Journal of Microscopy, VoL 196, Pt 2 November 1999, Pp. 194-202 Received 9 June 1998: accepted 15 Apri 1999 Transmission electron microscopy of interfaces in structural ceramic composites K M. KNOWLES. S. TURAN. t A. KUMAR.$ S.-. ChEN#& W.. CLEGG* University of Cambridge, Department of Materials Science and Metallurgy, Pembroke Street. fAnadolu Universites, Iki Eyliil Kampiisii, Seramik Miihendisligi BOliimii, 26470 Eskisehir BM Storage System Division, San Jose, CA 95193. U.s.A. #National Center for Electron Microscopy, Lawrence Berkeley Laboratory, Berkeley, CA 94720. Key words. Ceramics, chemical analysis, composites, fibres, interfaces. laminates. particles, transmission electron microscopy. slow. It is only recently that Si-C-O fibre-reinforced glass components have found a cost-effective use as replacement Ceramic composites based either on a particulate, fibre or a materials for handling hot glass during glass manufactur- lamellar architecture are potentially useful as damage tolerant high-temperature engineering materials. The ing operations (Beier Markman, 1997). For these ability of the interfaces in such systems to deflect cracks is applications, the tolerance of the materials to thermal vital to the damage tolerance of these materials. Transmis- shock and to modest heat of the order of 500C is being sion electron microscopy techniques enable the chemical exploited, rather than any long-term, high temperature and physical characterization of these interfaces, providing damage-tolerant ability. Further uses are anticipated information on interlayer thicknesses, chemical species exploiting the tribological properties of fibre-reinforced local bonding and the microstructural features which give lasses. There is an interesting parallel in the development rise to the interfacial properties, thereby enabling a full of fibre-reinforced ceramics and the development of understanding not only of composites after processing. but structural ceramic nanocomposites, for which applications also after exposure to aggressive environments such as air as wear parts are also envisaged at present for systems such at high temperature. Examples of the application of as SiC nanoparticles in an Al,O3 matrix(Davidge et aL. transmission electron microscopy to all three composite 1997: Sternitzke. 1997). rather than as high temperature. architectures are described damage-tolerant materials, despite claims in the literature of increases in both strength and toughness for ceramic nanocomposites relative to monolithic material. Indeed, it is Introduction probably true to say that, as with many other materials One of the major goals of research in structural ceramics research initiatives. the attempt to develop high tempera- in the latter half of this century has been the production ture, damage-tolerant ceramics has generated a knowledge of a range of high temperature, damage-tolerant, ceramic areas of ceramics technology. while failing to deliver components. Ceramic composites based either ona particulate, fibre or a lamellar architecture are potentially materials for the original intended end use(Campbell attractive materials in this regard, as toughening mechan 1997) isms such as crack deflection and fibre pull-out, more potent Central to the damage-tolerant behaviour of ceramic than those mechanisms available in monolithic ceramics composites at room temperature and elevated temperatures e able to impart damage tolerance(Evans Marshall is the chemical stability of interfaces between different 1989: Clegg, 1992). However, progress in the development phases formed du of technologically useful material has been disappointingly processing heat treatments. Interface layers can be pro duced either by in-situ chemical reactions during proces- spondence to: Dr K. M. Knowles. Tel: 44(0)1223334312: fax: +4 sing, or as coatings deposited by a separate processing step 23334567: e-mail: kmk10acam ac uk such as bn coatings on Si-C-O fibres(Sheldon et aL., 1996 194 0 1999 The Royal Microscopical Society
Transmission electron microscopy of interfaces in structural ceramic composites K. M. KNOWLES,* S. TURAN,† A. KUMAR,§ S.-J. CHEN# & W. J. CLEGG* *University of Cambridge, Department of Materials Science and Metallurgy, Pembroke Street, Cambridge CB2 3QZ, U.K. †Anadolu U¨ niversitesi, Iki Eylu¨l Kampu¨su¨, Seramik Mu¨hendisligˇi Bo¨lu¨mu¨, 26470 Eskis¸ehir, Turkey §IBM Storage System Division, San Jose´, CA 95193, U.S.A. #National Center for Electron Microscopy, Lawrence Berkeley Laboratory, Berkeley, CA 94720, U.S.A. Key words. Ceramics, chemical analysis, composites, fibres, interfaces, laminates, particles, transmission electron microscopy. Summary Ceramic composites based either on a particulate, fibre or a lamellar architecture are potentially useful as damagetolerant high-temperature engineering materials. The ability of the interfaces in such systems to deflect cracks is vital to the damage tolerance of these materials. Transmission electron microscopy techniques enable the chemical and physical characterization of these interfaces, providing information on interlayer thicknesses, chemical species, local bonding and the microstructural features which give rise to the interfacial properties, thereby enabling a full understanding not only of composites after processing, but also after exposure to aggressive environments such as air at high temperature. Examples of the application of transmission electron microscopy to all three composite architectures are described. Introduction One of the major goals of research in structural ceramics in the latter half of this century has been the production of a range of high temperature, damage-tolerant, ceramic components. Ceramic composites based either on a particulate, fibre or a lamellar architecture are potentially attractive materials in this regard, as toughening mechanisms such as crack deflection and fibre pull-out, more potent than those mechanisms available in monolithic ceramics, are able to impart damage tolerance (Evans & Marshall, 1989; Clegg, 1992). However, progress in the development of technologically useful material has been disappointingly slow. It is only recently that Si–C–O fibre-reinforced glass components have found a cost-effective use as replacement materials for handling hot glass during glass manufacturing operations (Beier & Markman, 1997). For these applications, the tolerance of the materials to thermal shock and to modest heat of the order of 500 8C is being exploited, rather than any long-term, high temperature, damage-tolerant ability. Further uses are anticipated exploiting the tribological properties of fibre-reinforced glasses. There is an interesting parallel in the development of fibre-reinforced ceramics and the development of structural ceramic nanocomposites, for which applications as wear parts are also envisaged at present for systems such as SiC nanoparticles in an Al2O3 matrix (Davidge et al., 1997; Sternitzke, 1997), rather than as high temperature, damage-tolerant materials, despite claims in the literature of increases in both strength and toughness for ceramic nanocomposites relative to monolithic material. Indeed, it is probably true to say that, as with many other materials research initiatives, the attempt to develop high temperature, damage-tolerant ceramics has generated a knowledge base which has enabled significant advances in a number of areas of ceramics technology, while failing to deliver materials for the original intended end use (Campbell, 1997). Central to the damage-tolerant behaviour of ceramic composites at room temperature and elevated temperatures is the chemical stability of interfaces between different phases formed during processing and subsequent postprocessing heat treatments. Interface layers can be produced either by in-situ chemical reactions during processing, or as coatings deposited by a separate processing step, such as BN coatings on Si–C–O fibres (Sheldon et al., 1996; Journal of Microscopy, Vol. 196, Pt 2, November 1999, pp. 194–202. Received 9 June 1998; accepted 15 April 1999 194 q 1999 The Royal Microscopical Society Correspondence to: Dr K. M. Knowles. Tel:þ 44 (0)1223 334312; fax: þ44 (0)1223 334567; e-mail: kmk10@cam.ac.uk
