2.1 Partial sintering

Partial sintering is the simplest, frequently used and most conventional method to fabricate porous SiC ceramics. Particles of powder compact are bonded by partial sintering due to surface diffusion, evaporation–condensation, recrystallization or a solution–reprecipitation process. We can see a homogeneous pore structure when sintering is terminated before fully densified (Fig. 1(a)). Full densification is often retarded or prohibited by reducing the sintering potential. Reduced sintering potential is achieved by low sintering temperature, a constrained network of coarse powders, sintering without additives, and recrystallization. Pore size and porosity are controlled by the size of starting powders and degree of partial sintering. Generally, the powder has to be two to five times larger than the pore size required [35]. It has been previously postulated that the elastic modulus is directly dependent on the neck radius to particle radius ratio and such neck formation allows the elastic behavior to be directly related to the sintering kinetics. Examples of partial sintering method reported in the previous studies are shown in Table 1.

Lin and Tsai [51] processed macroporous SiC ceramics by slip casting of submicron SiC powder-alumina additive mixtures and sintering at 1450–1800 °C. By varying both sintering temperature and alumina content (3–8 wt%), both porosity and pore size have been tailored within a range of 29–39% and 0.10–2.33 μm, respectively. Suwanmethanond et al. [5] processed porous SiC ceramics by pressureless sintering of SiC powder compacts with various additives including Al₂O₃, B₄C, phenol resin, and carbon black. Fukushima et al. [14,52] also processed porous SiC ceramics by cold isostatic pressing (CIP) of submicron β-SiC powder and alumina additive mixtures and subsequent sintering at 1500–1800 °C. By varying sintering temperature, CIP pressure, and alumina additive, both the microstructure and pore size have been tailored. The pore size and particle size of the porous SiC membrane is found to be increased when there is no addition of alumina. By an addition of 4% alumina, it showed the retention of small pore sizes and fine particles. This was mainly due to the formation of SiO₂–Al₂O₃ liquid phase from a thin SiO₂ layer that exists on the surface of SiC particles and alumina. Treatment with higher CIP pressure reduced the pore size of the porous SiC ceramics effectively. Fukushima and coworkers [53] prepared porous SiC membrane supports which are sintered at 1500–1800 °C by using different additive contents and weight ratios of Al₂O₃/Y₂O₃. For the membrane support with 1 wt% additives, the pore size and grain size increased with increasing sintering temperature. However, the linear shrinkage and porosity did not change during heating because grain growth took place through surface diffusion and evaporation–condensation, resulting in an increase of pore size without shrinkage. Eom et al. [54] processed porous SiC ceramics from a carbon-filled polysiloxane without any templates by carbothermal reduction and subsequent liquid-phase sintering with Al₂O₃–Y₂O₃ as an additive and the obtained porosity ranged from 39% to 54% and pore size ranged from 0.003 to 30 μm. The grain morphology of the SiC grains transformed from equiaxed to platelet grains with increasing sintering temperature as a result of the β → α phase transformation of SiC and subsequent grain growth. Zhao and his colleagues [55] processed porous SiC ceramics by pressureless sintering silicon-resin coated SiC powder compacts containing Al₂O₃–SiO₂–Y₂O₃ additives at 1350–1750 °C. The obtained porosity was in the range of 66–70%.

Ihle and his colleagues [56,194,195] have intensively studied the structure and phase formation of porous liquid-phase sintered SiC (LPS-SiC) containing Y₂O₃ and Al₂O₃ additives. Thermodynamic calculations and sintering experiments revealed that silicides or carbides can be formed in addition to stable oxides. The main parameters controlling the formation of the different reaction products are free carbon content, the oxygen activity, and the sintering temperature. They suggested that decomposition of oxide additives can be effectively suppressed and stable porous LPS-SiC can be produced by using CO containing atmospheres. Owens and Ruppel [62] fabricated porous SiC ceramics by forming SiC powder mixtures with a multimodal distribution of SiC particles and subsequent sintering. Kim et al. [44] processed porous SiC ceramics by adding large (65 μm) SiC particles or SiC whiskers into submicron SiC particles to restrain the densification. Yttrium aluminum garnet (Y₃Al₅O₁₂, YAG) was added as a sintering additive. By controlling the content of large SiC whiskers or particles and the applied pressure during sintering, porous SiC ceramics with porosity ranging from 0.3% to 39% were fabricated. Fukushima et al. [63] added nano-sized SiC powder into submicron-sized SiC powder and sintered the formed body prepared from the powder mixtures at 1500–1800 °C. A higher CIP pressure treatment was effective in reducing pore size. The pore size and grain size increased with increasing sintering temperature. Enlarged neck regions were observed, due to the enhanced mass transfer from the nano-sized particles to the neck area. The increased neck area improved the compressive strength. Jang et al. [65,66] fabricated porous SiC ceramics using a submicron-sized and micron-sized SiC powder mixture without sintering additives by vacuum sintering at 1700–2000 °C. The obtained porosity was in the range of 30–40% and the average grain size increased with increasing sintering temperature. Kawamura et al. [67] processed porous SiC ceramics by synthesizing SiC from pelletized powder mixtures of Si and fullerene or Si and amorphous carbon at a temperature as low as 1000 K in sodium vapor. The obtained porosity was in a range of 66–71%. Hotta et al. [60] processed porous SiC ceramics with 1 wt% Al₂O₃ or 1 wt% AlN–Y₂O₃ additives by spark plasma sintering. A flexural strength of 95 MPa at 42% porosity was obtained in the porous ceramics and the obtained porosity was in a range of 32–39%. Zhou et al. [61] fabricated porous SiC tubes by extrusion forming and partial sintering methods. The tubes sintered at 1800–1900 °C had porosities of 40–46% and pore sizes of 0.19–1.70 μm.

Liu et al. [68] fabricated wood-like porous SiC ceramics without a template via high temperature recrystallization process by mimicking the formation mechanism of cellular structure of woods. The decomposition of SiC produces the gas mixture of Si, Si₂C and SiC₂ above 1800 °C. The possible reactions are given as follows:(1) ${\text{SiC}}_{\text{(}s\text{)}}\to {\text{Si}}_{\text{(}g\text{)}}\text{+}{\text{C}}_{\text{(}s\text{)}}$(1) (2) ${\text{2SiC}}_{\text{(}s\text{)}}\to {\text{Si}}_{\text{(}g\text{)}}\text{+}{\text{SiC}}_2{}_{\text{(}g\text{)}}$(2) (3) ${\text{Si}}_{(g)}+{\text{SiC}}_{(s)}\to {\text{Si}}_{\text{2}}{\text{C}}_{(g)}$(3)

The Si-carrying gas species play a role as a transport medium for carbon and SiC. The directional flow of gas mixture in porous green body induces the surface ablation, rearrangement, and recrystallization of SiC grains, which leads to the formation of the aligned columnar fibrous SiC crystals and tubular pores in the axial direction. Representative morphologies of wood-like porous SiC ceramics are shown in Fig. 2 [68]. The pressure and concentration gradients are two driving forces for the transport of the produced gases. The orientation degree of SiC crystals and the pores in the axial direction strongly depend on the temperature and pressure so that it increases with increasing temperature while it decreases with increasing furnace pressure. The obtained porosity was in range of 53–61% and it increased with increasing the sintering temperature within a range of 1950–2300 °C.

Processing and properties of macroporous silicon carbide ceramics: A review

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Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

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18 April 2018

Fig. 2 Typical morphologies of wood-like porous SiC ceramic fabricated by partial sintering: (a) the radial section surface and (b) the axial section surface [68].

Reproduced with the permission of Elsevier.

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Kim et al. [69] processed porous SiC ceramics with interconnected huge plate-like grains via recrystallization from oxidized β-SiC powder with 1 wt% B₄C. β-SiC powders were oxidized at 650 °C for 2 h to remove the free carbon present in it. When β → α phase transformation occurred at 2100 °C, rapid grain growth of α-SiC consumed the unstable β-SiC matrix resulting in an interconnected network structure with huge plate like grains. Both the oxidation of β-SiC powder and the addition of B₄C are essential for rapid grain growth via a gas phase reaction, such as an evaporation–condensation reaction. Typical microstructure of porous SiC ceramics with interconnected huge plate-like grains is shown in Fig. 3 [69]. The porosity and mean pore size of the porous SiC ceramics were 47% and 6–7 μm, respectively.

