Preparation of aluminosilicate fiber-SiBOC composite foams for thermal protection applications

Abstract

In this study, aluminosilicate fiber-embedded SiBOC matrix composite foams are realised by a simple, cost-effective and easily scalable method. Aluminosilicate wool is transformed into a compressible preform using polyvinyl alcohol as a binder. The preforms are impregnated with methylvinylborosiloxane (MVBS), followed by heat treatment at 1400°C in an inert atmosphere. The composite foam's density, compressive strength and thermal conductivity are modulated in the ranges 0.71–0.46 g/cc, 0.79–0.39 MPa and 0.21–0.13 Wm⁻¹K⁻¹, respectively, by varying MVBS concentration from 29 to 54 wt.%. While the incorporation of aluminosilicate fibers was observed to lower the thermal conductivity of the composite foams, it did not contribute to their compressive strength. Exposure of the foams to air at 1300°C for 90 min didn't change their density, while it increased the compressive strength by up to 46% due to improved fiber-matrix adhesion and rounding of cracks.

Keywords

Ceramic fiber reinforced ceramic matrix composite foam aluminosilicate fiber preform borosiloxane oxidation resistance thermal protection system

Introduction

Ceramic foams offer a plethora of several unique properties which are normally extremely difficult in dense ceramics, like high thermal shock resistance, ultra-low thermal conductivity, very low thermal expansion and ease of machining [1 -6]. These properties are tailored by varying the foam density and cell size. Tailoring the cell size and structure of foams to some extent is possible by selecting appropriate forming processes such as replica technique, sacrificial template method, self-foaming, freeze-drying and additive manufacturing [5, 7 10]. Adding fillers to the preceramic polymers or the ceramic powder suspensions used to prepare these foams is another simple but efficient way of tailoring their properties [11 16]. Recently, some studies have explored the idea of using short fiber as passive fillers, i.e. the fillers which do not react with the matrix but contribute to improving the strength and thermal insulation. Wu et al. studied short YSZ fiber reinforcements in porous γ-Y₂Si₂O₇ ceramic by a foam-gel casting followed by freeze drying, obtained high compressive strength (1.35 MPa) and extremely low thermal conductivity (0.09 Wm⁻¹k⁻¹) [17]. Dong et al., in their study on short YSZ fiber-reinforced porous YSZ ceramics, observed that YSZ fiber obstructed densification during the sintering process leading to higher porosity. The reinforcing effect reached its optimum with 10 wt.% YSZ fiber addition and the fiber addition significantly changed the fracture mode of the porous ceramics from brittle to quasi-ductile due to fiber pull-out and crack deflection. Lang et al. reported porous Yttria-stabilised zirconia (YSZ) ceramics reinforced by four kinds of short fibers (mullite, aluminosilicate, Al₂O₃ and YSZ fibers) and observed significant improvement in the short Al₂O₃ fiber-reinforced porous YSZ ceramics, in terms of decrease in their thermal conductivity with increase in compressive strength [11, 18]. In these studies, short fibers are processed along with the slurry constituents by ball-milling, such that the fiber length is typically smaller than the thickness of the cell walls or struts. On the other hand, continuous fiber-reinforced ceramic matrix composites have been extensively studied. They are well-known for exhibiting the high toughness and strength required for thermo-structural applications [19, 20]. However, there is not much literature on using relatively long fibers as fillers in ceramic foams. The silica tile used in Space Shuttle is one such example of the success of long fiber-based thermal protection systems [21, 22].

Among the ceramic precursors investigated worldwide, siloxanes and their derivatives are gaining attention due to their versatile synthesis processes, cost-effectiveness and resistance to high-temperature oxidative degradation (at least up to 1300°C) [4, 23 -25]. Compared to vitreous silica-based materials, primarily used for thermal protection systems, borosiloxane-derived SiBOC materials demonstrate better strength, stiffness and hardness, as well as superior glass transition and crystallisation temperatures [26 30]. The glass phase obtained by heat treatment of borosiloxane between 900 and 1600°C has a complex structure comprising nanocrystals of β-SiC and turbostratic graphite embedded in an amorphous SiBOC matrix [27, 30, 31]. In view of this, the SiBOC glass has also been referred to as SiC/SiBOC nanocomposite [28, 32, 33].

