Effect of oxidation on the compressive strength of sintered SiC-fiber bonded ceramics

Materials Science and Engineering A 534 (2012) 394–399

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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea

Effect of oxidation on the compressive strength of sintered SiC-fiber bonded ceramics Joaquín Ramírez-Rico a,∗ , Julián Martínez-Fernández a , Mrityunjay Singh b a b

Dpto. Fisica de la Materia Condensada – ICMS, Universidad de Sevilla – CSIC, Avda. Reina Mercedes S/N, 41012 Sevilla, Spain Ohio Aerospace Institute, Cleveland, OH 44142, USA

a r t i c l e

i n f o

Article history: Received 17 August 2011 Received in revised form 24 November 2011 Accepted 25 November 2011 Available online 6 December 2011 Keywords: Strength Oxidation Microstructure Fiber bonded ceramic Silicon carbide

a b s t r a c t The compressive strength of SiC-fiber bonded ceramics obtained from hot-pressed amorphous Si–Al–C–O fibers and its degradation by high temperature exposure to an oxidizing environment was studied. Compressive strength was measured at room temperature as a function of strain rate, orientation, and oxidation temperature. Weight loss was monitored as a function of exposure time in atmospheric air at temperatures ranging from 800 to 1600 ◦ C, for times ranging from 0.5 to 5 h. Room-temperature compressive strength had a moderate decrease after exposures at 800 ◦ C associated to carbon burnout; increased for exposures in the range 1000–1500 ◦ C due to a defect-blunting action of the silica scale; and decreased significantly at 1600 ◦ C due to extensive surface recession. © 2011 Elsevier B.V. All rights reserved.

1. Introduction Ceramic Matrix Composites (CMCs) are subject of great interest due to their numerous potential applications as structural elements in high-temperature environments, such as gas turbine components [1]. Other applications range from thermal management systems [2], high-temperature gas filtration in coal gasification processes or nuclear reactors [3–6]. The wide scale interest in ceramic materials in general and CMCs in particular for aerospace applications is due to their high fracture toughness and good specific mechanical strength and creep resistance, allowing higher thrustto-weight ratios than metal-based systems [7]. For these reasons, CMCs are currently being designed, developed, and tested for various applications such as gas turbines components [8,9], thermal protection systems of reentry vehicles [10], and advanced friction systems [11]. Among them, SiC/SiC composites have received the highest amount of interest and attention due to their high operating temperature in oxidizing environments, and excellent thermal shock resistance. SiC/SiC CMCs have evolved in parallel with the development of novel ceramic fibers with enhanced properties, and there exists now a wide range of available SiC fibers with which to fabricate CMCs [12,13]. Among them, the Tyranno-SA fibers are being

∗ Corresponding author. E-mail address: [email protected] (J. Ramírez-Rico). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.11.085

considered for numerous applications due to their high crystallinity and low oxygen content, which imparts them with enhanced creep resistance. These fibers are obtained by sintering an amorphous, polymer derived Si–Al–C–O fiber. The incorporation of Al enhances their resistance to oxidation and chemical attack by glassy and alkaline deposits [14–16]. CMCs containing Tyranno-SA fibers have shown better properties than those obtained from other SiC fibers such as Nicalon or Hi-Nicalon [3,5,17]. Sintered SiC fiber bonded ceramics, obtained through hot pressing of stacked amorphous Si–Al–C–O fiber mats [15,16,18–21], have several advantages over both conventional CMCs and sintered fibrous monoliths. For instance, the manufacturing cost is lower than CMCs since its fabrication involves fewer processing steps, and it can be made into complex shapes by laying up the fiber mats into shapers, molds, or preforms. Full densities are nearly achieved and the absence of pores and defects results in good mechanical properties and excellent oxidation resistance at temperatures up to 1600 ◦ C, while the presence of a very high fiber-fiber interface density imparts high fracture toughness. In spite of its promising qualities, few studies exist for this material in the literature, and several aspects of its mechanical behavior and response in oxidizing environments are still unknown. Oxidation of SiC/SiC and SiC/C CMCs is a complex process involving several steps which are more or less relevant depending on temperature, oxygen partial pressure and the presence of water vapor typical of combustion environments. Polycrystalline SiC materials and CMCs form a protective silica scale that limits

