Journal of Materials Science & Technology  2019 , 35 (12): 2785-2798 https://doi.org/10.1016/j.jmst.2019.08.004

Orginal Article

Ablation characteristics of mosaic structure ZrC-SiC coatings on low-density, porous C/C composites

Yonglong Xua, Wei Suna*, Xiang Xionga, Fuqun Liub, Xingang Luanc

a.State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China
b. China Academy of Launch Vehicle Technology, Beijing 100076, China
c. Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an 710072, China

Corresponding authors:   *Corresponding author.E-mail address: sunweimse@csu.edu.cn (W. Sun).

Received: 2019-01-20

Revised:  2019-07-17

Accepted:  2019-07-22

Online:  2019-12-05

Copyright:  2019 Editorial board of Journal of Materials Science & Technology Copyright reserved, Editorial board of Journal of Materials Science & Technology

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Abstract

Mosaic structure ZrC-SiC coatings were fabricated on low-density, porous C/C composites via thermal evaporation and an in-situ method. ZrC was packed in a typical lamellar mode, and the mosaic structure was formed by the deposition of Zr and Si atoms on the shallow surface of the porous C/C composites. Ablation analysis showed that the defects in the coatings originate from the boundary between the ZrC and holes created by the consumption of SiC at 2500 °C. After ablation for 200 s at 3000 °C, a dense ZrO2 layer formed on the coating surface, and the defects were sealed owing to the continuous supply of ablative components. The mass and line ablation rates of the ZrC-SiC coatings were -0.46 ± 0.15 mg cm-2·s-1 and -1.00± 0.04 μm s-1, respectively.

Keywords: ZrC ; SiC ; Coating ; Mosaic structure ; Porous C/C composites ; Ablation

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Yonglong Xu, Wei Sun, Xiang Xiong, Fuqun Liu, Xingang Luan. Ablation characteristics of mosaic structure ZrC-SiC coatings on low-density, porous C/C composites[J]. Journal of Materials Science & Technology, 2019, 35(12): 2785-2798 https://doi.org/10.1016/j.jmst.2019.08.004

1. Introduction

For the development of hypersonic aerospace vehicles, vital components, such as nozzles (or noses) and leading edges, need to endure severe heat fluxes and oxidation at temperatures above 2000 °C during flight [1,2]. Carbon fiber-reinforced carbon matrix composites (C/C composites) are promising materials for such critical components owing to their low weight, good thermal shock resistance, and good high-temperature mechanical properties [3,4]. However, rapid oxidation and ablation failure at ultrahigh temperatures hinder their application in advanced aerospace systems [5].

The use of protective coatings is an effective way of improving the ablation resistance of C/C composites. Ultrahigh-temperature ceramic materials (UHTCs) are considered excellent coating components and have attracted much attention in recent years owing to their exceptional hardness, high melting point, chemical inertness, and low vapor pressure at high temperatures [[6], [7], [8]]. In particular, ZrC, which is an excellent ablative coating material, has a melting point of approximately 3540 °C and a significantly lower density than HfC and TaC. Moreover, its oxidation product ZrO2 has a high melting point as well as low thermal conductivity, high impact resistance, and good thermal corrosion resistance [9,10].

However, a ZrC coating has two limitations when it comes to protecting C/C composites: a significant difference in the thermal expansion coefficients (CTE) between ZrC (6.7 × 10-6  K-1) and C (1 × 10-6  K-1) and the formation of a loose and porous state during ablation [11]. Much attention has been given to adding SiC, LaB6, or C to solve the aforementioned problems. Sun et al. [12] deposited a ZrC/C gradient layer using the chemical vapor deposition (CVD) technique to improve the bonding between ZrO2 and a C/C composite. Yao et al. [13] fabricated a ZrC-SiC coating by supersonic plasma spraying; a SiC (4.5 × 10-6  K-1) transition coating was prefabricated to relieve the thermal expansion stress between the ZrC outer layer and the carbon matrix. Jia et al. [14] fabricated a LaB6-stabilized ZrC coating on a prefabricated SiC coating by combining the atmospheric plasma spraying technique and a packing method.

