Chemical Vapor Infiltration of the Hexagonal Boron Nitride Interphase for SiC Fiber-Reinforced HP-Si3N4 Matrix Composite
2018 Volume 59 Issue 8 Pages 1310-1316
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2018 Volume 59 Issue 8 Pages 1310-1316
This work describes the processing of SiC fiber-reinforced Si3N4 matrix composites with boron nitride (BN) interphase. The BN interphase was processed by chemical vapor infiltration (CVI) with BF3/NH3 gaseous precursors. The BN interphase modification involved the continuous treatment of Hi-Nicalon SiC fibers. The relations between (i) the processing parameters, (ii) the mechanisms controlling the kinetics of the CVI of the BN, and (iii) the structure of the deposited BN are presented. A single- or multi-layer BN interphase can be produced depending on the CVI conditions imposed on the fibers during the continuous process. A surface reaction mechanism controlling the CVI promotes a smooth, isotropic BN coating. An anisotropic BN coating can be produced when the CVI kinetics are controlled by a mass transport mechanism. With a controlled temperature gradient, the BN interphase is then made by stacking successive isotropic and anisotropic layers.
Ceramics are attractive for applications requiring a low-weight temperature- and corrosion-resistant material. However, the low toughness and poor reliability of monolithic ceramics strongly limit their applications. Ceramic fiber-reinforced ceramic matrix composites (CMCs) have been developed to overcome the brittle features of monolithic ceramics, which make them excellent potential candidates for high-temperature structural applications.1,2) The improved toughness of CMCs is attributed to the incorporation of fibers inside the brittle matrix and the presence of an interphase located between the fibers and the matrix. The interphase allows the optimization of the load transfer conditions between the fibers and the matrix. During thermomechanical loading, the interphase controls the balance between the two major damage mechanisms: matrix cracking and fiber–matrix debonding. These damage mechanisms are responsible for the non-brittle failure behavior of CMCs and are comparable to dislocations in metal materials. When the damage is initiated in the composite, an optimized interphase allows multiple matrix cracks and partial fiber–matrix debonding.
Pyrolytic carbon is one of the most efficient interphase materials used in CMCs because of the anisotropy of its structure and properties. When the graphite planes are oriented parallel to the fiber surface, this interphase allows cracks that propagate inside the matrix to be deflected near the surface of the fibers.3,4) In this condition, the interphase acts as a “mechanical fuse” because it allows the cracks to propagate through the fibers, which delays final composite rupture. The initiation and extension of the damage is a critical step in the mechanical behavior of CMCs. Indeed, an interphase that is too-strong induces brittle rupture of the composite, whereas a too-weak interphase promotes fiber overloading and their premature failure. Several interphase parameters can affect the load transfer conditions between the fibers and the matrix, for example, the thickness and the degree of the anisotropy of the carbon interphase.5)
Although pyrolytic carbon constitutes the most interesting interphase material in terms of the mechanical properties of CMCs, this is not the case in terms of oxidation resistance. Indeed, carbon oxidizes at temperatures as low as 500°C. When the composite is damaged, oxygen diffuses through the cracks of the matrix and oxidizes the carbon interphase, which results in strong degradation of the load transfer conditions between the fiber and the matrix as well as premature rupturing of the composite.6)
A large number of studies on boron nitride (BN) interphases have been carried out with the aim of increasing the oxidation resistance of damaged CMCs. Indeed, BN is characterized by a graphite-like structure, which is an advantage considering the anisotropy of its mechanical properties, accompanied by a higher oxidation resistance than that of carbon. The oxidation of BN starts at around 800°C. Contrary to carbon, the oxidation of BN is passive (i.e., it gives a condensed oxide product) below 1100°C at atmospheric pressure. The latter property can be used to promote the sealing of the matrix and the interphase cracks by the liquid oxide produced in the form of B2O3. However, for temperatures higher than 1100°C, active oxidation (giving gaseous oxide products) of BN occurs, as in the case of carbon oxidation. Under these conditions, a sealing process can be achieved only when the boron oxide products are combined with other oxides, such as silica.
