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Application of diffusion barriers in composite materials
Materials Science and Engineering A242 (1998) 235 – 247
Application of diffusion barriers in composite materials P. Mogilevsky *, A. Werner, H.J. Dudek Institute of Materials Research, German Aerospace Research Establishment (DLR), D-51140 Cologne, Germany Received 27 December 1996; received in revised form 15 July 1997
Abstract Application of diffusion barriers in composite materials is discussed from the point of view of thermodynamic and kinetic stability of the reinforcement/matrix interface. Two types of interface protective coating, i.e. inert diffusion barriers and reactive diffusion barriers can be distinguished. The application of one coating or the other depends on the type of damage the material suffers as a result of the interface reactions. A double-coating concept for SiC fiber-reinforced titanium matrix composites (TMC) is discussed from the viewpoint of this classification of diffusion barriers. Application of the C/TiB2 coating system for SiC fiber-reinforced TMC is considered in detail, and some approaches to improve the performance of this coating were suggested and examined experimentally with TiB2 coatings produced by magnetron sputtering deposition from a TiB2 target. The effect of the composition and the microstructure on barrier properties of TiB2 layers was examined. A crystalline microstructure with relatively large grains and of stoichiometric composition was found to be optimal for TiB2 protective coatings in SiC fiber-reinforced TMCs. A strong impact of layer stoichiometry on the diffusion of metal atoms through a crystalline TiB2 layer was found. However, stoichiometry did not significantly affect the diffusion of carbon. Crystalline TiB2 of proper stoichiometry can provide excellent protection for carbon coated (SCS6) SiC fibers in titanium aluminide based metal matrix composites (MMC). © 1998 Elsevier Science S.A. All rights reserved. Keywords: Titanium matrix composites; SiC fibers; Interface reactions; Diffusion barriers; Fiber coating
1. Introduction Ceramic-metal composite materials such as metal matrix composites (MMC), in particular titanium matrix composites (TMC), are high-performance materials for many advanced applications, e.g. aerospace high temperature applications. The performance of these materials is, however, often limited by the interface phenomena. The matrix/reinforcement interface can suffer from the coefficient of thermal expansion (CTE) mismatch, mechanical incompatibility, and thermodynamic instability. The last phenomenon is due to the fact that at high temperatures these materials often represent highly non-equilibrium systems, which suffer from spontaneous chemical reactions between the components. In many cases this is the major factor affecting the performance of the composite. Thus, SiC fiber* Corresponding author. Present address: University of Illinois at Urbana-Champaign, Materials Research Laboratory, 104 South Goodwin Avenue, Urbana, IL 61801-2985, USA. Tel.: + 1 217 3332367; fax: + 1 217 2442278. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 1 0 - 8
reinforced TMCs have received a great deal of attention for advanced aerospace applications. At the same time, their use is hindered by chemical reactions at the fiber/ matrix interface during processing (high temperature consolidation) and at service conditions [1–14]. In addition, high temperature post-consolidation heat treatment, which often can potentially improve the mechanical properties of a composite, is prevented by the rapid chemical reaction at the fiber/matrix interface [15,16]. Commercial SiC fibers used for the reinforcement of titanium alloys have a carbon-rich protective coating, which, given the required time of service of about 10 000 h, effectively prevents the reaction with the SiC fibers at temperatures up to 600°C [17]. Carbon coating also eliminates chemical bonding between SiC fibers and the matrix (‘weak interface’) and thus prevents crack propagation from the matrix into the fibers. In addition, carbon coating protects the fibers from surface abrasion [18]. At temperatures \ 600°C, however, the reaction between the carbon coating and Ti matrix becomes too fast to protect the fibers for a reasonable period of time. Further improvement of
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thermal stability of SiC fiber-reinforced TMC may be achieved using advanced coating systems, such as TiB2/ C [3–13,19,20] or TiC/C [14] double coatings. In such coating systems the carbon coating protects the fibers from surface abrasion and crack propagation from the matrix, while the additional coating protects the carbon layer from fast reaction at temperatures \ 600°C. In this paper we analyze the application of different types of protective coating (diffusion barriers) in composite materials and consider the use of a C/TiB2 coating system for SiC/Ti MMC. Application of the phases of the Ti– B system in protective coatings for SiC fibers in TMC as well as a number of ways to improve the performance and thermal stability of these coating systems are discussed. Recent experimental results on the performance of different types of C/Ti–B coatings for SiC fiber-reinforced TMC are reviewed.
2. Background
2.1. Application of diffusion barriers in composite materials The primary goal of any diffusion barrier is to separate two components and to prevent or slow down the interdiffusion and/or possible chemical reaction between them. Let us consider an uncoated reinforcement (F) introduced into a matrix material (M) (Fig. 1(a)). Two different types of diffusion barrier can be distinguished. The first class represents a layer of a material which is chemically inert with respect to both components to be separated (Fig. 1(b)). For example, a TiC protective layer on carbon coated SiC fibers in a Ti matrix is inert with respect for both carbon coating of the fibers and the matrix. Generally speaking, any naturally grown diffusion barrier (i.e. formed at the interface as a result of chemical reaction between the components) falls into this category. This is because the
Fig. 1. Interface reaction: (a) no diffusion barrier; (b) inert diffusion barrier; (c) reactive diffusion barrier. M, matrix; F, reinforcement; B, diffusion barrier; R, reaction zone.
reaction products are usually in thermodynamic equilibrium with the initial materials (uncommon cases of the formation of non-equilibrium phases during interface reactions is omitted here for simplicity). The performance of this type of diffusion barrier, either artificially created or formed during interface reaction, depends on the respective diffusion characteristics of each of the components in the barrier layer. The second class of diffusion barriers represents a coating which can react with one of the components (Fig. 1(c)). Carbon coating on SiC fibers (such as on SCS6 Textron fibers, SiC fibers coated with a carbon protective layer containing small additions of SiC) in TMC is a well-known example of this class of diffusion barriers. The kinetics of the interaction in this case is determined by the effective diffusion rate of the components in the resulting reaction zone. Accordingly, the performance of the diffusion barrier may strongly depend on the composition of the matrix. Thus, applicability of SCS6 fibers in Ti-6-4 alloy is practically limited to a temperature of 600°C. At higher temperatures the interdiffusion in TiC formed between the carbon coating and the matrix becomes too fast [17]. However, a very thin (100 nm) layer of a phase (Ti,Zr)x Siy which forms at the interface between SCS6 fibers and the IMI 834 titanium alloy provides sufficient stabilization of the interface at temperatures up to 700°C [21,22]. The principal difference between the two types of diffusion barrier and the resulting difference in their applications can be understood if we analyze the kinetics of the growth of the reaction zone at the composite interface in both cases. Let us assume that we have a composite consisting of a matrix and an uncoated reinforcement, and our task is to prevent or slow down the chemical attack of the reinforcement. Let us assume, for simplicity, that only matrix components diffuse (Fig. 1). Fig. 2 shows the kinetics of the development of the reaction zone for both types of the diffusion barrier as compared to an uncoated reinforcement. Typically, parabolic diffusion-controlled kinetics are observed for uncoated reinforcements (dashed line, Fig. 2). The same kinetics is most often observed experimentally for the case of inert diffusion barriers, with the important difference that an apparent ‘incubation’ period t%0 for the reaction is observed (dotted line, Fig. 2). Closer examination of the kinetics of the process reveals, however, that the apparent ‘incubation’ period in this case results from the parabolic approximation of the experimental data obtained on the advanced stage of the interaction, when the process is controlled by the interdiffusion through the growing reaction product. In reality, nevertheless, the reaction already takes place during the ‘incubation’ period as well. At this stage, when the thickness of the reaction zone is small, the process is controlled by the interdiffusion of the compo-
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Fig. 2. Kinetics of the reinforcement consumption and the reaction zone growth during interface reaction: A, inert diffusion barrier; B, C, reactive diffusion barrier (reinforcement consumption and overall reaction zone, respectively).