Processing and properties of macroporous silicon carbide ceramics: A review

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Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

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18 April 2018

Fig. 3 Typical microstructure of porous SiC ceramics with interconnected huge plate-like grains produced by partial sintering method [69].

Reproduced with the permission of Elsevier.

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In summary, partial sintering is the simplest and easiest way to fabricate porous SiC ceramics with porosity ranging below 65% and with pore size ranging from 0.1 to 10 μm. In this technique, particles of a SiC powder compact are bonded by partial sintering due to surface diffusion, evaporation–condensation, recrystallization or a solution–reprecipitation process. Full densification is retarded or prohibited by reducing the sintering potential. Reduced sintering potential is achieved by low sintering temperature, a constrained network of coarse powders, sintering without additives, and recrystallization. A very homogeneous pore structure with narrow pore size distribution has been successfully prepared through recrystallization process. One of the advantages of this method in the processing of porous SiC ceramics is an easiness of dimensional control, owing to the occurrence of grain growth without any shrinkage and densification, which can be obtained due to the strong covalent bonding characteristic of SiC.

2.2 Replica method

The replica method is based on the copy of the original foams concerning its pore shape and strut structure. Three different techniques were suggested to produce macroporous SiC ceramics: (1) impregnation of polymer foams with SiC suspension or precursor solution, (2) chemical vapor deposition (CVD) of SiC on polymeric foams, and (3) infiltration of natural wood-derived or artificial carbon foams with Si sources. Macroporous SiC ceramics with open cells of high volume porosity and with cell sizes ranging from 100 μm to millimeters are frequently fabricated by the replica method. Examples of replica method reported in the literature are shown in Table 2. The original invention of using a sponge as a template was from Schwartzwalder and Somers [196]. In early 1960s, they used a polymeric sponge to prepare ceramic cellular structure of various pore sizes, porosities, and chemical compositions. Since then the sponge replica technique has become the most popular method to produce macroporous ceramics with high porosity and open cells. The most frequently used synthetic template is polymeric sponge such as polyurethane (PU) [197^[198]–203]. The method is being widely used to produce filters or light weight structures in various applications. Porosity range of more than 90% can be produced using this method. SiC filter for molten metals is a typical example produced through the replica technique.

The first method based on impregnation of polymer foams with a SiC suspension or precursor solution involves (i) impregnating the porous or cellular structure with SiC suspension or precursor solution, (ii) removing the excess suspension by passing through rollers or centrifuges, (iii) drying the SiC or precursor coated polymer sponge, (iv) heat treatment through careful heating to remove the polymer sponge by decomposition, and (v) densification of the SiC coating by sintering in an appropriate temperature [70^[71]–73,79,80]. The impregnation of a cellular structure with a SiC suspension or precursor solution produces macroporous SiC ceramics which exhibit facsimile morphology of original porous material. Reticulated porous SiC ceramics processed by the above method contains triangular pores inside struts after the organic sponge is burned out, making them sensitive to structural stresses and limiting their structural applications. Zhu et al. [74] developed a recoating method to improve mechanical strength of the reticulated porous SiC ceramics. In the method, the coating process is composed of two stages. In the first stage, a thicker slurry is used to coat the sponge preform uniformly. The green body is preheated to produce a reticulated preform with sufficient handling strength after the sponge has been burned out. In the second stage, the preform is recoated by a thinner slurry using centrifugal method. By this method mechanical strength of the reticulated porous SiC ceramics has been improved significantly. The authors also suggested that the flexural strength of reticulated porous SiC ceramics improved remarkably as the solid content of SiC slurries increased [70,71].

Mouazer and coworkers [72] fabricated porous SiC ceramics using SiC powder, B₄C–graphite as a sintering aid, gelatin as a gelling agent, and polyurethane foam as a template for replica technique. In the study, pores of polyurethane foam were filled with aqueous slurry of ceramic powders with appropriate amount of gelling, anti-foaming and dispersing agents. The sponge is compressed to ensure that it is completely filled with slurry and rolled on to remove excess of slurry. The porous SiC foams with a porosity ranging from 78% to 88% were obtained after drying, calcination, and sintering at 2120 °C for 1 h in argon. Boron-doped SiC has a good electrical conductivity and making it suitable for electrically conductive foam. Yao et al. [75] processed reticulated porous SiC ceramics with polyurethane sponges by using MgO–Al₂O₃–SiO₂ additives at 1000–1450 °C. The sintering additives are employed from alumina, kaolin and talc powders. By adding the sintering additives, the authors could lower the optimum sintering temperature to 1300 °C, which was at least 100 °C lower than that with Al₂O₃–SiO₂ additives. Soy et al. [76] fabricated reticulated porous SiC ceramics using a ceramic slurry consisting of SiC powders and bentonite (montmorillonite). Bentonite was used as both a binder for ceramic slurry at room temperature and a sintering additive at elevated temperatures. Increasing bentonite content from 1 wt% to 10 wt% decreased the porosity from 86% to 84% and increased mechanical strength.

Due to the covalent nature of Si–C bonds, porous SiC ceramics normally needed to be sintered at high temperatures or/and with the addition of sintering additives, which have limited the application of porous SiC ceramics. New processing routes that overcome these problems are the preceramic polymer processes, during which the polymer precursors convert into ceramic materials. This process has an advantage of requiring usually low temperature of 1000–1200 °C [72]. Some examples of the processing route using precursor solutions are described here. Bao et al. [79] processed reticulated SiC ceramics using polysilane polymeric precursors. Polyurethane foams were immersed in polysilane precursor solutions and these were heated in nitrogen at 900–1300 °C. The SiC foams produced in this manner showed well-defined open-cell structures and the struts were free of voids. Nangrejo and Edirisinghe [80] used two types of polymeric precursors for SiC and dichloromethane as a solvent for processing reticulated SiC ceramics. PU foams were soaked in precursor solutions and subsequently pyrolyzed in nitrogen to produce SiC foams. It was possible to control the sintering shrinkage by varying the concentration of polymeric precursor and/or second phase fillers [204,205]. A representative microstructure of the reticulated SiC ceramics fabricated from SiC slurry consisted of well-defined open-cell structures as shown in Fig. 4(a) [73]. Vogt et al. [21] processed macroporous SiC foams by dipping PU foams several times into the polysiloxane (MK polymer) precursor slurry. After first coating cycle, the PU foam was decomposed, leaving hollow struts inside the ceramic structure, and simultaneously polysiloxane was pyrolyzed. Highly porous SiC ceramics with porosities higher than 96% were obtained by the process and the struts were polycrystalline and contain voids which were formed after removal of the unreacted carbon. The polysiloxane-derived phase acts as an additional binding phase at elevated temperatures and allows an appropriate handling for reinfiltration of the foam structures before nitridation at 1400 °C or recrystallization at 1850 °C [82]. Aoki and McEnaney [110] developed a modified replication method for producing SiC foams by (i) coating of polymer foams with phenolic resin, (ii) carbonization of the resin by pyrolysis, (iii) reaction of the resin-derived carbon foam with Si vapor at high temperatures and (iv) removal of the unreacted carbon by oxidation at 800 °C.

Processing and properties of macroporous silicon carbide ceramics: A review

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Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

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18 April 2018

Fig. 4 Typical microstructures of macroporous SiC ceramics produced via the replica technique: (a) SiC foams produced by impregnation of polyurethane sponges with SiC slurry [73], (b) SiC foams produced by CVD process [85], (c) SiC foams produced by infiltration of wood-derived carbon foams with Si vapor [88] and (d) SiC foams produced by infiltration of wood-derived carbon foams with Si vapor [111].