Much work has been done in the field of processing and characterization of BN interphase for CMCs.7–13) Isothermal/isobaric chemical vapor deposition and infiltration based on BCl3/NH3 and BF3/NH3 systems have been studied for the processing of turbostratic BN. Interesting results have been obtained with both systems, especially concerning the relationship between the kinetics of chemical vapor infiltration (CVI) and the processing parameters.
Prouhet7) showed that the temperature and gas pressure domains related to the chemical and physical mechanisms controlled the deposition rate of BN in the BF3/NH3 system. Hexagonal CVD/CVI BN has been produced by Naslain9) and Leparoux8) for the BF3/NH3 and BCI3/NH3 systems, respectively. These studies concerned the static treatments of fibrous preforms in order to produce a BN interphase for CMCs.
The present study investigates the continuous treatments of ceramic fibers. The fibers are first coated with a BN layer before incorporation into a Si3N4 ceramic matrix produced by a hot-pressing route. The aim of this work is to determine appropriate processing conditions for the BN interphase and the SiC/Si3N4 composites. The relationships between the processing parameters, the mechanisms controlling the kinetics of the BN CVI, and the morphology of the boron nitride layer are here presented for the first time and discussed. Then, the hot-pressing processing conditions for the Si3N4 matrix were also investigated. Finally, some properties of SiC/Si3N4 composites with a BN interphase are presented.
The ceramic fibers used in the present work were Hi-Nicalon*SiC fibers (from Nippon Carbon to NGS Advanced Fibers Co. Ltd., Japan). The BN interphase was processed by CVI. The CVI reactor is a hot wall furnace that heats a graphite susceptor placed inside a silica tube by induction. Hi-Nicalon*SiC fibers are wound on a spool and then move through the graphite susceptor up to a second spool. The experimental conditions controlled were temperature, gas pressure, gas flow rate, and rate of fiber displacement. Control of the fiber displacement allows determination of the duration of the CVI treatments and, consequently, the thickness of the deposits. BF3/NH3 was the gas source for BN and argon was used as the gas vector to dilute the reactive species because of the corrosive conditions of the system. The processing parameters analyzed in the present work are temperature, total pressure, and gas phase composition. The composition of the reactive species was determined by the following parameters: A = QNH3/QBF3 and B = QAr/(QBF3 + QNH3), where Q is the gas flow rate measured in sccm (standard cubic centimeters). Then, the Si3N4 matrix was processed in a second reactor using CH3SiCl3/H2 precursor gases at 1100°C. The ratio QH2/CH3SiCl3 was 0.75, the total gas flow rate was 70 cm3 min−1, and the pressure was 5 kPa. The duration of matrix infiltration was 5 h. Finally, some properties of SiC/Si3N4 composites with a BN interphase are presented.
2.1 CharacterizationThe morphology of the coatings was analyzed by scanning electron microscopy (SEM; Hitachi S800 equipment). The cross-section of the composite was studied by transmission electron microscopy (TEM; Topcon 002B, Japan) using bright field, high resolution, and selected area diffraction. The structure and the degree of anisotropy of BN were studied by Raman spectrometry and X-Ray diffraction, which are complementary techniques. X-ray diffraction gives information about the size of the coherent domains (denoted as Lc and La) present in the BN structure by measuring the full width at half maximum (FWHM) of the peaks observed at 2θ = 26.7° for Lc (corresponding to 002 reflections) and 2θ = 41.6° for La (corresponding to 100 reflections) and using the Scherrer equation: L = (kλ)/(FWHM cos θ). Raman spectrometry allows the highly accurate characterization of the crystallite size in the basal plane (La) by measuring the FWHM of the 1367 cm−1 line, which corresponds to a high-frequency active mode of hexagonal BN. Indeed, the FWHM of this peak is proportional to the inverse of La. This relation has been already determined for BN by Nemanich et al.14)
The influence of temperature and pressure on the structure of BN have been analyzed by means of successive CVI experiments in isothermal and isobaric conditions and without fiber displacement (in static conditions) in order to avoid temperature gradients during deposition. Transitions in the BN structure were studied by means of experiments using temperature or pressure gradients. The structure and degree of anisotropy of BN are presented in Fig. 1. The experimental conditions used for the CVI of BN are summarized in Table 1.