nents through the inert barrier of constant thickness and, therefore, has linear kinetics (A-I, Fig. 2). Later, with the increase of the reaction zone thickness, the control of the whole process is gradually overtaken by the diffusion through the reaction zone and the process finally gets the parabolic kinetics (A-II, Fig. 2). The apparent incubation period t%0 depends on the thickness of the barrier and on the coefficients of interdiffusion in the diffusion barrier and the reaction zone: t%0 :
2D0 rh 2b D0 2b
(1)
where D0 r and D0 b are the effective interdiffusion coefficients of the reaction zone and of the diffusion barrier, respectively, and hb is the diffusion barrier thickness. Due to typically small thickness of the reaction zone, the stage of linear growth is difficult to observe and may be missed in kinetic measurements. If such a thin reaction zone does not harm the properties of the composite, it can be ignored. More often, however, even a minor reaction with the reinforcement may strongly affect the properties of the composite, and the interpretation of the kinetic measurements by a simple parabolic kinetics with an incubation period may be misleading. Let us now consider the kinetics of the interaction in case of a reactive diffusion barrier. In this case a true incubation period t0 may be observed (B-I, Fig. 2). This is the period when the diffusion barrier reacts with the matrix. It was observed experimentally for carbon coated SiC fibers (SCS6) in Ti matrix, that unless all of the carbon coating is consumed, the tensile strength of the fibers (but not of the composite, however) remains unchanged [17]. At the same time it is known that even
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a minor reaction on the surface of SiC fibers results in a catastrophic drop in the fibers strength [23]. Therefore, we can conclude that no chemical attack of the fiber takes place unless all (or almost all) of the carbon coating (a reactive diffusion barrier) is consumed. Intuitively it can be explained by the strong affinity of the diffusing species to the material of the diffusion barrier (as is the condition for the barrier to be reactive). In this situation any atom diffusing through the barrier will be ‘trapped’ by the reaction with it as long as it still has to travel through the remaining barrier to reach the reinforcement. Contrary, in the case of an inert barrier where the diffusion species by definition have no affinity to the barrier material, they can diffuse through the barrier and attack the reinforcement already in the early stage of the process. After the reactive barrier has been consumed, the process will involve the interaction between the matrix and the reinforcement. The initial reactive barrier is usually transformed into the reaction products which are inert with respect to both the reinforcement and the matrix (for example, stoichiometric TiC formed between carbon coating of SCS6 fibers and Ti matrix is unreactive with respect to both SiC and Ti). Accordingly, the kinetics of the following interaction between the matrix and the reinforcement will be very similar to the case with an inert barrier (B-II and B-III, Fig. 2). The advantage of a reactive diffusion barrier can be understood from the above consideration. Unlike the inert one, it provides a true incubation period, during which the reinforcement can be kept completely out of the interaction. This may be critically important for such reinforcements as SiC fibers (apparently as well as other fiber reinforcements): as has been already mentioned, even minor a reaction on their surface results in a drastic decline in the fiber properties [23]. In this case it is simply not enough to slow down the reaction with the fiber, it should be completely prevented. Apparently, the only way to achieve this goal is the application of a reactive diffusion barrier. However, reactive barriers have their own drawbacks. Let us consider now the kinetics of the growth of the overall reaction zone (not to confuse with the consumption of the reinforcement) at the composite interface for both inert and reactive diffusion barriers. In the case of an inert barrier the thickness of the overall reaction zone follows the graph of the reinforcement consumption, since no other process is involved (curve A, Fig. 2). However, in the case of a reactive diffusion barrier the situation is different. During the incubation period, when the reinforcement does not react, the thickness of the overall reaction zone will follow the parabolic kinetics of the reaction between the matrix and the diffusion barrier (C-I, Fig. 2). After the barrier has been consumed, the reaction zone will continue to grow due to the reaction with the reinforce-
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ment, and the overall process will follow the kinetics of the reinforcement consumption (C-II and C-III, Fig. 2). The problem here is that the kinetics of the reaction between the matrix and a reactive diffusion barrier may be even faster than between the matrix and the reinforcement. Thus, from the kinetic measurements [24], the reaction constants k= 4.7 ×10 − 15 m2 s − 1 at 1000°C and k = 1.9 ×10 − 16 m2 s − 1 at 800°C can be evaluated for the interaction between polycrystalline SiC and pure Ti. At the same time for the interaction between pure carbon (graphite) and Ti in identical experimental conditions the respective values of k= 1.3 ×1 − 14 m2 s − 1 and k =1.6 ×l0 − 15 m2 s − 1 were obtained [25]. Accordingly, in this situation the growth of the overall reaction zone with such a diffusion barrier will be faster than without it (Fig. 2). The growth at the composite interface of a reaction zone, often consisting of brittle phases, results in the stress concentration, crack initiation, and the development of the internal transformation stresses, thus promoting the material failure [21]. It has been reported that the tensile strength of SCS6/IMI 834 composite initially remains unchanged during the heat treatment at 800°C, but drops about 50% when the thickness of the reaction zone between carbon coating and the IMI 834 matrix reaches a critical value of about 3 mm [26]. At this stage, the carbon coating of the SCS6 fibers is still not consumed, and the observed property decline can not be attributed to the fiber damage, but to the development of a brittle reaction zone. Therefore, we can summarize the advantages and disadvantages of both types of diffusion barrier as follows. An inert coating may have unlimited service time and is capable of slowing down the growth of the reaction zone, but it is unable to completely prevent the chemical attack of the reinforcement. In contrast to this, a reactive diffusion barrier has a limited service time and may cause an accelerated growth of the reaction zone, but it is capable to completely prevent (for a certain period of time) the chemical attack of the reinforcement. Accordingly, the application of one coating or the other will depend on the type of damage the material suffers as a result of the interface reactions. If the material is highly sensitive to reinforcement damage (as is the case for SiC and, apparently, other types of fiber reinforcements), the application of a reactive diffusion barrier is necessary. For other types of composite materials, e.g. particle-reinforced composites, the use of a proper inert diffusion barrier may suffice.
2.2. Fiber protecti6e coating for SiC fiber-reinforced TMC
the carbon coating as a means of protection of the fiber from surface abrasion, crack propagation, and immediate chemical attack from the matrix should remain as an inner layer of the coating. As has been already mentioned, a Ti-based composite reinforced with carbon coated SiC fibers fails not due to the fiber damage, as in the case of uncoated fibers, but due to the development of the brittle reaction zone at the coating/ matrix interface. Therefore, an ideal coating system for this type of materials would be a double coating consisting of one layer of a reactive coating (carbon) to protect the fiber from an immediate chemical attack, and another layer of an inert coating to slow down the growth of the reaction zone. With this coating system, the advantages of both types of diffusion barrier can be combined. The material of this second diffusion barrier of the coating system must, ideally, have high barrier properties for carbon and the components of the matrix, be thermodynamically stable with respect to both the carbon coating and the matrix, and its components should not dissolve and/or diffuse into the matrix. Taking into account numerous additional criteria which are not in scope of this paper (adhesion, CTE mismatch, mechanical compatibility), it is clear that the task of finding a material fulfilling all aforementioned requirements may be quite difficult or impossible, and that a compromise may have to be made. Different materials have been considered as an additional protective coating for carbon coated SiC fibers in Ti matrices, such as TiC [5,9,14], TiN [11,12], and TiB2 [3–13]. It is, however, hardly probable that the TiC layer can provide the necessary level of the interface stabilization for carbon coated SiC fibers in the Ti matrix. TiC forms naturally during the interface reaction between the carbon coating and the matrix, and the very fact that an additional protection at temperatures \600°C is needed means that at these temperatures the interdiffusion through TiC is too fast. The possible effect of an additional TiC layer on the kinetics of the reaction at the carbon/matrix interface can be estimated as follows. Assume that the growth of TiC at carbon/Ti matrix interface follows normal parabolic kinetics, and that normally the reaction zone of the critical thickness hc develops during the time tc (Fig. 3). If we place an initial TiC layer of a thickness h0 at the C/Ti interface, the reaction will follow the same kinetics, but starting from the point 0% (Fig. 3). Accordingly, we can calculate the time t%c during which a reaction zone of the critical thickness hc will develop on top of the initial TiC:
t%c = tc 1+ In light of the above discussion, the concept of a protective coating for SiC fibers in high temperature ( \600°C) TMCs may be suggested as follows. First,
2h0 hc
(2)
According to [26], the critical thickness of the reaction zone, after which a sharp decline in the mechanical
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properties of SCS6/IMI 834 composite occurs, is hc 3 mm. Assuming the thickness of the initial TiC layer in the range of 3–6 mm (a reasonable value), we obtain from Eq. (2) that such a layer would prolong the service time of the material by a factor of 3 – 5, respectively. This is obviously not enough to compensate the effect of rising temperature. According to [27], the activation energy of interdiffusion in TiC is : 330 920 kJ mol − 1, which means an increase of the coefficient of interdiffusion by a factor of :4900 with the temperature rising from 600 to 800°C (the target service temperature of modern TMC for aerospace applications). Application of TiN as an additional diffusion barrier for carbon coated SiC fibers and its barrier properties has not been studied in detail. It should be pointed out, however, that the relatively high solubility and diffusion rate of nitrogen in Ti [28 – 31] may have a negative effect on the performance of such coating system. Titanium diboride TiB2 has recently received much attention as a potential second coating for carbon coated SiC fibers in TMC applications. The choice of TiB2 is substantiated primarily by the low solubility (0.05 and 0.03 wt.% in a- and b-Ti, respectively, at 88695°C [28]) and the low rate of diffusion (: 9× 10 − 13 m2 s − 1 at 1000°C in b-Ti [29]) of B in Ti alloys. As follows from the Ti– C – B phase diagram [32], TiB2 is apparently thermodynamically stable with respect to carbon. According to the crystallographic data, TiB2 has a closed-packed crystallographic structure consisting of alternating Ti and B atomic layers [33], which may be indicative of good barrier properties of the compound. Indeed, it has been reported that a TiB2 coating is capable of preventing the diffusion of carbon, thus minimizing the consumption of the carbon protective coating [11]. TiB2 was reported to be the most promising choice for the protective coating in SiC/Ti composites, also from the point of view of its effect on
Fig. 3. The effect of an additional TiC coating on the kinetics of the reaction between the carbon coating of the SiC fiber and Ti matrix.
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the fiber strength [9,12]. However, reactive growth of TiB needles at the Ti matrix/TiB2 interface has been reported [3–9]. The formation of these needles enhances stress concentration and interface failure and is the major drawback of TiB2 as a protective coating for SiC fiber-reinforced TMC. Needle-like growth seems to be a general feature of the diffusion growth of TiB. It was previously observed in arc-melted Ti–Al–B alloys [33], during B4C–Ti interaction [34], in binary Ti–B planar diffusion couples, as well as at a Ti/B-fiber interface [35], at a Ti/TiB2-particle interface [36], and in planar Ti–TiB2 diffusion couples [37]. Thus, the interface phenomenon in a C/TiB2 coated SiC fiber-reinforced TMC is very complex and may involve multi-component diffusion in combination with thermodynamic and kinetic instabilities of the TiB2/matrix interface. It should be noted here, that from this point of view TiB2 is, in effect, a reactive diffusion barrier. Therefore, C/TiB2 coating system does not fulfill the requirements of an ideal coating system which, as discussed in the previous section, should consist of a combination of a reactive and an inert diffusion barrier. Let us now consider the physicochemical processes that can potentially take place at the interface between a C/TiB2 coated SiC fiber and a Ti matrix. As was already discussed, as long as the carbon coating has not been consumed, the fiber components are not involved in the interface processes. Taking also into account that the C/TiB2 interface is thermodynamically stable, we may consider the following processes at the interface between a C/TiB2 coated SiC fiber and a Ti matrix (Fig. 4): 1. Reaction of the TiB2 coating with the Ti matrix with formation of TiB. 2. Diffusion of carbon through the TiB2 layer and its reaction with the matrix. 3. Diffusion of the matrix’ elements through the TiB2 layer and their reaction with the carbon coating. All three processes lead to the degradation of the interface and deterioration of the properties of the entire composite. The first process reflects the thermodynamic and kinetic instability of the TiB2/Ti matrix interface, which leads to the transformation of the initially homogeneous TiB2 layer into the needle-like TiB. The second and the third processes are concerned with the proper barrier properties of the TiB2 coating and lead to the consumption of the carbon protective layer and to the formation of titanium and/or ternary carbides at both carbon/TiB2 and TiB2/matrix or TiB/ matrix interfaces. Accordingly, the problem with the C/TiB2 coatings is two-fold: 1. The development of TiB needles at TiB2/matrix interface should be eliminated. 2. High barrier properties of TiB2 coating against possible diffusion of carbon and/or matrix elements have to be ensured.
