Microstructural characterization of interafaces in fibre-reiforced ceramics

Composites Science and Technology 42 (1991) 221-249

Microstructural Characterisation of Interfaces in FibreReinforced Ceramics

M. H. Lewis & V. S. R. Murthy Centre for AdvancedMaterials Technology,Universityof Warwick, Coventry CV4 7AL, UK (Received 7 December1990; accepted 8 February 1991)

ABSTRACT A survey is presented of the principles underlying the design of model interface microstructures which have the required micromechanical response and therntal or environmental stability within fibre-reinforced ceramic matrix composites. Microstructural and compositional characterisation of interfaces, based on electron imaging and spectroscopic techniques, are reviewed with illustrative examples of interfaces formed by in-situ reaction during fabrication or by fibre precoating.

1 INTRODUCTION: THE DEVELOPMENT OF CERAMIC MATRIX COMPOSITES During the past five years there has been a rapid escalation in research activity on structural ceramics which are artificial combinations of two or more phases. These ceramic matrix composites (CMCs) are being developed in response to the failure of monolithic ceramics, in which microstructures evolve during natural processing reactions, to meet the reliability criteria for 'high-risk' engineering applications. Although there was an initial wave of enthusiasm for CMCs in the early 1970s involving theoretical modelling and fabrication studies, the absence of fibres appropriate for high-temperature service and fabrication in reactive ceramic matrices inhibited further 221 Composites Science and Technology 0266-3538/91/$03.50 © 1991 ElsevierSciencePublishers Ltd, England. Printed in Great Britain

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M. H. Lewis, V. S. R. Murthy

development. There was also an inhibiting influence in the promise of more easily fabricated monolithic ceramics (especially those based on Si3N~) in achieving moderate fracture toughness (Kc). During the early 1980s research and development on monolithic sinterable ceramics gained momentum with the introduction of the varied range of zirconia-toughened ceramics. However, the Si-based ceramics (Si3N4 and sialons) have reached refinement limits of < 10 M P a , j m , for K~ and ZrO2 toughening is confined mainly to oxides and has a severe temperature limitation. The wealth of experience gained in understanding toughening and deformation mechanisms in monolithic microstructures has subsequently been applied to the development of dispersion or whiskertoughened 'composites'. Although there are some commercial developments, such as SiC whiskers/A1203 matrices, the toughening increments are generally small and unlikely to reach theoretically projected limits above 15 MPax/-m. For many complex engineering applications the concept of a tolerance of service overstressing during load transfer to high-strength fibres in a microcracked matrix is essential. The experimental demonstration of this concept together with the development of new ceramic fibres, such as Nicalon SiC, has motivated the current level of interest in theoretical modelling and the evolution of varied matrices and fabrication methods, A well-established key to fibrous CMC performance is the chemistry and structure of the fibre/matrix interface. Although there have been earlier developments in composite fabrication technology and interface design for polymer and metal matrix composites, the micromechanical requirements (interface cohesion, etc.) generally differ from those of CMCs and the fabrication and temperature constraints are more severe. In this paper a survey of idealised interface behaviour is followed by a review of interface constitution and microstructure in real systems which illustrate currently available characterisation techniques.

2 PRINCIPLES OF I N T E R F A C E DESIGN Parameters which define the micromechanical behaviour of fibre/matrix interfaces in CMCs are Gt/G f, the critical strain energy release rate (fracture energy) for the interface relative to that for the fibre, and 3, the shear resistance of the interface after debonding. A prerequisite for interracial debonding, in preference to matrix microcrack propagation through fibres, is that Gj < Gff4.1'2 With this condition satisfied (Fig. l(a)) it is probable that, in unidirectionally stressed CMCs (parallel to the fibre axis) with small T, a steady state matrix

Microstructural characterisation of interfaces

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microcracking will occur at a stress o"m given by 2

60mr L" where O m is the fracture energy of the matrix, r is the fibre radius, and E m, E t and E are the Young's moduli of the matrix, the fibre and a composite of the fibre volume fraction, Ft. The ultimate tensile fracture stress, au, and increment in composite fracture energy, AG c, are given by 2'3

au = VtSf[z, r,m]

AGe = $ 3 + Vt(1- Vr)2E2m/6rEtE 2

where m is the fibre Weibull modulus and S is the fibre strength. Although similar expressions were obtained in the 1970s, notably by Aveston, Cooper and Kelly (ACK)'* based on energy balance arguments, the more rigorous fracture mechanics derivations also include the influence of fibre failure statistics via the function f(m), which contains the Weibull modulus for fibre samples. The influence of m on a~, etc., relates to the debond length, d, over which fibre/matrix shear occurs and the increased probability of fibre failure (below the mean stress) with increased debond length or fibre volume. An interesting effect is the inversion of dependence on ~/r for very low m, thus making a m and Gc synergistic. An important earlier conclusion (from the A C K theory) was the conflicting requirement on for tr= and Go, although for anticipated values of m (>5) this conflict remains.