Structure with anisotropy of BN interphase.
These experiments were performed in order to obtain information about the effects of each parameter on the structure of the deposits. As the effects of the experimental parameters on the kinetics of the CVI are already known,7) it was interesting to determine their relationship with the structure of the BN obtained in the present work.
3.1 Kinetics–structural relationships (a) Influence of total gas pressureBN was chemical vapor deposited on SiC fibers under static conditions at 1250°C in the pressure range of 0–60 kPa (see conditions 1 to 6 in Table 1). After an hour of deposition, the deposits were analyzed. Under the present CVI conditions, the anisotropy of the BN increased with the gas pressure.
A CVI of BN was performed with a pressure gradient (see condition 7 in Table 1). During this experiment, the pressure was continuously changed in the range of 0–20 kPa. The morphology of the deposit observed by SEM is shown in Fig. 2. As can be seen in the figure, a transition in the structure of BN was evident. At low pressure, the BN deposited seemed to be smooth and isotropic whereas at high pressure strong anisotropy was observed. Higher magnification of this zone evidenced a lamellar structure, which corresponded to the turbostratic form of BN.
Effects of pressures of 10 kPa (a) and 20 kPa (b) on the texture of the deposit at T = 1000°C.
BN was deposited in the temperature range of 1000–1300°C (see conditions 8 to 10 in Table 1). When BN was deposited under dynamic conditions (i.e., with the fiber moving through the reactor), a temperature gradient occurred through the graphite susceptor. BN was deposited, considering the temperature gradient, at a total gas pressure of 20 kPa (conditions 11 in Table 1). The cross-section of the deposit as observed by SEM (Fig. 3) and three parts could be seen: (1) a thin region located near the fiber that was isotropic; (2) a large zone where the BN was turbostratic; and (3) a large region outside of the deposit corresponding to isotropic BN again. Such a morphology of the deposit was evidence of two transitions in the BN structure related to the temperature gradient. At low temperatures, the BN deposited in the present condition was isotropic and weakly organized. At high temperatures, the turbostratic form of BN was observed. Finally, when the temperature decreased again, the isotropic form was obtained.
Gradient of texture in different temperature domains. SEM image with a succession of layers associated to granular isotropic, anisotropic, and smooth isotropic morphologies.
The effect of the composition of the gas mixture on the BN structure was studied in a range of α compositions of 0.5 to 4 at a temperature of 1250°C (conditions 12 to 15 in Table 1). The deposit was analyzed by X-ray diffraction; their La and Lc determinations are presented in Fig. 4 and their morphologies are shown in Fig. 2. Figure 4 shows that, for the considered temperature and gas pressure, the size La of the crystallites strongly decreases when the ratio a increases from 0.5 to 2, which means when ammonia became preponderant in the gaseous precursor. No significant change was observed when a is higher than 2. Moreover, the size Lc of the crystallites remained unchanged in the considered range of a. The SEM analyses performed on two cross sections of BN deposits at α = 0.5 and α = 4 (Fig. 5) showed a change in the structure of the deposited BN. For α = 4, the structure seemed to be isotropic, in contrast to that observed for α = 0.5, where a laminated structure of BN was evidenced.
Evolution of La and Lc versus α composition at 1250°C.
Evidence of the structure: with α = 4 the structure seemed to be isotropic (a), in contrast to that observed for α = 0.5, where a laminated structure of BN was evidenced (b).
The above results can be classified into three domains of temperature, gas pressure, and gas composition, for which specific structures of BN were observed (Fig. 6). Using a previous study of the kinetics of BN deposition based on the same gaseous precursors, it was possible to uncover a relationship between the mechanism controlling the kinetics of the CVI and the structure of BN observed in the present work.7)
Variation in the structure of BN with pressure and temperature.
