Effects of pyrocarbon interphase on microstructure and properties of C/SiBCN composites

C/SiBCN composites are expected to be widely used in aerospace applications because of their excellent high-temperature stability. However, the interfacial reactions have significantly limited their practical application. A pyrocarbon (PyC) interphase can improve the interfacial reactions of C/SiBCN composites. In this study, PyC interphases of different thicknesses (0.1 μm, 0.25 μm, and 0.5 μm) were introduced via chemical vapor deposition (CVD) process. The interface bonding of C/SiBCN composites with 0.1 μm and 0.25 μm thick interphases was relatively weak and the composites with 0.5 μm interphase exhibited strong interface bonding. After heat treatment at 1600 °C, the mechanical properties of the C/SiBCN composites with the 0.5 μm thick interphase was maintained at 131 MPa, and it was maintained at 105 MPa even after heat treatment at 1900 °C, indicating their excellent high-temperature mechanical properties. In short, 0.5 μm thick PyC interphase can effectively improve the interfacial reaction of the C/SiBCN composites, facilitating their application in high-temperature environments.


Introduction
SiBCN ceramics prepared via the pyrolysis of precursors have attracted considerable attention owing to their excellent high-temperature stability (up to 2000°C), oxidation resistance and creep resistance [1][2][3][4]. However, because of their brittleness, these ceramics cannot be utilized as structural components in aerospace systems [5]. Reinforcement of the SiBCN ceramic matrix with continuous carbon fibers is expected to improve its toughness, and the resulting carbon-fiber-reinforced SiBCN matrix composites (C/SiBCN composites) are expected to be widely used in aerospace applications.
Precursor infiltration and pyrolysis (PIP) is a promising method for the preparation of C/SiBCN composites [6]. However, a high-temperature environment significantly influences the properties of C/SiBCN composites prepared via PIP because of the occurrence of interfacial reactions between the carbon fibers and the SiBCN matrix at high temperatures; this drawback has significantly limited the practical application of these composites in the aerospace field [7,8]. Lee [9] prepared a C/SiBCN composite via PIP and found its flexural strength to be 255 MPa at 1500°C in argon; the strength decreased significantly (to approximately 70 MPa) at 1700°C. Ding et al [10,11] reported that a three-dimensional (3D) C/SiBCN composite fabricated by PIP had a flexural strength and modulus of 371 MPa and 31 GPa, respectively. As the heat treatment temperature increased beyond 1600°C, the flexural strength and modulus decreased greatly (to 92 MPa and 12 GPa, respectively). Luan et al [12] reported C/SiC-SiBCN composites with a flexural strength and modulus of 628.7 MPa and 74 GPa. After heat treatment at 1400°C, the flexural strength and modulus can be maintained at 292 MPa and 31.8 GPa.
Studies have rarely been conducted on an improvement in the interfacial reactions of C/SiBCN composites during the high-temperature treatment. Investigations of improvements in the properties of C/SiBCN composites after high-temperature treatment should focus on the interphase between carbon fibers and the matrix, because this interphase has a greatly influence on the properties of the composites [10]. Carbon fibers are typically coated with C [13,14], SiC [8,15], BN [16,17], etc as the interphase. The most common coating for optimizing the interphase are pyrocarbon (PyC). Nevertheless, the effects of the PyC interphase on the mechanical and high-temperature properties of C/SiBCN composites have rarely been reported thus far.
In this study, the PyC interphase was deposited onto carbon fiber perform by chemical vapor deposition (CVD), and then, C/SiBCN composites were fabricated using a SiBCN precursor via PIP. The effects of the PyC interphase on the mechanical properties, high-temperature properties, and microstructures of the C/SiBCN composites were investigated.

Materials
Plain-weave carbon cloth (1 k, T300, Toray) with dimensions of 180 mm × 180 mm was stacked to a thickness of 4.5 mm and a carbon preform was obtained by Z-stitching the stacked carbon cloth with carbon fibers (3 k, Toray) in a 5 mm × 5 mm space. The fiber volume fraction was ∼44.4%. The synthesis of the polyborosilazane (PBSZ) precursor has been described elsewhere [6,18]. The synthesis and use of PBSZ precursor were carried out using the Schlenk method [19].