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The second method based on CVD of SiC involves CVD coating of SiC on polyurethane foams. Duocell^® SiC foams (Oakland, California) are produced by this method. The cellular materials presented an open cell microstructure consisting of interconnected SiC struts which defined polyhedral cells with 12–14 faces (Fig. 4(b)) [85].

The third method based on the infiltration of natural wood-derived carbon foams with Si sources involves (i) pyrolysis of woods for making carbon foams and (ii) infiltration of the carbon foams with gaseous or liquid Si, SiO₂, or TEOS to form SiC or Si-SiC ceramics [8,38,39,48,71,82^[83]–106]. Ota and coworkers [82] were probably first team to use wood as a natural template to produce cellular ceramics with alkoxide solutions via sol–gel chemistry. Based on this work, Greil and coworkers [39,83^[84]–88] processed cellular SiC ceramics with anisotropic pore structures by infiltration of liquid Si into carbonized wood and subsequent reaction to SiC. The procedure was to heat the wood in non-oxidizing atmosphere at temperatures between 800 and 1800 °C which resulted in decomposition of the polyaromatic constituents to form carbon preforms. The biological carbon preforms were used as templates for infiltration with gaseous or liquid Si to form cellular SiC and Si/SiC ceramics. Depending on the initial cellular microstructure of the various kinds of wood, SiC ceramics of different anisotropic pore structures in the form of pseudomorphs of the original wood were obtained [83]. According to Greil’ data [86], the porosities ranging from 32% to 91% were obtained depending on the wood. Maximum diameters for the tracheal cells were in the range of 50 μm for pine up to 350 μm for oak. The anisotropic nature of cellular SiC ceramics produced using wood as template might be advantageous in applications that require open and highly oriented porous structures, such as in catalytic supports and in the filtration of hot gases [86]. Lopez-Alvarez and coworkers [106] used juncus maritimus which is also known as sea rush as a template to fabricate porous SiC scaffolds. Sea rush is thermally decomposed and subsequently infiltrated by molten Si in order to fabricate 3D porous SiC scaffolds that preserve the original microstructure of the plant. The average pore diameter and porosity of the SiC scaffolds were 12 μm and 48%, respectively. Qian et al. [107] processed porous SiC ceramics by carbothermal reduction of charcoal–SiO₂ composites in argon atmosphere, which was fabricated by infiltrating silica sol into porous wood-derived carbon foams. Many other researchers [89,91,93,94,96,99,100,104,108,109] processed porous SiC ceramics with a variety of wood-derived templates, but their strategies are basically same with Greil’ approach. Typical microstructure of SiC foams produced by infiltration of wood-derived carbon foams with Si vapor is shown in Fig. 4(c) [88].

Herzog et al. [97] developed modified wood-based process by using wood-based fiberboards as a template for overcoming the inhomogeneity problem that hinders the use of natural bulk wood as a template for SiC components. In the process, wood-based fiberboards were prepared by pressing technique and the boards were transformed into carbon foams by pyrolysis. The carbonized wood-based material was infiltrated with SiO₂ sol and transformed into porous SiC by the carbothermal reduction process. The obtained structural properties of the materials were isotropic and porosities were in the range of 55–74%. Qian et al. [92] processed macroporous SiC ceramics by (i) pressing phenol resin impregnated wood powders, (ii) pyrolysis, (iii) infiltration of Si at 1450–1600 °C, and (iv) removal of residual Si at 1700 °C. Pore structure of the obtained materials was homogeneous and porosities were ranged between 65% and 75%.

The artificial carbon foams for processing porous SiC ceramics by replication method involve resin-derived carbon foams [110], paper-derived carbon foams [111], and mesophase pitch [112,116]. SiC foams were produced by siliciding the carbon foams with Si vapor sources, such as Si₃N₄, vapor or liquid Si, and CH₃SiCl₃. A typical microstructure of the porous SiC ceramics fabricated from artificial carbon foams is shown in Fig. 4(d) [111]. Xue and Wang [113] processed porous SiC ceramics from cotton and phenol resin-derived carbon foams by liquid Si infiltration. The obtained porosity was in range of 10–35%. Cho et al. [114] fabricated porous SiC ceramics by dip-coating porous reaction-bonded SiC supports with slurries consisting of fine SiC powders and phenol resin and by slicing the coated layer with liquid Si. By the process, porous SiC ceramics with multi-layered pore structure were produced for hot gas filter applications. Shimada and his colleagues [115] processed porous SiC ceramics from a clutch lining waste-derived carbon foams and Si sources such as tetraethoxysilane and SiO by reacting them at 1500–1600 °C. The obtained porosity was about 80%.

In summary, the replica method is the most appropriate technique to produce open-cell SiC ceramics with pore sizes ranging from 10 μm to 5 mm at porosity levels between 60% and 95%. The adhesion of an impregnating suspension on a polymeric sponge is the most crucial step in the polymer replica technique. While this technique benefits from its overall simplicity, the mechanical strength and reliability of porous structures produced with this method can be substantially degraded by the formation of hollow struts and cracked struts during pyrolysis of the polymer sponge [32]. Wood-derived carbon foams have also been used as templates to produce porous SiC ceramics. Wood-derived porous SiC ceramics have an anisotropic pore structure, resulting in anisotropic properties. The several steps required to convert the wood structures into macroporous SiC ceramics might excessively increase the cost of the product [32]. Based on various processing parameters in replica method, the available data indicate a minimum flexural strength of 0.7 MPa with 86% porosity [70] and a maximum flexural strength of 205 MPa with 9% porosity [113] for the porous SiC ceramics.

2.3 Sacrificial template method

The sacrificial template method is one of the techniques to fabricate porous SiC ceramics by mixing appropriate amounts of sacrificial template as a pore forming agent with SiC powders or SiC precursors, forming a biphasic composite using the template-SiC powder or precursor mixtures, and removing the template before or during sintering process to create pores (Fig. 1(c)) [31,43,46,50,117,136]. This method leads to porous materials displaying a negative replica of the original sacrificial template, as opposed to the positive morphology obtained from the replica technique [32]. The way that the sacrificial template is extracted from the biphasic composites depends primarily on the type of pore former employed. A wide variety of sacrificial materials have been used as templates, including synthetic and natural organics, liquids, salts, metals, and ceramics. Table 3 illustrates some examples of possible sacrificial templates. Synthetic and natural organics are often extracted through decomposition by applying thermal treatments at temperatures between 200 and 1000 °C [43,45^[46]–47,118^[119]–124,134,135]. Liquids are often extracted through freeze drying or sublimation after freezing [137^[138]–139]. Carbon and SiO₂ templates are removed by oxidation [117,141^[142]–143] and chemical leaching [117,146], respectively. Salts are often extracted by leaching using water [140].

Polymer microbeads and microspheres have been frequently employed for sacrificial template method. For example, Kim and his colleagues [43,46,47,118] fabricated porous SiC ceramics by (i) mixing SiC powders, sintering additives, and polymer microbeads, (ii) forming by dry pressing, (iii) subsequent heat-treatments for decomposing the templates, and (iv) sintering the formed body. By controlling the microbead content and sintering temperature, it was possible to achieve tailored porosity ranging from 16% to 69% [43]. Eom et al. [47] processed porous SiC ceramics with various microstructures by controlling α/β ratio in the starting SiC powders. A typical microstructure obtained from 1% α-SiC and 99% β-SiC powders is shown in Fig. 5(a) [47]. Polymer precursors and artificial templates have also been used for producing porous SiC ceramics via sacrificial template method. Kim and his colleagues [45,119^[120]–123,125^[123]–128] developed a new processing method for fabricating macroporous SiC ceramics from carbon-filled polysiloxane by carbothermal reduction and subsequent sintering process. The strategy adopted for making porous SiC ceramics involves the following: (i) fabricating a formed body from a mixture of polysiloxane, carbon source such as phenol resin or carbon black, SiC powder as an optional inert filler, sacrificial templates, and sintering additives; (ii) pyrolysis of polymeric materials; (iii) synthesizing SiC by carbothermal reduction; and (iv) sintering the materials. The sacrificial templates used were polymer microbeads, hollow microspheres or expandable microspheres. In the process, SiC can be synthesized from the polysiloxane, if an additional carbon source is added and the mixture is heated at higher temperatures (>1400 °C) [119,206]. By controlling the template, inert filler, additive contents, and sintering temperature, it was possible to adjust the porosity so that it ranged from 40% to 95%. Kotani et al. [131] fabricated porous SiC ceramics using allylhydrido-polycarbosilane as a SiC precursor and polymer microbead as a template. The final product was formed without critical crack initiation and it was possible only, when the polymer was sufficiently contained to make a strong skeletal structure. Eom et al. [124] and Kumar et al. [133] fabricated highly porous SiC ceramics using poly(ether-co-octane) and hollow microsphere with more than 78% of porosity. Jin and Kim [50] processed highly porous SiC ceramics by pyrolyzing a cross-linked body consisting of polycarbosilane (PCS) and polymer microbeads at 1100–1400 °C in an argon atmosphere. By controlling the microbead content and pyrolysis temperature, it was possible to achieve tailored porosity ranging from 56% to 88% [50]. A typical microstructure obtained by the process is shown in Fig. 5(b) [50].

