SiC/(SiC+glass)/glass coating for carbon-bonded carbon fibre composites

Abstract

An oxidation resistant coating for highly porous carbon-bonded carbon fibre composites (CBCFs) was prepared by slurry method. The compositionally and structurally gradient coating consists of a carbon fibre reinforced SiC inner layer, a SiC with CaO–MgO–Al₂O₃–SiO₂ (CMAS) glass inter layer, and a dense CMAS glass outer layer. The effectiveness of the S–SG–G coating in protecting the CBCFs from degradation at 1000 °C in air has been evaluated. The results indicated that the coating could protect CBCFs from oxidation for 100 min with the weight loss of 20.7%. In contrast, the uncoated CBCFs were completely combusted after 50 min. Good oxidation resistance of the coating is mainly attributed to high temperature stability of SiC, CMAS glass sealant, and good adhesion within the coating layers and with the substrate.

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1 Introduction

Carbon-bonded carbon fibre composites (CBCFs) are a special kind of carbon/carbon (C/C) composites with low densities of 0.1–0.5 g cm⁻³ and high porosities of 70–90%.^1,2 Investigations into the microstructures,^1,3 mechanical properties,^2,4 and thermal properties³ of the CBCFs have been reported. Compared to the traditional C/C composites with densities of 1.4–1.8 g cm⁻³,^2,5 their lightweight structure, low thermal conductivity, and high temperature capability can better satisfy the development of thermal insulation for new spacecrafts.^6–8 However, application of the CBCF materials is significantly limited for the oxidation of carbon above 450 °C.^9–11

Up to now, most research has been focused on improving oxidation resistance of traditional C/C composites by coating method.^12–14 Preparation of porous-free coating on the highly porous CBCF substrate is difficult. The coating must provide complete surface protection from penetration of oxidizing species, and remain adherent to the substrate during prolonged high temperature exposures. F. Smeacetto, et al.¹⁵ have prepared a silica-based glass containing Y₂O₃ and Al₂O₃ coating on the surface of Si-infiltrated CBCFs, but oxidation behaviours of the coated composites are unknown for the lack of oxidation test. B. Xu, et al.⁸ have prepared a coating consisted of ZrB₂–SiC whiskers-borosilicate inner layer and ZrB₂–MoSi₂–SiC whiskers-borosilicate outer layer on the surface of ZrB₂-modified CBCFs with the density of 0.77 g cm⁻³. Oxidation resistance was achieved at the expense of added weight. Therefore, the protective coating system for the CBCFs needs to be further investigated.

Application of a multilayer coating with different function of each layer is considered as an effective way to prevent carbon from oxidation.^16–18 SiC is an ideal inner layer candidate due to its good physical and chemical adaptability with carbon, and excellent oxidation resistance.¹⁹ It is difficult to fabricate dense SiC using normal sintering method due to low vacancy diffusion for such covalent compounds. A promising technique is reaction-sintered method. During the sintering process, liquid Si infiltrates into a porous green mixture of SiC and C by capillary action. An in situ reaction between the Si and the C occurs, producing a secondary SiC phase that then bonds to the original SiC.²⁰ Among all the thoughts rise two problems: (i), the SiC layer is difficult to be fully densified on such a highly porous surface of CBCFs; and (ii), stress cracks are often produced in the layer due to the mismatched thermal expansion coefficient (CTE) between different phases. Application of a glass sealant has been proved to be an effective method for addressing these issues.¹⁷ CaO–MgO–Al₂O₃–SiO₂ (CMAS) glass has meet most of the requirements as an ideal defect sealant due to its high melting point, appropriate wetting properties and viscosity, high corrosion resistance, and good oxidation resistance at elevated temperatures.^21,22 The CMAS glass with the CTE of 3.9 × 10⁻⁶ °C⁻¹, similar to that of SiC 4.5 × 10⁻⁶ °C⁻¹,²³ was chosen.

In the present work, a multilayer coating system composed of a SiC (S) inner layer, a SiC–CMAS glass (SG) inter layer, and a CMAS glass (G) outer layer for CBCFs was designed. The S bond layer with a carbon fibre reinforced SiC structure is intended to improve the bonding between the outer layer coating and the CBCF substrate. The SG layer serves as a diffusion barrier which can delay transport of oxidizing species to the substrate. The G layer serves as a dense structure as well as a glass sealant for defects (holes or cracks). The S, SG and G layer was prepared by slurry painting followed by a heat treatment. The microstructure and oxidation behaviours of the S–SG–G coated CBCFs were investigated in detail.

2 Experimental procedures

Cubic CBCF specimens with the dimension of 15 mm × 15 mm × 15 mm used as substrates were cut from bulk anisotropic composites with densities of 0.23 ± 0.05 g cm⁻³ and porosities of 82–85%. The specimens were ultrasonically cleaned with ethanol and dried at 100 °C for 6 h. Subsequently, they were hand polished with 800 grit SiC paper to form a thin carbon debris layer on the surface to prevent over impregnation of coating slurry into the highly porous CBCF substrate.