INTERFACES IN STRUCTURAL CERAMIC COMPOSITES 19 Sun et al. 1996)or layered oxides with weak shear planes intergranular films. Such a even if partially as coatings in either fibrous or lamellar oxide composites recrystallised, will contribute te gth degradation at (Cinibulk Hay, 1996). An interesting recent development high temperatures by viscous The distribution of in this regard is a tough, thermally conductive, high volume remnant liquid phases in ceramic composites is therefore of fraction Si-Al-C-O fibre-based ceramic in which a thin particular interest interfacial carbon layer is created between adjacent fibres The existence of thin( l nm wide)amorphous films at during processing(Ishikawa et aL, 1998) general high-angle grain boundaries arizing from silica-rich Transmission electron microscopy (TEM) is a particularly liquid phases present during high temperature heat treat powerful research tool for the chemical and physical ments is now well established experimentally( Kleebe et al characterization of interfaces in ceramic composites. Such 1992; Turan Knowles, 1995a)and can be understood interfaces tend to be much more complex chemically than theoretically in terms of the competition between attractive in metallic materials. TEM techniques can provide informa- dispersion forces between grains forcing out the films from tion on interlayer thicknesses, chemical species, local the boundaries and repulsive disjoining forces from steric and the nature of crystalline and amorphous forces and electrical double-layer forces enabling films to roducts, thereby enabling a full understanding not only of be retained at the boundaries(Clarke, 1987: Clarke et al., composites after processing, but also after exposure to 1993) aggressive environments, such as air at high temperature. Our recent work in this area has concentrated on Thus, for example, the nce of chemical reactions particulate Si3 N4-SiC composites in which small h-BN which take place at a sub-micrometre level at fibre-matrix inclusions occur as a by-product of the processing route interfaces in oxidizing environments, and which lead to a (Turan& Knowles, 1995b). This multiphase material has decrease in damage tolerance of fibre-reinforced ceramics, proven to be very profitable for examining a range of can be specified by TEM procedures. Ultimately, such ceramic interphase boundaries. We have demonstrated information should enable more oxidation-resistant and experimentally an orientation dependence on the thickness damage-tolerant interfaces to be engineered in such of the silica-rich film seen at h-BN-SiC interphase bound materials aries, such as those shown in Fig. 1(Turan Knowles Although TEM has been used extensively to investigate 1997a). This orientation dependence can be understood m四 interfaces in particulate and fibre-reinforced ceramics, there theoretically in te isotropy of the Hamaker arizing from the anisotropy in the f h-Bn parallel and perpendicular to technique. In this review paper we give examples of the microstructural characterization of all three composite architectures, and show how an understanding of the microstructure helps to explain observed mechanical behaviour. as a function of heat treatment. and hence h-BN service lifetimes, of the current generation of ceramic Interphase boundaries in particulate engineering 3C Sic One way of trying to induce modest increases in toughness in conventional monolithic ceramics such as silicon carbide is to incorporate particles of a second phase such as silicon nitride. An alternative to this is to produce ceramic h-BN nanocomposites. in which both the matrix and the reinforcing phase are nanosized (Sternitzke, 1997). In order to densify powder compacts of such materials by either route, sintering aids are normally added at low concentra- tion levels. These sintering aids form a liquid at high 3C SIC temperature and can enable compacts of nearly theoretical density to be produced. However, this is at the expense of Fig. 1. Two differently misoriented interphase boundaries between eaving behind residual liquid which invariably cools to an h-BN and 3C Sic grains showing the effect of misorientation on amorphous phase present at triple junctions and as thin film thickness e 1999 The Royal Microscopical Society, Journtl of Microscop 196. 19+-202
Sun et al., 1996) or layered oxides with weak shear planes as coatings in either fibrous or lamellar oxide composites (Cinibulk & Hay, 1996). An interesting recent development in this regard is a tough, thermally conductive, high volume fraction Si–Al–C–O fibre-based ceramic in which a thin interfacial carbon layer is created between adjacent fibres during processing (Ishikawa et al., 1998). Transmission electron microscopy (TEM) is a particularly powerful research tool for the chemical and physical characterization of interfaces in ceramic composites. Such interfaces tend to be much more complex chemically than in metallic materials. TEM techniques can provide information on interlayer thicknesses, chemical species, local bonding and the nature of crystalline and amorphous products, thereby enabling a full understanding not only of composites after processing, but also after exposure to aggressive environments, such as air at high temperature. Thus, for example, the sequence of chemical reactions which take place at a sub-micrometre level at fibre–matrix interfaces in oxidizing environments, and which lead to a decrease in damage tolerance of fibre-reinforced ceramics, can be specified by TEM procedures. Ultimately, such information should enable more oxidation-resistant and damage-tolerant interfaces to be engineered in such materials. Although TEM has been used extensively to investigate interfaces in particulate and fibre-reinforced ceramics, there are few reported TEM observations of interfaces in lamellar ceramic composites, for which scanning electron microscopy has to date been the preferred microscopical technique. In this review paper we give examples of the microstructural characterization of all three composite architectures, and show how an understanding of the microstructure helps to explain observed mechanical behaviour, as a function of heat treatment, and hence service lifetimes, of the current generation of ceramic composites. Interphase boundaries in particulate engineering ceramics One way of trying to induce modest increases in toughness in conventional monolithic ceramics such as silicon carbide is to incorporate particles of a second phase such as silicon nitride. An alternative to this is to produce ceramic nanocomposites, in which both the matrix and the reinforcing phase are nanosized (Sternitzke, 1997). In order to densify powder compacts of such materials by either route, sintering aids are normally added at low concentration levels. These sintering aids form a liquid at high temperature and can enable compacts of nearly theoretical density to be produced. However, this is at the expense of leaving behind residual liquid which invariably cools to an amorphous phase present