Processing and properties of macroporous silicon carbide ceramics: A review

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Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 5 Typical microstructures of macroporous SiC ceramics produced via the sacrificial template technique: (a) SiC foams produced using polymer microbead as a template and β-SiC powders [47], (b) SiC foam obtained using polymer microbead as a template and polycarbosilane as a SiC precursor [50], and (c) SiC foams produced by gelation-freezing method parallel to the freezing direction [137] and (d) perpendicular to the freezing direction [137].

(a), (c) and (d) are reproduced with the permission of Elsevier and (b) is reproduced with the permission of Springer.

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Eom et al. [47,207,208] suggested three different strategies for controlling microstructure of porous SiC ceramics: (1) control of post annealing temperature, (2) control of additive content, and (3) control of initial α-SiC content in the starting powder. In the first strategy, the grain size increased with an increase of the annealing time and the morphology of SiC grains changed from equiaxed to cube after an annealing at 1750 °C and to hexagonal platelets after an annealing at 1950 °C [207]. In the second strategy, a toughened strut microstructure was obtained in porous SiC ceramics by the carbothermal reduction of polysiloxane-derived SiOC containing hollow microspheres, followed by sintering with Al₂O₃–Y₂O₃–CaO, as a result of enhanced grain growth due to the bimodal distribution of SiC, low viscosity of the liquid phase, and the β → α phase transformation of SiC. In the last strategy, by adjusting the initial α-SiC content in the processing of porous SiC ceramics, the grain morphology can be controlled from equiaxed to large platelet grains. A microstructure consisting of large platelet α-SiC grains were obtained from β powder or a mixture of α/β powders containing small (≤10%) amounts of α powders by sintering at 1950 °C for 4 h. Song et al. [132] fabricated porous SiC by a novel processing route with a duplex pore structure which is developed by using a polysiloxane, carbon black, SiC, Al₂O₃, Y₂O₃, and two kinds of pore former (expandable microspheres and PMMA spheres). It is fabricated by carbothermal reduction of a polysiloxane-derived SiOC and a subsequent sintering process. Two kinds of pore formers, expandable microspheres and PMMA spheres, were added simultaneously to create a duplex pore structure with high permeability. The duplex pore structure consists of large pores derived from the expandable microspheres and small windows in the struts that were replicated from the PMMA spheres. The presence of these small windows in the strut area improved the permeability of the porous ceramics.

Eom et al. [209] investigated the effect of forming methods and fabricated porous SiC from blends of carbon-filled polysiloxane. They used three different plastic forming methods like compression molding, injection molding and extrusion method. Each method gave different microstructures and porosities. The compression molding process led to more homogenous microstructure and lower porosity (72%) than the other forming methods. The extrusion and injection molding methods resulted in the production of porous SiC ceramics with 84% and 74% of porosities, respectively. Additive composition also found to result in different porosities and microstructures [129,130]. The obtained porosity of porous SiC ceramics fabricated from polysiloxane-derived SiOC by carbothermal reduction and subsequent sintering ranged from 56% to 72% after sintering at 1750 °C, depending on the additive composition.

Natural organics such as wax, dextrin, yeast, and agar were also used as templates for producing SiC foams. Roy et al. [134] processed porous SiC ceramics using SiC powder and wax as a template. Ying et al. [135] and Chi et al. [136] used dextrin and yeast as a template, respectively, to fabricate porous SiC ceramics. Yeast belongs to microorganism and is mainly composed of C, H, O, N, P and S elements. The surface of yeast has many pores with a porosity of 29.7%, which was measured from mercury porosimetry. Therefore, some slurry is adsorbed on the surface of yeast. Yeast completes its pyrolysis at higher temperature than general pore former, which relieves the closing of pores without deleterious residual and reaction with the green body. The sintering behavior of porous SiC ceramics showed a great difference with the change in content of sintering aid (Al₂O₃) and the optimal content of Al₂O₃. In addition, the sintering temperature and hold time had great influence on sintering properties of porous SiC ceramics [136]. Glycerol was adopted as stabilizing agent to decrease the sedimentation of the raw materials during processing. Glycerol has a high viscosity of 1499 mPa s and can bind SiC particles in the pores on the surface of the yeast.

Fukushima and his coworkers [137] produced porous SiC by a gelation-freezing method. In the method, raw powder was firstly dispersed into water (which acts as a pore former) with a gelation agent. After gelation, the wet gel is immersed into cold ethanol to form ice crystals. A vacuum freeze drier is used to dry the frozen gel, so that the ice is sublimated without shrinkage of green body and then sintered to obtain porous SiC ceramics. Using the gelation-freezing technique SiC foams with high porosity (∼87%) and unidirectional oriented micrometer-sized cylindrical pores were prepared [32,137]. Typical microstructures of SiC foams produced by the gelation-freezing method are shown in Fig. 5(c) and (d) [137]. Yoon et al. [138,139] processed porous SiC ceramics from PCS/camphene solution by a freezing method. In the technique, camphene is used as a pore former. The solution prepared at 60 °C was cast into a mold at a temperature ranging from 20 to −196 °C. It results in a bi-continuous structure with regularly patterned interconnection. After the freeze casting solidification, camphene grows dendritically toward the center of the cast body. Pore size decreases by lowering freezing temperature. The sample even frozen at liquid nitrogen shows a highly porous structure, implying the effectual growth of camphene dendrites during freezing. After the removal of the frozen camphene network by sublimation, the samples showed highly porous structure. Pyrolysis of porous PCS object at 1400 °C for 1 h in argon atmosphere produced porous SiC ceramics with oriented porosity. Fitzgerald et al. [140] used sodium chloride compact for producing fine open-cell SiC foams. In this process, PCS is first pressure infiltrated into a porous sodium chloride compact, having a controlled particle size and density. The salt is subsequently leached away using water; the resulting PCS foam is cured by heating in air; and finally pyrolyzed into a SiC ceramic. Homogeneous open-cell SiC foams were produced with controlled cell sizes varying roughly between 10 and 100 μm and controlled relative densities. Wang et al. [117] used carbon fiber, carbon nanotube, and nylon filter as templates and polymethylsilane as a precursor for SiC. Template derived porous SiC plates, SiC nano-net, fiber-inverse and bead-inverse porous SiC ceramics were prepared from polymethylsilane as a SiC precursor and the above templates. The fabrication procedure involves the infiltration of the templates with appropriate concentration of the polymethylsilane, curing of the precursor, precursor pyrolysis and subsequent template removal by oxidation or chemical leaching. Park and his coworkers [142] processed porous SiC ceramics by hot-pressing a SiC nano powder-carbon nano powder-sintering additive mixture and subsequent oxidation at 700 °C in air. By adjusting the carbon content it was possible to control the porosity within a range of 30–50%. Bereciartu et al. [144] fabricated porous SiC with different pore size distributions using mesocarbon microbeads as a sacrificial template from with/without using Y₂O₃ and Al₂O₃ as sintering additives. Different sizes of SiC and carbon powders are used to investigate their effect on porosity and mechanical properties. The experimental result shows that the usage of nanometric sized SiC and C powders resulted with the porous SiC ceramics with higher porosity than micrometer sized powders. Yamane et al. [145] fabricated porous SiC from Si powder and carbon black at 900 °C for 24 h in Na vapor. They used grains of Si powder as a source of Si and template for pore forming. The porosity of porous SiC ceramic was around 55–59%. Sung et al. [147] processed macroporous SiC with a highly ordered pore array by infiltrating polymethylsilane solution into a colloidal silica template, pyrolysis of polymethylsilane, and subsequent silica template leaching using HF solution.