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oxygen diffusion, but volatilizes in the presence of water vapor at very high temperatures, leading to recession [22–25]. If a carbon interphase is present, it can react with oxygen that ingresses the material through microcracks, leading to a reduction in fracture toughness. The kinetics of carbon burnout can be reactionor diffusion-limited, depending on the thickness of the interface [26,27]. If microcracks exist in the matrix, carbon burnout can be stress assisted as crack-opening displacement increases can lead to a transition from diffusion to reaction limited, too, and thermal cycling under load in oxidizing environments can lead to a reduction in stiffness and fracture strength [28]. Moreover, the formation of silica can have the effect of sealing cracks and pores limiting further carbon oxidation, but this enhancement is again stress dependent [29]. In this work, we study the high temperature oxidation of SiC-fiber bonded ceramics in atmospheric air under static conditions and its ex situ effect on room temperature compressive strength, under different temperature regimes. 2. Materials and methods The material studied in this work was a monolithic SiC-fiber bonded ceramic manufactured by Ube Industries, Ltd. (Yamaguchi, Japan) using a previously described process [15,17,21,30]. In brief, amorphous Si–Al–C–O fibers obtained through pyrolysis of meltspun polyaluminocarbosilane were woven into mats that were subsequently stacked together and hot-pressed at 1800 ◦ C and 50 MPa to yield a fibrous monolith. During the hot-pressing process, O is released as CO gas, and the fibers are sintered yielding a highly crystalline SiC-material. Some excess carbon diffuses outward and a carbon interphase is formed in situ between the fibers. SiC-fiber bonded ceramic billets measuring 1 cm × 2 cm × 4 cm were cut into parallelepiped shape with approximate dimensions of 3.5 mm × 3.5 mm × 5 mm using a low speed diamond saw. These samples were used for high-temperature oxidation and compressive strength measurements. Specimens for scanning electron microscopy (SEM) observation were prepared by conventional metallographic techniques. Pristine samples were tested at room temperature under compressive loading on an Instron 8562 machine using hydraulic grip platens and a SS 316 die for loading. Load was applied using steel rods to the longest dimension of the parallelepipeds. Constant crosshead velocities ranging from 5 ␮m min−1 to 500 ␮m min−1 were applied and corresponded to nominal strain rates of 1.66 × 10−5 s−1 to 1.66 × 10−5 s−1 . For oxidation resistance, thermal treatments were done in air at 800, 1000, 1300, 1400, 1500, and 1600 ◦ C for 5 min, 30 min, 2 h, and 5 h. The heating rate was 600 ◦ C h−1 and the cooling rate was 1000 ◦ C h−1 . The specimen weight before and after thermal treatment was measured with a precision of ±0.0001 g, and sample dimensions were measured with a precision of ±0.01 mm. More than 60 samples were studied. As-received and oxidized samples were observed by SEM on a Philips XL30 and JEOL 840 electron microscopes (CITIUS, University of Seville, Spain). To evaluate the impact of high-temperature oxidation on the mechanical properties of SiC-fiber bonded ceramic, oxidized samples were tested at room temperature using the previously described setup. Load was applied using steel rods to the longest dimension of the parallelepipeds at a crosshead velocity of 500 ␮m min−1 that corresponded to an initial strain rate of 1.66 × 10−3 s−1 . 3. Results and discussion 3.1. Microstructure Fig. 1a and b shows a representative cross-section of the asfabricated material in backscattered electrons contrast, in a plane

Fig. 1. Backscattered electron (BSE) SEM image of as-received sintered SiC-fiber bonded ceramic. (a) Low magnification view showing the fiber mat stacking. (b) Detail of the fiber stacking within a layer.