Commercial oxide fibers are primarily available as silica, alumina and their combinations with varying proportions and crystalline fractions. With the technological advancements in the development of mixed polycrystalline oxide fibers, there is renewed interest in using them as reinforcements for oxide-oxide ceramic matrix composites, metal-matrix composites and porous fiber-reinforced ceramics [19, 20, 34 -36]. These high-performance fibers, most popular among them NEXTEL™ 720 and NEXTEL™ 650, have good strength, microstructural stability and long-term creep resistance at temperatures above 1100°C, which make them suitable for applications such as in thermal protection systems for combustion chambers in aircraft engines [35]. The composite foams are realised by forming a reversibly compressible PVA-bonded ceramic fiber preform and subsequent impregnation with MVBS, followed by heat treatment at 1400°C in an inert atmosphere. The method is novel, simple and capable of realising composite foams of tailorable density.

Experimental

Materials

Aluminosilicate wool (Superwool® XTRA Unifelt), having a melting temperature of 1650°C and a bulk density of ∼0.22 gcm⁻³, was obtained from M/s Morgan Advanced Materials, Chennai, India. Poly(vinyl alcohol) (99%, Mw 89,000–98,000, Sigma Aldrich) was used as a binder to realise the ceramic fiber preform. MVBS, the precursor for SiBOC, was synthesised in-house from analytical reagent grade triethoxy(vinyl)silane (C₈H₁₈O₃Si, assay ≥ 97%, Evonik Industries, Germany), triethoxy(methyl)silane (C₇H₁₈O₃Si, assay ≥ 99%, Sigma-Aldrich Chemie GmbH, Germany) and boric acid (assay ≥ 99.5%, Merck Specialties, Mumbai, India) using the procedure reported earlier [30, 37 -40]. To synthesise MVBS, the three raw materials, viz., boric acid, triethoxy(vinyl)silane and triethoxy(methyl)silane were added in a glass reactor in the mole ratio of 1:0.5:1.5 and stirred for 5 h under reflux in an oil bath maintained at 80–85°C. The relevant reaction scheme is represented in Figure 1.

Figure 1

Reaction scheme of synthesis of methyl vinylborosiloxane (MVBS).

During the synthesis of MVBS, ethanol is produced as a byproduct, and thus the as-synthesised MVBS is an ethanol solution containing 37 wt.% MVBS oligomers, having a solution viscosity of ∼3 cps, and weight average molecular weight ( ) of ∼2000 gmol⁻¹ [26, 27]. The MVBS solutions of higher concentrations were obtained by distilling off some quantity of ethanol at 80^oC under reduced pressure, while MVBS solutions of lower concentrations were obtained by diluting the as-synthesised solution with distilled ethanol. The MVBS content was estimated by drying ∼50 g of the solution in an air oven at 175°C for 2 h.