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Fig. 4. Physicochemical processes at the interface between C/TiB2 coated SiC fibers and Ti matrix.
Below we suggest a number of approaches to meet these requirements and to improve the performance of the C/TiB2 coating system.
3. Experimental Samples for investigation of the interaction at TiB2/ TiAl and TiB2/Ti interfaces were prepared by magnetron sputtering deposition of TiB2 on TiAl and on Ti substrates. To model the interdiffusion at the interface between C/TiB2 coated SiC fibers and Ti-matrix C/ TiB2/TiAl ‘sandwich’ samples were prepared by magnetron sputtering deposition of TiB2 layers on graphite substrates, followed by a magnetron deposition of :30 mm thick g-TiAl layers. Heat treatments were performed in a vacuum of 10 − 4 Pa. Microstructure and composition of the as-deposited materials and annealed samples were studied by XRD, SEM, and EDX at 10 keV using pure Ti and Al standards. Boron content was evaluated by difference. Polished cross-sectional samples cut perpendicular to the interface were used for SEM and EDX examinations. For SEM examinations some of the samples were etched in a mixture of HF and HNO3.
4. Results and discussion
4.1. TiB needle formation As follows from the Ti – B phase diagram [30], the TiB2/Ti interface is thermodynamically unstable and TiB naturally forms at this interface. As seems from the experimental data, this interface is also kinetically unstable. This kinetic instability (the needle-like growth of TiB) is originated from the anisotropy of the boron diffusion in the lattice of TiB [33]. With the superimposed requirements for plastic accommodation of the
transformation stresses, this results in the growth of the needles in specific crystallographic directions of the Ti matrix (111 in b-Ti) [37]. It may be concluded, therefore, that the TiB2/Ti interface is intrinsically unstable both thermodynamically and kinetically, and the formation of TiB needles will always occur to smaller or larger extent whenever the formation of the monoboride takes place at the TiB2/matrix interface. Accordingly, a TiB2 coating is not an inert coating with respect to the matrix, and, therefore, the C/TiB2 coating represents a double coating, both the inner and the outer layers of which are reactive with respect to the matrix. As discussed in the previous section, the best results may be achieved with a combination of a reactive coating in the inner layer and an inert coating in the outer layer. The situation is additionally complicated by the nonhomogeneous nature of the reaction between TiB2 and Ti matrix (needle formation). There may be several possible ways to reduce or to eliminate the formation of TiB needles at the TiB2 coating/Ti matrix interface. According to the above discussion, the solution may be either of a thermodynamic or kinetic nature, aiming to completely prevent the formation of TiB, or to only eliminate the formation of the needles during TiB growth.
4.1.1. Composition of the matrix A natural way to radically eliminate the formation of TiB needles at TiB2/matrix interface is to use Ti matrices which are thermodynamically stable with respect to TiB2. Titanium aluminides are considered as prospective matrix materials for advanced applications due to their high melting points, good specific strength and modulus retention to high temperatures, resistance to oxidation, and good creep properties [38,39]. According to [32], TiB2 is thermodynamically compatible with all titanium aluminides. It also follows from thermodynamic calculations of equilibrium in the Ti–Al–B system per-
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formed using the TerQuat program [40] (Fig. 5) (though due to incomplete thermodynamic data on this system, this schematic diagram should be considered as tentative). This conclusion is corroborated by the study of the internal interfaces in a TiB2/TiAl composite [41] and by the investigation of arc-melted Ti – Al –B alloys [33]. The thermodynamic compatibility of the TiB2 coating with respect to the g-TiAl matrix was studied using TiB2 layers magnetron sputtered on g-TiAl substrates and model C/TiB2/g-TiAl sandwich-like samples. No new phases was detected by XRD after a heat treatment of the TiB2/TiAl samples at 1100°C for 50 h. At the same time, strong TiB peaks were observed in the diffraction pattern from a similar TiB2/Ti sample already after a heat treatment at 1000°C for 12 h (Fig. 6). This means that with respect to the carbon coating and the TiAl matrix, a TiB2 layer represents a true inert diffusion barrier, and the whole C/TiB2 coating will satisfy the requirements of an ideal doublecoating system. Accordingly, the performance of such a coating system will primarily depend on the barrier properties of TiB2 against the diffusion of the carbon and matrix elements. This topic is addressed in a separate section below. Thus, the first thermodynamic option to overcome the problem of TiB needle formation at the TiB2/matrix interface is to use the advanced matrices based on titanium aluminides, or other matrices which are thermodynamically stable with respect to TiB2. The feasibility of this approach is additionally supported by the investigation of TiB2/C coated SiC/Super-a2 TMC,
Fig. 5. Calculated isothermal section of the Ti–Al–B phase diagram.
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Fig. 6. Stability of the TiB2/TiAl and TiB2/Ti interfaces (XRD).
where no TiB needles were observed at the TiB2/matrix interface [8].
4.1.2. C/TiB double coating As has been mentioned above, TiB2/Ti interface is intrinsically unstable both thermodynamically and kinetically. In Ti matrices which are reactive with respect to TiB2, such as near-a or (a+ b) alloys, the growth of TiB (in particular in its needle-like form) will always occur at this interface. An alternative way to thermodynamically stabilize the coating/matrix interface is to alter the composition of the coating itself. Up to now, the only phase of the Ti–B system that has been considered as an additional protective coating for carbon coated SiC fibers in TMC was titanium diboride, TiB2 [3–12,19,20]. However, the problem of interface stability in non-aluminide Ti matrices can be potentially solved if the titanium monoboride, TiB, is used instead of TiB2 as a protective coating. The obvious advantage of this approach is that no chemical reaction, and, therefore, no needle formation can be expected at the TiB/matrix interface. Though TiB coating can dissolve in the Ti matrix, the kinetics of this dissolution is determined by the rate of boron diffusion and the solubility limit of boron in the matrix. From both points of view the TiB coating seems to be advantageous with respect to other prospective coatings, such as TiC or TiN, since both carbon and nitrogen have significantly higher diffusion rates and solubility in titanium than boron [28–31]. As evident from the experimental data [42] where TiB/Ti composites were investigated, the rate of TiB dissolution in Ti is quite low, and hence the dissolution of a TiB coating may probably be neglected.
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Fig. 7. C/TiB2/TiB triple coating for SiC fibers in Ti matrix.
However, this approach may have its own limitations. A serious problem in the use of a C/TiB double coating may be that, unlike TiB2, TiB is thermodynamically unstable with respect to carbon. As follows from the Ti–C –B phase diagram [32], a TiB coating should react with the underlying carbon coating according to the reaction: 2TiB+C TiB2 +TiC
(3)
Accordingly, the initial C/TiB interface may undergo a chemical reaction, in which both carbon and TiB coatings will be gradually consumed. Therefore, the overall interface process can be controlled either by the kinetics of the reaction (3), or by the diffusion of carbon and/or matrix elements through the products of this reaction, as well as through the TiB coating itself. It must be noted that the lattice of TiB is much more open than the lattice of TiB2 [33]. From this point of view, in such an assembly TiC and TiB2 formed at the C/TiB interface might probably play the main role as a diffusion barrier, while the remaining TiB would protect TiC from the dissolution in, and TiB2 from the reaction with, the matrix. The kinetics of these processes has not been yet studied in detail. The investigation of the C/amorphous TiB coatings in SiC/IMI 834 composites at temperatures 800 – 900°C has revealed that very fast diffusion of carbon and metal atoms through the TiB coating dominates the whole process. Ti atoms arriving at the C/TiB interface immediately react with the carbon coating with formation of TiC, thus effectively preventing the reaction (3) between the carbon and TiB coating [26]. Accordingly, though TiB is not thermodynamically stable with respect to carbon coating, the TiB layer in this situation plays the role of an inert diffusion barrier, but with very poor barrier properties. Preliminary results have indicated that in the case of crystalline TiB the situation is different: the diffusion of metal atoms through the crystalline TiB coating is not as fast as for the amorphous phase, and
the reaction (3) can be observed at the C/TiB interface (A. Werner, P. Mogilevsky, H.J. Dudek, unpublished results). The study of such C/TiB double coatings is now in progress.
4.1.3. C/TiB2 /TiB triple coating A compromise solution for the problem of stability of the TiB2 coating/matrix interface in nonaluminide matrices might be a triple coating with an additional TiB layer between the TiB2 layer and the matrix (Fig. 7). In such a structure with the sequence of titanium borides following the Ti–B phase diagram, an artificially created uniform TiB layer would separate the TiB2 from the matrix. All interfaces of the resulting triple C/TiB2/TiB coating will be in local thermodynamic equilibrium, so it can be expected that the growth of TiB needles may be greatly suppressed, if not prevented. It should be noted, however, that in this case the gradual transformation of TiB2 into TiB would proceed anyway, so that in the end the entire coating will eventually transform into the C/TiB coating system which has just been discussed. The results of our study of the kinetics of the TiB2 transformation are shown in Fig. 8. It gives the time of full transformation (lifetime) of a 10 mm TiB2 coating in pure Ti matrix. These results show that while at 1100°C such coating is transformed into TiB within only : 4.5 h, it will survive for about 4500 h at 800°C. It must be noted, that the reactions with Ti alloys such as Ti-6-4 and others are always much slower than with pure Ti [16], so a reasonable life-time of the coating can be achieved. Additional monoboride layer would also contribute into longer service time of the coating. Applying the same considerations as for deriving Eq. (2), we can write:
t%0 = t0 1+
2h hR
(4)
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where t0 is the time of full transformation of the initial TiB2 layer of a thickness h0 in the Ti matrix; t %0 is the time of full transformation of a similar TiB2 layer in the presence of an initial TiB layer of a thickness h at the Ti/TiB2 interface; and hR is the thickness of the TiB layer which results from the full transformation of the initial TiB2 layer. Taking into account the densities of TiB2 and TiB, rTiB2 =4.38 g cm − 3 and rTiB =5.09 g cm − 3 [31] and the molar weights MTiB2 =69.5 and MTiB = 58.7 g mol − 1, we can calculate hR :1.5h0. Accordingly, from Eq. (4) we can obtain that application at the TiB2/matrix interface of an additional TiB layer of a thickness, e.g. h =1.5h0 would increase the service time of the coating by a factor of three.