In summary, the sacrificial template method is the most appropriate technique to tailor the porosity, pore size distribution, and pore morphology of the final porous SiC ceramics through appropriate choice of the sacrificial template, because pore size, pore shape, and porosity are controlled by the size, shape, and content of the template used. The pore size and porosity attained in this method ranged from 1 to 700 μm and from 15% to 95%, respectively, depending on the template content and processing conditions. The replica method generally produces open-cell macroporous structure, whereas the sacrificial template method is preferred to produce micro or macrocellular structures with interconnected porosity [34,121]. A gelation-freezing process using water as a fugitive template can produce SiC foams with excellent compressive strength (17 MPa) at high porosity (∼86%) because of the formation of unidirectional oriented micrometer-sized cylindrical pores [137].

2.4 Direct foaming method

Direct foaming usually generates bubbles inside ceramic slurries containing SiC powders or inside polymeric precursor solutions to create SiC foams. In this technique, ceramic or precursor suspension is foamed with blowing agents to create stabilized foam, dried and subsequently sintered to obtain macroporous SiC ceramics (Fig. 1(d)). The blowing agents can be volatile liquids, gas evolving solids, or gases that can be evolved in situ by chemical reactions or can be added to the liquid mixture by mechanical stirring or bubbling (gas injection). Nucleation and growth of the gas bubbles inside the precursor liquids or ceramic slurries are influenced by the presence of suspended particles and the viscosity of the slurries. The blowing agents are classified into physical and chemical blowing agents. Chemical blowing agent produces gaseous products by a chemical reaction whereas the physical blowing agent does not accompany chemical reactions and the bubble/foam-making process is reversible. Table 4 illustrates some examples of possible blowing agents for producing porous SiC ceramics by direct foaming technique.

Several routes have been developed in last decades to prepare porous SiC ceramics using direct foaming method. Porous SiC ceramics can be produced by using thermosetting properties of Si-based polymers in combination with in situ blowing agents, either in the presence of surfactants [42,148] or by applying the pressure-drop technique [151,152] or by applying steam chest molding [154]. Mouazer et al. [42] introduced foaming-gel casting process with surfactants (Triton X-114 and Tergitol TMN-10) as a blowing agent for the production of porous SiC ceramics. They investigated the effect of the amount of agar (gelling agent) on the final porosity of the ceramic foams. The porosity increased with a decrease in the amount of agar. The suspension viscosity affects the foam volume and consequently the final density of the samples. By changing the agar concentration, it was possible to fabricate SiC foams with a wide range of porosity from 78% to 88% without cracks. Fukushima and Colombo [148] fabricated porous SiC foams from direct blowing of PCS with a chemical blowing agent (azodicarbonamide). This process was simple and efficient to produce porous SiC foams with tailored pore architecture and porosity. Porosity ranging from 59% to 85%, and cell size ranged from 416 to 1455 μm was obtained from this method. A typical microstructure of SiC foams produced by the direct blowing of PCS is shown in Fig. 6(a) [148]. Bao et al. [149] processed SiC foams by using volatiles, as a pore former, which were generated by the decomposition of the polymeric precursors. It was possible to control the porosity of the SiC ceramics by controlling the gas evaporation by tailoring the composition and structure of the polymeric precursors. Good shape retention was obtained in the polysilane precursors with higher functionalities and pyrolytic yields between 50 and 60 wt%. Qiu et al. [150] processed SiC foams with nano-sized grains by combustion synthesis of gel-cast Si/C foam containing n-octylamnie as a blowing agent. The as-synthesized SiC foam has a high porosity in the range of 70–90%.

Kim et al. [151] introduced a new route in direct foaming to fabricate porous SiC from preceramic precursors using CO₂ as a blowing agent. The subsequent steps involves: (i) saturation of preceramic polymers using gaseous, liquid, or supercritical CO₂, (ii) nucleation and growth of a large number of bubbles using thermodynamic instability via a rapid pressure drop and/or heating, and (iii) transforming the microcellular preceramic polymers into microcellular ceramics by pyrolysis and (optional) subsequent sintering [151^[152]–153,210]. The experimental result reveals that if the microcellular structures are preserved during the transformation of the preceramic to ceramic, porous ceramics with an extremely fine and homogeneous cell structure could be fabricated. Microcellular SiC ceramics having cell density above 10⁹ cells/cm³ and cell size below 10 μm were fabricated using PCS and/or polysiloxane precursors [151]. The same group also introduced a new processing route in direct foaming by applying steam chest molding [154]. The fabrication process involved the following steps: (i) blending a mixture of polysiloxane, carbon, hollow microspheres, and sintering additives, (ii) steam chest molding of the mixture, (iii) transforming the polysiloxane by pyrolysis into silicon oxycarbide, and (iv) fabricating SiC foams by carbothermal reduction of silicon oxycarbide and subsequent sintering. Porosity of the foams could be controlled by controlling the initial packing density in the steam chest molding step. A representative microstructure of SiC foams produced by the steam chest molding is shown in Fig. 6(b) [154]. It shows a uniform cellular structure with open-cells.

Processing and properties of macroporous silicon carbide ceramics: A review

All authors

Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 6 Typical microstructures of macroporous SiC ceramics produced via the direct foaming technique using (a) azodicabonamide as a chemical blowing agent [148] and (b) water vapor as a physical blowing agent [154].

(a) is reproduced with the permission of Elsevier and (b) is reproduced with the permission of John Wiley and Sons.

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In summary, the direct foaming technique is an easy and fast way to prepare both open and closed-cell structures with a wide range of pore size and porosity, up to 95% [32,34,35]. The structures have well interconnected porosity with unique permeability that allows finer adjustment of fluid transport within the structure [34,211]. The total porosity of directly foamed ceramics is proportional to the amount of gas incorporated into the suspension or liquid medium during the foaming process. The main advantage of this method is the possibility of producing dense ceramic struts and defect-free struts, providing a stronger foam in comparison with the replica method. One drawback of this method is the difficulty in producing a material with a narrow pore size distribution [28,32,151]. As long as PCS or polysiloxane is used as a precursor to SiC, the preceramic polymers should be blown first before cross-linking. Otherwise, direct foaming of the preceramics polymers would be difficult in the processing of porous SiC ceramics via direct foaming.

2.5 Bonding technique

Porous SiC ceramics can be produced mainly by partial sintering [212], replica [213,214], sacrificial template [215], and direct foaming [66,216] methods. Different processing routes for porous SiC ceramics have been developed for specific applications to satisfy the associated requirements of porosity, pore size, and degree of interconnectivity. However, all the above methods require high temperatures of ≥1500 °C for the processing if SiC powder is used as a starting material because of its strong covalent bonding characteristic. Thus, one of the key issues in the processing of porous SiC ceramics is how to produce porous SiC ceramics at relatively low temperatures. One way to lower the processing temperature is to use preceramic polymers as a precursor for SiC. Another way to lower the processing temperature of porous SiC ceramics is the use of bonding materials. The in situ reaction bonding process can realize the low-temperature fabrication of porous SiC ceramics (Fig. 1(e)) [155,156]. Alkali, cordierite (2MgO·2Al₂O₃·5SiO₂), mullite (3Al₂O₃·2SiO₂), silica (SiO₂), Si, SiC, silicon nitride (Si₃N₄), silicon oxycarbide (SiOC), and frit phases were investigated as bonding materials for porous SiC ceramics. Porous SiC ceramics containing a bonding phase and moderate porosity possess a unique set of characteristics, such as high thermal shock resistance, chemical stability, high specific strength, and cost-effectiveness. Table 5 illustrates some examples of possible bonding materials for producing porous SiC ceramics by the bonding technique.

