Mechanical Properties and Degradation Mechanism of SiC Fibers Exposed to Oxidative Environment up to 1600 °C

Abstract

In order to investigate the microstructure evolution and the degradation mechanism of SiC fiber in a high-temperature oxidative environment, the SiC fiber was thermally exposed at temperature up to 1600 °C in air. The morphologies of the surface and fracture surface were characterized by scanning electron microscopy. The consisting phase and crystallite size were analyzed by X-ray diffractometer. The mechanical properties of SiC fiber was characterized by a single-fiber tensile test technique. It was found that an obvious grain coarsening occurred at temperature above 1400 °C. A visible silica layer was formed at 1300 °C, and the morphology of silica layer was dependent on the exposure temperature. At 1400 °C, fiber surface formed a thick silica layer with cracks, while the silica layer exhibited a multilayered structure at 1600 °C. As for the tensile strength of fiber, it firstly decreased to about 1 GPa at 1200 °C, then the strength was maintained at 1400 °C. After thermal exposure at 1500 °C and 1600 °C, the strength decreased again. The degradation of mechanical properties was attributed to the grain coarsening and the decomposition of amorphous phase in fiber. Particularly, the decomposition of amorphous phase would damage the structure integrity of fiber. The current work would provide a valuable reference for research and application of SiC fiber.

1. Introduction

Silicon carbide (SiC) fiber has emerged as an important reinforcing material for high-temperature ceramic matrix composites (CMCs), attributed to its lightweight, exceptional mechanical properties and excellent compatibility with ceramic matrix.

As core thermal structural components in next generation of aerospace engineering, CMCs are put in a high priority in advancing high-temperature technological applications [1,2,3,4]. However, the critical demands for environment-adaptive performance of thermal structural materials have sparked intensive research focus. Notably, under complex and severe service conditions, the oxidation resistance of SiC fiber is closely dependent on both service temperature and exposed atmosphere, which exerts a decisive influence on its mechanical properties and microstructural stability. The oxidation behavior of SiC fibers is strongly associated with not only themselves, but also the formation and evolution of their oxidation products: at moderate temperature, a dense and continuous silica (SiO₂) layer is typically generated on the fiber surface in ambient atmosphere, which acts as an environmental barrier to inhibit further oxygen diffusion; however, when the temperature exceeds 1400 °C, the SiO₂ layer tends to undergo viscous flow, forming cracks or even volatilizing [5,6,7], for example, in a wet oxidation atmosphere, via reactions with trace moisture to produce Si(OH)₄, thereby losing its protective effect. This temperature-dependent microstructural evolution, variation of oxidation products directly governs the degradation mechanisms of SiC fibers, making the clarification of such processes essential. Meanwhile, the new SiC fiber demonstrated an increase of strength in as-fabricated state, but it still faced the same challenges in environmental resistance as the traditional SiC fibers do [8]. Thus, in-depth investigations into the performance evolution of SiC fiber under extreme environments have a great of significance for the development of CMCs.

In previous studies, the behavior of SiC fibers in high-temperature oxidative environments has been extensively studied [9,10,11,12,13]. Multiple analytical techniques were employed to characterize their physical and chemical performance under various service conditions, such as X-ray diffraction (XRD), scanning electron microscopy (SEM), and X-ray photoelectron spectroscopy (XPS). The crystallinity of SiC in the fibers played an important role in the formation and protection of the oxide layer. Sha et al. [9] reported that the mechanical properties (creep resistance) of SiC fibers were mainly related to the grain size of β-SiC and the composition of grain boundary. Wu et al. [10] reported that the high-temperature resistance of SA type SiC fibers was excellent, which was attributed to the high crystallization, large crystal grains and low carbon and oxygen content in the fiber. You et al. [13] found that the residual stress between SiC grains increased during the thermal shock of SiC fibers at 1200 °C, and the thermal expansion mismatch between the oxide layer and the fiber core led to oxide layer peeling and crack formation, resulting in defects in the fibers. These investigations consistently indicated: the oxidation behavior of SiC fibers, including the composition, thickness and structure of oxidation products, is strongly correlated with fiber type (e.g., precursor-derived SiC fiber, chemical vapor deposition-derived SiC fiber); secondly, the oxidation process, especially the breakage or volatilization of the SiO₂ protective layer, induces substantial degradation of the fibers’ mechanical properties.