Fabrication of ceramic fiber SiBOC composite foams

The scheme of realising the ceramic fiber-SiBOC composite foam is given in Figure 2(a). Firstly, the aluminosilicate ceramic fibers, received in the form of a woolen blanket, were converted into a disc-shaped preform. For each preform, 75 g ceramic fibers, as lumps of length 2–3 cm, were manually pulled from the woolen blanket and dispersed in 3 L of 30 g/L aqueous polyvinyl alcohol (PVA) solution. The ceramic wool was dispersed in the PVA solution by mechanical stirring and was poured into a filter-pressing set-up shown in Figure 2(b). Further, it was consolidated by filtration of the PVA solution through a carbon cloth placed on a Buchner funnel. A gentle hand pressing with a piston was given towards the end of the consolidation process. The PVA solution collected in the filter flask was reused to prepare subsequent samples. The filter-pressed bodies were dried in ambient condition for 24 h, followed by heating in an air oven at 55°C for ∼12 h, and at 0.5°C/min up to 120°C. Each disc-shaped preform thus obtained has dimensions of 135 mm diameter, 30 mm thickness, ∼80 g weight, and a bulk density of ∼0.19 gcm⁻³. The ceramic fiber preforms were immersed in methylvinylborosiloxane (MVBS) solutions of different concentrations for 1 h and compressed uniaxially in a hydraulic press up to 19.5 mm (35% of the preform thickness) to obtain the precursor-impregnated preforms. These impregnated preforms were dried in an air oven at 55°C for ∼12 h, followed by curing by heating at 0.5°C min⁻¹–175°C. Finally, the cured foams were heated under argon purging at 1400 °C for 1 h at a heating rate of 1 ^oC min⁻¹ up to 900^oC and thereafter 3°C min⁻¹ up to 1400^oC. The samples were unloaded after cooling the furnace at a rate of 3°C min⁻¹. The resulting foams were obtained as discs of ∼130 mm diameter and are referred to here as ceramic fiber-SiBOC composite foam (CF).

Figure 2

(a) Flowchart for the preparation of ceramic fiber - SiBOC composite foams (CF); Photograph of (b) filter-pressing set-up, (c) Ceramic fiber preform, and (d) CF before (circular) and after machining (square).

Characterisation

Microstructure of the ceramic fiber preforms - dry, with MVBS and after heating to 1400°C (CF), were observed using a scanning electron microscope (Hitachi S-2400, Japan). X-ray diffraction data (XRD, Brukers D8-Advance, USA) were collected using CuKα radiation 40 kV, 40 mA; step scan of 0.051, counting time of 5 s/step and λ = 1.54060Ǻ. The density of the ceramic foams was determined by their weight and volume. The skeletal density was measured using a helium pycnometer. The CF samples were thoroughly crushed to a fine powder using an agate mortar and pestle before the skeletal density measurement. Thermogravimetric analysis (TGA) in the air was carried out at a heating rate of 10°C/minute using a thermogravimetric analyzer (SDT2960, TA Instruments, USA).

The samples for compressive strength and thermal conductivity measurements were prepared by cutting the CF with a hack saw and polishing with 200-grit sandpaper. The CF samples of size 50 mm × 50 mm × 15 mm were used for the thermal conductivity measurements using the transient plane source method (Hot Disk, TPS 2500S, Sweden) at room temperature. The CF samples were evaluated for compressive strength in a Universal Testing Machine (Instron 5050, Instron, USA) at a crosshead speed of 0.5 mm/minute. For each value of compressive strength, four CF samples of size 24 mm × 24 mm × 12 mm (ASTM standard C365/C365M-05) were used.

Results and discussion

Composite foam preparation

The ceramic wool consists of isolated fibers of diameter 4–10 µm, as evaluated from the SEM micrograph depicted in Figure 3. The ceramic fibers in the blanket are mostly stacked as layers without any binder. Attempt to use the woolen blanket as a preform, without any treatment, suffers from two major limitations: (i) on immersing in the precursor (MVBS) and subsequent squeezing, the majority of porous space between fibers collapses due to wool's non-resilient nature, leading to denser bodies and (ii) the porous bodies realised from the blanket are anisotropic in properties, due to the unidirectional layered alignment of the fiber. Thus, in this study, the fibers are dispersed in an aqueous solution of polyvinyl alcohol (PVA), pressed in the filter pressing set-up, and dried to obtain disc-shaped ceramic fiber preforms of density ∼0.19 g/cc. These preforms have a more random orientation of fibers and are reversibly compressible due to their resilient nature. The resilient nature is evidenced by the restoration of the shape and dimension when a fiber preform compressed to 50% in the thickness direction using a hydraulic press is decompressed.