In our previous work, ZrC-SiC mixed coatings were prepared by thermal evaporation and in-situ methods on the graphite surface [15]. These kinds of coatings possess a ZrC/SiC/C multilayer structure that can relieve the CTE between ZrC and C. Furthermore, liquid SiO2 formed by the SiC can heal surface defects of ZrO2 at the early stage of ablation. In addition, Zr(g) and Si(g) can infiltrate into the graphite matrix and form a ZrC/SiC pinning structure. For this reason, we believe that this method has significant advantages for fabricating ZrC-SiC coating on porous C/C composites. There are more porous channels on the shallow surface of the porous C/C matrix, and a ZrC-SiC mosaic structure can be formed by the adsorption and reaction between Zr, Si vapor, and carbon in these porous channels, which can effectively improve the bonding force between the coating and matrix. Furthermore, matrices with high porosity can contain more ablative SiC and ZrC components, which is conducive to the improvement of long-term ablative resistance [16].

In this study, ZrC-SiC coatings were fabricated on low and higher-density C/C composites via thermal evaporation and an in-situ method, respectively. The formation mechanism, microstructure, and ablation properties of the multilayer ZrC-SiC coatings were investigated.

2. Experimental

The matrix was fabricated from bulk 2.5D needled integral C/C composites with a density of 1.42 and 1.65 g/cm3 densified by chemical vapor infiltration and then be cut into φ 28 × (5-10) mm specimens [17]. The initial open porosity of with a density of 1.42 and 1.65 g/cm3 were 22.03% and15.3%, respectively. The Zr and Si powders were mixed in a mass ratio of Zr:Si = 1:1, and then placed at the bottom of a graphite crucible, following a previously reported method [15]. The C powder was not added to the mixed powders, so as to avoid overreaction before evaporation and the subsequent formation of high-melting-point carbides. The crucible was then placed in a high-frequency induction furnace, which was firstly heated to 2400 °C under a heating rate of 15 °C/min and a micro-positive argon atmosphere. Then, during the insulation stage, the pressure was set at 200 Pa at 2400 °C, and argon was flowed into the furnace; the temperature was maintained for 2 h. An argon flow rate of 15 L/min was used to maintain a micro-positive pressure, and finally, the furnace was cooled down to room temperature. The evaporation temperature and low pressures were set to ensure high volatility of the molten Zr atoms. In our previous study [15,16], the carbon atoms in the C/C matrix had excellent activity, and could rapidly diffuse and react with the deposited Zr and Si atoms, thus forming a large amount of carbide in a short time. For convenience, the coating fabricated on the surface of the C/C composites with a density of 1.42 g/cm3 was termed “coating CL”, and that with a density of 1.65 g/cm3 was termed “coating CH”.

Phase analysis of the ZrC-SiC coatings before and after ablation was performed with an X-ray diffraction (XRD) analyzer (D/max 2550 vb +18 kW, Rigaku Co., Japan). The microstructure of the coating was analyzed by scanning electron microscopy (SEM, NanoSEM230, Novtma, Holland). The distribution of Zr, Si and C element was investigated by using an electron probe micro-analyzer (EPMA, JXA-8530 F, JEOL, Japan). Raman spectra were obtain using a Raman spectrometer with argon ion laser (λ = 532 nm, LabRAM Hr800).

2.1. Ablation test

To study the ablation behaviors of the ZrC-SiC coatings, an oxyacetylene torch test was performed according to the Chinese standard GJB323A-86 [18]. The pressure and flux were 0.095 MPa and 0.696 L/s for acetylene, and 0.4 MPa and 1.960 L/s for oxygen, respectively. The inner diameter of the gun tip used was 2 mm. During the test, the highest temperature on the ablation center was about 2500 and 3000 °C at distances of 20 and 10 mm from the torch nozzle to the sample surfaces, respectively [2].

The mass ablation rate was calculated according to the expression

$ r_{m}=Δm/t $

where rm refers to the mass ablation rate, Δm refers to the change in mass, and t is the ablation time.