Although SiC fibers have been commercially available for over three decades, considering the cost-efficiency and high performance, the research and development are still undergoing to fulfill the diverse application requirements. Moreover, clarifying the degradation mechanism of SiC fibers and establishing the quantitative relationship between their performance, microstructure and the service environments were necessary. Therefore, for a better understanding of the microstructural evolution and mechanical properties of fiber under thermal and oxidative coupling environment, the present work subjected the domestic SiC fiber to thermal exposure in ambient atmosphere at temperatures ranging from 1200 to 1600 °C. After thermal exposure, single-fiber tensile tests were conducted to assess the mechanical properties of the fibers, such as tensile strength and elastic modulus, and to analyze the potential degradation mechanisms caused by the high-temperature exposure and the oxidation process. Furthermore, the surface morphology and fracture surface characterization were also performed to establish the correlation between mechanical performance and microstructural evolution. Based on these comprehensive analyses, the degradation mechanism of SiC fibers in oxidative environments was systematically clarified. The combination of mechanical testing, microstructural analysis and fracture surface observation allowed for a comprehensive understanding of how thermal exposure and oxidation affect the integrity of SiC fibers. This would lay a foundation for the further improvement of thermal and mechanical stability of SiC fibers under extreme service environment.

2. Experimental Procedure

The domestic SiC fibers used in this study were prepared by polycarbonsilane-derived technology and crosslinked by electron beam irradiation. The diameter of fiber was about 12 µm, and the density was about 2.7 g/cm³. For the chemical composition, the mole ration of C to Si was 1.4, and the oxygen content was about 1.3 wt%.

The SiC fiber was placed in a muffle furnace (KSL-1700X, Kejing, Hefei, China) and subjected to a thermal exposure in air. The exposure temperature was varied from 1200 to 1600 °C. The heating rate was 10 °C/min and the fibers were isothermally heated at each temperature for 1 h. Subsequently, the fibers were cooled to room temperature and taken for tensile testing and microstructural characterization.

The crystallite structure of SiC fiber was characterized by X-ray diffractometer (XRD, D8ADVANCE, Bruker, Berlin, Germany) with Cu-Kα irradiation. The diameter, surface morphology and fracture surfaces of the fibers were characterized by using the field-emission scanning electron microscope (FE-SEM, KYKY2800B, Zhongke, Beijing, China). Due to the poor electrical conductivity of SiO₂, prior to the SEM observation, the SiC fibers were mounted on a substrate with the carbon adhesive tape and sputtered with a gold layer. Following this process, the high-quality SEM images could be obtained.

Due to the formation of viscous silica layer on the surface of fiber at high temperature in air, the individual fiber within a bundle were bonded together by the silica. For the tensile test of fiber, a single fiber should be extracted from the bundle. To minimize the possible damage during the extraction of the single fiber, the oxidized fibers were immersed in HF solution for a short time to remove the surface SiO₂ layer. Tensile testing was performed on an Instron testing machine (Instron3345, Instron, Boston, MA, USA) with a load cell of 5 N. During the tensile test, the load was applied by the displacement control mode with a speed of 0.1 mm/min. The gauge length was 25.0 mm for tensile test. For each condition, at least 25 valid tests were carried out for the data analysis.

3. Results and Discussion

3.1. XRD Analysis

Figure 1 shows the XRD patterns of the SiC fibers after thermal exposure at various temperatures. It was observed that the as-received fiber exhibited no detectable phase corresponding to β-SiC, indicating a typical amorphous phase of SiC_xO_y [5]. For the fiber exposed at 1200 °C, the characteristic peak of β-SiC appeared at 35.7°. This indicated that the obvious decomposition of amorphous SiC_xO_y phase started at 1200 °C [6], as described by the following reaction:

$${\mathrm{S}\mathrm{i}\mathrm{C}}_{\mathrm{x}}{\mathrm{O}}_{\mathrm{y}}\left(\mathrm{s}\right)\to \mathrm{S}\mathrm{i}\mathrm{C}\left(\mathrm{s}\right)+\mathrm{C}\left(\mathrm{s}\right)+\mathrm{C}\mathrm{O}\left(\mathrm{g}\right)+\mathrm{S}\mathrm{i}\mathrm{O}\left(\mathrm{g}\right)$$