4.1.4. Matrix microstructure It has been found that the TiB needles grow in parallel groups in specific crystallographic directions within single grains of the parent Ti matrix [37]. The length of the needles is rather limited by the size of the matrix grains. Therefore, it can be concluded that the growth of TiB needles at the TiB2/Ti matrix interface can be suppressed if a matrix with fine microstructure is used. To check this premise, the growth of TiB at the TiB2/Ti interface was studied in TiB2 – Ti diffusion couples with Ti produced from fine (B 10 mm) high purity Ti powder via the cold sintering technique and compared to the results obtained with cast commercially pure Ti [37]. Rapid grain growth typical for pure Ti [43] is not observed in cold sintered Ti. This is probably due to the very fine oxide films which are present on the surface of Ti powder particles. These oxide films are incorporated into the interparticle boundaries in cold sintered specimens, thus hampering the grain growth [44]. SEM micrographs of the interface for both types of diffusion couples are shown in Fig. 9. They show that
Fig. 9. Reactive formation of TiB at TiB2/Ti interface for cast Ti (a) and cold sintered Ti (b). 1000°C, 4 h (SEI).
while long (up to 100 mm) needles develop in the conventional Ti, no needles are formed in the cold sintered Ti, where the reaction zone consists of a rather uniform TiB layer. These findings confirm that the development of the TiB needles at TiB2/matrix interface can indeed be effectively suppressed through the grain refinement of the matrix microstructure. It should be noted, however, that the fine microstructure of the matrix can lead to low creep resistance of the composite, so the possible compromise between the two conflicting requirements should be examined.
4.2. Interdiffusion through the barrier coating
Fig. 8. Lifetime of a 10 mm TiB2 coating in pure Ti matrix.
Let us now consider the question of interdiffusion of the components of the composite (matrix elements and carbon from the carbon coating) through the additional (TiB2) coating. The barrier properties of a given compound for the diffusion of specific elements are determined by the thickness of the layer, its microstructure (defect density, grain size, texture, etc.), and composition. Since the possible thickness of the coating is limited by mechanical considerations and can hardly
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extend beyond a few micrometers [45], the proper design of the coating microstructure and composition is a key factor to improve the performance of the coating. For example, the interdiffusion coefficient of TiC is known to vary by a factor of : 100 over the homogeneity range [27]. In many cases, however, the microstructure and composition of the TiB2 coatings used in SiC/Ti composites have not been properly defined. In some cases the coating was either non-stoichiometric [9], or graded with partly amorphous microstructure [13]. In general, to fully utilize the potential of a material as a diffusion barrier, crystalline coatings of proper stoichiometry and, preferably, coarse microstructure avoiding fast grain boundary diffusion should be designed. The latter, however, is not always possible to achieve. Thus, in thin-film microelectronic devices amorphous diffusion barriers are often preferable [46]. This is due to the fact that unavoidable fine microstructure of crystalline thin-film materials promotes the fast grain boundary diffusion, and an amorphous material having no grain boundaries will have better barrier properties than a crystalline one. The matter may be different for diffusion barriers used in composite materials. A typical thickness of the barrier layer in these cases can be about a few micrometers, allowing for relatively large grains. Thus, TiB2 layers produced by magnetron sputtering had lateral grain size of :1 – 2 mm [19]. In such a situation crystalline microstructure of a diffusion barrier seems to be advantageous over amorphous. The effect of composition of crystalline TiB2 layers on their barrier properties was studied on model C/ TiB2/TiAl sandwich samples as suggested in [19,20]. g-TiAl was chosen as a model matrix material to avoid the formation of TiB at the matrix/TiB2 interface as suggested in the previous discussion. For a crystalline TiB2 of proper stoichiometry (349 0.5 at.% of Ti), no interdiffusion through the layer and no reactions at both C/TiB2 and TiB2/TiAl interfaces was found after the heat treatment at 1000°C for 50 h, as illustrated by the EDX line-scan (Fig. 10). For B-rich (29.69 0.7 at.% Ti) crystalline TiB2 layers both Ti and Al were found to be diffusing, and a reaction zone of about 5 mm thickness was observed at the C/TiB2 interface after the heat treatment at 1000°C for 10 h (Fig. 11). According to the XRD data, the reaction zone consisted of TiC and Ti2AlC (Fig. 12) which is similar to the reaction between uncoated SCS6 fibers and g-TiAl matrix [47]. Apparently, stoichiometry has a strong effect on the barrier properties of TiB2 layers, at least as far as the diffusion of matrix elements is concerned, and proper stoichiometry has to be ensured when this material is used as a diffusion barrier. It can be added that proper stoichiometry can not only result in better barrier prop-
Fig. 10. Qualitative EDX line scan across a C/TiB2/TiAl sample with stoichiometric TiB2. 1000°C, 50 h.
erties, but can also contribute to better coating/matrix interface stability. It was reported, for example, that harmful TiB needle formation was greatly enhanced in the case of B-rich TiB2 [9].
Fig. 11. SE image (a) and EDX line scan (b) across a C/TiB2/TiAl sample with non-stoichiometric TiB2. 1000°C, 10 h.
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nal lattice of TiB2 consists of alternating layers of B and close packed Ti layers parallel to the (001) plane. It is reasonable to assume, that the diffusion of both Ti and B within the corresponding atomic layers may be faster than the diffusion across the alternating Ti and B layers, and that (001) preferred orientation may, therefore, provide for better barrier properties across the layer. It was reported that TiB2 coating having a (110) preferred orientation could be produced via magnetron sputter deposition depending on sputtering conditions [48]. Therefore, it may be interesting to investigate the influence of preferred orientation on the barrier properties of crystalline TiB2 layers. Fig. 12. XRD of the reaction zone formed at the C/TiB2 interface of a C/TiB2/TiAl sample with nonstoichiometric TiB2. 1000°C, 2 h.
It should be mentioned, that though B-rich TiB2 apparently did not stop the diffusion of Ti and Al, it still could provide a sufficient protection against the diffusion of carbon: in both cases no reaction with carbon was observed at the TiB2/TiAl interface. This may be due to different mechanisms of diffusion of carbon and of metal atoms through the lattice of TiB2. Taking into account very similar atomic radii of ˚ , respectitanium and aluminum (1.45 and 1.43 A tively), both elements can diffuse by vacancy mechanism through the titanium sublattice of TiB2. Taking into account the large difference between atomic radii ˚ , respectively), of carbon and boron (0.77 and 0.91 A diffusion of carbon by vacancy mechanism through the boron sublattice is unlikely and diffusion through the interstitials is apparently the only possible mechanism for carbon diffusion. In more defective structures the diffusion of carbon can be enhanced. In graded amorphous TiBx coatings, for example, all the elements including carbon were found to be diffusing [13]. This is apparently due to the high defect concentration in the nonstoichiometric, amorphous phase. Other microstructural features, which potentially may have an impact on barrier properties of a protective coating, are morphology and crystallographic orientation. In our research on TiB2 coatings we have studied magnetron sputtered TiB2 layers [19,20,37]. It was found that the as-deposited layers showed columnar microstructure and preferred orientation with the c-axis of the hexagonal lattice of TiB2 parallel to the columns’ axis (Fig. 13). Such microstructure, in principle, can promote the fast diffusion along the intercolumnar boundaries oriented parallel to the diffusion direction. On the other hand, (001) preferred orientation associated with this microstructure may be beneficial with regard to lattice diffusion. The hexago-
5. Summary and conclusions Two types of diffusion barriers, i.e. inert coatings and reactive coatings, can be used at composite interfaces depending on the type of damage the material suffers as a result of the interface reactions. While an
Fig. 13. As-deposited TiB2 layer produced via magnetron sputtering deposition: (a) SEI; (b) XRD.
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inert coating may have unlimited service time and is capable of slowing down the formation/growth of the reaction zone, it is unable to completely prevent the chemical attack of the composite reinforcement. In contrast to this, a reactive diffusion barrier has a limited service time and may even accelerate the growth of the reaction zone at the interface, but it is capable to completely prevent (for a certain period of time) the chemical attack of the reinforcement. Accordingly, if the material is highly sensitive to the reinforcement damage (as is the case for, e.g. fiber reinforcements), the application of a reactive diffusion barrier is necessary. For other types of composite materials, e.g. particle-reinforced composites, the use of a proper inert diffusion barrier may be sufficient. A double-coating concept for SiC fiber reinforced TMC was discussed from the point of view of maximum thermodynamic and kinetic stability. Thermodynamic and kinetic aspects of the application of titanium boride coatings as a diffusion barrier for carbon coated SiC fibers in TMC were discussed. It was shown that the TiB2/matrix interface may be unstable both thermodynamically and kinetically. Kinetic instability manifests itself in the growth of TiB needles at the TiB2/matrix interface. Therefore, two factors can critically affect the performance of the coating: the thermodynamic and kinetic instability of the TiB2/matrix interface and the diffusion rate of the matrix components and carbon through the coating. Stabilization of the TiB2/matrix interface can be achieved through: 1. Matrix composition (e.g. unreactive Ti3Al and TiAl matrices). 2. Triple C/TiB2/TiB coatings. 3. Matrix microstructure (grain refinement). The diffusion rate of the components through the coating is determined by the microstructure and the composition of the coating. A crystalline microstructure with relatively large grains and of stoichiometric composition seems to be optimal for diffusion barrier coatings in composite materials. The effect of the composition on the barrier properties of crystalline TiB2 layers was examined. A strong impact of composition on the diffusion of metal atoms through TiB2 was found. At the same time, no diffusion of carbon through both stoichiometric and non-stoichiometric crystalline TiB2 coatings was observed. A conclusion can be made that TiB2 of proper stoichiometry can provide excellent protection for carbon coated (SCS6) SiC fibers in titanium aluminide based MMC. Further investigation is needed to study the feasibility and diffusion barrier properties of triple C/TiB2/TiB and double C/TiB coatings for application in non-aluminide titanium matrices, as well as the effect of the orientation and morphology on the diffusion barrier properties of crystalline TiB2 coatings.
Acknowledgements One of the authors (P.M.) would like to thank the MINERVA Committee for Scientific Cooperation between Germany and Israel for financial support. Cold sintered Ti specimens were supplied by E.Y. Gutmanas of the Technion–Israel Institute of Technology. The authors wish to thank R. Leucht and C. Leyens for magnetron depositions, H. Mangers for heat treatments and W.A. Kaysser for support of this research.