Preparation of C/SiBCN composites
The PyC interphase was deposited onto the fiber preform via CVD using C 3 H 6 at 960°C for 5 h, 15 h, and 30 h. Detailed information about this deposition can be found elsewhere [13]. The fiber preforms became more rigid and difficult to compress as the PyC content increased, making it unable to maintain the volume fraction of fibers in the preform at 44.4%. As a result, deposition over extended periods (more than >30 h) was not carried out. The PyC-coated preform was densified using the PBSZ precursor via the PIP process to obtain C/SiBCN composites. The preparation process for the C/SiBCN composites has been described in detail elsewhere [6]. Composites with 0.1-, 0.25-, and 0.5 μm thick PyC were labeled C0.1, C0.25, and C0.5, respectively, and the composite without the PyC interphase was labeled C0.

Characterizations
The microstructures of the PyC interphase and C/SiBCN composite samples were characterized via scanning electron microscopy (SEM; JSM-7900F, Japan). The phase compositions of the samples after heat treatment were analyzed using x-ray diffraction (XRD; Cu K α radiation; D8 Advance, Germany). The apparent densities and open porosities of the C/SiBCN composites were measured via Archimedes' method. The flexural strength (σ f ) and modulus (E f ) of the C/SiBCN composites were measured via three-point bending tests of 3.0 mm × 4.0 mm × 60 mm specimens with a 50 mm span and 0.5 mm min −1 crosshead speed, in accordance with ASTM C1341-06. Four specimens were used for each test.

Results and discussion
3.1. Microstructure of PyC coating Figure 1 shows micrographs of the PyC coating on the carbon fibers. The thicknesses of the PyC coating on the fibers after deposition times of 5 h, 15 h, and 30 h were ∼100 nm, ∼250 nm, and ∼500 nm, respectively (figure 1). The PyC interphases had rough surfaces that were strongly bonded with carbon fiber, which is conducive to improving the interface bonding strength between fiber and matrix [13,14]. After deposition for 5 h, carbon fibers were covered with a rough PyC coating with a thickness of 100 nm (figures 1(a), (b)); this rough surface is advantageous for increasing the frictional sliding resistance of pullout fibers in C/SiBCN composites [13]. As the deposition time increased to 15 h, the thickness of the PyC coating increased to ∼250 nm (figures 1(c), (d)). The PyC coating was evenly deposited on the surface of the carbon fibers. The grooves on the surface of the coating were similar to those on the surface of the original fibers. After deposition for 30 h, the grooves on the PyC surface were shallow and the thickness of PyC was increased to ∼500 nm (figures 1(e), (f)).

Mechanical properties and microstructures of C/SiBCN composites
The mechanical properties of C0.1, C0.25, and C0.5 are presented in table 1. The density of the C/SiBCN composites was approximately 1.6 g·cm −3 , which is lower than those of the other composites listed in table 2, because the density of the SiBCN matrix was only 1.82 g·cm −3 . The composites had a lower open porosity (∼6%-9%), indicating that the precursor effectively infiltrated the fiber preforms with coatings of different thicknesses. The flexural strengths of C0, C0.1, C0.25, and C0.5 were higher than those of composites prepared by other groups (table 2), with the exception of the composite prepared by Ding et al [11]; these differences are attributed to the different fiber weaving methods.
The strength and modulus of C0.5 were slightly higher than those of C0, C0.1, and C0.25 (table 1), indicating that the thickness of the PyC interphase did not significantly affect on the mechanical properties of the composites. However, C0.1, C0.25, and C0.5 exhibited different fracture behaviors, as shown in figure 2. C0.1 exhibited large fiber pullout lengths (∼50-80 μm) and fiber debonding (figures 2(a), (b)), indicating some degree of pseudo-plastic fracture. Fiber pullouts were also observed for C0.25 ( figure 2(c)), but the fiber pullout length  for C0.25 was smaller than that for C0.1. The fracture surface of C0.5 was relatively flat and smooth, exhibiting brittle fracture behaviors (figure 2(e)). A relatively strong interphase can transfer loads effectively in fiber-reinforced composites, whereas a relatively weak interphase can improve the toughness of the composites [22]. The PyC coating of C0.1 was tightly bonded to the carbon fibers, and interface debonding occurred between the coating and the matrix ( figure 2(b)). Therefore, the interphase of C0.1 was relatively weak, and the composite exhibited large fiber pullout lengths. The PyC coating of C0.25 was peeled off along with the matrix and the carbon fibers ( figure 2(d)). The interphase of C0.25 could not transfer loads effectively; therefore, the strength of C0.25 was comparatively lower than C0.1 and C0.5. The PyC coating of C0.5 caused strong bonding between the matrix and the fibers (figure 2(f)), which indicated the high strength of the interphase bond in C0.5. The flexural strength of C0.5 was slightly higher than those of C0, C0.1, and C0.25, indicating that loads were transferred more effectively in C0.5.