2.5.1 Porous mullite-bonded silicon carbide

It has been reported that the oxidation resistance of recrystallized SiC was improved by the incorporation of a mullite coating on the SiC [217]. Similarly, the corrosion resistance of SiC in contact with coal slag was improved by introducing a chemically vapor-deposited mullite coating on the SiC [218]. Ando et al. [219] reported that mullite-SiC composites have the ability to heal semicircular cracks having diameters of up to 200 μm. A new approach that has received much attention is the incorporation of mullite as a bonding phase in SiC. Silicon carbide has good hardness, strength, and wear resistance, whereas mullite (3Al₂O₃·2SiO₂) has good oxidation resistance and high temperature stability. Furthermore, mullite has good chemical compatibility with SiC, and both materials have similar thermal expansion properties [220]. Thus, the combination of mullite and SiC to form mullite-bonded SiC is an attractive approach and the porous ceramics are expected to exhibit better properties than porous mullite or porous SiC for some applications in oxidizing atmosphere.

Porous mullite-bonded SiC ceramics are fabricated from SiC, Al sources such as Al, AlN, Al₂O₃, and Al(OH)₃, and templates using an in situ bonding technique [156^[157]–168,221]. It has been reported that the oxidation-derived SiO₂ and Al sources react in situ to fabricate porous mullite (3Al₂O₃·2SiO₂)-bonded SiC ceramics at 1400–1550 °C in air [156,157,159,162,165]. Meanwhile, the templates are decomposed or burned out, leaving pores inside the materials. She et al. [156] firstly developed a unique technique to produce mullite-bonded porous SiC ceramics by an in situ oxidation bonding technique. The surface of SiC particles is oxidized to form SiO₂ with high reaction activity and the oxidation-derived SiO₂ can easily react with the added Al₂O₃ and form a mullite to bond SiC particles at 1400–1500 °C in air. Ding and coworkers [158] processed porous mullite-bonded SiC by the same technique, in situ oxidation bonding technique with an addition of Y₂O₃. The addition of Y₂O₃ enhanced neck formation (bonding part between SiC particles), resulted in improved mechanical properties. They suggested reaction steps involved in the formation of porous mullite-bonded SiC ceramics:(4) $\text{SiC}\text{+}{\text{2O}}_{\text{2}}\to {\text{SiO}}_{\text{2}}\text{(amorphous)}\text{+}{\text{CO}}_{\text{2}}$(4) (5) $\text{Amorphous}{\text{SiO}}_2\to \text{Cristobalite}$(5) (6) $3{\text{Al}}_{\text{2}}{\text{O}}_{\text{3}}\text{+}{\text{2SiO}}_{\text{2}}\to 3{\text{Al}}_{\text{2}}{\text{O}}_{\text{3}}\text{.}{\text{2SiO}}_{\text{2}}\text{(mullite)}$(6) The formation of mullite phase in between SiC grains can be explained on the basis of in situ reactions. SiC was oxidized to amorphous silica that further crystallized to form cristobalite during sintering at higher temperatures in air. The obtained cristobalite reacts with alumina upon higher temperatures (beyond 1400 °C) to form mullite [155,158,160].

Liu et al. [161] investigated the effect of preheat-treated aluminosilicate (PHAS) addition on the mullitization process and pore structure in the processing of porous mullite-bonded SiC by in situ bonding technique. The addition of 5 wt% PHAS into SiC–Al₂O₃–graphite batches promoted the mullitization process and the process was almost completed at 1450 °C and resulted in the enhancement of neck growth. The open porosity was reduced by the addition of PHAS, whereas the average pore size from the burnout of graphite was enlarged by the addition of PHAS. Kumar et al. [164,166] processed porous mullite-bonded SiC ceramics with four different Al sources (Al/AlN/Al₂O₃/Al(OH)₃) at 1450–1550 °C for 1–6 h duration in air. They suggested Al₂O₃ formation reactions from various Al sources as follows:(7) $2\text{Al}+1.5{\text{O}}_2\to {\text{Al}}_{\text{2}}{\text{O}}_{\text{3}}$(7) (8) $\text{2AlN}\text{+}\text{1}{\text{.5O}}_{\text{2}}\to {\text{Al}}_{\text{2}}{\text{O}}_{\text{3}}\text{+}{\text{N}}_{\text{2}}$(8) (9) ${\text{2Al(OH)}}_3\to {\text{Al}}_{\text{2}}{\text{O}}_{\text{3}}+3{\text{H}}_{\text{2}}O$(9) The oxidation derived cristobalite reacts with Al₂O₃ or Al source-derived Al₂O₃ upon increasing the temperature (beyond 1400 °C) to form mullite according to Eq. (6) [158]. The mullite-silica grains were well connected and strongly adhered with large SiC grains in porous mullite-bonded SiC ceramics by using Al₂O₃, whereas poor adhesion and weak bonding was found from usage of AlN. The minimum porosity of 17% and maximum porosity of 42% were obtained for the samples prepared with Al₂O₃ and AlN, respectively, when sintered at 1550 °C for 6 h and 1450 °C for 1 h, respectively [164].

Choi et al. [163,165] fabricated porous mullite-bonded SiC ceramics at temperatures ranging from 1350 to 1450 °C for 2 h using SiC, Al₂O₃, alkaline-earth metal oxides, and polymer microbeads. The addition of alkaline-earth metal oxides was beneficial for lowering the mullitization temperature from 1450–1550 °C to 1350–1450 °C and for the formation of dense strut, which should lead to the improved mechanical strength. Dey et al. [221] processed porous mullite-bonded SiC ceramics by infiltrating a SiC–Al₂O₃ powder compact with a liquid precursor of SiO₂ and subsequent heat treatment at 1500 °C. The bond phase was composed of needle-shaped mullite which was grown from a siliceous melt formed during the heat-treatment. The obtained porosity was 30% and the pore size was ranged from 2 to 15 μm with an average pore size of 5 μm. Kayal et al. [167,168] fabricated porous mullite-bonded SiC ceramics by infiltrating SiC-petroleum coke powder compacts with a liquid precursor of mullite and subsequent sintering at 1300–1500 °C for 4 h in air. The porous mullite bonded SiC ceramic was composed of SiC, mullite, and cristobalite. The porosity could be varied in the range of 31–49%, depending on the amounts of infiltrate and the petroleum coke powder, and the average pore diameter (∼8 μm) was almost independent of porosity.

Processing of porous mullite-bonded SiC is one of the low-temperature processing technique for porous SiC ceramics and the processing temperature was in the range of 1300–1550 °C in air. A typical microstructure of the porous mullite-bonded SiC is shown in Fig. 7(a) [159]. The obtained porosity was in the range of 17–55%, depending on the template content, composition of Al source, and sintering temperature.

Processing and properties of macroporous silicon carbide ceramics: A review

All authors

Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 7 Typical microstructures of macroporous SiC ceramics produced via the bonding technique: (a) mullite-bonded SiC foam [159], (b) silica-bonded SiC foam [223], (c) self-bonded SiC foam, and (d) SiOC-bonded SiC foam [190]. (a) and (b) are reproduced with permission of Elsevier. (d) is reproduced with permission from the Springer.

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Fig. 7 Typical microstructures of macroporous SiC ceramics produced via the bonding technique: (a) mullite-bonded SiC foam [159], (b) silica-bonded SiC foam [223], (c) self-bonded SiC foam, and (d) SiOC-bonded SiC foam [190]. (a) and (b) are reproduced with permission of Elsevier. (d) is reproduced with permission from the Springer.