Further, the linear ablation rate was calculated according to the expression

$r_{l}=Δd/t$

where rl refers to the linear ablation rate and Δd refers to the change in thickness.

3. Results and discussion

3.1. Microstructure characterization

Fig. 1 shows the X-ray diffraction (XRD) patterns of the two ZrC-SiC coatings. Clearly, there is no residual Si, Zr, or ZrSi2. The phases in the ZrC-SiC coatings are consistent with those given in the PDF data files of ZrC (65-0973), SiC (89-2214), and 3C-SiC (75-0254). A small amount of C (99-0057) is found in the CH coating. The preferential orientations of the CH and CL coatings are along the (200) and (111) planes, respectively.

Fig. 1.   X-ray diffraction patterns of ZrC-SiC coatings.

Surface SEM images shows that the ZrC ceramic fabricated by thermal evaporation has a typical lamellar structure. Fig. 2(a) shows that the CL coating is mainly composed of large, black SiC particles and a white ZrC phase. The ZrC phase existed mainly in two forms: a zigzag-like laminated ZrC phase and a polygonal-groove-like boundary (Fig. 2(b)-(d)). Inside these ZrC boundaries, the ZrC particles are packed along a specific direction, which can be described using the terrace-ledge-kink model [19]. The boundaries between ZrC and ZrC and that between ZrC and SiC are closely linked. Unlike that observed in the CL coating, the boundaries of the ZrC phase in the CH coating are random and separable (Fig. 3(a) and (b)). However, inside the boundary, the ZrC particles appear to have grown in a similar manner (Fig. 3(c)). In addition to the zigzag-like ZrC phase, the CH coating exhibited an additional terraced growth of the ZrC phase, which was packed in a spiral growth model under the low supersaturation state of Zr in the later stages of evaporation [20,21]. The growth rate of ZrC along the axis of screw dislocations was greater than that in the distant areas. Therefore, the growth of the ZrC particles in the coating was controlled by both the terrace-ledge-kink and spiral growth models.

Fig. 2.   Surface morphology of the CL coating: (a) surface morphologies of ZrC and SiC, (b) ZrC groove, (c) and (d) zigzag-like laminated ZrC phase.

Fig. 3.   Surface morphology of the CH coating: (a) surface morphologies of ZrC and SiC, (b) ZrC groove, (c) laminated ZrC phase, (d) ZrC packed in the terrace-ledge-kink and spiral growth models, (e) and (g) the Zr, C, and Si element distributions shown in Fig. 3(b).

Table 1 shows surface EPMA results for the CL and CH coatings shown in Figs. 2(b) and 3 (b), respectively. These areas in the ZrC grooves of the two coatings are composed of ZrCX (X > 1). As a transition-metal carbide, zirconium carbide tends to exist in the form of non-stoichiometric ZrCx, where x is the Zr/C ratio and is in the range of 0.66-0.99. As such, the zirconium carbide on the coating surface was covered evenly by free carbon or graphite (Fig. 3(e)-(g)) [22]. This phenomenon is similar to that observed in ZrC-SiC coatings prepared on highly pure graphite and high-density, compact C/C composites [15,16].

Table 1   Electron probe microanalysis results of CH and CL coatings.

PositionElement (at.%)
ZrSiC
143.98040.063155.9564
242.97530.145256.8794
335.77490.033164.1919
433.98950.124065.8885
548.55360.075351.3711
644.69170.084255.2241

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An interesting multilayer structure was formed in the ZrC-SiC coatings. Fig. 4 shows that the multilayer structure has two parts. (1) The CH coating is composed of an outer ZrC-SiC mixed layer, a SiC-rich transition layer, and an inner inhomogeneous distribution of the SiC-ZrC infiltration layer (Fig. 4(a) and (b)). The SiC-rich transition layer is formed in situ between Si and the carbon diffused from the matrix (Fig. 4(c)). (2) A C-SiC-ZrC-SiC-C interfacial structure is formed spontaneously inside the coating, and the ZrC ceramic is separated from the carbon matrix by SiC because of the difference in the melting points between the evaporation masterbatches (Fig. 4(d)). Such a special multilayer structure could resolve the mismatch in the coefficient of thermal expansion between the C/C composites and the ZrC coatings and promote interfacial bonding [13].