(1)

Based on Equation (1), The carbon phase should be formed from the decomposition of amorphous SiC_xO_y phase. However, in the oxidative environment, the carbon was oxidized by oxygen from the atmosphere, as depicted by following equation:

$$\mathrm{C}\left(\mathrm{s}\right)+{\mathrm{O}}_2\left(\mathrm{g}\right)\to 2\mathrm{C}\mathrm{O}\left(\mathrm{g}\right)$$

(2)

On the other hand, as the thermal exposure temperature increased, the cristobalite phase became obvious. Particularly, at temperature beyond 1400 °C, a sharp peak with high intensity presented due to an intensive oxidation. Meanwhile, the intensity of β-SiC peak progressively increased with rising temperature. However, the intensity of β-SiC peak was weaker than that of the cristobalite peak. This was due to the fact that the fiber oxidized severely at high temperature, and generated a large amount of SiO₂. Meanwhile, high temperature enhanced the crystallization of the SiO₂ oxide layer. Both high crystallinity and large volume fraction of SiO₂ resulted in high diffraction intensity of cristobalite in comparing with that of SiC. The crystallite size of β-SiC was also estimated using the Scherrer equation, expressed as Equation (3):

$$L={\displaystyle \frac{K\lambda }{D\mathrm{c}\mathrm{o}\mathrm{s}\theta }}$$

(3)

where K is the Scherrer constant (taken as 0.9), λ the wavelength (0.154 nm) for Kα irradiation, and D the full width at half maximum (FWHM).

By fitting the β-SiC peak, the D was calculated. Applying Equation (3), the crystallite size of β-SiC was calculated. Figure 2 shows the crystallite size as a function of thermal exposure temperature. A significant increase of crystallite size at temperature above 1400 °C was observed. After thermal exposure at 1600 °C, the crystallite size of β-SiC was about 16 nm.

3.2. Cross-Sectional Microstructure Characterization

Figure 3 shows the cross-sectional SEM micrographs of the SiC fibers after thermal exposure in air. The thickness of the oxide layer on the fiber surface progressively increased with increasing the thermal exposure temperature. At 1400 °C, a thick silica layer with cracks was formed on the surface of fiber (Figure 3d). Notably, at 1500 °C, the debonding occurred at the interface between the SiO₂ layer and the SiC fiber, forming an interfacial gap (Figure 3e). Figure 3f revealed that the SiO₂ layer formed at 1600 °C displaying a multilayered structure. Meanwhile, the profile of the fiber cross section was not circular. Such phenomenon might be associated with the viscous flow of silica layer. It is well known that at temperatures above 1500 °C, the viscosity of silica is very low. In this case, the silica tended to flow around the fiber surface. The silica was formed based on the reactions [7]:

$$2\mathrm{S}\mathrm{i}\mathrm{C}\left(\mathrm{s}\right)+3{\mathrm{O}}_2\left(\mathrm{g}\right)\to 2{\mathrm{S}\mathrm{i}\mathrm{O}}_2\left(\mathrm{s}\right)+2\mathrm{C}\mathrm{O}\left(\mathrm{g}\right)$$

(4)

To some extent, the SiO₂ layer could encapsulate the fiber core, effectively inhibiting the direct oxygen invasion. In this situation, the oxidation of the fiber core was mainly governed by oxygen diffusion, establishing a concentration gradient within the surface of SiC fiber [14]. Meanwhile, the continuous decomposition of the amorphous SiC_xO_y phase generates a gas pressure at the interface, produced the defects at the interface or in the early formed SiO₂ layer. As described in the literature [15,16], under atmospheric thermal exposure, the partial pressure of O₂ exceeded the partial pressures of SiO and CO. In this situation, the thermodynamic driving force enabled O₂ permeation through defects, initiating the secondary oxidation with SiC and SiO. Consequently, a new SiO₂ layer formed beneath the original oxide layer. Such repetitive oxidation eventually formed a multilayered structure of SiO₂ (Figure 3f). Figure 4 illustrated the oxidation mechanism for the SiC fiber exposed in air at high temperature.