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Application of diffusion barriers in composite materials P. Mogilevsky *, A. Werner, H.J. Dudek Institute of Materials Research, German Aerospace Research Establishment (DLR), D-51140 Cologne, Germany Received 27 December 1996; received in revised form 15 July 1997
Abstract Application of diffusion barriers in composite materials is discussed from the point of view of thermodynamic and kinetic stability of the reinforcement/matrix interface. Two types of interface protective coating, i.e. inert diffusion barriers and reactive diffusion barriers can be distinguished. The application of one coating or the other depends on the type of damage the material suffers as a result of the interface reactions. A double-coating concept for SiC fiber-reinforced titanium matrix composites (TMC) is discussed from the viewpoint of this classification of diffusion barriers. Application of the C/TiB2 coating system for SiC fiber-reinforced TMC is considered in detail, and some approaches to improve the performance of this coating were suggested and examined experimentally with TiB2 coatings produced by magnetron sputtering deposition from a TiB2 target. The effect of the composition and the microstructure on barrier properties of TiB2 layers was examined. A crystalline microstructure with relatively large grains and of stoichiometric composition was found to be optimal for TiB2 protective coatings in SiC fiber-reinforced TMCs. A strong impact of layer stoichiometry on the diffusion of metal atoms through a crystalline TiB2 layer was found. However, stoichiometry did not significantly affect the diffusion of carbon. Crystalline TiB2 of proper stoichiometry can provide excellent protection for carbon coated (SCS6) SiC fibers in titanium aluminide based metal matrix composites (MMC). © 1998 Elsevier Science S.A. All rights reserved. Keywords: Titanium matrix composites; SiC fibers; Interface reactions; Diffusion barriers; Fiber coating
1. Introduction Ceramic-metal composite materials such as metal matrix composites (MMC), in particular titanium matrix composites (TMC), are high-performance materials for many advanced applications, e.g. aerospace high temperature applications. The performance of these materials is, however, often limited by the interface phenomena. The matrix/reinforcement interface can suffer from the coefficient of thermal expansion (CTE) mismatch, mechanical incompatibility, and thermodynamic instability. The last phenomenon is due to the fact that at high temperatures these materials often represent highly non-equilibrium systems, which suffer from spontaneous chemical reactions between the components. In many cases this is the major factor affecting the performance of the composite. Thus, SiC fiber* Corresponding author. Present address: University of Illinois at Urbana-Champaign, Materials Research Laboratory, 104 South Goodwin Avenue, Urbana, IL 61801-2985, USA. Tel.: + 1 217 3332367; fax: + 1 217 2442278. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 1 0 - 8
reinforced TMCs have received a great deal of attention for advanced aerospace applications. At the same time, their use is hindered by chemical reactions at the fiber/ matrix interface during processing (high temperature consolidation) and at service conditions [1–14]. In addition, high temperature post-consolidation heat treatment, which often can potentially improve the mechanical properties of a composite, is prevented by the rapid chemical reaction at the fiber/matrix interface [15,16]. Commercial SiC fibers used for the reinforcement of titanium alloys have a carbon-rich protective coating, which, given the required time of service of about 10 000 h, effectively prevents the reaction with the SiC fibers at temperatures up to 600°C [17]. Carbon coating also eliminates chemical bonding between SiC fibers and the matrix (‘weak interface’) and thus prevents crack propagation from the matrix into the fibers. In addition, carbon coating protects the fibers from surface abrasion [18]. At temperatures \ 600°C, however, the reaction between the carbon coating and Ti matrix becomes too fast to protect the fibers for a reasonable period of time. Further improvement of
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thermal stability of SiC fiber-reinforced TMC may be achieved using advanced coating systems, such as TiB2/ C [3–13,19,20] or TiC/C [14] double coatings. In such coating systems the carbon coating protects the fibers from surface abrasion and crack propagation from the matrix, while the additional coating protects the carbon layer from fast reaction at temperatures \ 600°C. In this paper we analyze the application of different types of protective coating (diffusion barriers) in composite materials and consider the use of a C/TiB2 coating system for SiC/Ti MMC. Application of the phases of the Ti– B system in protective coatings for SiC fibers in TMC as well as a number of ways to improve the performance and thermal stability of these coating systems are discussed. Recent experimental results on the performance of different types of C/Ti–B coatings for SiC fiber-reinforced TMC are reviewed.
2. Background
2.1. Application of diffusion barriers in composite materials The primary goal of any diffusion barrier is to separate two components and to prevent or slow down the interdiffusion and/or possible chemical reaction between them. Let us consider an uncoated reinforcement (F) introduced into a matrix material (M) (Fig. 1(a)). Two different types of diffusion barrier can be distinguished. The first class represents a layer of a material which is chemically inert with respect to both components to be separated (Fig. 1(b)). For example, a TiC protective layer on carbon coated SiC fibers in a Ti matrix is inert with respect for both carbon coating of the fibers and the matrix. Generally speaking, any naturally grown diffusion barrier (i.e. formed at the interface as a result of chemical reaction between the components) falls into this category. This is because the
Fig. 1. Interface reaction: (a) no diffusion barrier; (b) inert diffusion barrier; (c) reactive diffusion barrier. M, matrix; F, reinforcement; B, diffusion barrier; R, reaction zone.
reaction products are usually in thermodynamic equilibrium with the initial materials (uncommon cases of the formation of non-equilibrium phases during interface reactions is omitted here for simplicity). The performance of this type of diffusion barrier, either artificially created or formed during interface reaction, depends on the respective diffusion characteristics of each of the components in the barrier layer. The second class of diffusion barriers represents a coating which can react with one of the components (Fig. 1(c)). Carbon coating on SiC fibers (such as on SCS6 Textron fibers, SiC fibers coated with a carbon protective layer containing small additions of SiC) in TMC is a well-known example of this class of diffusion barriers. The kinetics of the interaction in this case is determined by the effective diffusion rate of the components in the resulting reaction zone. Accordingly, the performance of the diffusion barrier may strongly depend on the composition of the matrix. Thus, applicability of SCS6 fibers in Ti-6-4 alloy is practically limited to a temperature of 600°C. At higher temperatures the interdiffusion in TiC formed between the carbon coating and the matrix becomes too fast [17]. However, a very thin (100 nm) layer of a phase (Ti,Zr)x Siy which forms at the interface between SCS6 fibers and the IMI 834 titanium alloy provides sufficient stabilization of the interface at temperatures up to 700°C [21,22]. The principal difference between the two types of diffusion barrier and the resulting difference in their applications can be understood if we analyze the kinetics of the growth of the reaction zone at the composite interface in both cases. Let us assume that we have a composite consisting of a matrix and an uncoated reinforcement, and our task is to prevent or slow down the chemical attack of the reinforcement. Let us assume, for simplicity, that only matrix components diffuse (Fig. 1). Fig. 2 shows the kinetics of the development of the reaction zone for both types of the diffusion barrier as compared to an uncoated reinforcement. Typically, parabolic diffusion-controlled kinetics are observed for uncoated reinforcements (dashed line, Fig. 2). The same kinetics is most often observed experimentally for the case of inert diffusion barriers, with the important difference that an apparent ‘incubation’ period t%0 for the reaction is observed (dotted line, Fig. 2). Closer examination of the kinetics of the process reveals, however, that the apparent ‘incubation’ period in this case results from the parabolic approximation of the experimental data obtained on the advanced stage of the interaction, when the process is controlled by the interdiffusion through the growing reaction product. In reality, nevertheless, the reaction already takes place during the ‘incubation’ period as well. At this stage, when the thickness of the reaction zone is small, the process is controlled by the interdiffusion of the compo-
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Fig. 2. Kinetics of the reinforcement consumption and the reaction zone growth during interface reaction: A, inert diffusion barrier; B, C, reactive diffusion barrier (reinforcement consumption and overall reaction zone, respectively).