224

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This latter point raises the interesting question of design criteria in CMCs. There is a tendency to overemphasise the requirement for high work of fracture, whereas in many applications a m must be a design limit and this is reinforced by environmental factors discussed below. With increasing r, a mis raised so as eventually to meet a (slowly falling) UTS, with the onset of totally 'brittle' fracture, characteristic of monolithic ceramics. Hence there is a need, during processing, to control r within design criteria for G¢ and a m with an additional constraint of fibre properties (S, m) and size r. It is important to realise that the relatively refined modelling of fracture behaviour applies only to unidirectional fibre microstructures stressed axially. For multiaxial stressing, almost invariably encountered in real engineering components, 2D and 3D cross-plied and woven fibre architectures are essential. Thence an interaction between the different failure modes (transverse fracture, delamination and normal microcracking) will modify the relative weighting of the parameters discussed. The simplistic model of a single 'interface" with appropriate debond and shear properties has to be modified to cope with the problems of reactivity during fabrication between fibre, matrix and/or debond layer, and the 'inservice' reaction between the atmosphere and fibre or debond layer. To cope with such problems multilayer interfaces may be designed of the type shown in Fig. l(b). An inner debond/shear layer may be a ceramic 'layer' structure with largely van der Waals bonding, such as graphite (C) or hexagonal boron nitride (BN), a metallic layer which has low matrix cohesion or exhibits plastic flow at low stress for the band of service temperature. Such debond layers are susceptible to oxidation, the extreme problem being with carbon which oxidises to gaseous products above - 600°C. Although debond layer removal may not destroy the matrix microcracking response of the composite, the formation of solid "bridging' layers which have higher cohesion may produce a reversion to monolithic, brittle ceramic behaviour. The bridging layers may be metallic oxides (which have a larger atomic volume than the metallic interface layer) or silica oxidation layers on the underlying SiC-based fibre, as described below. Although oxidation resistance may be conferred naturally by low oxygen diffusion rates through a fully dense matrix, such as stoichiometric crystalline SiC, porous or high diffusivity ceramic matrices (such as silicate glass ceramics) may dictate the presence of a diffusion barrier as a second fibre coating (DB in Fig. l(b)). A third fibre coating may be required if the oxygen barrier is incompatible with matrix chemistry either in service or during the shorter, but normally higher temperature, fabrication process. A further complication of fibre/' debond layer reaction may similarly necessitate an inner reaction barrier (RBI and RBII in Fig. l(b)). Such complex multilayer interfaces are likely to

Microstructural characterisation of interfaces

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be prohibitively expensive, apart from ultra-high-temperature aerospace applications, with a probable limit of bi-layer (debond + reaction/diffusion barrier) systems. For small fibres such as Nicalon and Tyranno with thin interface layers, differential thermal expansion may not be a critical issue in requiring the additional complication of radial variations in chemistry. Although the principles of interface design remain at a simplistic level, an experimental modelling of these principles is constrained by an absence of data on cohesion and plasticity of these layers with varied chemistry and the problems of multilayer synthesis. Vapour deposition (PVD or CVD) of interface layers onto fibre tows or woven preforms offers the greatest flexibility in preparation. An alternative of in-situ reaction during processing is normally restricted to specific fibre/matrix combinations, described below, resulting in single interface layers with limited temperature capability.

3 INTERFACE CHARACTERISATION TECHNIQUES There are two main approaches to a definition of the constitution and structure of interfaces which apply to all forms of composite material. Firstly, the interface may be studied in-situ by means of microscopy and microanalysis of sectioned surfaces in a scanning electron microscope (SEM) or of thin sections in a transmission electron microscope (TEM). The latter is required for chemical and structural information at a resolution consistent with normal interface layer thicknesses. A second approach is to use surface-sensitive spectroscopies on fibres that are chemically or mechanically extracted from the matrix, most conveniently following fracture and 'pull-out'. A comparison of typical excitation or fluorescence volumes for the various analytical signals illustrates the restriction in spatial resolution for each technique (Fig. 2). The in-situ thin section TEM-based techniques are superior but impose more difficult specimen preparation problems associated with varied matrix, interface and fibre chemistries in CMCs. In the imaging mode SEM has a potential resolution of < 10 nm but in practice is rarely better than 100 nm owing to the back-scattered electron excitation volume or the signal/noise ratio for atomic number contrast within light element-containing interfaces. Hence imaging of fine fibre interfaces, with typical thicknesses between 10 and 50 nm exemplified by insitu carbon-rich layers between SiC fibres and silicate matrices, requires TEM. SEM is useful for thick interface layers formed by fibre precoating on the micron scale and at this level may be supplemented by low keV X-ray excitation to provide compositional data from spectra (EDS) or, at a lower