Fig. 4.   Cross-sectional morphologies of ZrC-SiC coatings: (a) cross-sectional morphology of coating CH, (b) morphology of the outer ZrC-SiC mixed layer and the SiC-rich transition layer in the CH coating, (c) SiC-C boundary, (d) C-SiC-ZrC-SiC-C boundary in the infiltration layer (depth of approximately 500 μm from the coating surface), (e) cross-sectional morphology of the CL coating, (f) morphology of the outer ZrC-SiC mixed layer and SiC-rich transition layer in the CL coating, (g) C-SiC-ZrC-SiC-C boundary in the infiltration layer of the CL coating (depth of approximately 1000 μm from the coating surface).

There was no change in the multilayer structure with a rapid increase in the ceramic content in the CL coating. Fig. 4(e)-(g) shows cross-sectional images of the CL coating after polishing. Similar to that in the CH coating, the thickness of the outer ZrC-SiC mixed layer is approximately in the range of 100-200 μm. However, low-density C/C composites exhibit higher porosity and multiple capillary pores, which are beneficial to the deposition of Zr and Si atoms on the shallow surface of the matrix. Notably, inside the matrix, there is a significant inhomogeneous distribution of the SiC-ZrC infiltration layer throughout the C/C composite, and a mosaic structure is formed with the ZrC/SiC ceramics inlaying the larger pores and holes in the matrix.

3.2. Formation mechanism of ZrC-SiC coating

In our previous studies, the formation mechanism of coatings was studied [15]. Si and then Zr powders were heated to the melting temperature, evaporated, and applied successively to the carbon matrix to formed a ZrC/SiC/C multilayer structure. However, it is interesting that there is no low melt point ZrSi2 or ZrSi on the coating, which is in contrast to other methods such as packing cementation. Combined with the phenomena observed in XRD and SEM, the formation of high purity ZrC-SiC coatings is likely to depend on the thermal decomposition of SiC at 2400 °C. A part of the SiC is thermally decomposes into Si(g) and graphite at 2300 °C [23,24]. A portion of the Si escapes from the coating surface, and the rest of the graphite can react with the subsequent evaporative Zr and Si for a second time. For the high density C/C with a lower pore content, the absorption of Zr(g) and Si(g) on the surface of matrix is limited. The residual graphite content is higher after Si escape. For low density C/C with a higher pore content, the matrix can accommodate more Zr and Si vapors and the decomposed graphite can react completely. This is reason why the C peak can be detected in the CH coating but not in the CL coating (Fig. 1).

To verify this hypothesis, Raman spectra and the graphitization degree of the matrix were obtained by performing Raman spectroscopy. If the content of graphite near the coating is higher than that far away from the coating, this indicates that there is thermal decomposition of SiC in the coating during fabrication. As shown in Fig. 5(a), a disorder-induced (D) peak and a graphite (G) peak are detected at 1332 and 1583 cm-1 in the CH and CL coatings, respectively. The G peak gradually widens and the D peak sharpens along the direction of the coating toward the matrix. The graphitization degree of the carbon substrates (g) was calculated using Eq. (3), where R is the intensity rate of ID/IG [25]. Fig. 5(b) shows that the graphitization degree of the carbon matrix in the two samples decreased from 50 to 200 nm. The graphitization degree of the matrix adjacent to the SiC inner layer was higher than that inside the matrix, which is far away from the coating. The graphite content decreases then increases from the boundary toward the inside of the matrix, unlike the typical concentration gradient of carbon, which deceases along the direction of diffusion. The thermal decomposition of SiC and evaporation of Si intensify the carbon diffusion on the surface of the coating, which is beneficial to the formation of carbide.

$g=1-2.05exp(-2.11R^{-1}) $

Fig. 5.   Raman spectra and graphitization degree of the C/C matrix near the CH and CL coatings: (a) Raman spectra of the carbon matrix adjacent to the SiC inner layer at an interval of 50 nm, and (b) graphitization degree of the carbon matrix shown in Fig. 5(a).