3.1 Room temperature flexural and compressive strengths

Mechanical properties of porous SiC ceramics depend on their porosity, pore size, microstructure of the strut, and compositions of bonding phase and additives [28,46,208]. Each processing method to produce porous SiC is best suited for producing a specific range of porosity, pore size, and microstructure of the strut [28,170,239]. The effects of porosity, pore size, microstructure of the strut, and compositions of bonding phase and additives on flexural and compressive strengths are discussed in this session.

The most influential parameter on the flexural strength of porous SiC ceramics is porosity. The collection of flexural strength data for various porous SiC ceramics produced via different processing strategies is shown as a function of porosity in Fig. 8. The porosity varied from 9% to 91% and the flexural strength varied from 0.7 to 205 MPa. From the plot, the following observations were made: (1) the flexural strength generally decreases with increasing porosity; (2) at a porosity range of 30–50%, the template technique yields porous SiC ceramics with better strength than the bonding technique; and (3) at a porosity range of 50–80%, the replica technique (mostly from wood templates) provides porous SiC ceramics with higher strength than both the partial sintering and the template techniques.

Processing and properties of macroporous silicon carbide ceramics: A review

All authors

Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 8 Flexural strength as a function of porosity of macroporous SiC ceramics produced via different processing strategies. Data points are labeled with the corresponding reference numbers.

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In porous SiC ceramics fabricated by partial sintering technique, a maximum flexural strength of 152 MPa at 65% porosity [58] and a minimum flexural strength of 5.2 MPa at 67% porosity [55] were obtained depending on the various processing parameters. Based on various processing parameters in replica technique, a minimum flexural strength of 0.7 MPa at 86% porosity [70] and a maximum flexural strength of 205 MPa at 9% porosity [113] were reported. In porous SiC ceramics processed by the sacrificial template technique, a maximum flexural strength of 170 MPa at 50% porosity [141] and a minimum flexural strength of 2.4 MPa at 77% porosity [45] were obtained. The flexural strength data of porous SiC ceramics processed by the direct foaming were quite limited and a maximum and a minimum strengths were 6.2 MPa at 86% porosity and 2.9 MPa at 88% porosity, respectively [42]. In porous SiC ceramics fabricated by bonding technique, a maximum flexural strength of 185 MPa at 31% porosity [169] and a minimum flexural strength of 3.6 MPa at 82% porosity [55] were obtained depending on both the bonding method and processing parameters.

According to the model proposed by Rice [240], the strength of a porous material is related to its porosity through the expression.(14) $\sigma ={\sigma }_{\text{o}}\exp (-bP)$(14) where σ and σ_o are strengths of a porous material with porosity P and nonporous material, respectively, and b is a constant that is dependent on pore characteristics. The model does not expect any dependence of strength on pore size, but a variation only with porosity of porous ceramics. However, the pore size affects the flexural strength of porous SiC ceramics. The effect of pore size on flexural strength of porous SiC ceramics has been investigated systematically by Eom and Kim [46]. They fabricated porous SiC ceramics with different pore sizes by using three different sizes of templates. The obtained porosity of the porous SiC ceramics almost did not change with the size of template at the same content of template. Thus, the porous ceramics have almost same porosity at the same content of template, but different pore sizes. Typical flexural strengths of the porous SiC ceramics with 45% porosity were 55 MPa for the porous SiC with a pore size of 7 μm, 47 MPa for the porous SiC with a pore size of 14 μm, and 20 MPa for the porous SiC with a pore size of 31 μm. Thus, the flexural strength of porous SiC ceramics increased with decreasing the pore size at the same porosity. The result proved that the flexural strength of porous SiC ceramics increases with decreasing the pore size, because a decrease in pore size by using a smaller template decreased the critical flaw size of the porous SiC ceramics. Spherical templates with a small diameter would be beneficial for preparing high-strength, porous SiC ceramics.

Effect of microstructure on the flexural strength of porous SiC ceramics has been investigated by Eom et al. [47,118,208] and Kennedy et al. [177]. Eom et al. [47,118] fabricated porous SiC ceramics with various microstructures by adjusting the initial α-SiC content in the starting composition by the template method. When β-SiC powder or a mixture of α/β powders containing a small amount (≤10%) of α-SiC powder was used, the microstructure consisted of large platelet grains (Fig. 5(a)). In contrast, when using α-SiC powder or α/β powders containing a large amount (>10%) of α powders, the microstructure consisted of small equiaxed grains. The microstructural development was insensitive to the additive composition (Al₂O₃–Y₂O₃ or Al₂O₃–Y₂O₃–CaO) and the additive content (5% or 7%) [47,118]. A maximum strength of 75 MPa was obtained in the porous SiC with a microstructure consisting of equiaxed grains, whereas a minimum strength of 30 MPa was obtained in the porous SiC with a microstructure consisting of large platelet grains [118]. Kennedy et al. [177] fabricated porous self-bonded SiC ceramics with different grain sizes by adjusting the initial SiC particle size. They used three different sizes of micron-size SiC powders (25 μm, 50 μm, 65 μm) and a submicron-size SiC powder (0.7 μm) as a bonding phase. The porous SiC ceramics fabricated from a 25 μm-size powder showed a maximum flexural strength of 35.5 MPa at 46% porosity, indicating that the porous SiC ceramics with smaller grain size is beneficial in improving the flexural strength of porous self-bonded SiC ceramics at the equivalent porosity.

Eom et al. [129,130] investigated the effect of additive composition on flexural strength of porous SiC ceramics, which were processed by the template method. Variation of additives led to different porosities in the range of 56–72% and different flexural strengths in the range of 8 MPa to 122 MPa. However, some specimens showed different flexural strengths at the same porosity. For example, the flexural strength of porous SiC ceramics sintered with 3 wt% Al₂O₃–2 wt% Y₂O₃ was 55 MPa at 62% porosity, whereas that of porous SiC ceramics sintered with 3 wt% Al₂O₃–1 wt% Y₂O₃–1 wt% CaO was 45 MPa at the same porosity [129]. Choi et al. [163] fabricated porous mullite-bonded SiC ceramics with 4 wt% alkaline earth oxide additives. The addition of alkaline earth oxides was beneficial for lowering the sintering temperature from 1450–1550 °C to 1350–1450 °C and for improving the flexural strength dramatically. The mullite-bonded SiC ceramics sintered with 4 wt% SrO showed a flexural strength of 44 MPa at 46% porosity. In contrast, the flexural strength of the mullite-bonded SiC ceramics sintered with no additives was 6 MPa at 49% porosity. The results suggest that judicious selection of sintering additive composition is an efficient way to improve the mechanical properties of porous SiC ceramics.

Since the bonding technique has a wide variation in flexural strength at the equivalent porosity (Fig. 8), the flexural strength data belong to each bonding material was plotted as a function of porosity in Fig. 9. From the plot, we can find the following observations: (1) at a porosity range of 10–30%, porous SiOC-bonded SiC ceramics showed greater flexural strength than the mullite-bonded, silica-bonded, and cordierite-bonded SiC ceramics; (2) at a porosity range of 30–40%, porous silica-bonded SiC ceramics showed better flexural strength than the other materials; and (3) at a porosity range of 40–50%, porous mullite-bonded SiC ceramics showed better flexural strength than the other materials.

Processing and properties of macroporous silicon carbide ceramics: A review

All authors

Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 9 Flexural strength as a function of porosity of macroporous SiC ceramics produced with different phases. Data points are labeled with the corresponding reference numbers.