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level, by X-ray mapping. Light element (windowless) X-ray detectors are essential for this work and are now available on TEMs as well as SEMs. Examples of SEM imaging and X-ray mapping are given in Section 5. T E M imaging provides a resolution of < 1 nm using absorption (atomic number) or phase contrast, exemplified in Section 4.1 for structure within the thickness of in-situ reaction interfaces in SiC/silicate matrix composites. This may be supplemented by lattice or structure images for crystalline interfacial/matrix phases, with <1 nm resolution, or by Fresnel fringe images taken in through-focal series in determining the width of interfacial

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layers on this scale. The latter technique, exemplified for a SiC whisker/Al203 matrix composite in Ref. 5. is also capable of providing constitutional data for the interface via quantitative Fresnel fringe profiling. However, these ultra-high resolution techniques have stringent electron optical and geometric requirements, with interfaces parallel to the electron beam within very thin sections. Analytical TEM also offers better spatial resolution as a consequence of a much smaller and more precisely defined excitation volume in thin sections. Electron energy loss spectroscopy (EELS) has the ultimate resolution dictated only by initial electron excitation volume, with a lateral extent (Fig. 2) not much greater than electron probe diameter, whereas EDS is subject to secondary X-ray fluorescence. However. the sensitivity (signal/background in the spectrum) is normally superior for EDS, even for light elements, and it is a more convenient technique, less sensitive to specimen/spectrometer geometry and with parallel element processing (although parallel EELS detectors have recently evolvedJ. Various surface sensitive spectroscopies have been used on extracted fibres or matrix 'pull-out' channels, for example Auger electron spectroscopy (AES), secondary ion mass spectroscopy (SIMS) and X-ray photoelectron spectroscopy (XPS). This approach to compositional analysis of interface sections requires sequential removal, by ion sputtering, of surface layers followed by spectral analysis. The spatial resolution of these spectroscopies normal to interface layer is potentially a few atom layers, owing to the limited emission/escape depths for the tow energy electrons and ions. However, non-uniform ion sputtering profiles may result in averaged compositions over much larger radial dimensions in the fibre. This is less probable for AES, in which small electron probes may sample the centre of the ion sputtered crater (Fig. 2(b)), whereas ion probes are normally of similar size. In XPS the irradiated area is normally comparable with the fibre diameter such that compositional profiling of interfaces is not possible. Examples of surface-sensitive spectroscopies are given in Fig. 3, illustrating composition profiling with AES and its good light element sensitivity,6,7- together with superposed XPS from SiC whiskers with different surface composition, s illustrating the capability of providing structural information via 'chemical shifts' for similar atoms in different coordination (in this case silicon in carbon or oxygen environment).

4 FIBRE CONSTITUTION AND STRUCTURE Interface behaviour in a stressed composite is a function of intrinsic mechanical properties of fibre and matrix relative to the interface structure. The latter is frequently related to fibre/matrix reaction during processing or

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to fibre/vapour reaction during fibre precoating. Hence a brief review of fibre chemistry and structure is a necessary prelude to an understanding of interface microstructure. Fine (10-20/~m) fibres produced by melt-spinning of organic precursors, e.g. Nicalon (Nippon Carbon Co., Japan) or Tyranno (Ube Industries, Japan), are most convenient for CMC fabrication but lack high-temperature stability (above --,1000°C). Nicalon and Tyranno both have a silicon oxycarbide constitution, being synthesised from polycarbosilane and polytitanocarbosilane, respectively. The fibres have similar constitutions, with Si/C/O ratios of 1/1-46/0-36 for Nicalon and 1/1.32/0.96 for Tyranno. Some of the excess carbon (above the SiC stoichiometry) is dispersed in the