at triple junctions and as thin intergranular films. Such a phase, even if partially recrystallised, will contribute to strength degradation at high temperatures by viscous flow. The distribution of remnant liquid phases in ceramic composites is therefore of particular interest. The existence of thin (< 1 nm wide) amorphous films at general high-angle grain boundaries arizing from silica-rich liquid phases present during high temperature heat treatments is now well established experimentally (Kleebe et al., 1992; Turan & Knowles, 1995a) and can be understood theoretically in terms of the competition between attractive dispersion forces between grains forcing out the films from the boundaries and repulsive disjoining forces from steric forces and electrical double-layer forces enabling films to be retained at the boundaries (Clarke, 1987; Clarke et al., 1993). Our recent work in this area has concentrated on particulate Si3N4–SiC composites in which small h-BN inclusions occur as a by-product of the processing route (Turan & Knowles, 1995b). This multiphase material has proven to be very profitable for examining a range of ceramic interphase boundaries. We have demonstrated experimentally an orientation dependence on the thickness of the silica-rich film seen at h-BN–SiC interphase boundaries, such as those shown in Fig. 1 (Turan & Knowles, 1997a). This orientation dependence can be understood theoretically in terms of the anisotropy of the Hamaker constants of interfaces arizing from the anisotropy in the refractive indices of h-BN parallel and perpendicular to [0001]. q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202 Fig. 1. Two differently misoriented interphase boundaries between h-BN and 3C SiC grains showing the effect of misorientation on film thickness. INTERFACES IN STRUC TURAL CERAMIC COMP OSITES 195
196 K. M. KNOWLES ET AL Fibre-matrix interfaces in fibre-reinforced glass ceramics While the existence of carbon-rich layers in tough Si-C-O fibre-reinforced lithium aluminosilicates has been known for some time(Brennan, 1986: Brennan. 1987). it is only recently that a relatively unambiguous identification of the structure and chemistry of interfacial layers in as-processed and oxidized Si-C-O fibre-reinforced magnesium alum silicates(MAS)has been made by transmission electron microscope techniques(Kumar Knowles, 1996a, b).The recent work by Le Strat et al. (1998)on Nicalon fibre-Pyrex h-BN glass matrix composites has shown the complexity of the 5 nm reaction products at the fibre-matrix interfaces using a variety of spectroscopic techniques, highlighting in parti- cular the difficulty in distinguishing via electron energy loss spectroscopy(EELS) between amorphous silica-rich Fig. 2. A regularly stepped interphase boundary between an and amorphous oxygen-rich silicon oxycarbide interphase h-Bn particle and a B-Si3 N4 grain where the beam direction is 120 h-BNII[0001 8- as-processed material, a simplification of the reaction events at the fibre-matrix interfaces during processing is that carbon and silica occur as a diphasic layer adjacent to By comparison, we have consistently observed non- the fibres at fibre-matrix interfaces through the oxidation random orientation relationships at h-BN-Si3N4 interface in the same samples, and a lack of evidence for amorphous Sic(s)+O2(g)- Sio2(g)+C(s) intergranular films at these interfaces, as in the example This chemical reaction occurs during composite pre in Fig. 2(Turan Knowles, 1997b). Such interphase even when hot pressing in graphite dies and is favor boundaries can be understood qualitatively on the basis of thermodynamically up to 1500C in comparison with other ear-coincidence lattice theory as being somehow special oxidation reactions for SiC (Cooper Chyung. 1987) in terms of low misfit and therefore presumably of low In addition to the diphasic layer, there can be a thin energy, enabling them to expel liquid phase during the diffusion layer of elements diffusing from the matrix into the processing heat treatment fibres, and there can also be a separate carbon layer. On the basis of geometrical near-coincidence lattice depending on the processing conditions used(Kumar theory arguments alone, there is no orientation relationship Knowles, 1996a). Le Strat et al.(1998) have shown via between h-bN and Sic which leads to low misfit either X-ray photoelectron spectroscopy(XPS) that the silica-rich two dimensions or in three dimensions, even though it is interfacial reaction product between the fibre and the found experimentally that it is common for the(0001) basal carbon layer adjacent to the matrix in the composites they plane of h-BN particles to be parallel to(111)3C SiC planes examined consisted of some silicon carbide, some carbon Turan Knowles, 1996). Thus, the clear observation of and at least two oxygen-rich silicon oxycarbides. It is likely amorphous material at h-BN-SiC interphase boundaries that other systems exhibit equally complicated reaction contrast to the absence of amorphous material at h-BN- products, for which XPS, rather than EELS, would appear to Si3n4 interphase boundaries can be rationalized on the be the preferred experimental technique for spectral basis of simple geometrical models, although the need for deconvolution into the various chemical species present further research on both direct measurements and computer The detail of the fibre-matrix interphase region deter- calculations of interfacial energies is evident. Given the mines the degree of fibre pull-out that a composite will interest in the development of orientation relationships and exhibit. By correlating interfacial microstructure with the presence or absence of amorphous material at interfaces composite mechanical properties, it can be shown that in the context of nanocomposites (Pan et aL. 1996). well-developed separate turbostratic carbon layers at the particularly in view of claims of enhanced strength and fibre-matrix interfaces, in which the basal planes are toughness of ceramic nanocomposites relative to more orientated parallel to the fibre axes. clearly facilitate conventionally processed materials, interfaces in particulate debonding in comparison with composites in which a engineering ceramics are likely to remain of continuing more intimately mixed carbon and silica diphasic layer is interest, benefiting in the future from the advances in produced at the fibre-matrix interfaces(Kumar& Knowles, interface computer modelling made possible by the latest 1996c) Composites subjected to oxidation heat treatments generation of supercomputers. are able to modify the diphasic layer to create silica-rich e 199y The Royal Microscopical Society, Jotarntl of Microscop 196, 19+-202