Commercially available Si powders (9–11 μm, 99.99%, Zhejiang Kaihua Yuantong Silicon Industry Company, China), SiC powders (6–8 μm, Showa Denko K. K, Japan), CMAS glass powders (200 mesh, Shaanxi University of Science and Technology, China), and phenolic resin (no. 2130, Xi'an Resin Factory, China) were used in this study. The slurry of S, SG and G layer was prepared by ball milling of the ingredients in Table 1 for 6 h, respectively. S slurry was painted twice on the surface of CBCFs using a hard brush. The painted samples were subsequently dried at 80 °C overnight to remove the residual ethanol. Considering high porosities of the substrate, phenolic resin with ethanol was chosen as the dispersant of S slurry. Phenolic resin acting as “glue” can bond Si and SiC particles together, enabling a slow and homogeneous infiltration into the porous fibre webs. Then the samples were heated up to 1450 °C for 2 h in Ar atmosphere to form S layer. SG slurry was painted on the surface of the S layer using the same painting method. The painted samples were heated in a tubular oven at 1250 °C for 1 h in Ar atmosphere. G layer was prepared on the surface of the SG layer using G slurry in the same way as the SG layer.

Table 1 Composition of S, SG and G layer slurry

Layer slurry	Composition (wt%)
	SiC	Si	Resin	Glass	Ethanol
S	35	17.5	17.5	—	30
SG	9	—	—	21	70
G	—	—	—	30	70

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The phase composition of the coating was investigated by X-ray diffractometry (XRD, X'Pert PRO). The morphologies and element distribution of the multilayer coating was analysed using scanning electron microscopy (SEM, VEGAII XMU) equipped with an energy dispersive spectroscopy (EDS) detector. The detailed differential scanning calorimetric (DSC) studies of the CMAS glass powders were carried out using the thermal analyser instrument SDTQ600 at a heating rate of 5 °C min⁻¹ in Ar atmosphere. Isothermal oxidation test was conducted to investigate oxidation behaviours of the coated samples at 1000 °C in an electrical furnace. After oxidation for a certain time at the set temperature, samples were taken out of the furnace and cooled naturally to room temperature for weighing. The samples were measured by a high accuracy electronic balance (BP221S) with a sensitivity of ±0.1 mg. Then they were put into the furnace again for the next thermal cycle. Weight changes of the samples were recorded as a function of oxidation time. Average value from five samples for each kind treated under the same condition was used to improve the accuracy and reliability of the data.

3 Results and discussion

Appropriate preparation temperature for SG and G layer containing CMAS glass was determined by DSC measurements of the glass powders (Fig. 1). During non-isothermal heating of the glass, a sequence of thermal effects, e.g., structural relaxation, glass transition, softening and nucleation, crystallization, and melting occur over the transformation temperature range. A sharp endothermic peak centred at 1182.4 °C in the DSC curve is associated with the softening of the glass. No apparent exothermic peaks related to the crystallization can be observed. The measuring results reveal a glass transition temperature (T_g) at 1141.7 °C, a softening temperature (T_s) at 1182.4 °C and a melting point (T_m) at 1315.1 °C. Therefore, the glass was heated up to 1250 °C that is above the softening point during the preparation of SG and G layer.

Fig. 1 DSC curve of CMAS glass powders.

Fig. 2 shows surface XRD patterns of the CBCF substrate, S, SG, and G layer. The diffraction peak at 25.8° identifies the graphite phase in the CBCFs. In the S layer, diffraction peaks of α-SiC and β-SiC are detected along with the graphite. Decrease of the graphite peak intensities indicates that most of the carbon fibres on the surface of CBCFs are covered. In the SG layer, no graphite phase is found except for the SiC crystals, showing that surface of the CBCFs is completely covered. A typical amorphous XRD pattern is observed in the curve of G layer, indicating the amorphous nature of CMAS glass. The results agree well with the designation of the coating system.

Fig. 2 Surface XRD patterns of the CBCF substrate, S, SG, and G layer.