3.3. Ablation properties

The ZrC-SiC coatings show excellent ablation properties after oxyacetylene ablation at 2500 °C. As listed in Table 2, the mass and linear ablation rates are 0.03 ± 0.01 mg·cm-2·s-1 and -0.15 ± 0.02 μm/s, respectively, for the CH coating and 0.04 ± 0.01 mg·cm-2·s-1 and -0.14 ± 0.03 μm/s, respectively, for the CL coating after ablation for 60 s.

Table 2   Mass and linear ablation rates of ZrC-SiC coatings.

SampleAblation temperature (°C)Ablation time (s)Mass ablation rate (mg·cm-2·s-1)Linear ablation rate (μm·s-1)
CH2500600.03 ± 0.01-0.16 ± 0.02
CL2500600.04 ± 0.01-0.14 ± 0.03
CL3000300.03 ± 0.01-0.13 ± 0.02
CL300060-0.10 ± 0.02-1.33 ± 0.04
CL3000200-0.46 ± 0.15-1.00 ± 0.04

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The XRD patterns shown in Fig. 6 indicate that SiO2 is in a glass stage and is difficult to detect and that the surface ZrC is completely oxidized to ZrO2 [26]. The mass loss comprises the mass gain from the oxidations of ZrC and SiC, mass loss from the evaporation of SiO2, and mechanical denudation of the oxide [5].

Fig. 6.   X-ray diffraction patterns of the ZrC-SiC coatings after ablation.

Fig. 7 shows the surface morphology of the CH coating after 60 s of ablation. The center ablation region is composed of large gray SiO2 and white ZrO2 particles. Moreover, the boundary of ZrC dehisced during oxidation and formed cracks. Fig. 7(b) and (c) shows magnified views of the cracks. Numerous micropores are observed near the crack. The cracks form paths for oxygen permeation, and the CO gas produced by the inner carbide diffuses and forms micropores (Eqs. (4)-(7)) [[27], [28], [29]].

$ZrC(s)+xO_{2}(g)=ZrC_{1-x}O_{x}(s)+xCO(g)$

$2ZrC_{1-x}O_{x}(s)+(2-x)O_{2}(g)=2ZrO(s)2+(2-2x)C(s)$

$2SiC(s)+3O_{2}(g)=2SiO^{2}(s)+2CO(g)$

$2C(s)+O_{2}(g)=2CO(g)$

Fig. 7.   Morphologies of central ablation surface regions of the CH coating after 60 s of ablation at 2500 °C: (a) central ablation region, and (b) and (c) morphology of ZrO2 near the crack.

Besides the formation of micropores due to CO gas emission, there is another reason for the pores on the surface of ZrO2. Fig. 8 shows the surface morphology of the CL coating ablated at 2500 °C for 60 s. After the evaporation and mechanical denudation of SiO2, holes are formed on the surface of the coating, and numerous micropores are found near the SiO2 holes. The dissipative evolution process of SiO2 (Fig. 8(c)-(e)) shows that ZrO2 particles are protected by the melting liquid SiO2 but are gradually exposed to the ablation flame with the consumption of SiO2. The ZrO2 near the SiO2 holes suffered from the combined effect of SiO2 volatilization and CO gas emission. These SiO2 holes were another important path for oxygen erosion.

Fig. 8.   Surface morphology of the CL coating after 60 s of ablation at 2500 °C: (a) central ablation region, (b) and (c) morphologies of ZrO2 near the SiO2 hole, and (d) and (e) transition region of the CL coating.

Fig. 9(a) and (b) shows the cross-sectional morphology of the CH coating after 60 s of ablation. The outer ZrC-SiC mixed layer is oxidized, forming a loose ZrO2 layer, and the inner SiC pinned with the ZrC layer acts as an actual oxygen erosion barrier in the later stages of ablation. However, SiC is largely consumed and continuously oxidized; thus, it melts and evaporates via the holes and cracks in the outermost porous ZrO2 layer. Under the combined effect of the surface acetylene flame and internal gas-liquid flow of SiO2 and CO, the CH coating is vulnerable to mechanical erosion and has limited protection capability.

