Journal of Ceramics
Volume 2014 (2014), Article ID 169748, 9 pages
http://dx.doi.org/10.1155/2014/169748

Research Article

Modeling of Thermal and Mechanical Behavior of ZrB₂-SiC Ceramics after High Temperature Oxidation

Jun Wei,¹ Lokeswarappa R. Dharani,¹ K. Chandrashekhara,¹ Gregory E. Hilmas,² and William G. Fahrenholtz²

¹Department of Mechanical and Aerospace Engineering, Missouri University of Science and Technology, Rolla, MO 65409-0050, USA
²Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, MO 65409-0340, USA

Received 31 July 2014; Accepted 18 October 2014; Published 11 November 2014

Academic Editor: Guillaume Bernard-Granger

Copyright © 2014 Jun Wei et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Abstract.

The effects of oxidation on heat transfer and mechanical behavior of ZrB₂-SiC ceramics at high temperature are modeled using a micromechanics based finite element model. The model recognizes that when exposed to high temperature in air ZrB₂-SiC oxidizes into ZrO₂, SiO₂, and SiC-depleted ZrB₂ layer. A steady-state heat transfer analysis was conducted at first and that is followed by a thermal stress analysis. A “global-local modeling” technique is used combining finite element with infinite element for thermal stress analysis. A theoretical formulation is developed for calculating the thermal conductivity of liquid phase SiO₂. All other temperature dependent thermal and mechanical properties were obtained from published literature. Thermal stress concentrations occur near the pore due to the geometric discontinuity and material properties mismatch between the ceramic matrix and the new products. The predicted results indicate the development of thermal stresses in the SiO₂ and ZrO₂ layers and high residual stresses in the SiC-depleted ZrB₂ layer.

1. Introduction

Ultrahigh temperature ceramics (UHTCs) such as zirconium diboride and hafnium diboride (ZrB₂ and HfB₂) have been proposed for thermal protection of hypersonic aerospace vehicles, which may be exposed to temperatures above 1500°C in oxidizing environments. These materials are chemically and physically stable above 1600°C and have melting points above 3000°C [1]. In particular, ZrB₂ because of its lower theoretical density is attractive for aerospace applications [2]. Exposure of solid zirconium boride (ZrB₂ (s)) to air at elevated temperatures results in its oxidation to solid zirconia (ZrO₂ (s)) and liquid boria (B₂O₃ (l)). The oxidation resistance of ZrB₂ (s) can be improved by adding SiC (s) to promote the formation of a silica-rich scale. At high temperature, above 1100°C, SiC (s) oxidizes by reaction to form SiO₂ (l) which has a lower volatility and a higher melting point and viscosity compared with B₂O₃ (l) [3–5]. Based on the experimental observations, Fahrenholtz [3] proposed a reaction sequence for the formation of the SiC-depleted layer during the oxidation of ZrB₂-SiC at 1500°C in air. The oxide scales that form on ZrB₂-SiC consist of an outer layer of SiO₂, a middle layer of porous ZrO₂, sometimes filled with SiO₂, and a layer of SiC-depleted ZrB₂ adjoining the unoxidized ZrB₂-SiC at around 1500°C [2–6]. Parthasarathy et al. [7] developed a chemical reaction model for the oxidation of ZrB₂-SiC ceramics to predict the thicknesses of the above three new productions. For temperatures below ~1600°C, an external glassy SiO₂ layer forms and completely fills in pores of the porous ZrO₂ scale whereas at higher temperatures, the glassy scale recedes due to evaporation of SiO₂ (l) so that it only partially fills the pores in the ZrO₂ layer.

The region of particular interest, from a mechanical perspective, is the interface between the pores and the corner of the pores in the ZrO₂ scale. The pore itself may or may not be filled with liquid SiO₂ (l). The interface therefore consists of three materials (ZrO₂ scale, solid/liquid SiO₂, and SiC-depleted ZrB₂ layer) of significantly different thermal and mechanical properties. This thermomechanical mismatch and geometric discontinuity would lead to residual stresses, and additional stress concentrations during the cool-down process from the processing temperature, thereby leading to potential cracking. Some researchers [8–12], having performed furnace oxidation and high velocity thermal shock tests on ZrB₂-SiC, have indeed shown cracking in the ZrO₂ scale.