containing the fiber mat stacking direction. Fibers from perpendicularly running tows are clearly visible. Fibers show polygonal cross sections characteristic of the sintering process, and some pores and voids are observed at the triple points. Darker contrast in the core of the fibers is indicative of a concentration gradient, and the cores were found to have higher carbon content than their surroundings. It is known that the microstructure of sintered Si–Al–C–O fibers consists of SiC grains of ≈200 nm in diameter, with a considerable amount of free turbostratic carbon at triple points which is homogeneously distributed in the fiber [31]. The observed concentration gradient probably occurs during the fiber sintering step and could be due to differences in diffusivity between the species involved. In recent studies, this carbon rich core and the interface has been shown to play a major role in the stress-rupture behavior of the sintered SiC-fiber bonded monolithic ceramic materials [30]. 3.2. Oxidation behavior Fig. 2 shows a plot of the weight change versus annealing time for all the temperatures studied. In all cases there was a weight loss due to the thermal treatment. However, the magnitude of the weight loss and its dependence with the time varied significantly with temperature. At 800 ◦ C, there was a clear increase of weight loss with time for all the conditions studied, although the weight loss rate was higher for shorter times. For temperatures ranging from 1000 ◦ C to 1500 ◦ C, the weight loss was almost constant for all the treatment lengths studied and smaller than that at 800 ◦ C, and the weight loss was mostly independent of temperature in that range. For that reason, average values of weight loss per unit area in the range 1000–1500 ◦ C are plotted in Fig. 2. At 1600 ◦ C, however a rapid linear decrease in weight was found and the weight loss was larger than for any other thermal treatment.

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To rationalize the previous observations, several mechanisms have to be taken into account; the relative importance of each of them will depend on temperature. At low temperatures, the oxidation behavior of the composite is dominated by carbon burnout, which is present both in the carbon interphase between the fibers as well as in the fibers’ core. Windisch et al. [27] showed that, in SiC/SiC composites with a thick carbon interphase, oxidation kinetics were limited by the rate of carbon oxidation reactions: 1 O2 → CO(g) 2 C + H2 O(g) → CO(g) + H2 (g) C + CO2 → 2CO(g) C+

Fig. 2. Weight loss per unit area as a function of exposure time in atmospheric air at temperatures from 800 ◦ C to 1600 ◦ C. Error bars are equal to one standard deviation and, when not plotted, are smaller than the symbol size.

(1)

Then, weight loss per unit area was linear with time for temperatures ranging from 800 ◦ C to 1000 ◦ C and O2 partial pressures ranging from 0.31 to 24 kPa. Other studies in CMCs with thin carbon interphases, for example those of Filipuzzi et al. [26,32], determined the reaction to be diffusion limited. In the case of the SiC-fiber bonded ceramic, the weight loss per unit area at 800 ◦ C showed a parabolic reaction rate, which is typical of a diffusion-controlled process. In this case the weight loss can be approximated as: −

dc 1 dm = −D A dt dx

(2)

Fig. 3. Backscattered electron (BSE) SEM images of sintered SiC-fiber bonded ceramic after exposures at high temperature. (a) After oxidation in air at 800 ◦ C for 5 h. (b) After oxidation in air at 1500 ◦ C for 5 h. (c) Outer surface of a sample exposed to atmospheric air at 1500 ◦ C for 5 h. (d) and (e) Sections of a sample exposed to atmospheric air at 1600 ◦ C for 5 h. The mark on micrographs represents 10 ␮m.

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Fig. 4. Schematics of the oxidation mechanism the SiC-fiber bonded ceramic monoliths studied in this work (see text for further discussion).

where A is the exposed surface area, D is the diffusion coefficient of the limiting reactant and dc/dx is the concentration gradient at the solid/gas interface. In the steady-state, this gradient can be approximated as: dc c − = w dx

(3)

This equation assumes as a simplification that the concentration profile is linear at a given time and forms a boundary layer of thickness w. The solution of the associated, √ one-dimensional diffusion problem allows us to estimate w = 4Dt and thus we can integrate Eq. (2) for constant exposed area, giving: −

m = A



2 cDt =



Kc2 t

(4)

where Kc is a rate constant for the diffusion-controlled process. At higher temperatures, SiC oxidation occurs, producing silica as the solid product and carbon monoxide as the gaseous product: SiC (s) +

3 O2 (g) → SiO2 (s) + CO (g) 2

(5)

The silica scale growth is controlled by oxygen molecular diffusion and its rate is given by [23]: Kp dx = 2x dt