By comparison, we have consistently observed nonrandom orientation relationships at h-BN–Si3N4 interfaces in the same samples, and a lack of evidence for amorphous intergranular films at these interfaces, as in the example in Fig. 2 (Turan & Knowles, 1997b). Such interphase boundaries can be understood qualitatively on the basis of near-coincidence lattice theory as being somehow ‘special’ in terms of low misfit and therefore presumably of low energy, enabling them to expel liquid phase during the processing heat treatment. On the basis of geometrical near-coincidence lattice theory arguments alone, there is no orientation relationship between h-BN and SiC which leads to low misfit either in two dimensions or in three dimensions, even though it is found experimentally that it is common for the (0001) basal plane of h-BN particles to be parallel to (111) 3C SiC planes (Turan & Knowles, 1996). Thus, the clear observation of amorphous material at h-BN–SiC interphase boundaries in contrast to the absence of amorphous material at h-BN– Si3N4 interphase boundaries can be rationalized on the basis of simple geometrical models, although the need for further research on both direct measurements and computer calculations of interfacial energies is evident. Given the interest in the development of orientation relationships and the presence or absence of amorphous material at interfaces in the context of nanocomposites (Pan et al., 1996), particularly in view of claims of enhanced strength and toughness of ceramic nanocomposites relative to more conventionally processed materials, interfaces in particulate engineering ceramics are likely to remain of continuing interest, benefiting in the future from the advances in interface computer modelling made possible by the latest generation of supercomputers. Fibre–matrix interfaces in fibre-reinforced glass ceramics While the existence of carbon-rich layers in tough Si–C–O fibre-reinforced lithium aluminosilicates has been known for some time (Brennan, 1986; Brennan, 1987), it is only recently that a relatively unambiguous identification of the structure and chemistry of interfacial layers in as-processed and oxidized Si–C–O fibre-reinforced magnesium aluminosilicates (MAS) has been made by transmission electron microscope techniques (Kumar & Knowles, 1996a, b). The recent work by Le Strat et al. (1998) on Nicalon fibre–Pyrex glass matrix composites has shown the complexity of the reaction products at the fibre–matrix interfaces using a variety of spectroscopic techniques, highlighting in particular the difficulty in distinguishing via electron energy loss spectroscopy (EELS) between amorphous silica-rich and amorphous oxygen-rich silicon oxycarbide interphase products. In as-processed material, a simplification of the reaction events at the fibre–matrix interfaces during processing is that carbon and silica occur as a diphasic layer adjacent to the fibres at fibre–matrix interfaces through the oxidation reaction SiCðsÞ þ O2ðgÞ → SiO2ðgÞ þ CðsÞ This chemical reaction occurs during composite processing even when hot pressing in graphite dies and is favourable thermodynamically up to 1500 8C in comparison with other oxidation reactions for SiC (Cooper & Chyung, 1987). In addition to the diphasic layer, there can be a thin diffusion layer of elements diffusing from the matrix into the fibres, and there can also be a separate carbon layer, depending on the processing conditions used (Kumar & Knowles, 1996a). Le Strat et al. (1998) have shown via X-ray photoelectron spectroscopy (XPS) that the silica-rich interfacial reaction product between the fibre and the carbon layer adjacent to the matrix in the composites they examined consisted of some silicon carbide, some carbon and at least two oxygen-rich silicon oxycarbides. It is likely that other systems exhibit equally complicated reaction products, for which XPS, rather than EELS, would appear to be the preferred experimental technique for spectral deconvolution into the various chemical species present. The detail of the fibre–matrix interphase region determines the degree of fibre pull-out that a composite will exhibit. By correlating interfacial microstructure with composite mechanical properties, it can be shown that well-developed separate turbostratic carbon layers at the fibre–matrix interfaces, in which the basal planes are orientated parallel to the fibre axes, clearly facilitate debonding in comparison with composites in which a more intimately mixed carbon and silica diphasic layer is produced at the fibre–matrix interfaces (Kumar & Knowles, 1996c). Composites subjected to oxidation heat treatments are able to modify the diphasic layer to create silica-rich Fig. 2. A regularly stepped interphase boundary between an h-BN particle and a b-Si3N4 grain where the beam direction is [1120]h-BN jj[0001] b-Si3N4 . 196 K. M. KNOWLES ET AL. q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202
INTERFACES IN STRUCTURAL CERAMIC COMPOSITES 197 morphologies in the surface regions of oxidized samples and of load transfer between the matrix and the fibres(Knowles complex multilayered carbon-containing morphologies in et aL, 1997). Thus, around this temperature. the advantage the centre of samples where the partial pressure of oxygen is gained by reinforcing the relatively weak brittle glass- low(Kumar& Knowles, 1996b). Such composites are still ceramic matrix with high strength fibres. thereby producing able to fail non-catastrophically even after 120 h in air in high strength, damage tolerant materials, is lost, and the the temperature range 1000-1200c because of the composite strength reduces towards that of the matrix. carbon interlayers developed in the interior of the compo- However. as not all the fibre-matrix interfaces will be sites during these heat treatments(Kumar Knowles. affected in the same way, because of local microstructural 996d factors affecting the oxidation kinetics, some fibres may still The interfacial microstructure of Si-C-O fibre-reinforced contribute to the strength of the composites, even after barium magnesium aluminosilicate containing 5 wt extended isothermal exposure. It is only at temperatures potassium borosilicate glass developed by Corning Incorpo- above 1000C that the voids formed by oxidation of the rated as a material which has an improved resistance to turbostratic carbon at the fibre-matrix interfaces are closed oxidation embrittlement over other fibre-reinforced glass by the formation of silica, enabling contact between the ceramics(D. C. Larsen, personal communication, 1996) is fibre and matrix to be resumed, and allowing most of the similar to the Si-C-O/MAS composites examined by Kumar room temperature strength and damage tolerance to be Knowles. In the example shown in Fig. 3, taken from retained after 100 h isothermal heat treatments(Knowles as-processed material, conventional bright field imaging of a et aL., 1997) fibre-matrix interface edge-on to the electron beam is able The significance of microstructural observations of fibre to distinguish a 90-nm thick interfacial zone containing two reinforced ceramics has been to establish microstructural interlayers 1 and 2, between the fibre and the matrix principles determining the toughness of fibre-reinforced (Knowles et al., 1997). Energy dispersive X-ray spectroscopy ceramics in the absence of separate costly processing steps (EDX) of the interlayers shown in Figs 3(b)and(c)confirms such as laying down carbon coatings on fibres prior to that the lighter contrast interlayer is carbon-rich and that embedding them in a ceramic matrix. It is therefore the darker layer is both silica-rich and barium-rich. somewhat ironic that these principles are now being used consistent with the contrast from the two layers to specify advanced fibre coating methodologies in an After 100 h in air at 704C, oxidation of the carbon-rich attempt to overcome the limitations of carbon as an interlayer produces debond cracks running parallel to the interfacial layer in high-temperature oxidizing environments axis of the fibres, as a result of which there will be a loss (see, for example, Cinibulk Hay, 1996) Fibre Matrix Interlayers 2 00 Fig. 3.(a)Example of a fibre-matrix interface seen in bright field in the TEM in an as-processed plate of Si-C-O fibre-reinforced MAS-5 (b) EDS spectrum from interlayer 1 and(c) EDS spectrum from interlayer 2 e 1999 The Royal Microscopical Society, Journtl of Microscop 196. 19+-202