Figure 5 shows the thickness of SiO₂ layer as a function of temperature. For the fibers exposed at temperatures below 1300 °C, the thickness of silica layer was less than 0.5 μm. When the temperature exceeded 1400 °C, the thickness of silica layer increased significantly, reaching 3 μm at 1600 °C. The increased silica thickness mainly caused by the fast decomposition of the SiC_xO_y phase above 1400 °C, which degraded the structure of fiber and consequently accelerated the oxidation.

Figure 6 shows the mass change of SiC fibers after thermal exposure at different temperatures in air. Based on the observations of the mass variation, the mass change should be strongly related to a competing mechanism: mass loss caused by the decomposition of the amorphous SiC_xO_y phase, and mass gain caused by the oxidation of SiC. From 1200 to 1300 °C, a slight mass gain was observed, remaining below 0.5%. This meant that the mass gain and the mass loss were in a balance state. Meanwhile, in this temperature range, both the oxidation and decomposition should not be serious. This was evidenced by a weak β-SiC peak in XRD patterns (Figure 1) and a thin silica layer in Figure 3b,c. The partial decomposition of SiC_xO_y phase at low temperature generated a limited SiC and SiO, which subsequently reacted with atmospheric oxygen, forming a small amount of SiO₂. As observed in Figure 3b,c, the silica layer is very thin.

At 1400 °C, the mass loss was about 2%, as shown in Figure 6. This indicated that the decomposition of fiber was a dominant mechanism for mass loss, while the oxidation of fiber was minor. This was due to that high temperature drove a fast decomposition of amorphous SiC_xO_y, leading to the release of the gaseous products. In this case, the mass loss exceeded the oxidation-induced mass gain. At 1500 °C, the intensified oxidation promoted the oxidation, forming massive SiO₂, as shown in Figure 3e. The mass change was higher than 8%, as shown in Figure 6. Consequently, the vapor-phase reactions between SiO and C from the decomposition products regenerated SiC at the surface, accompanied by a significant grain coarsening (Figure 2 and Figure 3e). A thick silica layer would induce the thermal stress accumulation during cooling, finally resulting in the generation of crack within the SiO₂ layer. At 1600 °C, the rapid formation of oxide layer could inhibit the both SiC_xO_y decomposition and oxygen invasion, leading to a mass loss comparing with that of 1500 °C. However, multilayered oxidation might compromise the fiber integrity (Figure 3f).

3.3. Morphology of Fiber Surface

Figure 7 shows the morphologies of SiC fibers after thermal exposure in air. The as-received fiber presented a smooth surface (Figure 7a). At 1200 °C, the surface exhibited no discernible changes; particularly, no observable oxide layer, as shown in Figure 7b. When the temperature increased to 1300 °C, an observable SiO₂ layer formed on the fiber surface. The oxide layer encapsulated the fiber (Figure 7c), to some extent, which could effectively inhibit the oxygen invasion, thereby suppressing further oxidation.

At 1400 °C, a relatively thick SiO₂ layer was formed on the fiber surface, accompanied by extensive cracking, as shown in Figure 7d. This phenomenon was probably caused by two factors: first, the rapid decomposition of the amorphous SiC_xO_y could generate a large amount of gases, resulting in an internal pressure at the interface between fiber core and the silica layer. When the pressure of gases exceeded the fracture stress of the silica layer, the silica layer would rupture. Secondly, the mismatch of coefficient of thermal expansion (CTE) between the silica layer and fiber core also generated stress during cooling, leading to the rupture of silica layer. This is quite possible because the CTE of silica is larger than that of SiC fiber [6]. During cooling, the SiO₂ layer with large CTE would experience a tensile stress, leading to the crack formation [17,18].