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free state (as graphitic microcrystals) and the remainder is contained within a non-crystalline network, together with some of the oxygen, in mixed coordination tetrahedra (SiOxC,~_~, 0 < x < 4). This co-ordination chemistry and the higher oxygen content for Tyranno is demonstrated by means of MAS NMR spectroscopy 9 in Fig. 4. High-temperature structural degradation of Nicalon or Tyranno fibres occurs by a progressive segregation of carbon and oxygen to their respective co-ordination tetrahedra (SIC4, SiO4) with an accompanying increase in long range order, evidenced by spectral line narrowing (Fig. 4). The ultimate structure is a microcrystalline mixture of fl-SiC and SiO2 (cristobalite). In vacuum or inert atmospheres oxygen is lost as SiO and CO, leaving a microporous fl-SiC crystalline fibre of low strength. Tyranno fibres also contain 2-4% titanium, which reacts with free carbon to precipitate TiC. Recently developed fine fibres which have improved high-temperature stability are based on non-crystalline carbonitrides. Fiberamic [RhonePoulenc, France) 1° and HPZ (Dow Coming, USA) are claimed to be stable to nearly 1400°C, maintaining their disordered Si(N, C),~-based tetrahedral network structure. An example of MAS NMR spectra together with TEM image and EDS analysis is given for the HPZ fibre (Fig. 5). The broad NMR spectral peak is based on SiNa co-ordination with an asymmetric peak broadening towards the SiC~ polytypic structures, indicative of mixed (C, N) co-ordination.tl'~2 The fibre also contains oxygen (Si/N/C/O = 57/28/10/4), detected in the EDS analysis, and the TEM image exhibits a phase contrast fine structure which may originate from excess carbon in precipitated graphitic t'orm. ~ Stability and interface reactions, within silicate matrices, have been the subject of a brief comparative study, described in Section 5. The limited temperature capability of the polymer precursor silicon oxycarbide fibres has encouraged the use of SiC monofilaments, synthesised by means of CVD, which have been developed largely for metal matrix composites. SCS6 monofilaments (Textron, USA), 140#m diameter, are produced by thermal decomposition of silane/hydrogen gas mixtures onto a 30-~tm carbon core (Fig. 6). The CVD SiC is a crystalline polytypic (largely 3C) form with ultra-fine (50-100 nm) radially oriented grains. Variations in deposition stoichiometry (from SiC) are intentionally introduced and accommodated by a fine silicon dispersion which results in a marked strength reduction and impaired creep resistance above 1200~C. ~3 Monofilaments are normally used in a precoated form; the SCS6 source contains a thick (4/~m) carbon-rich bi-layer which may survive hightemperature processing in certain matrix compositions to act as a debond layer, as exemplified in Section 5. Of the varied oxide fibres, many are based on crystalline e-AlzO3 (e.g. FP

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Fig. 6. Electron imaging (SEM and TEM), EDS analysis and NM R spectrum from an SCS6 silicon carbide monofilament (Textron--USA) illustrating the crystalline !mair~i) 3C polytype) ultra-line oriented grain structure, variable SiC ratio, carbon fibre core and carbon-rich surface coating. AIzO 3 from du P o n t a n d Nextel 312 f r o m 3M). T h e y are reactive with most matrices (nitrides, silicates, etc.), such that reaction barrier coatings are essential, a n d have little h i g h - t e m p e r a t u r e potential on a c c o u n t of their plasticity.

5 INTERFACE MICROSTRUCTURES 5.1 Fibre/silicate reaction mechanisms The controlled reaction between silicate matrices and silicon oxycarbide (Nicalon) fibres, identified originally by Prewo and Brennan,~'14".5 produces

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an interface structure which results in the necessary mechanical response of matrix microcracking and non-catastrophic failure. It has been established that the reaction is a fortuitous combination of silicate matrix chemistry and non-stoichiometric/non-crystallinefibre structure. A majority of research has been performed on Nicalon/Li20-A1203-SiO2 (LAS) systems in relation to the understanding of interface reactions and the modelling of mechanical behaviour. 15-19 Similar reactions have been identified in MgOAl203-SiO 2 (MAS), CaO-A1203-SiO2 (CAS) and BaO-A1203-SiO2 (BAS) systems, 2°-22 and for Tyranno as well as Nicalon fibres.22 The basic reaction required to produce the observed carbon-rich interface is the oxidation of SiC at the fibre surface: SiC + 02 --~ SiO2 + C The interracial carbon, normally in graphitic form, will act as a diffusion barrier either for gaseous oxygen (O2), if this is derived from dissolved matrix oxygen, or for silicon and oxygen down the activity gradient from fibre to matrix. Cooper and Chyung 2t have linked interface development with S i O 2 activity in the silicate matrix in explanation ofdifferences in reaction kinetics with matrix 'basicity'. The basicity is related to the degree of structural polymerisation in silicates; the addition of network modifying oxides (e.g. Li20 or CaO) increases the degree of ionic bonding (or 'non-bridging' oxygens in a glass) and decreases the S i O 2 activity. Hence matrices with higher modifier ion content should have enhanced interface reaction kinetics and those with alkali ions should have enhanced kinetics compared to the less basic calcium or magnesium silicates. This model assumes that carbon diffusion into the silicate matrix is relatively slow, consistent with the sharp carbon/matrix interface and the observation of interface growth into the fibre. A modified model 23 based on thermodynamic analysis of the stability of interfacial carbon indicates that CO formed as a result of carbon oxidation may diffuse to the C/SiC fibre interface, resulting in the alternative interface growth reaction: SiC + 2CO --* SiO 2 + 3C Hence CO diffusion, rather than Si + O diffusion, may be a rate-limiting step. They also suggest that the influence of multivalent oxide additions (such as Nb205 or As203) may be to increase the CO activity and hence kinetics of this reaction. However, the redox reactions involved in creating CO consume part of the interfacial carbon and in the case of transition metal additions this is observed in precipitated form (NbC, etc.) 15-ts although not always in the form of a continuous interfacial layer which might suppress reactions of the type described above.