(6)

where x is the silica scale thickness, and Kp is a rate constant for the diffusion-controlled process. Additionally, in presence of water vapor, the silica layer undergoes a reaction rate-controlled volatilization reaction according to: SiO2 (g) + 2H2 O(g) → Si(OH)4 (g)

(7)

Combining both effects, the silica scale thickness is given by Eq. [23]: Kp dx = − Kl 2x dt

(8)

This behavior is termed paralinear [22,25,33] and, once the steady-state is reached, Eq. (8) predicts a constant scale thickness and a linear recession rate. In this regime we can approximate: 1 dm ≈ Kl A dt

(9)

where Kl is the rate constant associated with silica scale volatilization. At 800 ◦ C there is a clear oxidation of the fiber cores and interfaces as shown in Fig. 3a, that is in agreement with the dependence of weight loss with time at temperature. The weight loss can be fit to a parabolic dependence with time typical of a diffusion limited reaction (Fig. 4a), and in this manner Kc can be calculated. At temperatures in the range of 1000–1500 ◦ C the weight loss per unit area was almost constant for holding times between 0.1 and 5 h. Fig. 3b shows a representative cross section of the material oxidized at 1500 ◦ C for 5 h. Long, cylindrical pores along the fiber interface were again observed (white arrows in the micrograph), but often these pores were closed through formation of silica (denoted as black arrows) that acted as a diffusive barrier inhibiting further oxidation of the remaining carbon in the fibers. Fibers on the sample surface were coated with SiO2 , both on the inside and the outside. Therefore, two competing processes are involved in the oxidation of SiC-fiber bonded ceramic in this temperature range: carbon burn-out, which is initially limited by gaseous oxygen diffusion through annular pores and hollow cores in the fibers, and silica formation through SiC oxidation (Fig. 3c). Despite the limited overall change in weight, these two processes are significant in the early stages of oxidation. Once the silica scale forms in the cylindrical pores and closes them, further carbon burnout must proceed by oxygen diffusion in the solid state, a much slower process. This socalled “pinching effect” (Fig. 4b) was previously observed in SiC/SiC composites [26,32]. In the time scale involved in our experiments these two effects cannot be discerned, as the formation of the silica scale occurs earlier than the lower exposure time considered in our experiments. At a temperature of 1600 ◦ C the linear relationship of weight loss and oxidation time is indicative of a paralinear regime with

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2500

Table 1 Kinetic parameters determined. See text for details.

◦

800 C 1000–1500 ◦ C 1600 ◦ C

Parabolic fit, Kc (mg cm−2 h−1/2 )

Linear fit, Kl (mg cm−2 h−1 )

5.12 – –

– <0.01 2.45

SiO2 volatilization. Indeed, severe surface recession was observed (Fig. 3d and e). Fibers were hollowed and thinned from the inside out by volatilization of the silica scale leaving a large amount of residual porosity. The mechanism involved is shown schematically in Fig. 4c. We have fitted the weight loss per unit area as a function of time for different temperatures according to the discussion above, and the results are included in Table 1. For the weight loss at 800 ◦ C the trend is clearly that described by Eq. (4) for a diffusion limited process, while for the weight loss at 1600 ◦ C the trend is linear and was fitted to Eq. (9). In all other cases, the weight loss was approximately constant and independent of exposure time and linear fits of the weight loss per unit area vs. time in this regime yield a slope that is practically zero (<10−2 mg cm−2 h−1 ). There is, however, an initial weight loss that occurs very quickly (in less than 0.5 h), which implies that some carbon is lost before the silica scale forms preventing further oxygen diffusion along the pores. 3.3. Compressive strength To establish the baseline behavior under compression, tests were carried out in directions parallel and perpendicular to the pressing direction in the material, and results are shown in Fig. 5. Marked strength anisotropy was observed, with samples perpendicular to the pressing direction showing ∼10% of the strength observed parallel to the pressing direction. This was attributed to shear stresses acting on the fibers’ carbon interphase, as cracks were observed to propagate mainly along interfaces in both cases. Strength was found to be independent of initial strain rate within the experimental scatter. Fig. 6 shows the room-temperature strength of the SiC-fiber bonded ceramic as a function of exposure time to atmospheric air at 800 ◦ C and 1600 ◦ C. In both cases, the weight loss is associated to a decrease in strength. After oxidation at 800 ◦ C the strength decreased approximately to a 60% of the as-received strength after 5 h as a consequence of the voids created by carbon oxidation,