Fig. 9.   Cross-sectional morphologies of the ZrC-SiC coatings: (a) and (b) cross-sectional morphologies of the CH coating, (c)-(e) cross-sectional morphologies of the CL coating.

However, the CL coating shows a superior anti-ablation property after 60 s of ablation. As shown in Fig. 9(c)-(e), the ablation erosion damage to the inner sides of the coating is limited. This can be attributed to the restricted penetration depth of oxygen owing to the high internal content of the ceramic. SiC continues to liquefy and form an oxide protective film, thereby establishing a dynamic balance between protection against liquefaction, volatilization, and mechanical denudation.

The ablation temperature (2500 °C) is lower than the melting point of ZrO2 (2700 °C) and much higher than that of SiO2 (1780 °C) [30]. Hence, in the early stages of ablation, only the liquefaction of SiO2 was effective, and the defects were sealed, compared with the loose scale of ZrO2. Therefore, higher-temperature oxyacetylene flame ablation (3000 °C) was used to further test the properties of the CL coating. It was found that ZrC, instead of SiC, acts as the actual oxidation barrier during a long-duration ablation process at 3000 °C.

Fig. 10 shows the structural evolution of ZrO2 on the coating surface after ablation for 60 s at 2500 °C and for 30 s and 60 s at 3000 °C. The laminated ZrC could have oxidized and formed a monolithic dense ZrO2 after ablation at 2500 °C; however, cracks are easily formed on its boundary (Fig. 10(a) and (b)). With the increase in the ablation temperature, the boundary of ZrO2 gradually softens, and the inner ZrO2 particles gradually melt and form monolithic ZrO2 with several micropores (Fig. 10(c) and (d)). After ablation for 60 s at 3000 °C, a fused and monolithic ZrO2 without obvious micropores formed at the ablation center, with the boundaries appearing to cohere with each other.

Fig. 10.   Structural evolution of ZrO2 on the surface of the CL coating after ablation: (a) and (b) ablation for 60 s at 2500 °C, (c) and (d) ablation for 30 s at 3000 °C, and (e) and (f) ablation for 60 s at 3000 °C.

The cross-sectional morphology of the CL coating after ablation for 30 s at 3000 °C (Fig. 11) shows that the ablative damage on the coating surface is limited. SiC on the coating surface is rapidly oxidized and consumed, forming holes for oxygen erosion. Micropores are easily formed near the SiO2 holes, while a dense melting ZrO2 region is formed away from the holes (Fig. 11(b)-(d)). This ablation behavior is similar to that observed at 2500 °C. Compared with the discontinuous distribution of the ceramics, shown in Fig. 4, a channel of evaporated Zr and Si atoms entering the pores and holes of the C/C matrix can be clearly observed in the cross-sectional morphology of the coating, parallel to the carbon fiber bundles (Fig. 12(a)). The SiC-C interface is formed and remains stable, even when large quantities of ceramics are introduced into the matrix (Fig. 11(e)); this confirms the uniqueness of the thermal evaporation method.

Fig. 11.   Cross-sectional morphology of the CL coating after ablation for 30 s at 3000 °C: (a) distributions of ZrC and SiC ceramics in the C/C matrix, (b) morphology of the outer ZrC-SiC mixed layer, (c) and (d) cross-sectional morphologies of SiO2 holes and distribution of micropores in ZrO2, and (e) C-SiC-ZrC-SiC-C boundary in the infiltration layer.

Fig. 12 shows the cross-sectional morphology of the CL coating after ablation for 200 s at 3000 °C. The mass ablation and linear ablation rates are only -0.46 ± 0.15 mg·cm-2·s-1 and -1.00 ± 0.04 μm·s-1, respectively, which is attributed to a dense ZrO2 layer on the coating surface. The thickness of the outer oxide layer is approximately 20 μm; this layer is closely attached to the inner carbide (Fig. 12(a) and (b)). Most of the holes and cracks heal during the long-duration, high-temperature ablation. Some residual holes are formed in the internal pores because of the high temperature, pressure, and velocity of the combustion flame.