The morphologies of the CBCFs, S, SG, G layer, and area scanning of the G layer are shown in Fig. 3. Fig. 3a and e show the quasi-layered carbon fibres with diameters of 9–13 μm are randomly distributed. A large number of holes constructed by the bridging fibres need to be densified. Fig. 3b and f show the images of S layer. The structure seems loose, but sizes of the holes in the surface are largely diminished. No visible cracks are observed in this layer because of the framework structure of carbon fibres and solid skeleton effect of the original SiC particles. Packed SiC particles with diameters of 8–17 μm have course surfaces due to the newly formed SiC from thermally evaporated Si reacted with carbon debris and pyrolyzed resin. After the application of SG layer, the surface becomes denser and smoother, as shown in Fig. 3c and g. SiC further deposits in the holes of the S layer. The melting glass flows and connects the solid particles to be an integral structure. From Fig. 3d and h, one can see a fully dense, continuous and defect-free surface. Defects in the SG layer are completely sealed by the sticky flow of the glass melted objects. No traces of carbon fibres and SiC particles can be observed. Fig. 3i–l shows area scanning of the G layer surface. It can be seen O, Si, Al, and Ca elements are uniformly distributed in the layer. Traces of Mg were also detected. B₂O₃ could be contained in the glass but element B can hardly be detected by EDS detector.

Fig. 3 Surface morphologies of (a) CBCF, (b) S, (c) SG, (d) G layer; (e), (f), (g) and (h) are high magnification, respectively; (i–l) elemental area scanning of (h).

The typical micrographs of cross-section and elemental line scanning of S–SG–G coated CBCFs are shown in Fig. 4. It can be seen the coating is 140–180 μm in thickness and no transverse cracks are observed (inset in Fig. 4a). The concentration of Ca, Mg, Al, Si and O is high from the surface to ∼57 μm, indicating the thickness of G layer. Based on the analysis of elemental distribution, the thickness of SG layer and S layer is estimated to be ∼49 μm and ∼47 μm, respectively. Actually, there is no obvious interface, and distinct composition boundary to distinguish these three layers. The density of G layer mainly depends on the molten glass, while that of the S layer depends on the compaction of SiC particles. G layer is more compact than the S layer. During heat treatment, the melted glass with appropriate viscosity in G and SG layer could partially infiltrate into the porous S layer, especially in the interfacial region, thus resulting increase of sintering densification and high combination interface between the S and SG layer. In the S layer, SiC powders are uniformly dispersed in the top webs of carbon fibres, as shown in Fig. 4d. This is a carbon fibre reinforced SiC ceramics structure, which could enhance the bonding between the coating and the substrate by various toughening mechanisms, including fibre–matrix debonding, fibre bridging and pull-out, and crack deflection to dissipate energy.²⁴ A gradient distribution of porous structure from G to SG layer, and a carbon fibre reinforced S layer along the thickness direction of the coating is achieved. Such a compositionally and structurally gradient coating has good adhesion within the coating layers and with the substrate.

Fig. 4 (a) Cross-sectional microstructure of S–SG–G coated CBCFs; (b) elemental line scanning of the coating; interface between (c) G and SG layer, (d) S layer and CBCF substrate.

Isothermal and cyclic oxidations of the uncoated and coated CBCFs at 1000 °C in static air were performed to evaluate effectiveness of the coating in protecting CBCFs from degradation, as shown in Fig. 5. CBCFs with no protective coatings were completely combusted in 50 min. S coated CBCFs showed improved oxidation resistance. But the weight loss has reached 70% after 100 min. Low protective abilities of the porous S layer coating can be explained by: (i) defects (pores and holes) in the layer acting as oxygen diffusion channels result in the degradation of carbon; (ii) oxidation of exposed carbon fibres on the surface leads to reaction spreading into internal substrate by creating new channels for oxygen diffusion. Therefore, a dense layer is necessary to isolate carbon from oxygen. After the application of SG and G layer, oxidation resistance of CBCFs was significantly improved. The coating can protect CBCFs for 100 min with a weight loss of 20.7 wt% accompanying with 4 thermal cycles between 1000 °C and room temperature. Diffusion rate of oxygen through the coating and the erosion rate of carbon were decreased significantly.

Fig. 5 Isothermal oxidation curves of the coated CBCFs composites at 1000 °C in static air.

To further investigate oxidation mechanisms of the coated samples, surface morphologies and EDS analysis of the oxidized coating are shown in Fig. 6. A glass layer containing many bubbled shapes and pores is observed on the surface of the composites. The bubbled shapes are formed from the volatilization of B₂O₃ contained in the glass, and gaseous CO and CO₂ generated from the oxidation of SiC and carbon. The gas pressure reaches a high value with oxidation time, exceeding the surface tension of the bubbles. Gases bubbling out results in the formation of pores (Fig. 6b). The oxidized G layer is composed of Ca, Mg, Al, Si and O elements, and the content of Ca and Mg is relatively low (Fig. 6d). Fig. 7a shows cross-sectional SEM image of coating after oxidation test, it can be seen the G layer presents coarse with microscopic collapsed holes, implying that damage of the coating is dominant during this process. Due to the oxidation of coating, some voids can be seen in the coating. The oxidized SG and G layer became less dense in comparison with that before oxidation test. Although the test temperature is not high enough for SiC particles oxidized to silica glass, the holes can hardly penetrate through the inner layer because the complex coating structure could stop crack propagation, and the fluid CMAS glass could fill some disfigurements in the coating. Fig. 7b shows line scanning analysis of the oxidized coating, Ca, Mg and Al elements can be detected in the SG and S layer, indicating that the glass has partially infiltrated into the inner layer. It is inevitable that the defects form due to the CTE mismatch between SiC (4.45 × 10⁻⁶ °C⁻¹) and CBCFs (1.19 × 10⁻⁶ °C⁻¹).²⁵ Some of the formed cracks could be self-sealed by the viscous flow of glass. With the increase in the number of thermal cycles, sizes and numbers of defects increase. The defects provide an oxygen diffusion path resulting in slight oxidation of carbon. The carbon fibre networks remain intact after oxidation (Fig. 7a), indicating good oxidation protective ability of the coating.