234

M. H. Lewis, V. S. R. Murthy

Quantitative studies of interface reaction kinetics with different matrix chemistries are difficult because of the number of interdependent variables, but a limited study of carbon interface thickening in LAS indicates diffusion control. 24 In this work typical interface thicknesses were shown by means of AES to increase to about 500nm in 4 h at ll00°C, with parabolic kinetics, and an activation energy of about 25 kcal/mol is claimed to be typical of that for gaseous diffusion in amorphous solids. The model for interface formation, based on oxidation of the SiC fibre component and subsequent rate control via diffusion through a carbon layer, is over-simplistic. In real silicate systems matrix constitution and structure are changing in parallel with the interface reaction. One of the benefits of using silicate glass frit as a matrix precursor is its rapid densification by viscous flow. However, fabrication must be conducted at temperatures below that for fibre degradation and high enough for interface development. For most silicates this may be below the liquidus (Te--Fig. 7), such that matrix crystallisation kinetics may compete with interface development and also inhibit final densification. The ideal thermal cycle in which densification and interface development precede a controlled matrix crystallisation (Fig. 7) is rarely achieved. 5.2 ln-situ reacted interfaces

The following examples of interface structure, taken largely from programmes at Warwick University, illustrate the principles and techniques outlined above and the current limitations of reaction routes to interface development. 5.2.1 Glass matrices

Borosilicate (Pyrex) glasses, which have low thermal expansion (3.3 x 10-6/°C) and softening temperature, have been extensively studied 25'26 as matrices for Nicalon fibres in generating high specific strength/stiffness composites with limited temperature capability (500°C). The development of interfaces with the required debond/shear properties for non-brittle failure occurs above --~800°C with optimal UTS values developed below the liquidus (,-~ll00°C). Unlike the higher temperature glass/ceramic compositions, borosilicates exhibit a limited in-situ interface reaction 20.27 with carbon precipitation barely detectable below 1000°C and preceded by a 50-rim band of cross diffusion (Fig. 8(a)). At higher hot pressing temperatures a continuous 10-20 nm carbon layer is developed at the original fibre/matrix interface and a diffusion band of greater width within the fibre (Fig. 8(b)). Tyranno fibres behave similarly except for an enhanced diffusion band and the precipitation of TiC particles beneath the

Microstructural characterisation of interfaces

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developing carbon layer (Fig. 8(c)). Tyranno fibres have greater oxygen content and degree of structural disorder which induces enhanced diffusivity for matrix elements. The reduction in carbon layer thickness is anticipated from the reduced free carbon in Tyranno and as a result of competition from the precipitation of TiC following titanium diffusion to the fibre surface. 28 Conventional Pyrex compositions present an example of undesirable matrix crystallisation when hot pressed below the liquidus. Heterogeneous nucleation ofcristobalite (SiO~) occurs on glass frit surfaces prior to porosity removal and on fibre/matrix interfaces (Fig. 8(d)). Nucleation kinetics are maximised in the 900-1000°C interval such that hot pressing in excess of 30 min produces a high volume fraction of cristobalite (thermal expansion 12 x 1 0 - 6 / ° C ) and extensive matrix cracking. 2v'28 Borosilicate glass compositions which do not crystallise during hot pressing over a range of temperature and time provide a good basis for studies of micromechanical properties of interfaces in varied structural states from initial carbon precipitation to continuous layers of variable thickness. These interfaces are structurally stable up to the intended application temperatures but very long

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(b) Fig. 8. Interface microstructure and microanalysis in borosilicate glass matrix/SiC (Nicalon or Tyranno) fibre CMCs: (a) Nicalon, 950°C fabrication temperature; (b) Nicalon, II00°C fabrication temperature.