morphologies in the surface regions of oxidized samples and complex multilayered carbon-containing morphologies in the centre of samples where the partial pressure of oxygen is low (Kumar & Knowles, 1996b). Such composites are still able to fail non-catastrophically even after 120 h in air in the temperature range 1000–1200 8C because of the carbon interlayers developed in the interior of the composites during these heat treatments (Kumar & Knowles, 1996d). The interfacial microstructure of Si–C–O fibre-reinforced barium magnesium aluminosilicate containing 5 wt.% potassium borosilicate glass developed by Corning Incorporated as a material which has an improved resistance to oxidation embrittlement over other fibre-reinforced glass ceramics (D. C. Larsen, personal communication, 1996) is similar to the Si–C–O/MAS composites examined by Kumar & Knowles. In the example shown in Fig. 3, taken from as-processed material, conventional bright field imaging of a fibre–matrix interface edge-on to the electron beam is able to distinguish a 90-nm thick interfacial zone containing two interlayers 1 and 2, between the fibre and the matrix (Knowles et al., 1997). Energy dispersive X-ray spectroscopy (EDX) of the interlayers shown in Figs 3(b) and (c) confirms that the lighter contrast interlayer is carbon-rich and that the darker layer is both silica-rich and barium-rich, consistent with the contrast from the two layers. After 100 h in air at 704 8C, oxidation of the carbon-rich interlayer produces debond cracks running parallel to the axis of the fibres, as a result of which there will be a loss of load transfer between the matrix and the fibres (Knowles et al., 1997). Thus, around this temperature, the advantage gained by reinforcing the relatively weak brittle glass– ceramic matrix with high strength fibres, thereby producing high strength, damage tolerant materials, is lost, and the composite strength reduces towards that of the matrix. However, as not all the fibre–matrix interfaces will be affected in the same way, because of local microstructural factors affecting the oxidation kinetics, some fibres may still contribute to the strength of the composites, even after extended isothermal exposure. It is only at temperatures above 1000 8C that the voids formed by oxidation of the turbostratic carbon at the fibre–matrix interfaces are closed by the formation of silica, enabling contact between the fibre and matrix to be resumed, and allowing most of the room temperature strength and damage tolerance to be retained after 100 h isothermal heat treatments (Knowles et al., 1997). The significance of microstructural observations of fibrereinforced ceramics has been to establish microstructural principles determining the toughness of fibre-reinforced ceramics in the absence of separate costly processing steps such as laying down carbon coatings on fibres prior to embedding them in a ceramic matrix. It is therefore somewhat ironic that these principles are now being used to specify advanced fibre coating methodologies in an attempt to overcome the limitations of carbon as an interfacial layer in high-temperature oxidizing environments (see, for example, Cinibulk & Hay, 1996). q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202 Fig. 3. (a) Example of a fibre–matrix interface seen in bright field in the TEM in an as-processed plate of Si–C–O fibre-reinforced MAS-5, (b) EDS spectrum from interlayer 1 and (c) EDS spectrum from interlayer 2. INTERFACES IN STRUC TURAL CERAMIC COMP OSITES 197
198 K. M. KNOWLES ET AL Crack-deflecting interlayers in laminar ceramic reaching to the end of the sample in three-point bend tests Thus, there is a distinct advantage to increasing the fracture energy of the crack-deflecting interlayers It has been shown As with fibre-reinforced ceramics, carbon interlayers have that such increases can be achieved by doping colloidal clearly been demonstrated to provide weak interfaces to graphite layers with chromium( Phillipps et aL., 1993) enable crack deflection in ceramic laminates(Clegg, 1992). although the reasons for this are not yet fully understood. Thus, these interlayers have the same advantages, but also In our tem work in this area sic laminates with carbon the same disadvantages, as those seen in fibre-reinforced rich interlayers fabricated by a plastic mixing and rolling ceramics. Tough interfaces can be produced at room method(Clegg, 1992) have been examined. Two different temperature, but above about 600C the interlayers suffer techniques of producing crack-deflecting interlayers were from oxidation, although recently other crack-deflecting used: (i) painting the Sic laminates with colloidal graphite interlayers have been suggested, such as porous interlayers with 20 vol. polyvinyl alcohol in solution and (ii) painting (Blanks et al., 1998: Clegg, 1998), glasses(Clegg, 1992) the Sic laminates with colloidal graphite doped with and weak shear structures such as hibonite(Filonenko& chromium nitrate, Cr(NO3)3 9H,O, with 20 vol. polyvinyl Lavrov, 1949) alcohol in solution. This latter procedure for producing As might be expected, the energy absorbed by layered interlayers increases the interfacial fracture energy of the ceramic structures is critically dependent on the fracture interlayers from 7.5 to 16 ]m(Phillipps et aL., 1993) energy of the interlayers(Phillipps et al., 1993). Interest- Laminates were stacked together, pyrolysed in a flowing ingly, provided crack deflection occurs, the fracture energy nitrogen atmosphere at 600C and then sintered in a of the interlayers does not appear to affect dramatically the graphite furnace under flowing argon at 2050C for total energy required to break samples of ceramic laminates 30 min, after which they were furnace cooled. Samples for (Phillipps et al., 1993). Samples with low fracture energy TEM could be prepared by mechanically polishing interlayers show substantial delamination, sometimes 500 um thick slabs cut from the laminates down to a Sic Sic Interlayer I um Fig. 4. General appearance of astandard carbon-rich interfacial layer between two SiC laminae in a damage-tolerant SiC/C laminar ceramic e 199y The Royal Microscopical Society, Jotarntl of Microscop 196, 19+-202
Crack-deflecting interlayers in laminar ceramic composites As with fibre-reinforced ceramics, carbon interlayers have clearly been demonstrated to provide weak interfaces to enable crack deflection in ceramic laminates (Clegg, 1992). Thus, these interlayers have the same advantages, but also the same disadvantages, as those seen in fibre-reinforced ceramics. Tough interfaces can be produced at room temperature, but above about 600 8C the interlayers suffer from oxidation, although recently other crack-deflecting interlayers have been suggested, such as porous interlayers (Blanks et al., 1998; Clegg, 1998), glasses (Clegg, 1992) and weak shear structures such as hibonite (Filonenko & Lavrov, 1949). As might be expected, the energy absorbed by layered ceramic structures is critically dependent on the fracture energy of the interlayers (Phillipps et al., 1993). Interestingly, provided crack deflection occurs, the fracture energy of the interlayers does not appear to affect dramatically the total energy required to break samples of ceramic laminates (Phillipps et al., 1993). Samples with low fracture energy interlayers show substantial delamination, sometimes reaching to the end of the sample in three-point bend tests. Thus, there is a distinct advantage to increasing the fracture energy of the crack-deflecting interlayers. It has been shown that such increases can be achieved by doping colloidal graphite layers with chromium (Phillipps et al., 1993), although the reasons for this are not yet fully understood. In our TEM work in this area, SiC laminates with carbonrich interlayers fabricated by a plastic mixing and rolling method (Clegg, 1992) have been examined. Two different techniques of producing crack-deflecting interlayers were used: (i) painting the SiC laminates with colloidal graphite with 20 vol.