At 1500 °C, SiO₂ layer was partially peel off. Meanwhile the fiber surface presented a rough morphology (Figure 7e). The silica layer peeling off should be also related to the thermal stress caused by CTE mismatch. At 1600 °C, extensive large cracks were formed within the silica layer, and the structural integrity of the fiber was destructed (Figure 7f). Previous work also found that the fracture of the silica layer above 1500 °C could induce the premature failure of the fiber core in air [19].

3.4. Mechanical Properties of Fibers

In order to elucidate the degradation mechanism of mechanical properties, the single-fiber tensile test was conducted. However, the challenge was that the individual fiber within a fiber bundle was bonded by the silica formed during thermal exposure in air. To avoid the possible damage of the fibers during the tensile specimen preparation, the oxidized fibers were soaked into the HF solution to remove the surface SiO₂ layer.

Figure 8a shows the typical stress–strain curves for SiC fibers exposed at elevated temperatures in air. It was clear that fibers presented a brittle fracture behavior. Furthermore, the slope of the stress–strain curve varied with the thermal exposure temperature.

Based on the fracture stress obtained from the single-fiber tensile test, the strength (σ_i) of fiber was calculated. The data of tensile strength was analyzed by the two-parameter Weibull statistical model [20,21]. The Weibull distribution function is expressed as:

$$F_i=1-\mathrm{e}\mathrm{x}\mathrm{p}[-L{\left({\displaystyle \frac{\sigma }{{\sigma }_0}}\right)}^m]$$

(5)

where F_i is the fracture probability, σ₀ is the scale parameter, L is the gauge length of fiber specimen, m is the Weibull shape parameter (Weibull modulus), and σ is the fiber tensile strength. Taking the natural logarithm of both sides of Equation (5) yields:

$$\ln \left[\ln \left(1/(1-F_i)\right)\right]=m\left(\ln {\displaystyle \frac{{\sigma }_i}{{\sigma }_0}}\right)+c$$

(6)

where c is a constant. The expression for F_i is defined as follows:

$$F_{\mathrm{i}}={\displaystyle \frac{\mathrm{n}}{N+1}}$$

(7)

where N represents the total number of samples and n represents the intensity ranking of the test samples. The average tensile strength is calculated by following equation:

$${\sigma }_{\mathrm{a}\mathrm{v}\mathrm{g}}={\sigma }_0\Gamma (1+{\displaystyle \frac{1}{m}})$$

(8)

where Γ is the gamma function.

Using Equation (8), the statistical data was plotted in Figure 8b, where ln[ln(1/(1 − F_i))] was as a function of ln(σ_i). Then, the data was fit by the general least square method. The slope of the fitting line corresponded to the Weibull modulus, m, which reflected the variability of the fiber strengths.

Figure 8c shows the temperature dependence of the average tensile strength for fibers exposed at elevated temperatures in air. The as-received fiber exhibited a tensile strength about 2 GPa. After 1200 °C exposure, the strength decreased to about 1 GPa. From 1300 to 1400 °C, the strength was almost unchanged within the statistical errors. After thermal exposure at 1500 °C and 1600 °C, the strength decreased again, as shown in Figure 8c. This was due to the rapid decomposition of amorphous phase and a cracked silica layer on fiber surface (Figure 7e,f). The elastic modulus is also a critical property of SiC fibers. The elastic modulus of the SiC fibers was calculated from the slope of stress–strain curves in Figure 8a. The elastic modulus as a function of exposure temperature was plotted in Figure 8d. The elastic modulus of the as-received fibers was about 115 GPa. After thermal exposure at 1300 °C, the modulus exceeded 200 GPa. This might be due to the fact that the original SiC fibers contain amorphous SiCxOy phases, which are usually composed of disordered atomic structures. Through appropriate thermal exposure treatment, the amorphous phases were transformed into SiC, resulting in high crystallinity and a near stoichiometric composition. Thereby, the binding force between grain boundaries was strengthened and the elastic modulus was increased. A similar phenomenon was also observed that the elastic modulus of SiC fiber is much higher than that of the Si-C-O fiber [22].