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TABLE 1 Constitution and Thermal Characteristics of Potential Matrix Glass Ceramtcs

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heat treatment at 550:C amplifies the diffusional segregation oF matrix elements in the fibre surface beneath the debond laver. Figure 8(a)-(d) provides examples of the analytical EM techniques outlined above, in particular the high spatial resolution light element analysis of interface structure and the selective debonding property associated with carbon enrichment, even in the initially precipitated state. 5.2.2 Glass ceramic matrices

Crystalline silicate matrices, formed from glass precursors, are based on low thermal expansion phases which are well established for monolithic glass ceramics in LAS, MAS, CAS and BAS svstems (Table 1L The comparatively high liquidus temperatures in these svstems dictate that fabrication is normally conducted in the 1000 1300:C interval for times consistent with full densification without severe fibre degradation. Even f'or nonstoichiometric compositions these temperatures are normally below the liquidus such that matrix crystatlisation accompanies densification and interface reaction. Typical interface structures in Nicalon-reinforced MAS and CAS matrix composites (Fig. 9(a,b)) have a well-developed graphitic layer form with minimal diffusion zone, indicative of rapid reaction kinetics and the influence of graphite as a diffusion barrier. Where transition metal oxides (e.g. TiO 2) occur as nucleating agents within the parent glass, some of the carbon reacts to form carbide particles on the matrix side of the interface (Fig. 9(a)). An extreme example of a graphite-containing interface occurs in BAS (Fig. 9(c)) driven by the requirement for higher temperature fabrication and probably by reduced matrix SiO2 activity. The interface layer is typically 250rim wide containing an interconnected array of graphite (00t) flakes within a SiO_, rich environment (Fig. 9(c)). A significant in-diffusion of metallic matrix elements (Ba, AI) occurs beyond the carbon layer at these temperatures, especially in Tyranno fibres.

Microstructural characterisation of interfaces

239

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Fig. 9. Interface microstructure (TEM) and microanalysis (EDS) in glass ceramic matrix/ Nicalon fibre CMCs: (a) MAS matrix (TiOz doped); (b) CAS matrix; (c) BAS matrix (inset enlargement of graphite-rich interface with electron diffraction pattern). Although temperature has a primary influence on interface reaction kinetics and carbon graphitisation, it is possible to compare the influence o f constitution for selected eutectic silicates. The increased matrix basicity for MAS results in a thicker interface layer than for borosilicates under identical fabrication conditions. For example, at 950°C for 1 h with Nicalon (Fig. 10(b)) the a m o r p h o u s carbon-enriched zone is 30 nm thick compared with 10 nm for the borosilicate matrix.

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Fig. 10. Oxidative degradation of interface microstructure: (a) SiO 2 bridging layer in MAS following carbon removal: (b) and (c) low-temperature MAS interface with gas bubble nucleation for a prolonged fabrication time.

Microstructural characterisation of interfaces

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A key question currently being pursued at Warwick University is the extent to which interface properties (especially z and am relative to UTS)may be controlled within the temperature/time/constitutional variables of fabrication. Preliminary data indicates that the 10-nm amorphous precipitated carbon types of interfaces have higher r and a mvalues than the 100-nm graphitic carbon layers (M. W. Pharaoh, 1990, unpublished). Carbon-rich interfaces result both from SiC oxidation reactions described above and from residual free carbon initially present in the fibre. It is unlikely that such interfaces can be developed by in-situ reactions with the nitride-based fibres (HPZ--Dow Corning, Fiberamic--Rhone Poulenc). An example for HPZ/MAS composites fabricated at 950°C shows a fine cellular structure, believed to be a phase-separated mixture of SiO2 with residual carbon (Fig. 11). The evolved nitrogen may be partially dissolved in the matrix glass prior to crystallisation. These high SiO2 interfaces have not been mechanically characterised but are unlikely to provide the low debond stresses associated with high carbon enrichment.

5.2.3 Interface stability An established interface degradation mechanism is the oxidation of carbon to gaseous (CO or CO 2) form followed by SiO 2 bridging due to SiC fibre oxidation. Atmospheric oxygen, which is necessary for this mechanism, may be diffused through the silicate matrix or be 'piped' longitudinally from exposed fibres at free surfaces or microcracks. There is a carbon removal threshold above ~600°C with the onset of SiO2 bridging at -,-800:C. An example following 5-h exposure at 1200:C (Fig. 10(a)) shows complete bridging and crystallisation of the SiO 2 layer, resulting in brittle composite behaviour. 2z A more subtle degradation mechanism has been detected on sustained exposure to fabrication conditions within the (normally reducing) graphite die atmosphere. Figure 10(c) shows the nucleation of gas bubbles near the surface of the carbon-enriched interface layer for MAS/Nicalon composites. They are probably carbon oxidation products which occur at a stage where SiO2 activity has been amplified locally by interface cross-diffusion and the inward transport rate for CO and 02 is reduced due to thickening of the interface layer.