Figure 3

SEM micrograph of aluminosilicate wool (photograph of a piece of as-received aluminosilicate blanket is shown in inset picture).

The ceramic fiber preform is immersed in MVBS solution, during which it absorbs the precursor throughout the bulk due to its high porosity. Before drying the fully infiltrated preform, it is essential to squeeze out a considerable amount of precursor solution. This prevents capillary movement of the liquid to the surface during drying, thus preventing the build-up of higher concentrations of the ceramic precursors on the surface than in the interior [30]. The density of the foams is further controlled by using MVBS solutions of different concentrations from 29% to 54% by weight. The MVBS-impregnated preforms on ambient drying and subsequent heating at 175°C become hard and lose all flexibility due to cross-linking. The subsequent heat treatment of the foams in an inert atmosphere was restricted to 1400°C in order to retain the properties of aluminosilicate wool. SEM micrograph of the cured precursor foam, depicted in Figure 4, indicates that MVBS coats on a majority of the ceramic fibers and fill the gaps between the fibers while leaving some fibers un-bonded.

Figure 4

SEM micrograph of ceramic fiber-MVBS precursor foam after curing at 175°C.

On heat treatment in inert atmospheres, at temperatures between 1200 and 1600°C, MVBS transforms to a SiBOC glassy phase [30, 37, 38]. Phase stability studies on ceramics derived from siloxanes and borosiloxanes, viz., SiOC and SiBOC, respectively, have revealed that the presence of boron influences the crystallisation [41 -43]. In this temperature regime, SiBOC is reported to exhibit a transition from completely amorphous to nano-crystalline glass-ceramic phases, with nano-domains of β-SiC and stacked nano-carbon layers surrounded by amorphous SiBOC [27, 30, 31]. The XRD spectra depicted in Figure 5 indicate that although the SiBOC phase and the as-received aluminosilicate wool are amorphous, on heating to 1400°C, the ceramic wool in the composite foam undergoes crystallisation. The XRD spectrum of the CF shows a majority of the mullite peaks, indicating the crystallisation of aluminosilicate fiber to mullite. A peak at 2θ close to 22^o corresponding to silica is also observed. However, the SEM micrographs of the CF (Figure 6) indicate that the fibers are intact, without any damage by reaction with the SiBOC matrix or by self-degradation due to high temperatures, as reported elsewhere [18]. From the SEM micrographs, it is also evident that the CF prepared at lower MVBS solution concentrations (29–34 wt.%) consists of many isolated fibers bonded only at certain locations by the matrix, while at higher MVBS solution concentrations (50–54 wt.%) the fibers are more or less completely embedded in the matrix. Moreover, the porosity is predominantly governed by the inter-fiber porosity at lower MVBS solution concentrations, while at higher MVBS concentrations, it is a combination of porosity of the matrix, and of those formed by bridging of the entangling fibers.

Figure 5

XRD spectra of neat SiBOC, as received aluminosilicate wool and the realised ceramic composite foam (CF).

Figure 6

SEM micrographs of CF prepared at MVBS solution concentrations of (a) 29 wt.%, (b) 41 wt.% and (c) 54 wt.%.

Effect of MVBS solution concentration

The effect of MVBS solution concentration on the bulk density of the CF is presented in Figure 7. The density of CF increases from 0.46 g/cc to 0.71 g/cc when the MVBS solution concentration increases from 29% to 54%. The increase in density is due to the progressive filling of inter-fiber porosity by the SiBOC matrix produced from the infiltrated MVBS. The skeletal density of the CF measured using a pycnometer is provided in Table S1, supplementary information. The porosity calculated from the bulk density of the CF and the corresponding skeletal density, decreases from 84% to 73% when the MVBS solution concentration increases from 29 to 54 wt.%. A rough estimate of fiber content in the CF is obtained from the weight of fiber preform, weight of impregnated MVBS and ceramic yield from MVBS. The MVBS gives a ceramic yield of 83.7% at 900^oC, as evidenced by the TGA in nitrogen atmosphere shown in Figure S1, supplementary information. The fiber content in the CF increases from 42 to 72 wt.% when the MVBS solution concentration decreases from 54 to 29 wt.%. (Table 1).