1600 oC 800oC

2000

Strength (MPa)

Temperature

1500

1000

500

0 0

1

2

3

4

5

6

Time (h) Fig. 6. Room-temperature compressive strength as a function of exposure time for temperatures of 800 ◦ C and 1600 ◦ C.

which can act as flaws to initiate fracture. After oxidation at 1600 ◦ C, the strength decreased approximately to a 15% of the strength measured in the as-received samples after 30 min. It is clear, in light on the results of Section 3.2, that carbon burn-out and silica recession significantly damage the microstructure of these materials. Samples oxidized at temperatures between 1000 ◦ C and 1500 ◦ C exhibited a compressive strength that was independent of exposure time within the experimental scatter. Fig. 7 shows the room-temperature strength of the SiC-fiber bonded ceramic as a function of oxidation temperature for 5 h exposures. Compressive strength decreased after exposure to atmospheric air at 800 ◦ C but increased after exposures to higher temperatures and for samples exposed to 1000–1500 ◦ C, being in fact higher than that of pristine materials. This increase in strength can be attributed to a defect-blunting action of the silica scale which has been observed extensively in silica forming ceramics such as Si3 N4 and SiC [34]. Moreover, the aforementioned “pinching effect” could have a crack-healing effect under static loading [35], although we cannot conclude this directly from our work. Strength decreased significantly at 1600 ◦ C because the extensive surface recession and the large amount of residual porosity.

1800 1600

Perpendicular 2000

1200 1000

Strength (MPa)

Strength (MPa)

1400

2500

Parallel

800 600 400 200 0 10 -5

1500

1000

500

10 -4

10 -3

10 -2

Strain rate (s-1) Fig. 5. Compressive strength in the SiC-fiber bonded ceramic at room temperature, as a function of orientation and strain rates. Circles and squares correspond to experiments where the compressive load was parallel and perpendicular to the pressing direction, respectively.

0 0

500

1000

Holding temperature

1500

(oC)

Fig. 7. Room-temperature compressive strength as a function of exposure temperature, for 5 h exposures.

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4. Conclusions Weight loss per unit area of SiC-fiber bonded ceramic monoliths after exposure to atmospheric air at high temperatures shows three different regimes. At 800 ◦ C, a parabolic behavior is observed. Oxidation of the fiber’s interphase and carbon-rich core is diffusion limited and the main mechanism contributing to weight loss. At intermediate temperature, formation of a silica scale takes place which closes the pores formed by carbon burn-out, blocking further oxygen diffusion and effectively halting the oxidation process, resulting in a weight loss per unit area that is constant with respect to exposure time. This process is confirmed by microstructural observations that show that the carbon oxidizes only until a silica scale thick enough to pinch off the channels is formed. At high temperatures (1600 ◦ C) significant recession was observed, which was attributed to silica volatilization, and the weight loss is linear with time in the time range observed. Compressive strength decreased after exposure to atmospheric air at 800 ◦ C, but increased after exposures to temperatures between 1000 ◦ C and 1500 ◦ C, being this effect attributed to a defect-blunting action of the silica scale. After exposure at 1600 ◦ C, the extensive surface recession and the large amount of residual porosity caused a marked decrease of strength. Acknowledgements The authors would like to thanks Dr. T. Ishikawa and T. Matsunaga for providing materials for testing as well as technical discussion. SEM was performed at CITIUS (Universidad de Sevilla) and was funded by project P09-TEP-5152(Junta de Andalucía). References [1] A.G. Evans, Philos. T. Roy. Soc. A 351 (1995) 511–527. [2] D.B. Marshall, B.N. Cox, Annu. Rev. Mater. Res. 38 (2008) 425–443. [3] M. Kotani, A. Kohyama, Y. Katoh, J. Nucl. Mater. 289 (2001) 37–41.

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