Fig. 12.   Cross-sectional morphology of the CL coating after ablation for 200 s at 3000 °C: (a) and (b) cross-sectional morphologies of the CL coating, (c) dense ZrO2 layer on the coating surface, and (d) inner holes in the oxide layer.

It can be speculated that in the densification process of ZrO2 at 3000 °C, a significant amount of oxide is volatilized and undergoes severe mechanical denudation. Fig. 13 shows the surface morphology of the CL coating after ablation for 200 s at 3000 °C. Although ZrO2 finally fused to form a dense oxygen protective film, the thickness of the coating decreases dramatically after the apparent mechanical denudation, and some of the C fibers near the coating are exposed. As shown in Fig. 13(b), the surface fibers are chemically eroded. Nevertheless, the original micropores and holes on the coating surface disappear and heal (Fig. 13(c)-(e)).

Fig. 13.   Surface morphology of the CL coating after 200 s of ablation at 3000 °C: (a) central ablation surface regions, (b) eroded carbon fibers exposed on the coating surface, and (c)-(e) dense ZrO2 surface with healing holes.

3.4. Model and mechanism of ablation

The ablation behavior of the ZrC-SiC coatings at 2500 °C can be attributed to two factors. (1) Oxidation and cracking in the boundary of ZrC owing to thermal stress. Micropores are easily formed near the cracks, and the coating is not dense overall because the ablation temperature is lower than the melting point of ZrO2 (Fig. 14(a)) [31]. Near the cracks, the carbide inside the coating was oxidized, and the CO gas diffused and hindered the densification process [[27], [28], [29]]. However, these cracks would be sealed by extending the ablation time or temperature. (2) Consumption of SiC. The oxidation and liquefaction of SiO2 are favorable to the reduction in the temperature and provide a positive oxide protective layer to prevent the oxidation of the inner carbide in the initial stage, thereby reducing the mass loss within a short ablation time [32,33]. However, with prolonged ablation, potholes would be formed; these potholes are the main cause for failure of the coating on the higher-density matrix under a limited supply of ablative components. Loose scales of ZrO2 are easily formed near the potholes and can undergo severe mechanical denudation (Fig. 14(b) and (c)).

Fig. 14.   Schematic of ablation model of the ZrC-SiC coatings: (a) structural evolution of ZrO2 during ablation, (b) formation of SiO2 holes, and (c) and (d) structural evolutions of the CL coating during long-duration ablation at 3000 °C.

Nevertheless, at higher temperatures, the ablation analysis of the ZrC-SiC coatings at 3000 °C shows the formation of a large number of porous oxide scales at the early stages of ablation because of the oxidation of carbide, liquefaction, volatilization of SiO2, and the mechanical denudation of porous ZrO2. The ablation properties of the CL coating strongly depend on its mosaic structure, which provides a continuous supply of ablative components. When the surface SiC is exhausted, the residual ZrO2 melts, densifies, and acts as a protective oxide layer. The cracks and holes on the surface and inside the coating formed at the early stages could be compressed and sealed because of the high temperature, pressure, and velocity of the combustion flame (Fig. 14(d).

4. Conclusions

(1)ZrC-SiC coatings were fabricated on low-density, porous C/C composites via thermal evaporation and an in-situ reaction. A C-SiC-ZrC-SiC-C multilayer structure was formed inside the matrix. Zr and Si atoms deposited onto the multiple capillary pores on the shallow surface of the porous C/C composites and formed a mosaic structure.

(2)The graphitization degree of the matrix increased because of the decomposition of SiC at 2400 °C and facilitated the formation and purification of the carbides.

(3)Severe dissipation of the carbides occurred on the surface of the ZrC-SiC coating during ablation at 3000 °C; this was accompanied by severe mechanical denudation. The residual ZrO2 could be densified to form a compact oxide layer because of the high temperature, pressure, and velocity of the combustion flame. Furthermore, the pores and holes on the surface and inside the coating were compressed and sealed.

Acknowledgements

This work was supported by National Science Foundation of China (No. 51405522) and the self-fund of State Key Laboratory for Powder Metallurgy (PM-CSU-2015-03).


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