This indicates that appropriate thermal exposure could improve the elastic modulus of fibers. However, if the temperature was too high, it would damage the structure integrity of fiber due to the rapid decomposition of amorphous phase [23]. Thus, the decrease of elastic modulus was primarily attributed to the damaged microstructure of fiber. Comparing with other studies, it was found that the strength of Hi-Nicalon fiber was 2.8 GPa in as-received state [24]. After thermal exposure at 1500 °C, the current SiC fibers presented a low strength, while Hi-Nicalon fibers still retained a strength higher than 76% [25]. This meant that the high crystallinity and stoichiometric composition were crucial factors for high thermal and mechanical stability of SiC fiber.

Figure 9 shows the surface morphologies of fibers without the silica layer. The high strength of the as-received fiber was due to its smooth surface and dense structure, as shown in Figure 9a. The thermal exposure at 1200 °C caused the crystallization of fiber and incomplete silica layer coverage. In this situation, the incomplete silica formation on fiber surface could not effectively inhibit the oxygen ingress (Figure 9b), accelerating the internal oxidation and resulting in the strength degradation, as shown in Figure 8c. From 1300 °C and 1400 °C, due to the rapid formation of protective SiO₂ layer, the fiber oxidation and generation of defects were alleviated, as shown in Figure 9c,d. After thermal exposure at 1500 °C and 1600 °C, due to the rapid decomposition of amorphous phase and a cracked silica layer on fiber surface (Figure 7e,f), the fiber surface presented an inhomogeneous oxidation and many pits were formed on the fiber surface (Figure 9e,f). This led to a more significant strength decrease, as shown in Figure 8c. Furthermore, the abnormal grain growth also degraded the fibers’ microstructure, as shown in Figure 2.

After the single-fiber tensile test, the microstructure observations on the fracture surfaces were also performed by SEM, as shown in Figure 10. The distinct feature was that the fracture surfaces of fibers displayed a brittle morphology, originating from the different defects. The fracture surface of the as-received fiber presented a glass state and no obvious fracture origin (Figure 10a). The fracture of fiber exposed at 1200 °C and 1300 °C mainly originated from the inner defects, as shown in Figure 10b,c.

When the exposure temperature increased to 1400 °C, the initiation site of fracture shifted to surface (Figure 10d). This should be attributed to the surface defects generated from the decomposition of amorphous phase. At 1500 °C (Figure 10e), the grain coarsening was obvious, and the microstructure became porous.

4. Conclusions

The domestic SiC fiber was thermally exposed at temperatures ranging from 1200 °C to 1600 °C in air. After thermal exposure, the mechanical properties were evaluated via a single-fiber tensile test technique. The microstructure evolution and degradation mechanism of SiC fiber in high-temperature oxidative environment were investigated. The main findings were summarized as follows:

• The as-received SiC fibers mainly compose of an amorphous SiC_xO_y phase. The decomposition of amorphous phase starts at 1200 °C, and is strongly dependent on the thermal exposure temperature. Meanwhile, the decomposition of amorphous phase enhances the crystallinity and oxidation of fiber. Particularly, at 1600 °C, the oxidation-formed silica layer on fiber surface presents a multilayered structure.
• The tensile strength of fiber firstly decreases at 1200 °C. Then, the strength almost remains unchanged within the statistical error from 1300 to 1400 °C. After thermal exposure at 1500 °C and 1600 °C, a significant strength reduction is observed. The degradation of fibers’ strength is attributed to the grain coarsening and the damage of structure integrity.
• Thermal exposure at moderate temperature improves the elastic modulus. For the fiber exposed at 1300 °C, the elastic modulus is the highest, exceeding 200 GPa. Above 1300 °C, the elastic modulus shows a decrease trend with increasing the temperature. The reduction of elastic modulus was mainly attributed to the high-temperature decomposition of the amorphous phase, which leads to a damage of structure integrity of fiber.
• SiC fiber, as the reinforcement of ceramic matrix composite, is the backbone for load bearing. To achieve the excellent thermal and mechanical thermal stability, the fiber with high crystallinity and near stoichiometric composition is crucial for the fabrication and application of ceramic matrix composite. This is a critical consideration for the long-term durability of SiC fibers in CMCs used in aerospace components like turbine blades and combustion chambers.