nents through the inert barrier of constant thickness and, therefore, has linear kinetics (A-I, Fig. 2). Later, with the increase of the reaction zone thickness, the control of the whole process is gradually overtaken by the diffusion through the reaction zone and the process finally gets the parabolic kinetics (A-II, Fig. 2). The apparent incubation period t%0 depends on the thickness of the barrier and on the coefficients of interdiffusion in the diffusion barrier and the reaction zone: t%0 :
2D0 rh 2b D0 2b
(1)
where D0 r and D0 b are the effective interdiffusion coefficients of the reaction zone and of the diffusion barrier, respectively, and hb is the diffusion barrier thickness. Due to typically small thickness of the reaction zone, the stage of linear growth is difficult to observe and may be missed in kinetic measurements. If such a thin reaction zone does not harm the properties of the composite, it can be ignored. More often, however, even a minor reaction with the reinforcement may strongly affect the properties of the composite, and the interpretation of the kinetic measurements by a simple parabolic kinetics with an incubation period may be misleading. Let us now consider the kinetics of the interaction in case of a reactive diffusion barrier. In this case a true incubation period t0 may be observed (B-I, Fig. 2). This is the period when the diffusion barrier reacts with the matrix. It was observed experimentally for carbon coated SiC fibers (SCS6) in Ti matrix, that unless all of the carbon coating is consumed, the tensile strength of the fibers (but not of the composite, however) remains unchanged [17]. At the same time it is known that even
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a minor reaction on the surface of SiC fibers results in a catastrophic drop in the fibers strength [23]. Therefore, we can conclude that no chemical attack of the fiber takes place unless all (or almost all) of the carbon coating (a reactive diffusion barrier) is consumed. Intuitively it can be explained by the strong affinity of the diffusing species to the material of the diffusion barrier (as is the condition for the barrier to be reactive). In this situation any atom diffusing through the barrier will be ‘trapped’ by the reaction with it as long as it still has to travel through the remaining barrier to reach the reinforcement. Contrary, in the case of an inert barrier where the diffusion species by definition have no affinity to the barrier material, they can diffuse through the barrier and attack the reinforcement already in the early stage of the process. After the reactive barrier has been consumed, the process will involve the interaction between the matrix and the reinforcement. The initial reactive barrier is usually transformed into the reaction products which are inert with respect to both the reinforcement and the matrix (for example, stoichiometric TiC formed between carbon coating of SCS6 fibers and Ti matrix is unreactive with respect to both SiC and Ti). Accordingly, the kinetics of the following interaction between the matrix and the reinforcement will be very similar to the case with an inert barrier (B-II and B-III, Fig. 2). The advantage of a reactive diffusion barrier can be understood from the above consideration. Unlike the inert one, it provides a true incubation period, during which the reinforcement can be kept completely out of the interaction. This may be critically important for such reinforcements as SiC fibers (apparently as well as other fiber reinforcements): as has been already mentioned, even minor a reaction on their surface results in a drastic decline in the fiber properties [23]. In this case it is simply not enough to slow down the reaction with the fiber, it should be completely prevented. Apparently, the only way to achieve this goal is the application of a reactive diffusion barrier. However, reactive barriers have their own drawbacks. Let us consider now the kinetics of the growth of the overall reaction zone (not to confuse with the consumption of the reinforcement) at the composite interface for both inert and reactive diffusion barriers. In the case of an inert barrier the thickness of the overall reaction zone follows the graph of the reinforcement consumption, since no other process is involved (curve A, Fig. 2). However, in the case of a reactive diffusion barrier the situation is different. During the incubation period, when the reinforcement does not react, the thickness of the overall reaction zone will follow the parabolic kinetics of the reaction between the matrix and the diffusion barrier (C-I, Fig. 2). After the barrier has been consumed, the reaction zone will continue to grow due to the reaction with the reinforce-
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ment, and the overall process will follow the kinetics of the reinforcement consumption (C-II and C-III, Fig. 2). The problem here is that the kinetics of the reaction between the matrix and a reactive diffusion barrier may be even faster than between the matrix and the reinforcement. Thus, from the kinetic measurements [24], the reaction constants k= 4.7 ×10 − 15 m2 s − 1 at 1000°C and k = 1.9 ×10 − 16 m2 s − 1 at 800°C can be evaluated for the interaction between polycrystalline SiC and pure Ti. At the same time for the interaction between pure carbon (graphite) and Ti in identical experimental conditions the respective values of k= 1.3 ×1 − 14 m2 s − 1 and k =1.6 ×l0 − 15 m2 s − 1 were obtained [25]. Accordingly, in this situation the growth of the overall reaction zone with such a diffusion barrier will be faster than without it (Fig. 2). The growth at the composite interface of a reaction zone, often consisting of brittle phases, results in the stress concentration, crack initiation, and the development of the internal transformation stresses, thus promoting the material failure [21]. It has been reported that the tensile strength of SCS6/IMI 834 composite initially remains unchanged during the heat treatment at 800°C, but drops about 50% when the thickness of the reaction zone between carbon coating and the IMI 834 matrix reaches a critical value of about 3 mm [26]. At this stage, the carbon coating of the SCS6 fibers is still not consumed, and the observed property decline can not be attributed to the fiber damage, but to the development of a brittle reaction zone. Therefore, we can summarize the advantages and disadvantages of both types of diffusion barrier as follows. An inert coating may have unlimited service time and is capable of slowing down the growth of the reaction zone, but it is unable to completely prevent the chemical attack of the reinforcement. In contrast to this, a reactive diffusion barrier has a limited service time and may cause an accelerated growth of the reaction zone, but it is capable to completely prevent (for a certain period of time) the chemical attack of the reinforcement. Accordingly, the application of one coating or the other will depend on the type of damage the material suffers as a result of the interface reactions. If the material is highly sensitive to reinforcement damage (as is the case for SiC and, apparently, other types of fiber reinforcements), the application of a reactive diffusion barrier is necessary. For other types of composite materials, e.g. particle-reinforced composites, the use of a proper inert diffusion barrier may suffice.
2.2. Fiber protecti6e coating for SiC fiber-reinforced TMC
the carbon coating as a means of protection of the fiber from surface abrasion, crack propagation, and immediate chemical attack from the matrix should remain as an inner layer of the coating. As has been already mentioned, a Ti-based composite reinforced with carbon coated SiC fibers fails not due to the fiber damage, as in the case of uncoated fibers, but due to the development of the brittle reaction zone at the coating/ matrix interface. Therefore, an ideal coating system for this type of materials would be a double coating consisting of one layer of a reactive coating (carbon) to protect the fiber from an immediate chemical attack, and another layer of an inert coating to slow down the growth of the reaction zone. With this coating system, the advantages of both types of diffusion barrier can be combined. The material of this second diffusion barrier of the coating system must, ideally, have high barrier properties for carbon and the components of the matrix, be thermodynamically stable with respect to both the carbon coating and the matrix, and its components should not dissolve and/or diffuse into the matrix. Taking into account numerous additional criteria which are not in scope of this paper (adhesion, CTE mismatch, mechanical compatibility), it is clear that the task of finding a material fulfilling all aforementioned requirements may be quite difficult or impossible, and that a compromise may have to be made. Different materials have been considered as an additional protective coating for carbon coated SiC fibers in Ti matrices, such as TiC [5,9,14], TiN [11,12], and TiB2 [3–13]. It is, however, hardly probable that the TiC layer can provide the necessary level of the interface stabilization for carbon coated SiC fibers in the Ti matrix. TiC forms naturally during the interface reaction between the carbon coating and the matrix, and the very fact that an additional protection at temperatures \600°C is needed means that at these temperatures the interdiffusion through TiC is too fast. The possible effect of an additional TiC layer on the kinetics of the reaction at the carbon/matrix interface can be estimated as follows. Assume that the growth of TiC at carbon/Ti matrix interface follows normal parabolic kinetics, and that normally the reaction zone of the critical thickness hc develops during the time tc (Fig. 3). If we place an initial TiC layer of a thickness h0 at the C/Ti interface, the reaction will follow the same kinetics, but starting from the point 0% (Fig. 3). Accordingly, we can calculate the time t%c during which a reaction zone of the critical thickness hc will develop on top of the initial TiC:
t%c = tc 1+ In light of the above discussion, the concept of a protective coating for SiC fibers in high temperature ( \600°C) TMCs may be suggested as follows. First,
2h0 hc
(2)
According to [26], the critical thickness of the reaction zone, after which a sharp decline in the mechanical
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properties of SCS6/IMI 834 composite occurs, is hc 3 mm. Assuming the thickness of the initial TiC layer in the range of 3–6 mm (a reasonable value), we obtain from Eq. (2) that such a layer would prolong the service time of the material by a factor of 3 – 5, respectively. This is obviously not enough to compensate the effect of rising temperature. According to [27], the activation energy of interdiffusion in TiC is : 330 920 kJ mol − 1, which means an increase of the coefficient of interdiffusion by a factor of :4900 with the temperature rising from 600 to 800°C (the target service temperature of modern TMC for aerospace applications). Application of TiN as an additional diffusion barrier for carbon coated SiC fibers and its barrier properties has not been studied in detail. It should be pointed out, however, that the relatively high solubility and diffusion rate of nitrogen in Ti [28 – 31] may have a negative effect on the performance of such coating system. Titanium diboride TiB2 has recently received much attention as a potential second coating for carbon coated SiC fibers in TMC applications. The choice of TiB2 is substantiated primarily by the low solubility (0.05 and 0.03 wt.% in a- and b-Ti, respectively, at 88695°C [28]) and the low rate of diffusion (: 9× 10 − 13 m2 s − 1 at 1000°C in b-Ti [29]) of B in Ti alloys. As follows from the Ti– C – B phase diagram [32], TiB2 is apparently thermodynamically stable with respect to carbon. According to the crystallographic data, TiB2 has a closed-packed crystallographic structure consisting of alternating Ti and B atomic layers [33], which may be indicative of good barrier properties of the compound. Indeed, it has been reported that a TiB2 coating is capable of preventing the diffusion of carbon, thus minimizing the consumption of the carbon protective coating [11]. TiB2 was reported to be the most promising choice for the protective coating in SiC/Ti composites, also from the point of view of its effect on
Fig. 3. The effect of an additional TiC coating on the kinetics of the reaction between the carbon coating of the SiC fiber and Ti matrix.
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the fiber strength [9,12]. However, reactive growth of TiB needles at the Ti matrix/TiB2 interface has been reported [3–9]. The formation of these needles enhances stress concentration and interface failure and is the major drawback of TiB2 as a protective coating for SiC fiber-reinforced TMC. Needle-like growth seems to be a general feature of the diffusion growth of TiB. It was previously observed in arc-melted Ti–Al–B alloys [33], during B4C–Ti interaction [34], in binary Ti–B planar diffusion couples, as well as at a Ti/B-fiber interface [35], at a Ti/TiB2-particle interface [36], and in planar Ti–TiB2 diffusion couples [37]. Thus, the interface phenomenon in a C/TiB2 coated SiC fiber-reinforced TMC is very complex and may involve multi-component diffusion in combination with thermodynamic and kinetic instabilities of the TiB2/matrix interface. It should be noted here, that from this point of view TiB2 is, in effect, a reactive diffusion barrier. Therefore, C/TiB2 coating system does not fulfill the requirements of an ideal coating system which, as discussed in the previous section, should consist of a combination of a reactive and an inert diffusion barrier. Let us now consider the physicochemical processes that can potentially take place at the interface between a C/TiB2 coated SiC fiber and a Ti matrix. As was already discussed, as long as the carbon coating has not been consumed, the fiber components are not involved in the interface processes. Taking also into account that the C/TiB2 interface is thermodynamically stable, we may consider the following processes at the interface between a C/TiB2 coated SiC fiber and a Ti matrix (Fig. 4): 1. Reaction of the TiB2 coating with the Ti matrix with formation of TiB. 2. Diffusion of carbon through the TiB2 layer and its reaction with the matrix. 3. Diffusion of the matrix’ elements through the TiB2 layer and their reaction with the carbon coating. All three processes lead to the degradation of the interface and deterioration of the properties of the entire composite. The first process reflects the thermodynamic and kinetic instability of the TiB2/Ti matrix interface, which leads to the transformation of the initially homogeneous TiB2 layer into the needle-like TiB. The second and the third processes are concerned with the proper barrier properties of the TiB2 coating and lead to the consumption of the carbon protective layer and to the formation of titanium and/or ternary carbides at both carbon/TiB2 and TiB2/matrix or TiB/ matrix interfaces. Accordingly, the problem with the C/TiB2 coatings is two-fold: 1. The development of TiB needles at TiB2/matrix interface should be eliminated. 2. High barrier properties of TiB2 coating against possible diffusion of carbon and/or matrix elements have to be ensured.