5.3 Presynthesised interfaces 5.3.1 Monofilament/nitride matrix systems The higher temperature potential of non-oxide matrices, such as silicon nitride, has dictated the use of SiC-based monofilaments which are structurally stable above 1000°C although they suffer a significant reduction

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Interface reaction layer in a MAS HPZ nitride-based fibre composite. with EDS and ELS analysis of the C SiO,-rich phase-separated structure.

in strength and creep resistance above 1300°C. Earlier matrices were fabricated by reaction sintering (RBSN}, with the benefit of minimal shrinkage but the problem of high porosity and hence of atmospheric interaction with fibrematrix interfaces, More recently, dense silicon nitride matrices produced by hot pressing with liquid sintering aids have been fabricated within SCS6 monofilament arrays, z9"3° Interfaces for both matrix types consist of thick (4 pm) carbon-rich precoatmgs (see Section 4l which have the dual function of reaction barrier and debond layer. SEM sectional imaging and analysis techniques have adequate spatial

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Fig. 12. Textron CVD monofilaments within a silicon nitride matrix unidirectional composite fabricated by hot pressing tape-cast laminates of SRBSN constitution. 3~ The SEM back-scattered electron image (a) shows atomic number contrast from carbon core and carbon-rich interface layers which have survived the high-temperature fabrication. Some reaction of the outer carbon layer is detected via EDS and X-ray mapping itbi and {c)l.

resolution for studies of the relatively coarse interfacial microstructure (Fig. 12j. Back-scattered electron imaging shows the low atomic number carbon core and carbon-rich coatings in dark contrast, together with a matrix reaction layer, of variable thickness. X-ray line-scanning and mapping define the mixed chemistry of this reaction layer, derived from both matrix elements and carbon (Fig. 12). An essential structural element for composite performance is survival of the inner carbon-rich interlace layer, which performs a debond function, following 1700°C hot pressing with appropriate sintering additives. 3°

5.3.2 Nicalon fibre precoat&gs Of the various attempts to replicate the mechanical response of carbonenriched interfaces produced by in-situ reactions, the majority have used single-layer CVD coatings on Nicalon. Elemental carbon, boron, silicon and compound BN coatings have been studied with variable thickness and their influence on mechanical response measured, normally in simple flexure tests. 6 Some comparisons of interface shear stress have been made by means of indentation techniques with a general conclusion that measured z values of less than 2-5 MPa are required for non-brittle response, which is only observed for carbon or BN coatings. Matrices studied most frequently are SiC fabricated by CVI which, in low porosity form, should act as an effective oxidation barrier.

Fig. 13. Exam pies ofinterfaces produced by CVD precoating of Nicalon or Tyranno fibres: (a) titanium nitride crystalline single-layer interface in a glass-ceramic matrix: (bi and !el bilayer interface of the type modelled in Fig. 1. The potential ofstoichiometric crystalline SiC as a reaction barrier for high-temperature composite fabrication is illustrated by its nonreactive interface with a BAS ~lass ceramic Idt.

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The implication of these studies is that non-reactive layered (vander Waals) interface structures, such as graphite or hexagonal BN, are necessary for low debond and Tvalues. However, there have been few studies involving more varied chemistry with control of substrate temperature and coating state. Metallic or intermetaUic compound coatings may offer low-energy debond interfaces with oxide matrices or fibres but will normally require reaction barriers with non-oxide fibres. Figure 13 provides examples of single-layer interfaces, and of the less frequently studied multilayer interfaces, introduced by Nicalon fibre coating. The TiN layer, produced by plasma-assisted CVD, exemplifies the problems of coating uniformity and some reaction with fibre and matrix during fabrication, resulting in poor mechanical response. The bi-layer coating exemplifies a thin van der Waals layer protected by an oxidation reaction barrier of SiC (1/~m), applied to a Nicalon fibre by conventional CVD. 31 The fibre debond normally occurs at the Nicalon interface during isolated fibre fracture (Fig. 13(c)) and within a glass-ceramic matrix composite (Fig. 13(b)). The potential of stoichiometric crystalline SiC as a reaction barrier for high-temperature glass ceramics has been assessed by using dispersions of SiC whiskers in refractory silicate matrices (e.g. BaOA1203-SiO 2 near to the stoichiometric celsian composition) fabricated at temperatures up to 1600°C (Fig. 13(d)). Composite fabrication or application in this regime must await the development of new stable fibres.