Fig. 6 S–SG–G coating after oxidation at 1000 °C for 60 min: (a) surface morphologies; detail showing (b) defects, (c) glass; (d) EDS spectra indicated in (c).

Fig. 7 (a) Cross-sectional SEM image of the S–SG–G coated CBCFs after oxidation at 1000 °C for 100 min; (b) line scanning elemental analysis marked in (a).

Surface XRD pattern of the S–SG–G coated CBCFs after oxidation test is shown in Fig. 8. Diffraction peak of carbon can be detected, which may result from the volatilization of G layer and concomitant decrease of the coating thickness, leaving some carbon fibres near under the surface. Combined with the EDS analysis (Fig. 6d), it can be determined that anorthite (CaAl₂Si₂O₈) and diopside (MgCaSi₂O₆) crystalline phases formed during the oxidation process. Anorthite and diopside are thermally stable phases with good oxidation resistance at high temperature. Their existence could improve the stability of glass and oxidation resistance of the coating. In addition, the formed glass–ceramics have a similar CTE with the value of that before oxidation for the CMAS glass contains low content of CaO and MgO, which are known to increase the thermal expansion.²⁶ This indicates the coating has good ability to avoid thermal stress during repeated temperature variations.

Fig. 8 Surface XRD patterns of S–SG–G coating after oxidation test.

Based on the above analysis, oxidation mechanism of the S–SG–G coating can be summarized as follows:

B₂O₃(s) → B₂O₃(l) → B₂O₃(g) (1)

SiC(s) + 2O₂(g) → SiO₂(s) + CO₂(g) (2)

2SiC(s) + 3O₂(g) → 2SiO₂(s) + 2CO(g) (3)

SiC(s) + O₂(g) → SiO(g) + CO(g) (4)

2C(s) + O₂(g) → 2CO(g) (5)

C(s) + O₂(g) → CO₂(g) (6)

G layer is damaged owing to the thermal evaporation of B₂O₃ in the CMAS glass (reaction (1)), forming defects for oxygen entering the inner SG and S layer. SiC would consume some oxygen by forming SiO₂ and SiO (reaction (2)–(4)), while the S layer could hardly become stable and densified silica glass phase because the test temperature is not high enough for complete oxidation of SiC. Viscous flow of the glass could heal some defects in a short time. As the oxidation time and thermal cycles increased, various kinds of gaseous byproducts formed such as CO, CO₂ and SiO (reaction (4)–(6)). These gases would be suppressed by the coating layers first, existing in the form of bubbles. Vapour pressure increases with oxidation time and finally leads to the bursting of the bubbles. This results in the weight loss and formation of pores on the surface of the coating. However, the multiple layers with complex structure offer difficult and tortuous paths for oxygen to diffuse into the internal CBCF substrate. Only a small amount of oxygen finally attack carbon, so no obvious oxidation region can be seen in the CBCF substrate, which indicates good protective ability of the coating.

4 Conclusions

This work proposed a simple and low-cost method to produce compositionally and structurally gradient oxidation resistant S–SG–G coating for the highly porous CBCF materials. The refractoriness of the coating was enhanced by SiC crystals with excellent oxidation, erosion resistance, and similar CTE to that of CMAS glass. The glass of suitable properties was applied as a sealant for the preparation of pore-free layer by heating at the temperature above the glass softening point. The multiple layer coating possesses good oxidation protective ability, which could protect CBCFs for 100 min with a weight loss 20.7 wt% accompanying with 4 thermal cycles between 1000 °C and room temperature. The carbon fibre networks remain intact after the oxidation test. Overall, oxidation resistance of the CBCFs is significantly improved by the multilayer coating, indicating the potential application of the CBCFs used as thermal insulation in the air.

Acknowledgements

The authors appreciate the financial support of the National Natural Science Foundation of China (51472198), the National Basic Research Program of China (973 Program) granted No. 2015CB655200 and the Fundamental Research Funds for the Central Universities (No. xkjc2014009).

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