% polyvinyl alcohol in solution and (ii) painting the SiC laminates with colloidal graphite doped with chromium nitrate, Cr(NO3)3·9H2O, with 20 vol.% polyvinyl alcohol in solution. This latter procedure for producing interlayers increases the interfacial fracture energy of the interlayers from 7·5 to 16 J m¹2 (Phillipps et al., 1993). Laminates were stacked together, pyrolysed in a flowing nitrogen atmosphere at 600 8C and then sintered in a graphite furnace under flowing argon at 2050 8C for 30 min, after which they were furnace cooled. Samples for TEM could be prepared by mechanically polishing , 500 mm thick slabs cut from the laminates down to a Fig. 4. General appearance of a ‘standard’ carbon-rich interfacial layer between two SiC laminae in a damage-tolerant SiC/C laminar ceramic composite. 198 K. M. KNOWLES ET AL. q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202
INTERFACES IN STRUCTURAL CERAMIC COMPOSITES 19 10 nm Fig. 5. Detail of the carbon interlayer in a SiC/C laminar ceramic composite showing small graphite flakes embedded within an amorphous Crack Crack 0.1um Fig. 6. Kink band formation at a crack in a graphite flake in the carbon layer of a SiC/C laminar ceramic composite e 1999 The Royal Microscopical Society, Journtl of Microscop 196. 19+-202
q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202 Fig. 5. Detail of the carbon interlayer in a SiC/C laminar ceramic composite showing small graphite flakes embedded within an amorphous carbonaceous matrix. Fig. 6. Kink band formation at a crack in a graphite flake in the carbon layer of a SiC/C laminar ceramic composite. INTERFACES IN STRUC TURAL CERAMIC COMP OSITES 199
200 K M. KNOWLES ET AL hickness of 100 um, mounting these on 3 mm diameter they have a high degree of preferred orientation. in contrast copper grids with epoxy resin and ion beam thinning to o the flakes in the centre of the interlayers. transparency at 5 kV using argon ions Examination of the sic lamellae close to the carbon The general appearance of the microstructure of the interlayers shows evidence for turbostratic carbon in pores standard interlayers without chromium nitrate doping is at triple junctions, such as the example shown in Fig. 7. in shown at low magnification in Fig 4. The layers can be seen which wrinkled(0002)layers characteristic of turbostratic to be quite distinct from the adjacent SiC laminae, and there carbon are evident. The source of this carbon is the colloida is no obvious interdiffusion between the SiC laminae and graphite- the silicon carbide sheets formed by the plastic the carbon interlayer. The differential thinning of the mixing and rolling method are porous, enabling th interlayer and the Sic laminae and the higher electron graphite paint to penetrate the sheets prior to sintering. scattering power of the Sic laminae both contribute to the The ability of the carbon-rich interlayer to deflect cracks relatively dark appearance of the SiC laminae in Fig. 4 in in the Sic laminae perpendicular to the interlayers can comparison with the carbon interlayer. The carbon inter- therefore be rationalized straightforwardly from these layer is composed of a number of relatively small graphite transmission electron microscope observations, on the basis flakes embedded in an amorphous carbonaceous matrix of a suitable texture arising in the graphite-rich interlayer (Fig. 5). These flakes have a tendency to be contorted. during processing which will deflect cracks parallel to the particularly within the centre of the interlayer, and can basal planes, between which the van der Waals bonding exhibit clear kink bands, such as the one shown bridging a provides only modest resistance to shear stresses crack in Fig. 6. The texture of the graphite in the interlayer corporation of chromium nitrate in the colloidal varies as a function of perpendicular distance away from th graphite paint solution causes the formation within the SiC lamellae. Not surprisingly, the graphite aggregates close interlayer of a relatively uniform distribution of metallic to the lamellae appear to use the lamellae as flat templates chromium-rich nanoparticles(Fig. 8). These particles also against which to align themselves with their(0001)basal form as a thin -400 nm layer on the surfaces of the Sic planes parallel to the lamellae. As a consequence of this laminae. In comparison with the chromium-free carbon sicl Sic Turbostratic carbon 10 nm Fig. 7. Turbostratic carbon at triple junctions within a SiC lamina arising from the colloidal graphite used for painting the carbon interlayers e 199y The Royal Microscopical Society, Jotarntl of Microscop 196, 19+-202
thickness of , 100 mm, mounting these on 3 mm diameter copper grids with epoxy resin and ion beam thinning to transparency at 5 kV using argon ions. The general appearance of the microstructure of the standard interlayers without chromium nitrate doping is shown at low magnification in Fig. 4. The layers can be seen to be quite distinct from the adjacent SiC laminae, and there is no obvious interdiffusion between the SiC laminae and the carbon interlayer. The differential thinning of the interlayer and the SiC laminae and the higher electron scattering power of the SiC laminae both contribute to the relatively dark appearance of the SiC laminae in Fig. 4 in comparison with the carbon interlayer. The carbon interlayer is composed of a number of relatively small graphite flakes embedded in an amorphous carbonaceous matrix (Fig. 5). These flakes have a tendency to be contorted, particularly within the centre of the interlayer, and can exhibit clear kink bands, such as the one shown bridging a crack in Fig. 6. The texture of the graphite in the interlayer varies as a function of perpendicular distance away from the SiC lamellae. Not surprisingly, the graphite aggregates close to the lamellae appear to use the lamellae as flat ‘templates’ against which to align themselves with their (0001) basal planes parallel to the lamellae. As a consequence of this they have a high degree of preferred orientation, in contrast to the flakes in the centre of the interlayers. Examination of the SiC lamellae close to the carbon interlayers shows evidence for turbostratic carbon in pores at triple junctions, such as the example shown in Fig. 7, in which wrinkled (0002) layers characteristic of turbostratic carbon are evident. The source of this carbon is the colloidal graphite – the silicon carbide sheets formed by the plastic mixing and rolling method are porous, enabling the graphite paint to penetrate the sheets prior to sintering. The ability of the carbon-rich interlayer to deflect cracks in the SiC laminae perpendicular to the interlayers can therefore be rationalized straightforwardly from these transmission electron microscope observations, on the basis of a suitable texture arising in the graphite-rich interlayer during processing which will deflect cracks parallel to the basal planes, between which the van der Waals bonding provides only modest resistance to shear stresses. Incorporation of chromium nitrate in the colloidal graphite paint solution causes the formation within the interlayer of a relatively uniform distribution of metallic chromium-rich nanoparticles (Fig. 8). These particles also form as a thin , 400 nm layer on the surfaces of the SiC laminae. In comparison with the chromium-free carbon Fig. 7. Turbostratic carbon at triple junctions within a SiC lamina arising from the colloidal graphite used for painting the carbon interlayers onto the laminae. 200 K. M. KNOWLES ET AL. q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202