The purpose of this study is to develop a thermal and mechanical simulation model for ZrB₂-SiC ceramics after oxidation. A steady-state heat transfer analysis was conducted using finite element analysis (FEA) modeling. An adpative remeshing technique is employed in both heat transfer and thermal stress analysis. A “global-local modeling” technique is used to combine finite element with infinite element for the thermal stress and the stress concentration analysis near a pore. Temperature, thermal, and residual stress distributions will be presented.

2. FEA Model and Simulation Procedure for ZrB₂-SiC after Oxidation

To simplify the problem, the ZrO₂ scale was assumed to be of uniform thickness with regularly distributed pores. The pores were assumed to be straight, columnar in structure without tortuosity. A cylindrical representative volume unit (CRVU) was constructed and further treated as a two-dimensional (2D, pseudo-3D) axisymmetric problem subjected to local heating as shown in Figure 1. It is assumed that the body is stress free prior to heating.

Figure 1: A schematic of a model with cylindrical representative volume unit (CRVU) subjected to local heating.

The oxide scale that forms on ZrB₂-SiC consists of an outer layer of SiO₂, a middle layer of porous ZrO₂, and a layer of SiC-depleted ZrB₂ next to the unoxidized ZrB₂-SiC [3]. Based on the chemical oxidation models [7] and the experimental observation [3], the FEA models for ZrB₂-SiC ceramic after oxidation at high temperature were created as shown in Figure 2. An adaptive remeshing zone was created to cover the SiO₂, ZrO₂, ZrB₂ (SiC-depleted) and part of the ZrB₂-SiC base near the pore. The temperature dependent dimensions of the ZrO₂ scale (crystalline oxide), glassy SiO₂, and SiC-depleted ZrB₂ layer in ZrB₂-20 vol% SiC were obtained from the chemical oxidation model [7]. The unoxidized ZrB₂-SiC ceramic is treated as a macroscale continuous solid with properties of a predetermined ratio of 4 : 1 of ZrB₂ to SiC (ZrB₂-20 vol% SiC).

(a)

(b)

Figure 2: FEA models for ZrB₂-SiC ceramic after oxidation. (a) Lower temperature (external layer of SiO₂), and (b) higher temperature (partial evaporation of the glassy SiO₂).

The heat conduction equation for an axisymmetric problem can be expressed as

where is the time, is the temperature, and are polar axis and longitudinal axis, is the mass density, is the specific heat, and is the thermal conductivity. The thermoelastic model is given by

where is the elasticity matrix, the coefficient of thermal expansion, the temperature increment, the stress vector, and the strain vector.

The heat flux condition is given by

where, is the heat flux, is the surface film coefficient, and is the sink temperature. All simulations were conducted using ABAQUS finite element code.

The temperature dependent thermal and mechanical properties of the solid phases, needed for the heat transfer and mechanical analyses, can be found in the literature or databases [13, 14]. However, the thermal conductivity of liquid phase of SiO₂ (l) and the elastic constants cannot be found. As such, the temperature dependence of the thermal conductivity of liquid SiO₂ (l) and the elastic constants have to be predicted based on thermodynamics and some available test data. The predictive methods used for calculating the above properties are outlined in the next section. The cylindrical representative volume unit with equivalent pore diameter was treated as a 2D axisymmetric model (pseudo-3D). The modeling involves a steady-state heat transfer analysis representing local heat-up to calculate the temperature distribution and then a transient heat transfer analysis for 30 minutes representing a cool-down event to calculate the residual temperature distribution. The resulting temperature distributions were then applied to a thermomechanical finite element model to calculate the thermal stress distribution in the material. Adaptive remeshing technique was employed for the heat transfer analysis to improve accuracy. A “global-local modeling”, along with the adaptive remeshing technique, is used to combine finite element with infinite element for thermal stress analysis. The procedure is summarized in Figure 3.

Figure 3: Flow diagram showing the procedure and steps for thermal and mechanical analyses.

