higher ultimate strengths than the uncoated control specimens. In a another study, porous oxide (ZrO2–SiO2 mixture) and monazite were evaluated in Nextel 720TM-reinforced Blackglas.107 BNcoated and uncoated fibers were used as controls for comparison. Fig. 11. Weibull plots of the strengths of minicomposites using dense Blackglas as the matrix show that porous (c) ZrO2–SiO2 and (d) monazite coatings on Nextel 720 fibers are as effective as the (b) BN-coated fibers. (a) Control is significantly weaker than the fiber, showing that Blackglas might be a good model matrix to evaluate interface coatings.107 2608 Journal of the American Ceramic Society—Kerans et al. Vol. 85, No. 11