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Fig. 4. Physicochemical processes at the interface between C/TiB2 coated SiC fibers and Ti matrix.
Below we suggest a number of approaches to meet these requirements and to improve the performance of the C/TiB2 coating system.
3. Experimental Samples for investigation of the interaction at TiB2/ TiAl and TiB2/Ti interfaces were prepared by magnetron sputtering deposition of TiB2 on TiAl and on Ti substrates. To model the interdiffusion at the interface between C/TiB2 coated SiC fibers and Ti-matrix C/ TiB2/TiAl ‘sandwich’ samples were prepared by magnetron sputtering deposition of TiB2 layers on graphite substrates, followed by a magnetron deposition of :30 mm thick g-TiAl layers. Heat treatments were performed in a vacuum of 10 − 4 Pa. Microstructure and composition of the as-deposited materials and annealed samples were studied by XRD, SEM, and EDX at 10 keV using pure Ti and Al standards. Boron content was evaluated by difference. Polished cross-sectional samples cut perpendicular to the interface were used for SEM and EDX examinations. For SEM examinations some of the samples were etched in a mixture of HF and HNO3.
4. Results and discussion
4.1. TiB needle formation As follows from the Ti – B phase diagram [30], the TiB2/Ti interface is thermodynamically unstable and TiB naturally forms at this interface. As seems from the experimental data, this interface is also kinetically unstable. This kinetic instability (the needle-like growth of TiB) is originated from the anisotropy of the boron diffusion in the lattice of TiB [33]. With the superimposed requirements for plastic accommodation of the
transformation stresses, this results in the growth of the needles in specific crystallographic directions of the Ti matrix (111 in b-Ti) [37]. It may be concluded, therefore, that the TiB2/Ti interface is intrinsically unstable both thermodynamically and kinetically, and the formation of TiB needles will always occur to smaller or larger extent whenever the formation of the monoboride takes place at the TiB2/matrix interface. Accordingly, a TiB2 coating is not an inert coating with respect to the matrix, and, therefore, the C/TiB2 coating represents a double coating, both the inner and the outer layers of which are reactive with respect to the matrix. As discussed in the previous section, the best results may be achieved with a combination of a reactive coating in the inner layer and an inert coating in the outer layer. The situation is additionally complicated by the nonhomogeneous nature of the reaction between TiB2 and Ti matrix (needle formation). There may be several possible ways to reduce or to eliminate the formation of TiB needles at the TiB2 coating/Ti matrix interface. According to the above discussion, the solution may be either of a thermodynamic or kinetic nature, aiming to completely prevent the formation of TiB, or to only eliminate the formation of the needles during TiB growth.
4.1.1. Composition of the matrix A natural way to radically eliminate the formation of TiB needles at TiB2/matrix interface is to use Ti matrices which are thermodynamically stable with respect to TiB2. Titanium aluminides are considered as prospective matrix materials for advanced applications due to their high melting points, good specific strength and modulus retention to high temperatures, resistance to oxidation, and good creep properties [38,39]. According to [32], TiB2 is thermodynamically compatible with all titanium aluminides. It also follows from thermodynamic calculations of equilibrium in the Ti–Al–B system per-
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formed using the TerQuat program [40] (Fig. 5) (though due to incomplete thermodynamic data on this system, this schematic diagram should be considered as tentative). This conclusion is corroborated by the study of the internal interfaces in a TiB2/TiAl composite [41] and by the investigation of arc-melted Ti – Al –B alloys [33]. The thermodynamic compatibility of the TiB2 coating with respect to the g-TiAl matrix was studied using TiB2 layers magnetron sputtered on g-TiAl substrates and model C/TiB2/g-TiAl sandwich-like samples. No new phases was detected by XRD after a heat treatment of the TiB2/TiAl samples at 1100°C for 50 h. At the same time, strong TiB peaks were observed in the diffraction pattern from a similar TiB2/Ti sample already after a heat treatment at 1000°C for 12 h (Fig. 6). This means that with respect to the carbon coating and the TiAl matrix, a TiB2 layer represents a true inert diffusion barrier, and the whole C/TiB2 coating will satisfy the requirements of an ideal doublecoating system. Accordingly, the performance of such a coating system will primarily depend on the barrier properties of TiB2 against the diffusion of the carbon and matrix elements. This topic is addressed in a separate section below. Thus, the first thermodynamic option to overcome the problem of TiB needle formation at the TiB2/matrix interface is to use the advanced matrices based on titanium aluminides, or other matrices which are thermodynamically stable with respect to TiB2. The feasibility of this approach is additionally supported by the investigation of TiB2/C coated SiC/Super-a2 TMC,
Fig. 5. Calculated isothermal section of the Ti–Al–B phase diagram.
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Fig. 6. Stability of the TiB2/TiAl and TiB2/Ti interfaces (XRD).
where no TiB needles were observed at the TiB2/matrix interface [8].
4.1.2. C/TiB double coating As has been mentioned above, TiB2/Ti interface is intrinsically unstable both thermodynamically and kinetically. In Ti matrices which are reactive with respect to TiB2, such as near-a or (a+ b) alloys, the growth of TiB (in particular in its needle-like form) will always occur at this interface. An alternative way to thermodynamically stabilize the coating/matrix interface is to alter the composition of the coating itself. Up to now, the only phase of the Ti–B system that has been considered as an additional protective coating for carbon coated SiC fibers in TMC was titanium diboride, TiB2 [3–12,19,20]. However, the problem of interface stability in non-aluminide Ti matrices can be potentially solved if the titanium monoboride, TiB, is used instead of TiB2 as a protective coating. The obvious advantage of this approach is that no chemical reaction, and, therefore, no needle formation can be expected at the TiB/matrix interface. Though TiB coating can dissolve in the Ti matrix, the kinetics of this dissolution is determined by the rate of boron diffusion and the solubility limit of boron in the matrix. From both points of view the TiB coating seems to be advantageous with respect to other prospective coatings, such as TiC or TiN, since both carbon and nitrogen have significantly higher diffusion rates and solubility in titanium than boron [28–31]. As evident from the experimental data [42] where TiB/Ti composites were investigated, the rate of TiB dissolution in Ti is quite low, and hence the dissolution of a TiB coating may probably be neglected.
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Fig. 7. C/TiB2/TiB triple coating for SiC fibers in Ti matrix.
However, this approach may have its own limitations. A serious problem in the use of a C/TiB double coating may be that, unlike TiB2, TiB is thermodynamically unstable with respect to carbon. As follows from the Ti–C –B phase diagram [32], a TiB coating should react with the underlying carbon coating according to the reaction: 2TiB+C TiB2 +TiC
(3)
Accordingly, the initial C/TiB interface may undergo a chemical reaction, in which both carbon and TiB coatings will be gradually consumed. Therefore, the overall interface process can be controlled either by the kinetics of the reaction (3), or by the diffusion of carbon and/or matrix elements through the products of this reaction, as well as through the TiB coating itself. It must be noted that the lattice of TiB is much more open than the lattice of TiB2 [33]. From this point of view, in such an assembly TiC and TiB2 formed at the C/TiB interface might probably play the main role as a diffusion barrier, while the remaining TiB would protect TiC from the dissolution in, and TiB2 from the reaction with, the matrix. The kinetics of these processes has not been yet studied in detail. The investigation of the C/amorphous TiB coatings in SiC/IMI 834 composites at temperatures 800 – 900°C has revealed that very fast diffusion of carbon and metal atoms through the TiB coating dominates the whole process. Ti atoms arriving at the C/TiB interface immediately react with the carbon coating with formation of TiC, thus effectively preventing the reaction (3) between the carbon and TiB coating [26]. Accordingly, though TiB is not thermodynamically stable with respect to carbon coating, the TiB layer in this situation plays the role of an inert diffusion barrier, but with very poor barrier properties. Preliminary results have indicated that in the case of crystalline TiB the situation is different: the diffusion of metal atoms through the crystalline TiB coating is not as fast as for the amorphous phase, and
the reaction (3) can be observed at the C/TiB interface (A. Werner, P. Mogilevsky, H.J. Dudek, unpublished results). The study of such C/TiB double coatings is now in progress.
4.1.3. C/TiB2 /TiB triple coating A compromise solution for the problem of stability of the TiB2 coating/matrix interface in nonaluminide matrices might be a triple coating with an additional TiB layer between the TiB2 layer and the matrix (Fig. 7). In such a structure with the sequence of titanium borides following the Ti–B phase diagram, an artificially created uniform TiB layer would separate the TiB2 from the matrix. All interfaces of the resulting triple C/TiB2/TiB coating will be in local thermodynamic equilibrium, so it can be expected that the growth of TiB needles may be greatly suppressed, if not prevented. It should be noted, however, that in this case the gradual transformation of TiB2 into TiB would proceed anyway, so that in the end the entire coating will eventually transform into the C/TiB coating system which has just been discussed. The results of our study of the kinetics of the TiB2 transformation are shown in Fig. 8. It gives the time of full transformation (lifetime) of a 10 mm TiB2 coating in pure Ti matrix. These results show that while at 1100°C such coating is transformed into TiB within only : 4.5 h, it will survive for about 4500 h at 800°C. It must be noted, that the reactions with Ti alloys such as Ti-6-4 and others are always much slower than with pure Ti [16], so a reasonable life-time of the coating can be achieved. Additional monoboride layer would also contribute into longer service time of the coating. Applying the same considerations as for deriving Eq. (2), we can write:
t%0 = t0 1+
2h hR
(4)
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where t0 is the time of full transformation of the initial TiB2 layer of a thickness h0 in the Ti matrix; t %0 is the time of full transformation of a similar TiB2 layer in the presence of an initial TiB layer of a thickness h at the Ti/TiB2 interface; and hR is the thickness of the TiB layer which results from the full transformation of the initial TiB2 layer. Taking into account the densities of TiB2 and TiB, rTiB2 =4.38 g cm − 3 and rTiB =5.09 g cm − 3 [31] and the molar weights MTiB2 =69.5 and MTiB = 58.7 g mol − 1, we can calculate hR :1.5h0. Accordingly, from Eq. (4) we can obtain that application at the TiB2/matrix interface of an additional TiB layer of a thickness, e.g. h =1.5h0 would increase the service time of the coating by a factor of three.