3. Thermal and Mechanical Properties

As mentioned earlier, the thermal conductivity and elastic constants of liquid phases of SiO₂ are not readily available. In an earlier work [15], the authors developed a method for calculating the thermal conductivity of liquid SiO₂ at a given temperature. The following thermal conductivity equation for a liquid by Hirschfelder et al. [16] is used in the method [15]:

In the above equation, is Boltzmann’s constant, is molecules per unit volume for the liquid, is Avogadro’s number, is molar mass, and is the temperature dependent bulk density of the liquid, and are the specific heats at constant pressure and at constant volume, respectively, and is speed of sound in the liquid. The temperature dependent specific heat at constant pressure, , for liquid SiO₂ was reported in [14]. Then, the specific heat at constant volume, , is calculated using the following relationship [17]:

where is the coefficient of thermal expansion. The temperature dependence of density and coefficient of thermal expansion of liquid SiO₂ were given in [18]. The speed of sound in liquids of SiO₂ was found in [19]. With all the needed parameters, the thermal conductivity of liquid SiO₂ was calculated using (4).

Using the temperature dependent values for density and speed of sound in liquid of SiO₂, the bulk and shear moduli of liquid SiO₂ were calculated using the Newton-Laplace equation [20–24]:

where and are the sound velocity of longitudinal and transverse wave, respectively. To simplify the problem, the temperature dependent elastic properties were used for the liquid phase of SiO₂ instead of the viscous properties because the stress state in liquid phase of SiO₂ was not of interest in the present study.

4. Results and Discussions

A 2D (pseudo-3D) 4-node linear axisymmetric heat transfer quadrilateral element was used in the thermal analysis. Heat flux was used as an error indicator variable to control the adaptive remeshing rule [25]. The dimensions of the new products after oxidation were taken from the chemical oxidation model [7].

Two steps were used in the heat transfer analyses. The first step in the heat transfer analysis was a steady-state analysis representing local heating at the top surface to calculate the temperature distribution. The second step was a transient heat transfer analysis for 30 minutes representing a cooling event to room temperature to predict residual temperature distribution. The surface heating temperature was set as (1780 K or 2240 K) during heating and 293 K during cooling. Outside the local heating area, the sink temperature was set at 293 K. The initial temperature of the material was 293 K. The heating, cooling and sink temperature conditions are summarized in Figure 4. The surface film coefficient, , was set as 2500 W/(m²·K) during the heating representing a high speed fluid flow and 100 W/(m²·K) during cooling assuming a cooler fluid flow next to a solid boundary in air. The surface film coefficient was set as 100 W/(m²·K) at all other boundaries during both heating and cooling.

Figure 4: The heating, cooling, and sink temperature conditions used in the analysis.

4.1. Results of Heat Transfer Analysis

The calculated temperature distributions in the body after surface heating temperatures of 1780 K and 2240 K are shown in Figures 5 and 6, respectively. In the following results, the temperatures shown in parenthesis correspond to the case of 2240 K. The maximum temperature at the top surface of the outer SiO₂ layer is 1168 K (1492 K) which is less than the applied heating temperature of 1780 K (2240 K). This is due to the effect of the surface film coefficient on heat transfer between a fluid and a solid and the thermal conduction at the boundaries. The temperatures at the interface between the outer SiO₂ layer and the ZrO₂, and at the interface between the oxide scale and ZrB₂, are 1160 K (1432 K) and 1148 K (1404 K), respectively. The temperature at the bottom surface is 1124 K (1370 K) which is much less than the heating temperature applied at the top surface. The temperatures at locations shown in Figures 5 and 6 were also calculated for different heating temperatures. The predicted temperatures in various materials and at the interfaces are linearly dependent on heating temperature, except in ZrO₂ layer. This deviation could be due to the increase in ZrO₂ layer thickness accompanied by a decrease in SiO₂ layer thickness. Figure 7 shows the predicted heat flux distribution of ZrB₂-SiC after steady-state analysis for heating to 2240 K. It is seen that a heat flux concentration occurs at the pore corner due to the geometric discontinuity and thermal conductivity mismatch.

Figure 5: Predicted temperature distribution of ZrB₂-SiC after steady-state analysis at heating to 1780 K.

Figure 6: Predicted temperature distribution of ZrB₂-SiC after steady-state analysis for heating to 2240 K.

Figure 7: Predicted heat flux distribution of ZrB₂-SiC after steady-state analysis for heating to 2240 K.