INTERFACES IN STRUCTURAL CERAMIC COMPOSITES 201 Cr-rich particles 0.2um Fig. 8. General appearance of a chromium-doped carbon-rich interfacial layer between two SiC laminae in a damage-tolerant SiC/C laminar ceramic composite showing the formation of chromium-rich nanoparticles. interlayers, the carbon in the chromium-doped interlayers is ceramics, TEM is able to show unambiguously the presence less crystalline, forming amorphous carbon and highly or absence of thin amorphous films at interfaces after corrugated local regions of turbostratic carbon at the processing at high temperature and the presence or absence expense of more fully developed graphitic carbon of favoured orientation relationships. Such data are On the basis of these observations, the greater interfacial important for particulate-reinforced ceramics in the context fracture energy from the chromium-doped carbon inter- of microstructural development during high temperature layers in comparison to chromium-free interlayers can be consolidation processes and the effect on mechanical rationalized in terms of the chromium providing a greater properties such as creep resistance and fracture toughness degree of bonding within the interlayer, and also possibly causing a more tortuous crack path through the presence of turbostratic. rather than graphitic, carbon. For low Acknowledgements temperature damage-tolerant lamellar ceramic composites, We would like to thank the Turkish Government. the United there is clearly scope for further interfacial engineering States Air Force European Office of Aerospace Research and using this experimental methodology of modifying the Development, London, U.K.,the Materials Directorate chemical constitution of the carbon-rich crack-deflecting Wright Patterson Air Force Base, Dayton,Ohio, U.S.Aand Interlayers. the european Community Brite/ Euram programme (project no BE 7414)for financial support for various aspects of this Concluding remarks The examples presented and discussed here of interfaces in ceramic composites examined by TEM provide valuabl References microstructural information For fibre-reinforced ceramics Beier W.& Markman, S(1997) Fiber-reinforced glass. Adv. Mater and ceramic laminates this can be related directly to the Process.152(6),37-40 ability of interfaces to deflect cracks. In particulate-reinforced Blanks, K.S., Kristoffersson. A, Carlstrom, E& Clegg. W.(1998) e 1999 The Royal Microscopical Society, Journtl of Microscop 196. 19+-202
interlayers, the carbon in the chromium-doped interlayers is less crystalline, forming amorphous carbon and highly corrugated local regions of turbostratic carbon at the expense of more fully developed graphitic carbon. On the basis of these observations, the greater interfacial fracture energy from the chromium-doped carbon interlayers in comparison to chromium-free interlayers can be rationalized in terms of the chromium providing a greater degree of bonding within the interlayer, and also possibly causing a more tortuous crack path through the presence of turbostratic, rather than graphitic, carbon. For low temperature damage-tolerant lamellar ceramic composites, there is clearly scope for further interfacial engineering using this experimental methodology of modifying the chemical constitution of the carbon-rich crack-deflecting interlayers. Concluding remarks The examples presented and discussed here of interfaces in ceramic composites examined by TEM provide valuable microstructural information. For fibre-reinforced ceramics and ceramic laminates this can be related directly to the ability of interfaces to deflect cracks. In particulate-reinforced ceramics, TEM is able to show unambiguously the presence or absence of thin amorphous films at interfaces after processing at high temperature and the presence or absence of favoured orientation relationships. Such data are important for particulate-reinforced ceramics in the context of microstructural development during high temperature consolidation processes and the effect on mechanical properties such as creep resistance and fracture toughness. Acknowledgements We would like to thank the Turkish Government, the United States Air Force European Office of Aerospace Research and Development, London, U.K., the Materials Directorate, Wright Patterson Air Force Base, Dayton, Ohio, U.S.A. and the European Community Brite/Euram programme (project no. BE 7414) for financial support for various aspects of this work. References Beier, W. & Markman, S. (1997) Fiber-reinforced glass. Adv. Mater. Process. 152 (6), 37–40. Blanks, K.S., Kristoffersson, A., Carlstro¨m, E. & Clegg, W.J. (1998) q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202 Fig. 8. General appearance of a chromium-doped carbon-rich interfacial layer between two SiC laminae in a damage-tolerant SiC/C laminar ceramic composite showing the formation of chromium-rich nanoparticles. INTERFACES IN STRUC TURAL CERAMIC COMP OSITES 201
202 K. M. KNOWLES ET AL Crack deflection in ceramic laminates using porous interlayers. intermediate temperature embrittlement in a Si-C-O fibre- Eur Ceram Soc. 18.1945-1951 reinforced barium magnesium aluminosilicate. Microscopy of Brennan, J.J. (1986) Interfacial characterization of glass and glas Oxidation 3(ed. by S B Newcomb and J. A. Little). pp. 695-707 ceramic matrix/Nicalon SiC fiber composites. Tailoring Multiphase Book 675. Institute of Materials. London, U.K and Composite Ceramics (Proc. 21st University Conference on Kumar, A.& Knowles. K M.(1996a)Microstructure-property Ceramic Science, July 17-19, 1985. Pennsylvania State University relationships of SiC fibre-reinforced magnesium aluminosilicates niversityy Park, PA), Materials Science Research, Vol. 20(ed. by R. I. Microstructural characterisation. Acta Materialia. 44. 2901 E. Tressler, G. L. Messing. C G. Pantano and R E. Newnham) 2921 pp 549-560. Plenum Press, New York, U.S.A. Kumar, A& Knowles, K M. (1996b)Oxidation behavior of a Si-C- Brennan, J].(1987) Interfacial chemistry and bonding in fiber- O-fiber-reinforced magnesium aluminosilicate. Am. Ceram Soc. reinforced glass and glass-ceramic matrix composites. Ceramic 79.2364-2374 Microstructures86 Materials Science Research, Vol 21(ed by J. A. Kumar, A.& Knowles, K.M.(1996c) Microstructure-property Pask and A G. Evans), pp. 387-399 Plenum Press, New York relationships of SiC fibre-reinforced magnesium aluminosilicates U.S.A IL. Mechanical properties and failure characteristics. Acta Campbell, (1997)Opportunities for ceramics industr Mater.44.2923-2934. 96.237-246 Kumar, A.& Knowles, K M.(1996d) Effect of oxidation heat Cinibulk. M K& Hay. R.S. (1996) Textured magnetoplumbite fiber- treatments on the mechanical behavior of a Si-C-O-fiber- matrix interphase derived from sol-gel fiber coatings. J. Am reinforced magnesium aluminosilicate. Am. Ceram. Soc. 79 Cerar.Soc.79,1233-1246. 2375-2378. Clarke, D R (1987)On the equilibrium thickness of intergranular Le Strat. E, Lancin, M. Fourches-Coulon N.& Marhic, C(1998) glass phases in ceramic materials. Am. Ceram. Soc. 70. 15 SiC Nicalon fibre-glass matrix composites: interphases and their nechanism of formation. Philos. Mag. A. 78. 189-202 Clarke, D... Shaw, T M, Philipse. A P. Horn, R.G.(1993) Pan, X. Mayer. ], Ruhle. M.& Niihara, K.