Figure 7

Variation of density of CF with MVBS solution concentration.

Table 1

Properties of the ceramic composite foams prepared at various MVBS solution concentrations before and after exposure to air atmosphere at 1300°C.

Exp. No.	MVBS solution concentration (%)	Ceramic Fiber content (Wt. %)	Before exposure			After exposure
Exp. No.	MVBS solution concentration (%)	Ceramic Fiber content (Wt. %)	Density (g/cc)	Thermal conductivity (Wm⁻¹ K⁻¹)	Comp. Strength (MPa)	Density (g/cc)	Thermal conductivity (Wm⁻¹K⁻¹)	Comp Strength (MPa)
1	54	42	0.71	0.21	0.794 ± 0.04	0.71	0.12	1.01 ± 0.03
2	50	47	0.68	0.19	0.683 ± 0.06	0.67	0.104	1 ± 0.1
3	41	52	0.58	0.16	0.56 ± 0.09	0.57	0.088	0.582 ± 0.12
4	34	60	0.51	0.15	0.481 ± 0.04	0.52	0.029	0.379 ± 0.01
5	29	72	0.46	0.13	0.395 ± 0.04	0.46	0.017	0.335 ± 0.08

The representative compressive stress–strain curves of CF prepared at various MVBS solution concentrations are represented in Figure 8. On comparison, it is observed that the stress–strain graph of CF at lower MVBS solution concentration (lower density) is smoother compared to that prepared at higher MVBS solution concentration (higher density). Further, the stress–strain graphs of the CF are slightly positively inclined with fewer fluctuations compared to that of monolithic SiBOC foams realised in our earlier study, which showed a highly serrated plateau region before densification [30]. This can be attributed to the toughening effect of the embedded aluminosilicate fiber by a more uniform distribution of compressive loads and crack-bridging. Moreover, from the morphology of the CF, it can also be proposed that the fibers might have a role in preventing the sudden breakage of the cells by connecting the cell walls through flexible fibers, which are in higher quantity in the CF with lower density (Figure 6). The effect of foam density on the thermal conductivity and compressive strength of the CF is depicted in Figures 9(a,b), respectively. With an increase in density from 0.46–0.71 g/cc, the thermal conductivity increases from 0.13 to 0.21 Wm⁻¹k⁻¹. In our earlier study on monolithic SiBOC foams, a thermal conductivity of 0.13 Wm⁻¹k⁻¹ was obtained for a density of 0.39 gcm⁻³, while in this study, similar thermal conductivity is obtained at an 18% higher density of 0.46 gcm⁻³ [30]. This can be attributed to the lesser thermal conductivity of the aluminosilicate fiber as compared to that of the SiBOC matrix. On the other hand, in this study, the compressive strength was found to be in the range of 0.395–0.794 MPa, which is comparatively lesser than that obtained in our earlier study, viz., 0.3–0.87 MPa at foam density in the range of 0.18–0.39 g.cm⁻³ [28]. The lower compressive strength of the CF in this study, as compared to monolithic foams reported earlier, can be due to poor fiber-matrix adhesion, which limits the matrix to fiber stress transfer. Similar observations were made by other researchers, even for short fiber-based foams [44, 45]. Lang et al., in their study of ceramic fiber-reinforced YSZ porous ceramics, have concluded that the fiber-reinforcing effect is observed below a certain porosity [18]. As an example, for Al₂O₃ fiber or YSZ fiber-reinforced porous YSZ ceramics, fiber addition showed a considerable reinforcing effect only at a porosity lower than 60%. Apart from improving the compressive strength, the fiber addition can improve the toughness of the otherwise brittle ceramic foams by restricting small cracks from growing by crack-bridging. Thus, the role of fibers in composite foams can be multifold.