4.2. Results of Thermal Stress Analysis

In the thermal stress analysis, the layout of infinite elements and finite elements, as well as the displacement constraints for the stress analysis shown in Figure 8, are used. The distribution of maximum principal stresses for the steady-state heating at 1780 K is shown in Figure 9(a). The maximum principal stress distribution after cooling from 1780 K is shown in Figure 9(b). The maximum value of the maximum principal stresses of 946 MPa occurs at the top surface of SiO₂ layer. The temperature at this location is about 1166 K (Figure 5), which is below the glass melting point. The brittle glassy SiO₂ is sensitive to tensile stress with an average tensile strength of  MPa [26]. Therefore, a tensile stress of 946 MPa may induce cracking in the SiO₂ layer. The maximum value of the maximum principal stresses of 568 MPa occurs at the upper corner of the pore in the ZrO₂ layer and is less than the flexural strength (900 MPa) of ZrO₂ [27]. The maximum stress in the ZrB₂ is 451 MPa and occurs near the lower corner of the pore. This is higher than the measured bend strength of ZrB₂ [28]. The largest principal stress in the ZrB₂-SiC is 191 MPa, located near the lower corner of the pore. The flexural strength of ZrB₂-SiC is 1000 MPa [29–31]. For the cool-down case, the residual stresses in Figure 9(b) vary between 153 and 360 MPa.

Figure 8: The infinite element, the finite element, and the boundary conditions for the thermal stress analysis.

(a) Heating-up

(b) Cool-down

Figure 9: Maximum principal stresses distribution in the enlarged area near pore after steady-state thermal analysis at 1780 K.

The distribution in the maximum principal stresses near the pore for steady-state heating to 2240 K and cool-down from 2240 K to 293 K are shown in Figures 10(a) and 10(b), respectively. The stresses for heating vary between 182 and 2702 MPa while the corresponding stresses for cool-down vary between 281 and 529 MPa. Once again, the maximum values occur in the SiO₂ (2702 MPa) and ZrO₂ (2224 MPa) near the pore. This may initiate tensile cracking. These results are consistent with the experimental observations by Levine et al. [8] who performed furnace oxidation and high velocity thermal shock of ZrB₂ + 20 vol.% SiC ceramic tests. Their results show that both pores and cracks appeared in the ZrO₂ when oxidized in air at 1927°C for ten 10-min cycles. The highest maximum principal stress near the lower corner of the pore in the ZrB₂ layer is 657 MPa shown in Figure 10(a) which is higher than the bending strength of the ZrB₂ as indicated above. For the cool-down case, the maximum principal stresses shown in Figure 10(b) are less than the respective material strengths except in the ZrB₂ layer.

(a) Heating-up

(b) Cool-down

Figure 10: Maximum principal stresses distribution in the enlarged area near pore after steady-state thermal analysis at 2240 K.

The results were also obtained for additional temperatures. The variation of the maximum principal stress at indicated locations in Figures 9 and 10 with temperature is shown in Figure 11(a) for heating up to and in Figure 11(b) for cool-down from to 293 K. The thermal stress at heating temperature in the ZrB₂-SiC matrix is relatively small and does not vary much with heating temperature (Figure 11(a)). The residual stress at 293 K (Figure 11(b)) in the ZrB₂-SiC decreases with the increasing heating temperature . Watts et al. [32] measured thermal residual stresses in ZrB₂-30 vol% SiC composites using neutron diffraction. Their results indicated that stresses begin to accumulate at about 1673 K during cool-down from the processing temperature of 2172 K. The stress increased to an average compressive stress of 880 MPa in the SiC phase and to an average tensile stress of 450 MPa in the ZrB₂ phase. By using the rule of mixtures for 34 vol% SiC, the stress in the SiC (880 MPa) converts to an equivalent stress of 453 MPa which is very close to the measured stress of 450 MPa [32].

(a) Thermal stress

(b) Residual stress

Figure 11: Thermal and residual maximum principal stress at indicated locations for different heating temperatures.