6 CONCLUSIONS (i) There is now a moderate understanding of the interracial properties that are necessary for non-brittle behaviour in CMCs based on theoretical modelling of unidirectional fibrous microstructures. Experimental confirmation of the quantitative relationship between interracial parameters and mechanical properties, such as the matrix microcracking stress, is limited by the small range of microstructures which have been prepared, with little control of interracial debond and shear stresses. (ii) Microstructural and compositional characterisation of interfaces is possible at the nanometre level of resolution using a combination of electron beam imaging and spectroscopic techniques which have been exemplified in this paper. Transmission electron microscopy combined with 'windowless' X-ray energy dispersive spectroscopy and electron energy loss spectroscopy offers the best spatial resolution but the analytical techniques have a limited quantitative capability for light elements.

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(iii) The most frequently studied microstructures are those with polymer precursor fibres (such as Nicalon or Tyranno SiC) in silicate glass ceramic matrices. This has been motivated by the ability to fabricate these CMCs at temperatures below that for fibre degradation, the control of matrix thermal expansion and the fortuitous development of carbon-rich interfaces via in-situ reaction. (iv) The formation of carbon-rich interfaces in Nicalon/silicate matrix CMCs is based on the non-crystalline/non-stoichiometric (oxygencontaining) fibre structure. The precipitation of carbon originates from a SiC oxidation reaction, with the fibre surface, driven by differences in SiO 2 activity across the interface and rate-controlled by gaseous (O2 or CO) diffusion through the reaction layer. Carbon formed by oxidation of SiC, supplemented by free carbon in the original fibre, exhibits a gradation in structure from the continuous graphitic layer formed at high temperatures to the amorphous particulate state at lower temperatures (typically <1000°C). This range of interface structure may exhibit the necessary debond property with a significant variation in interface shear stress. (v) The control of interfacial reactions in silicate matrix CMCs is not always compatible with fibre stability, densification kinetics and matrix microstructure. Fabrication is typically conducted at temperatures below the silicate liquidus such that undesirable crystatlisation may occur (as in Pyrex glasses) or there is little control over crystal nucleation and growth in glass ceramic matrices. (vi) Carbon-rich interfaces, whether formed by fibre/matrix reaction or by fibre precoating, are susceptible to oxidative degradation above 500°C Carbon removal modifies the interface debond and shear length adjacent to a matrix microcrack with consequent change in composite response. The ultimate problem of reversion to brittle failure is due to a subsequent bridging of the matrix/fibre gap by the SiC surface oxidation product of SiO 2. (vii) A suppression of interface oxidation reactions may be achieved bv using matrices with low oxygen diffusivity such as stoichiometric SiC or Si3N,, but these are not easily fabricated in the absence of fibre degradation. Non-stoichiometric SiC may be introduced by CVI at temperatures which avoid Nicalon degradation but hot pressed silicon nitride matrices may only be fabricated with SiC monofilaments which have higher temperature stability. In both cases the preformed CVD graphitic interfaces survive both fabrication and high-temperature oxidation treatments. (viii) For matrices with high oxygen diffusivity, such as silicates, an oxygen barrier may be provided as a CVD fibre coating overlying the debond

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layer. A C/SiC bi-layer is one example, the SiC acting as an oxygen barrier during service and a reaction barrier during fabrication. (ix) Although interface instability problems during fabrication and service are potentially soluble for refractory matrices by using a range of bilayer chemistry and innovative deposition processes, the ultimate temperature limit is that of fibre degradation. Current limits for small fibres approach 1300°C for the more reactive S i - N - C - b a s e d compositions. For higher temperatures, stoichiometric crystalline SiC has potential but is not available in small-diameter, weaveable form. Diffusional creep in such fine-grained fibres will present the ultimate temperature ceiling for composite application.

A C K N O W L E D G E M ENTS The authors wish to thank the SERC, MOD, Rolls-Royce plc and BP Research for supporting research programmes at Warwick University, from which some of the material reviewed in this paper has been derived.

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28. Murthy, V. S. R., Pharaoh, M. W. & Lewis, M. H., Interface microstructure and matrix crystailisation in SiC-borosilicate (Pyrex) composites. Materials Lett., 10 (1990) 161-4. 29. Foulds, W., LeCostaouec, J. F., Landry, C., DiPietro, S. & Vasilos, T., Tough silicon nitride matrix composites using Textron silicon carbide monofilaments. Ceram. Eng. Sci. Proc., 10 (1989) 1083-99. 30. Razzell, A. G. & Lewis, M. H., Silicon carbide monofilament reinforced silicon nitride ceramic composites. Proc. 15th Annual Conf. on Composites and Advanced Ceramic Materials, Cocoa-Beach, 1991 (Amer. Ceram. Soc., in press). 31. Cain, M. G. & Lewis, M. H., Presynthesised interfaces for ceramic matrix composites via fibre coating. University of Warwick (in prep.).