4.1.4. Matrix microstructure It has been found that the TiB needles grow in parallel groups in specific crystallographic directions within single grains of the parent Ti matrix [37]. The length of the needles is rather limited by the size of the matrix grains. Therefore, it can be concluded that the growth of TiB needles at the TiB2/Ti matrix interface can be suppressed if a matrix with fine microstructure is used. To check this premise, the growth of TiB at the TiB2/Ti interface was studied in TiB2 – Ti diffusion couples with Ti produced from fine (B 10 mm) high purity Ti powder via the cold sintering technique and compared to the results obtained with cast commercially pure Ti [37]. Rapid grain growth typical for pure Ti [43] is not observed in cold sintered Ti. This is probably due to the very fine oxide films which are present on the surface of Ti powder particles. These oxide films are incorporated into the interparticle boundaries in cold sintered specimens, thus hampering the grain growth [44]. SEM micrographs of the interface for both types of diffusion couples are shown in Fig. 9. They show that
Fig. 9. Reactive formation of TiB at TiB2/Ti interface for cast Ti (a) and cold sintered Ti (b). 1000°C, 4 h (SEI).
while long (up to 100 mm) needles develop in the conventional Ti, no needles are formed in the cold sintered Ti, where the reaction zone consists of a rather uniform TiB layer. These findings confirm that the development of the TiB needles at TiB2/matrix interface can indeed be effectively suppressed through the grain refinement of the matrix microstructure. It should be noted, however, that the fine microstructure of the matrix can lead to low creep resistance of the composite, so the possible compromise between the two conflicting requirements should be examined.
4.2. Interdiffusion through the barrier coating
Fig. 8. Lifetime of a 10 mm TiB2 coating in pure Ti matrix.
Let us now consider the question of interdiffusion of the components of the composite (matrix elements and carbon from the carbon coating) through the additional (TiB2) coating. The barrier properties of a given compound for the diffusion of specific elements are determined by the thickness of the layer, its microstructure (defect density, grain size, texture, etc.), and composition. Since the possible thickness of the coating is limited by mechanical considerations and can hardly
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extend beyond a few micrometers [45], the proper design of the coating microstructure and composition is a key factor to improve the performance of the coating. For example, the interdiffusion coefficient of TiC is known to vary by a factor of : 100 over the homogeneity range [27]. In many cases, however, the microstructure and composition of the TiB2 coatings used in SiC/Ti composites have not been properly defined. In some cases the coating was either non-stoichiometric [9], or graded with partly amorphous microstructure [13]. In general, to fully utilize the potential of a material as a diffusion barrier, crystalline coatings of proper stoichiometry and, preferably, coarse microstructure avoiding fast grain boundary diffusion should be designed. The latter, however, is not always possible to achieve. Thus, in thin-film microelectronic devices amorphous diffusion barriers are often preferable [46]. This is due to the fact that unavoidable fine microstructure of crystalline thin-film materials promotes the fast grain boundary diffusion, and an amorphous material having no grain boundaries will have better barrier properties than a crystalline one. The matter may be different for diffusion barriers used in composite materials. A typical thickness of the barrier layer in these cases can be about a few micrometers, allowing for relatively large grains. Thus, TiB2 layers produced by magnetron sputtering had lateral grain size of :1 – 2 mm [19]. In such a situation crystalline microstructure of a diffusion barrier seems to be advantageous over amorphous. The effect of composition of crystalline TiB2 layers on their barrier properties was studied on model C/ TiB2/TiAl sandwich samples as suggested in [19,20]. g-TiAl was chosen as a model matrix material to avoid the formation of TiB at the matrix/TiB2 interface as suggested in the previous discussion. For a crystalline TiB2 of proper stoichiometry (349 0.5 at.% of Ti), no interdiffusion through the layer and no reactions at both C/TiB2 and TiB2/TiAl interfaces was found after the heat treatment at 1000°C for 50 h, as illustrated by the EDX line-scan (Fig. 10). For B-rich (29.69 0.7 at.% Ti) crystalline TiB2 layers both Ti and Al were found to be diffusing, and a reaction zone of about 5 mm thickness was observed at the C/TiB2 interface after the heat treatment at 1000°C for 10 h (Fig. 11). According to the XRD data, the reaction zone consisted of TiC and Ti2AlC (Fig. 12) which is similar to the reaction between uncoated SCS6 fibers and g-TiAl matrix [47]. Apparently, stoichiometry has a strong effect on the barrier properties of TiB2 layers, at least as far as the diffusion of matrix elements is concerned, and proper stoichiometry has to be ensured when this material is used as a diffusion barrier. It can be added that proper stoichiometry can not only result in better barrier prop-
Fig. 10. Qualitative EDX line scan across a C/TiB2/TiAl sample with stoichiometric TiB2. 1000°C, 50 h.
erties, but can also contribute to better coating/matrix interface stability. It was reported, for example, that harmful TiB needle formation was greatly enhanced in the case of B-rich TiB2 [9].
Fig. 11. SE image (a) and EDX line scan (b) across a C/TiB2/TiAl sample with non-stoichiometric TiB2. 1000°C, 10 h.
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nal lattice of TiB2 consists of alternating layers of B and close packed Ti layers parallel to the (001) plane. It is reasonable to assume, that the diffusion of both Ti and B within the corresponding atomic layers may be faster than the diffusion across the alternating Ti and B layers, and that (001) preferred orientation may, therefore, provide for better barrier properties across the layer. It was reported that TiB2 coating having a (110) preferred orientation could be produced via magnetron sputter deposition depending on sputtering conditions [48]. Therefore, it may be interesting to investigate the influence of preferred orientation on the barrier properties of crystalline TiB2 layers. Fig. 12. XRD of the reaction zone formed at the C/TiB2 interface of a C/TiB2/TiAl sample with nonstoichiometric TiB2. 1000°C, 2 h.
It should be mentioned, that though B-rich TiB2 apparently did not stop the diffusion of Ti and Al, it still could provide a sufficient protection against the diffusion of carbon: in both cases no reaction with carbon was observed at the TiB2/TiAl interface. This may be due to different mechanisms of diffusion of carbon and of metal atoms through the lattice of TiB2. Taking into account very similar atomic radii of ˚ , respectitanium and aluminum (1.45 and 1.43 A tively), both elements can diffuse by vacancy mechanism through the titanium sublattice of TiB2. Taking into account the large difference between atomic radii ˚ , respectively), of carbon and boron (0.77 and 0.91 A diffusion of carbon by vacancy mechanism through the boron sublattice is unlikely and diffusion through the interstitials is apparently the only possible mechanism for carbon diffusion. In more defective structures the diffusion of carbon can be enhanced. In graded amorphous TiBx coatings, for example, all the elements including carbon were found to be diffusing [13]. This is apparently due to the high defect concentration in the nonstoichiometric, amorphous phase. Other microstructural features, which potentially may have an impact on barrier properties of a protective coating, are morphology and crystallographic orientation. In our research on TiB2 coatings we have studied magnetron sputtered TiB2 layers [19,20,37]. It was found that the as-deposited layers showed columnar microstructure and preferred orientation with the c-axis of the hexagonal lattice of TiB2 parallel to the columns’ axis (Fig. 13). Such microstructure, in principle, can promote the fast diffusion along the intercolumnar boundaries oriented parallel to the diffusion direction. On the other hand, (001) preferred orientation associated with this microstructure may be beneficial with regard to lattice diffusion. The hexago-
5. Summary and conclusions Two types of diffusion barriers, i.e. inert coatings and reactive coatings, can be used at composite interfaces depending on the type of damage the material suffers as a result of the interface reactions. While an
Fig. 13. As-deposited TiB2 layer produced via magnetron sputtering deposition: (a) SEI; (b) XRD.
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inert coating may have unlimited service time and is capable of slowing down the formation/growth of the reaction zone, it is unable to completely prevent the chemical attack of the composite reinforcement. In contrast to this, a reactive diffusion barrier has a limited service time and may even accelerate the growth of the reaction zone at the interface, but it is capable to completely prevent (for a certain period of time) the chemical attack of the reinforcement. Accordingly, if the material is highly sensitive to the reinforcement damage (as is the case for, e.g. fiber reinforcements), the application of a reactive diffusion barrier is necessary. For other types of composite materials, e.g. particle-reinforced composites, the use of a proper inert diffusion barrier may be sufficient. A double-coating concept for SiC fiber reinforced TMC was discussed from the point of view of maximum thermodynamic and kinetic stability. Thermodynamic and kinetic aspects of the application of titanium boride coatings as a diffusion barrier for carbon coated SiC fibers in TMC were discussed. It was shown that the TiB2/matrix interface may be unstable both thermodynamically and kinetically. Kinetic instability manifests itself in the growth of TiB needles at the TiB2/matrix interface. Therefore, two factors can critically affect the performance of the coating: the thermodynamic and kinetic instability of the TiB2/matrix interface and the diffusion rate of the matrix components and carbon through the coating. Stabilization of the TiB2/matrix interface can be achieved through: 1. Matrix composition (e.g. unreactive Ti3Al and TiAl matrices). 2. Triple C/TiB2/TiB coatings. 3. Matrix microstructure (grain refinement). The diffusion rate of the components through the coating is determined by the microstructure and the composition of the coating. A crystalline microstructure with relatively large grains and of stoichiometric composition seems to be optimal for diffusion barrier coatings in composite materials. The effect of the composition on the barrier properties of crystalline TiB2 layers was examined. A strong impact of composition on the diffusion of metal atoms through TiB2 was found. At the same time, no diffusion of carbon through both stoichiometric and non-stoichiometric crystalline TiB2 coatings was observed. A conclusion can be made that TiB2 of proper stoichiometry can provide excellent protection for carbon coated (SCS6) SiC fibers in titanium aluminide based MMC. Further investigation is needed to study the feasibility and diffusion barrier properties of triple C/TiB2/TiB and double C/TiB coatings for application in non-aluminide titanium matrices, as well as the effect of the orientation and morphology on the diffusion barrier properties of crystalline TiB2 coatings.
Acknowledgements One of the authors (P.M.) would like to thank the MINERVA Committee for Scientific Cooperation between Germany and Israel for financial support. Cold sintered Ti specimens were supplied by E.Y. Gutmanas of the Technion–Israel Institute of Technology. The authors wish to thank R. Leucht and C. Leyens for magnetron depositions, H. Mangers for heat treatments and W.A. Kaysser for support of this research.
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