(1996) Silicon nitride Possible electrical double-layer contribution to the equilibrium based ceramic nanocomposites. J. Am. Ceram Soc. 79, 585-590 thickness of intergranular glass films in polycrystalline ceramics. Phillipps, A J. Clegg. W.J.& Clyne. Tw(1993) Fracture behaviour of . Am. Ceram Soc. 76, 201-204 ceramic laminates in bending-IL. Comparison of model predictions Clegg. w. (1992) The fabrication and failure of laminar with experimental data. Acta Metall. Mater. 41. 819-827 opposites. Acta Metall. Mater. 40. 3085-3093. Sheldon, B W. Sun, EY. Nutt. S.R.& Brennan. J.(1996) Clegg, w.].(1998)Hard and fast ceramic composite Oxidation of BN-coated SiC fibers in ceramic-matrix composites. Worl.6.215 J. Am. Ceram. Soc. 79, 539-543 Cooper. R. . Chyung K(1987)Structure and chemistry of fibre- Sternitzke, M.(1997) Structural ceramic nanocomposites. J. Eur matrix interfaces in silicon carbide fibre-reinforced glass-ceramic Cerar.Soc.17,1061-1082. omposites: an electron microscopy study. J. Mater. Sci. 22, Sun. E.Y. Nutt, S.R.& Brennan, J.(1996) High-temperature 3148-3160. tensile behavior of a boron nitride- coated silicon carbide-fiber Davidge, R W. Brook, R J, Cambier, E. Poorteman, M. Leriche, A. glass-ceramic composite. J. Am. Ceram. Soc. 79. 1521-1529 OSullivan, D. Hampshire. S& Kennedy. T(1997)Fabrication. Turan, S.&Knowles, K M. (1995a)A comparison of the properties, and modelling of engineering ceramics reinforced microstructure of silicon nitride-silicon carbide composites with nanoparticles of silicon carbide. Br. Ceran. Trans. 96. 121 made with and without deoxidized starting material. J. Microsc. 77.287-304 Evans. A.G.& Marshall, D B(1989)The mechanical behaviour of Turan, S.& Knowles, KM.(1995b) Formation of boron nitride ceramic matrix composites. Acta Metall. 37. 2567-2583 inclusions in hot isostatically pressed silicon nitride-silicon Filonenko, N E.& Lavrov, I V.(1949)Calcium hexaluminate in carbide composites. Am. Ceram Soc. 78, 680-684 the system CaO-Al2O3-SiO2. Dokl. Akad. Nauk SSSR, 66. 673- Turan, S& Knowles, K M.(1996) Effect of boron nitride on the 676 phase stability and phase transformations in silicon carbide. J. wa. T, Kajii, S, Matsunaga, K, Ogami, T, Kohtoku, Y.& Am. Ceran.Soc.79,3325-3328 gasawa, T.(1998)A tough, thermally conductive silicon Turan, S.&Knowles. KM.(1997a) Orientation-dependent carbide composite with high strength up to 1600.C in air. equilibrium film thickness at interphase boundaries in ceramic- Science.282,1295-1297. ceramic composites. Inst. Phys. Conf. Ser. 153, 483-486 Kleebe, H- ] Hoffmann. M J.& Ruhle, M. (1992) Influence of Turan, S& Knowles, K M (1997b) Interphase boundaries between secondary phase chemistry on grain boundary film thickness in hexagonal boron nitride and beta silicon nitride in silicon nitride silicon nitride. Z. Metallk. 83. 610-617 ilicon carbide particulate composites. J. Eur. Ceram. Soc. 17 Knowles, K.M. Kumar. A& Fox, AG.(1997) Investigation of 1849-1854. e 199y The Royal Microscopical Society, Jotarntl of Microscop 196, 19+-202
Crack deflection in ceramic laminates using porous interlayers. J. Eur. Ceram. Soc, 18, 1945–1951. Brennan, J.J. (1986) Interfacial characterization of glass and glassceramic matrix/Nicalon SiC fiber composites. Tailoring Multiphase and Composite Ceramics (Proc. 21st University Conference on Ceramic Science, July 17–19, 1985, Pennsylvania State University, University Park, PA), Materials Science Research, Vol. 20 (ed. by R. E. Tressler, G. L. Messing, C. G. Pantano and R. E. Newnham), pp. 549–560. Plenum Press, New York, U.S.A. Brennan, J.J. (1987) Interfacial chemistry and bonding in fiberreinforced glass and glass-ceramic matrix composites. Ceramic Microstructures 0 86 Materials Science Research, Vol. 21 (ed. by J. A. Pask and A. G. Evans), pp. 387–399. Plenum Press, New York, U.S.A. Campbell, J. (1997) Opportunities for ceramics industry. Br. Ceram. Trans. 96, 237–246. Cinibulk, M.K. & Hay, R.S. 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(1997) Fabrication, properties, and modelling of engineering ceramics reinforced with nanoparticles of silicon carbide. Br. Ceram. Trans. 96, 121– 127. Evans, A.G. & Marshall, D.B. (1989) The mechanical behaviour of ceramic matrix composites. Acta Metall. 37, 2567–2583. Filonenko, N.E. & Lavrov, I.V. (1949) Calcium hexaluminate in the system CaO–Al2O3–SiO2. Dokl. Akad. Nauk. SSSR, 66, 673– 676. Ishikawa, T., Kajii, S., Matsunaga, K., Hogami, T., Kohtoku, Y. & Nagasawa, T. (1998) A tough, thermally conductive silicon carbide composite with high strength up to 16008C in air. Science, 282, 1295–1297. Kleebe, H.-J., Hoffmann, M.J. & Ru¨hle, M. (1992) Influence of secondary phase chemistry on grain boundary film thickness in silicon nitride. Z. Metallk. 83, 610–617. Knowles, K.M., Kumar, A. & Fox, A.G. (1997) Investigation of ‘intermediate’ temperature embrittlement in a Si–C–O fibrereinforced barium magnesium aluminosilicate. Microscopy of Oxidation 3 (ed. by S. B. Newcomb and J. A. Little), pp. 695–707. Book 675, Institute of Materials, London, U.K. Kumar, A. & Knowles, K.M. (1996a) Microstructure-property relationships of SiC fibre-reinforced magnesium aluminosilicates – I. Microstructural characterisation. Acta Materialia, 44, 2901– 2921. Kumar, A. & Knowles, K.M. (1996b) Oxidation behavior of a Si–C– O-fiber-reinforced magnesium aluminosilicate. J. Am. Ceram. Soc. 79, 2364–2374. Kumar, A. & Knowles, K.M. (1996c) Microstructure-property relationships of SiC fibre-reinforced magnesium aluminosilicates – II. Mechanical properties and failure characteristics. Acta Mater. 44, 2923–2934. Kumar, A. & Knowles, K.M. (1996d) Effect of oxidation heat treatments on the mechanical behavior of a Si–C–O-fiberreinforced magnesium aluminosilicate. J. Am. Ceram. Soc. 79, 2375–2378. Le Strat, E., Lancin, M., Fourches-Coulon, N. & Marhic, C. (1998) SiC Nicalon fibre-glass matrix composites: interphases and their mechanism of formation. Philos. Mag. A, 78, 189–202. Pan, X., Mayer, J., Ru¨hle, M. & Niihara, K. (1996) Silicon nitride based ceramic nanocomposites. J. Am. Ceram. Soc. 79, 585–590. Phillipps, A.J., Clegg, W.J. & Clyne, T.W. (1993) Fracture behaviour of ceramic laminates in bending – II. Comparison of model predictions with experimental data. Acta Metall. Mater. 41, 819–827. Sheldon, B.W., Sun, E.Y., Nutt, S.R. & Brennan, J.J. (1996) Oxidation of BN-coated SiC fibers in ceramic-matrix composites. J. Am. Ceram. Soc. 79, 539–543. Sternitzke, M. (1997) Structural ceramic nanocomposites. J. Eur. Ceram. Soc. 17, 1061–1082. Sun, E.Y., Nutt, S.R. & Brennan, J.J. (1996) High-temperature tensile behavior of a boron nitride-coated silicon carbide-fiber glass-ceramic composite. J. Am. Ceram. Soc. 79, 1521–1529. Turan, S. & Knowles, K.M. (1995a) A comparison of the microstructure of silicon nitride-silicon carbide composites made with and without deoxidized starting material. J. Microsc. 177, 287–304. Turan, S. & Knowles, K.M. (1995b) Formation of boron nitride inclusions in hot isostatically pressed silicon nitride-silicon carbide composites. J. Am. Ceram. Soc. 78, 680–684. Turan, S. & Knowles, K.M. (1996) Effect of boron nitride on the phase stability and phase transformations in silicon carbide. J. Am. Ceram. Soc. 79, 3325–3328. Turan, S. & Knowles, K.M. (1997a) Orientation-dependent equilibrium film thickness at interphase boundaries in ceramicceramic composites. Inst. Phys. Conf. Ser. 153, 483–486. Turan, S. & Knowles, K.M. (1997b) Interphase boundaries between hexagonal boron nitride and beta silicon nitride in silicon nitride – silicon carbide particulate composites. J. Eur. Ceram. Soc. 17, 1849–1854. 202 K. M. KNOWLES ET AL. q 1999 The Royal Microscopical Society, Journal of Microscopy, 196, 194–202