5. Conclusion

A “global-local modeling” technique is used combining finite element with infinite element for thermal stress analysis for the oxidation effects on heat transfer and mechanical behavior of ZrB₂-SiC ceramics at high temperature. Thermal conductivity was calculated for the liquid phase of SiO₂ based on a theoretical formulation. The predicted temperature at the top surface of the outer SiO₂ layer is less than the applied heating temperature due to the surface film coefficient effect on the heat transfer between a fluid and a solid and the thermal conduction at the boundaries. An increase in ZrO₂ layer thickness, accompanied by a decrease in SiO₂ layer thickness, during oxidation will affect heat transfer in the body. Heat flux concentration occurs at the pore corner due to the geometric discontinuity and the material property mismatch. Thermal and residual stress concentrations occur near the pore due to geometric discontinuity and the material properties mismatch between the ceramic matrix and the new products. Thermal stresses in the surface oxide layers consisting of SiO₂ and ZrO₂, are higher than their respective materials strengths. Thermal and residual stresses in the layer of a new oxidation product of SiC-depleted ZrB₂ layer for both heating and cooling cases are higher than the material strength. Therefore, it is expected that damage may initiate in the layers of new oxidation products.

Conflict of Interests

The authors declare that there is no conflict of interests regarding the publication of this paper.

Acknowledgments

This project was funded under subcontract 10-S568-0094-01-C1 through the Universal Technology Corporation under prime contract number FA8650-05-D-5807. The authors are grateful to the technical support on the program by the Air Force Research Laboratory and specifically to Dr. Mike Cinibulk at AFRL for both his collaboration and guidance.

References

M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, “Oxidation-based materials selection for 2000°C + hypersonic aerosurfaces: theoretical considerations and historical experience,” Journal of Materials Science, vol. 39, no. 19, pp. 5887–5904, 2004.

W. G. Fahrenholtz, G. E. Hilmas, A. L. Chamberlain, and J. W. Zimmermann, “Processing and characterization of ZrB₂-based ultra-high temperature monolithic and fibrous monolithic ceramics,” Journal of Materials Science, vol. 39, no. 19, pp. 5951–5957, 2004.

W. G. Fahrenholtz, “Thermodynamic analysis of ZrB₂-SiC oxidation: formation of a SiC-depleted region,” Journal of the American Ceramic Society, vol. 90, no. 1, pp. 143–148, 2007.

C. M. Carney, P. Mogilvesky, and T. A. Parthasarathy, “Oxidation behavior of zirconium diboride silicon carbide produced by the spark plasma sintering method,” Journal of the American Ceramic Society, vol. 92, no. 9, pp. 2046–2052, 2009.

W. G. Fahrenholtz, “The ZrB₂ volatility diagram,” Journal of the American Ceramic Society, vol. 88, no. 12, pp. 3509–3512, 2005.

A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, “Evolution of structure during the oxidation of zirconium diboride-silicon carbide in air up to 1500°C,” Journal of the European Ceramic Society, vol. 27, no. 6, pp. 2495–2501, 2007.

T. A. Parthasarathy, R. A. Rapp, M. Opeka, and M. K. Cinibulk, “Modeling oxidation kinetics of SiC-containing refractory diborides,” Journal of the American Ceramic Society, vol. 95, no. 1, pp. 338–349, 2012.

S. R. Levine, E. J. Opila, M. C. Halbig, J. D. Kiser, M. Singh, and J. A. Salem, “Evaluation of ultra-high temperature ceramics for aeropropulsion use,” Journal of the European Ceramic Society, vol. 22, no. 14-15, pp. 2757–2767, 2002.

J. C. Han, P. Hu, X. Zhang, S. Meng, and W. Han, “Oxidation-resistant ZrB2-SiC composites at 2200°C,” Composites Science and Technology, vol. 68, no. 3-4, pp. 799–806, 2008.

X.-H. Zhang, P. Hu, and J.-C. Han, “Structure evolution of ZrB₂-SiC during the oxidation in air,” Journal of Materials Research, vol. 23, no. 7, pp. 1961–1972, 2008.

P. Hu, W. Guolin, and Z. Wang, “Oxidation mechanism and resistance of ZrB₂-SiC composites,” Corrosion Science, vol. 51, no. 11, pp. 2724–2732, 2009.

P. Sarin, P. E. Driemeyer, R. P. Haggerty et al., “In situ studies of oxidation of ZrB2 and ZrB2–SiC composites at high temperatures,” Journal of the European Ceramic Society, vol. 30, no. 11, pp. 2375–2386, 2010.