Domain A corresponded to low temperatures (T < 1200°C), low pressures (P < 10 kPa), and a gaseous composition characterized by a > 1 (high ammonia content). The BN deposits were smooth and the structure was weakly organized (Fig. 6, left). Under these conditions, the surface adsorption and reactions control the kinetics of BN deposition.7) Because the mass transport of gaseous phases is not a limiting step, the density of sites for BN nucleation is suspected to be high. In this work, it resulted in a limited extension of the well-ordered structure of hexagonal BN, and the deposit tended to be isotropic. The fact that BN was not well crystallized is expected to be related to the gas composition, but the specific roles of BF3 and NH3 in the crystallization of BN have still not been determined.
Domain B corresponded to high temperatures (T > 1200°C), intermediate pressures (P = 20 kPa), and a gaseous composition characterized by a < 1 (low ammonia content). Under such conditions, the deposits are lamellar and correspond to the turbostratic form of BN (Fig. 6, bottom). The mass transport of the gaseous phase becomes the limiting mechanism, which is suspected to promote the extension of a well-ordered hexagonal structure.7) As previously mentioned, the relation between the diffusion mechanism controlling the CVI of BN and a well-organized structure of BN has been observed by several authors.8,9) Both studies noted that a diffusion screen that promotes kinetics controlled by mass transport promotes the organization of hexagonal BN. However, in the case of BN infiltration of large fibrous preforms, such a mechanism tends to increase the heterogeneity of the thickness of the deposit between the outer part and the inner part of the preform.5,8) In the present work, the substrate was composed of a thin fiber tow, which limited the problem of heterogeneity of the deposits. Since a diffusion screen was not used, it was necessary to reduce the total gas flow rate as much as possible in order to obtain a sufficiently large diffusion layer over the fiber substrate.
Domain C corresponded to high temperatures (T > 1200°C), high pressures (P > 20 kPa), and a gaseous composition determined by a < 1. The deposits were characterized by a higher ordered crystalline structure compared to Domain B. However, the morphology remained isotropic (Fig. 6, right). The BN coatings were porous and a strong heterogeneity was observed in the thickness of the deposits between the outer and the inner parts of the fiber tow. Considering these observations, nucleation in the gaseous phase is proposed to explain the morphology of the deposits and the low yield observed.
(d) Choice of BN structure for an interphase in CMCsAs noted above, the main function of an interphase in a CMC is to deflect the crack propagation near the fiber surface in order to avoid fiber failure and subsequent composite rupture. A laminated interphase is of strong interest because of the anisotropy of the structure and mechanical properties. BN can be used as a mechanical fuse in CMC if the structure of the deposit is sufficiently laminated to deflect crack propagation. The efficiency of BN for crack deflection was evidenced by thick coatings. TEM analyses of these deposits demonstrated the influence of the structure on the crack propagation through the BN (Fig. 7). In the case of isotropic BN deposited using Domain A conditions, the crack propagated through the deposit without deflection occurring.
The crack propagation of the turbostratic planes of the boron nitride.
When lamellar BN was deposited under conditions corresponding to Domain B, crack deflection was observed inside the interphase. The crack deflection (Fig. 7) followed the direction of the turbostratic planes of the BN. In Fig. 7, both the anisotropy of the structure and that of the mechanical properties were evident. Such micrographic evidence shows that the best BN interphase material for crack deflection in a CMC is likely under conditions corresponding to Domain B (Fig. 3).
The gaseous environment of the reactor can be strongly corrosive for Hi-Nicalon SiC fibers. It is well known that BF3 degrades silicon carbide, giving a SiF4 gaseous product. Such corrosion is active and the SiC fibers can be degraded in the same manner as when oxygen acts on the carbon fibers at high temperatures. For a CVI process involving BF3 gas content, the fiber is corroded at the beginning of the deposition process. Changing the geometry of the spire of the furnace can optimize the cycles with a maximum temperature (1250°C) corresponding to Domains A–C. A first protective coating has to be produced at the surface of the fiber in order to avoid reactions between the corrosive gas and the silicon carbide fiber.
In the present work, the sub-layer was made of isotropic BN produced using non-aggressive conditions before the deposition of the anisotropic layer. A multilayer BN coating is thus realized, which will help to promote crack deflection inside the interphase and consequently improve the lifetime of the interphase of CMCs in oxygen environments.