Microstructures and properties of composites after heat treatment
The properties of C0.1, C0.25, and C0.5 after heat treatment at 1600°C are presented in table 3. The weight loss rate, density, and mechanical properties of C0.1 after heat treatment at 1600°C were similar to those of C0.25 after heat treatment at 1600°C (table 3). However, C0.5 after heat treatment at 1600°C exhibited a higher flexural strength (131 MPa). Furthermore, the weight loss rate of C0.5 was lower than those of C0.1, and C0.25. The XRD patterns of the C/SiBCN composites are shown in figure 3. The SiBCN matrix of the as-prepared C0.1, C0.25, and C0.5 samples was amorphous, and peaks corresponding to carbon fibers appeared at 26°and 43°. After heat treatment at 1600°C, peaks of SiC grains appeared in the XRD patterns of C0.1, C0.25, and C0.5; the SiC peaks for C0.1 were sharper than those for C0.25 and C0.5. SiC grains formed after heat treatment at 1600°C because of the carbothermal reduction reaction of the volatile silicon and carbon [10,23]. The SiC peaks for C0.1 were sharper than those for C0.25 and C0.5, probably because of the more intense interfacial reactions of the carbon fibers in C0.1, which resulted in severe fiber damage and severe degradation of the mechanical properties of C0.1. Figure 4 shows the microstructures of polished cross-sections of C0.1, C0.25, and C0.5 after heat treatment at 1600°C. SiC rings appeared around the carbon fibers along the cracks in the composites [7], where volatile silicon elements could easily diffuse to the surface of the fibers. Because of the weak interface bonding and  thinner PyC coatings of C0.1 and C0.25, the interfacial reactions in these two composites were more violent than those in C0.5, which resulted in the formation of more SiC rings in C0.1 and C0.25. These interfacial reactions led to numerous cracks (figures 4(b), (d)) and higher weight loss rates and porosities (table 3). Therefore, the flexural strengths of both C0.1 and C0.25 after heat treatment at 1600°C were significantly reduced to 18.9-31.4 MPa. However, due to the strong interface bonding in C0.5, most carbon fibers were tightly bonded to the matrix after the heat treatment at 1600°C, and the volatile silicon could not directly react with the carbon fibers. In addition, because of the thicker PyC coating of C0.5, only a small amount of PyC underwent the carbothermal reduction reaction. Therefore, the strength of the fibers was well retained and C0.5 exhibited a high flexural strength of 131 MPa (table 3).   carrying capacity of the SiBCN matrix caused by the more numerous cracks and voids in these composites. C0.5 exhibited a typical fracture morphology (figure 5) that was similar to those of carbon-fiber-reinforced ceramic matrix composites [23,24]. This is because of the strong interlaminar bonding in the fiber cloth due to the higher density and lower open porosity of C0.5 after heat treatment at 1600°C (table 3).
For further investigation of the high-temperature properties C0.5, the heat treatment temperature was increased to 1900°C. The flexural strength and modulus of C0.5 after heat treatment at 1900°C were105 MPa and 39.7 GPa, respectively. Polished cross-sections of C0.5 after heat treatment at 1900°C are shown in figures 7(a), (b). Few SiC rings appeared in C0.5 after heat treatment at 1900°C, indicating that no significant interfacial reactions occurred even at 1900°C. The load-displacement curves of C0.5 after heat treatment at 1900°C are shown in figure 8. As-prepared C0.5 exhibited brittle fracture behavior, and the load decreased sharply to zero after reaching the maximum value (figure 8), which resulted in the flat fracture surface exhibited in figure 2(e). C0.5 showed pseudo-plastic fracture behavior after heat treatment at 1600°C, which could be ascribed to fiber pull-out and interface separation that occurred during the damage to the composite, as shown in figure 6(e). However, with heat treatment at 1900°C, the composite shrank and the number of cracks decreased, resulting in a strong interface bond. Therefore, the fracture surface of C0.5 after heat treatment at 1900°C was flat ( figure 7(c)), and it showed somewhat brittle fracture behavior ( figure 8).     mechanical properties can be maintained even at 1900°C, which is promising for applications involving microwave absorption [25], high-temperature ablation resistance [26], and structural components [27].