“Material Property Database (MPDB),” JAHM Software, Inc., http://www.jahm.com

I. Barin, Thermochemical Data of Pure Substances, Wiley-VCH, New York, NY, USA, 3rd edition, 1995.

J. Wei, L. R. Dharani, K. Chandrashekhara, G. E. Hilmas, and W. G. Fahrenholtz, “Modeling of oxidation effects on heat transfer behavior of ZrB₂ and ZrB₂-SiC Ceramics at high temperature,” in Proceedings of the 53rd AIAA/ASME/ASCE/AHS/ASC Structures, Structural Dynamics and Materials Conference, AIAA, Honolulu, Hawaii, USA, April 2012.

J. O. Hirschfelder, C. F. Curtiss, and R. B. Bird, Molecular Theory of Gases and Liquid, John Wiley & Sons, New York, NY, USA, 1954.

R. S. Hixson, M. A. Winkler, and J. W. Shaner, “Sound speed measurements in liquid lead at high temperature and pressure,” in Conference: 10. High Pressure Conference on Research in High Pressure Science and Technology, Amsterdam, The Netherlands, 1985, LA-UR-85-2363.

M. S. Ghiorso, “An equation of state for silicate melts. I. Formulation of a general model,” American Journal of Science, vol. 304, no. 8-9, pp. 637–678, 2004.

A. Polian, V.-T. Dung, and P. Richet, “Elastic properties of a-SiO₂ up to 2300 K from Brillouin scattering measurements,” Europhysics Letters, vol. 57, no. 3, pp. 375–381, 2002.

2013, https://en.wikipedia.org/wiki/Speed_of_sound

D. D. Joseph, A. Narain, and O. Riccius, “Shear-wave speeds and elastic moduli for different liquids. Part 1. Theory,” Journal of Fluid Mechanics, vol. 171, pp. 289–308, 1986.

D. D. Joseph, O. Riccius, and M. Arney, “Shear-wave speeds and elastic moduli for different liquids. Part 2: experiments,” Journal of Fluid Mechanics, vol. 171, pp. 309–338, 1986.

O. Riccius, D. D. Joseph, and M. Arney, “Shear-wave speeds and elastic moduli for different liquids Part 3. Experiments-update,” Rheologica Acta, vol. 26, no. 1, pp. 96–99, 1987.

M. S. Greenwood and J. A. Bamberger, “Measurement of viscosity and shear wave velocity of a liquid or slurry for on-line process control,” Ultrasonics, vol. 39, no. 9, pp. 623–630, 2002.

“ABAQUS 6.9 documentation,” http://www.simulia.com/support/documentation.html

W. N. Sharpe Jr., J. Pulskamp, D. S. Gianola, C. Eberl, R. G. Polcawich, and R. J. Thompson, “Strain measurements of silicon dioxide microspecimens by digital imaging processing,” Experimental Mechanics, vol. 47, no. 5, pp. 649–658, 2007.

http://www.ferroceramic.com/zirconia.htm

D. Kalish, E. V. Clougherty, and K. Kreder, “Strength, fracture mode, and thermal stress resistance of HfB₂ and ZrB₂,” Journal of the American Ceramic Society, vol. 52, no. 1, pp. 30–36, 1969.

A. L. Chamberlain, W. G. Fahrenholtz, G. E. Hilmas, and D. T. Ellerby, “High-strength zirconium diboride-based ceramics,” Journal of the American Ceramic Society, vol. 87, no. 6, pp. 1170–1172, 2004.

A. Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, “Effect of hot pressing time and temperature on the microstructure and mechanical properties of ZrB₂-SiC,” Journal of Materials Science, vol. 42, no. 8, pp. 2735–2744, 2007.

S. Zhu, W. G. Fahrenholtz, and G. E. Hilmas, “Influence of silicon carbide particle size on the microstructure and mechanical properties of zirconium diboride-silicon carbide ceramics,” Journal of the European Ceramic Society, vol. 27, no. 4, pp. 2077–2083, 2007.

J. Watts, G. Hilmas, W. G. Fahrenholtz, D. Brown, and B. Clausen, “Measurement of thermal residual stresses in ZrB₂-SiC composites,” Journal of the European Ceramic Society, vol. 31, no. 9, pp. 1811–1820, 2011.