3.2 Continuous dynamic CVI process studyThe continuous fiber treatments by CVI coatings included a temperature gradient. The temperature of the fiber moving through the reactor first increased to the programmed temperature before decreasing to room temperature. The temperature profile can be optimized in order to control the structure of the BN deposited on the fiber substrate.
Several temperature profiles, described in Fig. 3, were tested. The first temperature profile (Profile 1) corresponded to the case where the kinetics always remained in Domain A. In Profile 2, the temperature went through Domains A and B, then returned back to A. Finally, in Profile 3, by changing the geometry of the spire of the furnace, it was possible to impose successive heating–cooling cycles with a maximum temperature corresponding to Domains A, B, or C. The interest in exploring Profile 3 is to increase the temperature gradient in order to make the transition between the successive structures of the BN more pronounced. The processing parameters of the continuous treatments of the SiC fiber using these temperature profiles for BN deposition are described in Table 2.
Considering Profile 1, an isotropic monolayer BN is expected. This coating was produced in order to compare the mechanical properties of CMCs with a monolayer BN interphase to those obtained with a multilayer BN interphase. The interest in a multilayer coating was to obtain a laminated layer that promotes crack deflection. Such a laminated layer can be obtained in Domain B, as described previously. The anisotropy of the BN structure was then obtained if the gas pressure was increased (Profile 2 in Table 2) or if the gas phase was BF3-rich (Profile 3 in Table 2). The mechanical resistance of the fibers after BN coating was tested under tension and the results are summarized in Table 3. The BN was isotropic, as expected considering the processing conditions. When several temperature domains were crossed (Profiles B and C), multilayer BN was observed (Fig. 6, bottom and right). They are both characterized by the presence of a laminated BN layer located in between isotropic layers. However, the degree of anisotropy of this layer and the mechanical bond between each layer are probably different. This difference in the morphology of the BN interphase will be analyzed when considering the tensile properties of the SiC/Si3N4 composites.
To determine the tensile resistance of the fibers after the BN coating, the measured load to rupture is divided either by the section of the fiber or by the full section, that is, the fiber and BN coating (see Table 3). However, because the as-processed BN is supposed to be more compliant than the silicon carbide fiber, the stress to rupture has to be determined considering the fiber section only. As reported in Table 4, the tensile results showed that the resistance of the silicon carbide fibers was not significantly modified after the processing of the BN coating.
Finally, an adherence analysis of the of the interphase of BN with SiC fibers is presented in Fig. 8, where after a strength test (Fig. 8(a)) it was noted that there was a large amount of interphase attached to the fiber (Fig. 8(b)), and also that this adhesion is present in the SiC matrix (Fig. 8(c)).
Morphology of the failure surface: smooth pulled out fibers (a), BN interphase fixed on pulled out fiber (b), and the influence of the structure on crack propagation through the BN interphase.
An interphase of BN with a lamellar texture was optimized and deposited as a ceramic matrix reinforcement for Hi-Nicalon fibers. By making small changes to the temperature and pressure, it was possible to optimize the interphase’s multilayer gradient properties with transitions in the texture of the interphase BN. We were able to obtain coatings consisting of a succession of isotropic and anisotropic layers, with isotropic layers being protective and the anisotropic layers providing a mechanical fuse. This was corroborated by TEM micrographs, which showed that the deviation of fissures in the interphase protected the fibers. The CVI process rests on the existence of areas with different temperatures in the experimental reactor. When coated, most of the fibers have mechanical characteristics very similar to those of the initial Hi-Nicalon fibers. All the interphases developed contain a laminate underlayer and play the role of a mechanical fuse, acting to divert cracks away from other areas of the ceramic. The presence of a laminar layer does not guarantee good mechanical properties; the quality of the fiber–interphase and matrix–interphase contact surfaces is crucial. Characterization by TEM showed such crack diversion, notably in the fiber–surface BN and the matrix SiC–interphase BN. These behaviors and properties are linked to the synthesis conditions and, in particular, the feed rate in the reactor.