Figure 8

Compressive stress-strain graph of CF prepared at various MVBS solution concentrations.

Figure 9

(a) Thermal conductivity and (b) compressive strength of CF, as a function of density.

Effect of exposure to a high-temperature oxidising atmosphere

To evaluate the potential of the CF in oxidising atmospheres, the constituents, as well as the foam samples, were subjected to oxidising conditions at high temperature. Figure 10 shows the oxidation TGA up to 1300°C of CF and that of the as-received aluminosilicate wool, showing negligible weight loss. Similar results were obtained on heating the CF bodies in a muffle furnace at 1300 °C for 90 min. After exposure to air at 1300 ^oC, it was observed that in comparison to the CF samples before exposure, the density of the foams remained the same; compressive strength remained approximately the same for the CF samples of lower density of 0.46–0.57 gcm⁻³, while for higher CF density of 0.67 and 0.71 gcm⁻³, the compressive strength increases by as high as 46% (Table 1). The SEM micrographs (Figure 11) of the CF samples exposed to air atmosphere at 1300^oC show that the fibers are intact, and the edges of the micro-cracks and the pores are more rounded as compared to the CF samples before exposure (Figure 6). This change in the microstructure is apparently due to the partial oxidation of SiBOC to SiO₂ and B₂O₃, which leads to the expansion and softening of the matrix at this temperature [32]. The XRD spectrum of the SiBOC after heating in air at 1300^oC for 90 min shows peaks corresponding to SiO₂ and B₂O₃ as depicted in Figure S2, supplementary information. The resultant glass phases from SiO₂ and B₂O₃ are known to act as sealants for cracks and pores in the matrix [46]. Thus, the improvement in compressive strength of the oxidised samples can be attributed to the sharp cracks converting to blunt/rounded pores and the sealing of some of the micro-cracks. Moreover, the softening and volume expansion of the SiO₂-B₂O₃ glass phase at these high temperatures can facilitate better adhesion of the matrix with the fibers, and the joining of nearby fibers resulting in better load dispersion. However, these factors are observed to be effective only above threshold matrix content (higher foam density), below which the fibers behave as independent entities and do not contribute to compressive strength. The properties of the ceramic composite foams before and after exposure to air at 1300°C are given in Table 1. Thus, in this study, the use of ceramic fibers is found advantageous (a) in making the foam fabrication simple, cost-effective and easily scalable and (b) in decreasing the foam's thermal conductivity.

Figure 10

Oxidation TGA of ceramic wool and ceramic composite foam (CF) prepared at a MVBS solution concentration of 41 wt.%.

Figure 11

SEM micrographs of the CF samples prepared at MVBS solution concentrations of (a) 29 wt.% (b) 41 wt.% and (c) 54 wt.% after heating in air atmosphere at 1300°C for 90 minutes.

Conclusions

In this study, aluminosilicate fibers embedded SiBOC composite foams are realised for thermal protection applications up to 1300°C. MVBS solutions impregnated on a compressible aluminosilicate fiber preform were dried and subsequently heat-treated for cross-linking and ceramisation to produce ceramic composite foams (CF). The MVBS solution concentrations in the range of 29–54 wt.% modulated the density, compressive strength and thermal conductivity of CF in the ranges of 0.71–0.46 g/cc, 0.79–0.39 MPa, 0.21–0.13 Wm⁻¹K⁻¹, respectively. The incorporation of aluminosilicate fiber decreased the thermal conductivity by 18% compared to a monolithic SiBOC foam, whereas it did not contribute to the compressive strength. However, exposure of the CFs to air atmosphere at 1300°C for 90 min increased their compressive strength by up to 46%. The low thermal conductivity and ability to resist oxidation prove the high potential of the realised CFs for use as a thermal protection material up to 1300°C in oxidising atmospheres.

Footnotes

Acknowledgements

The authors are thankful to Director VSSC and Director IIST for their support and encouragement.

Disclosure statement

No potential conflict of interest was reported by the authors.

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