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The collection of compressive strength data for various porous SiC ceramics produced via different processing strategies is shown as a function of porosity in Fig. 10. The porosity varied from 30% to 90% and the compressive strength varied from 0.8 to 600 MPa. From the compressive strength versus porosity plot, we made the following observations: (1) the compressive strength generally decreases with increasing porosity; (2) at a porosity range of 30–40%, the partial sintering and the sacrificial template methods yield porous SiC ceramics with better strength than the other processing technique; and (3) at a porosity range of 40–55%, the sacrificial template method provides porous SiC ceramics with higher strength than the other methods. In porous SiC ceramics fabricated by partial sintering technique, a maximum compressive strength of 513 MPa with 39% porosity [63] and a minimum compressive strength of 61 MPa with 54% porosity [54] were obtained depending on the processing parameters. Based on various processing parameters in replica technique, a minimum compressive strength of 0.3 MPa with 90% porosity [73] and a maximum compressive strength of 70 MPa with 55% porosity [83] were reported. In porous SiC ceramics processed by sacrificial template technique, a maximum compressive strength of 600 MPa with 30% porosity [46] and a minimum compressive strength of 4.8 MPa with 78% porosity [139]. The compressive strength data of porous SiC ceramics processed by direct foaming were quite limited and a maximum and a minimum strengths were 77 MPa at 59% porosity and 51 MPa at 62% porosity, respectively [154]. In porous SiC ceramics fabricated by the bonding technique, a compressive strength of 200 MPa at 40% porosity in silica-bonded SiC [170] and a compressive strength of 0.9 MPa at 81% porosity [182] in alkali-bonded SiC were reported.

Processing and properties of macroporous silicon carbide ceramics: A review

All authors

Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 10 Compressive strength as a function of porosity of macroporous SiC ceramics produced via different processing strategies. Data points are labeled with the corresponding reference numbers.

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From Figs. 8–10, it is evident that the flexural and compressive strength generally decreases with increasing porosity in porous SiC ceramics. This tendency has also been observed in many other porous ceramics [241^[242]–244] and can be attributed to the pore coalescence under load at higher porosities. The pore coalescence increases the defect size and reduces the strength [34]. From the model proposed by Rice (Eq. (14)) [240], the value b was reported to be 6 for cubic stacking and 9 for rhombic stacking [245]. The strength data of porous SiC ceramics processed by various methods showed different b values. For example, the reported b values were 3.4 in porous SiC ceramics fabricated by the partial sintering [44], 3.19–3.24 in porous SiC ceramics fabricated by the template method [141,142], 4.36 in porous mullite-bonded SiC [156], 6.54–7.95 in porous silica-bonded SiC ceramics [169,170], 4.8 in porous cordierite-bonded SiC [140], and 6–16 in porous SiOC-bonded SiC [190]. The observed deviation in the b values suggests a more pronounced influence of density on the strength than that predicted by the model [34]. This can probably attributed to the following factors: (1) the model assumed a constant cell size, whereas the distribution of cell sizes are often found in porous SiC ceramics, especially porous ceramics processed by the bonding technique [156,170,185,190]; (2) the presence of both open and closed pores, which should have some effect on the stress distribution near pores [34]; (3) the variation of solid area for each specimen irrespective of the material porosity [190]; and (4) the presence of abnormally large pores or defects in struts, which lead to the unexpectedly low flexural strength [34].

In summary, from a collection of flexural and compressive strength data for various porous SiC ceramics produced via different processing strategies, the following observations were made: (1) both the flexural and compressive strengths generally decrease with increasing porosity; (2) at a specific porosity range, the strength of porous SiC ceramics are strongly dependent on the processing method; (3) an excellent flexural strength of 152 MPa at 65% porosity was obtained in porous SiC ceramics fabricated by the partial sintering technique; and (4) compressive strengths of 513 MPa at 39% porosity and 77 MPa at 59% porosity were reported in porous SiC ceramics processed by partial sintering and by direct foaming, respectively.

3.2 High temperature flexural strength

The collection of flexural strength data at high temperatures for various porous SiC ceramics is shown in Fig. 11. Singh and Salem [105] measured flexural strength of porous biomorphic SiC ceramics fabricated by the replica method at temperatures from RT to 1350 °C. Flexural strengths of SiC ceramics fabricated from maple wood precursors were 344 MPa and 230 MPa at RT and 1350 °C, respectively. In contrast, those of SiC ceramics fabricated from mahogany wood precursors were 144 MPa and 112 MPa at RT and 1350 °C, respectively. The RT-strength retention of maple-based and mahogany-based SiC ceramics was 67% and 78% at 1350 °C, respectively. The flexural strength of porous SiC ceramics, which were fabricated by CVD of SiC onto reticulated carbon skeletons obtained by pyrolysis of polyurethane foams, was 19 MPa at 200 °C and the material maintained its RT-strength up to 1400 °C [81]. The excellent strength retention of the porous CVD-SiC was attributed to its high purity, probably >99.9% pure. The flexural strength values of porous mullite-bonded SiC ceramics with 30% porosity were 47 MPa at RT, 49 MPa at 900 °C, and 58 MPa at 1100 °C [221]. The increase of flexural strength at 1100 °C was due to the crack healing of porous SiC ceramics by oxidation of SiC during flexural strength testing.

Processing and properties of macroporous silicon carbide ceramics: A review

All authors

Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 11 High temperature flexural strength of macroporous SiC ceramics produced via different processing strategies. Data points are labeled with the corresponding reference numbers.

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In summary, wood-derived porous SiC ceramics processed by the replica method maintained 67–78% of their room temperature strengths at 1350 °C, whereas porous SiC ceramics fabricated by CVD of SiC on reticulated carbon foams obtained by pyrolysis of polyurethane foams (replica method) maintained their RT-strength up to 1400 °C. The excellent strength retention of the porous CVD-SiC was attributed to its high purity (99.9% pure).

3.3 Fracture toughness

The collection of fracture toughness data for various porous SiC ceramics produced via different processing strategies is shown as a function of porosity in Fig. 12. The porosity varied from 8 to 95% and the fracture toughness varied from 0.2 to 4.3 MPa m^1/2. From the fracture toughness versus porosity data for porous SiC ceramics produced via different processing strategies, the following observations were made: (1) the fracture toughness generally decreases with increasing porosity; and (2) at a porosity range of 40–70%, the replica and partial sintering techniques yield porous SiC ceramics with higher toughness than the template method.

Processing and properties of macroporous silicon carbide ceramics: A review

All authors

Jung-Hye Eom, Young-Wook Kim & Santosh Raju

https://doi.org/10.1016/j.jascer.2013.07.003

Published online:

18 April 2018

Fig. 12 Fracture toughness as a function of porosity of macroporous SiC ceramics produced via different processing strategies. Data points are labeled with the corresponding reference numbers. The fracture toughness values obtained in Refs [37,47,81], Ref [58], and Ref. [64] were measured by a single-edge notched beam, a chevron-notched beam, and an indentation-strength method, respectively.

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The obtained fracture toughnesses of porous SiC ceramics processed by partial sintering and by replica methods were 0.9 MPa m^1/2 at 64% porosity and 1.4 MPa m^1/2 at 50% porosity, respectively [58,64]. In porous SiC ceramics fabricated by the replica method, the fracture toughness ranged from 0.2 MPa m^1/2 at 95% porosity [81] to 1.4 MPa m^1/2 at 50% porosity [38]. The fracture toughness data of porous SiC ceramics fabricated by template method were quite limited and the toughness values were 0.3–0.4 MPa m^1/2 at 57–58% porosity [47]. Comparison of the fracture toughness values obtained in template method with values obtained in the other methods indicates that the toughness values obtained in template method are lower than those of the other methods. A crack-tip blunting mechanism contributes to the toughening of porous ceramics [58]. However, the microstructure of the porous SiC ceramics fabricated by template method by Eom et al. [47] consisted of large platelet grains, as shown in Fig. 5(a). The growth of platelet SiC grains in the porous SiC ceramics formed sharp edges and V-shaped notches at the contacting region of two platelet grains in the pores. The formation of sharp edges and V-shaped notches in the pore surfaces hindered the contribution of the crack blunting mechanism in the porous ceramics, resulting in lower fracture toughness [47]. The above results suggest that the fracture toughness of porous SiC ceramics can be improved by taking advantage of crack blunting mechanism, which can be accomplished by controlling pore morphology and grain arrangement [58].

Generally, the fracture toughness values are dependent on the measurement methods. It should be noted that the fracture toughness values plotted in Fig. 12 were measured in different methods: the data from Refs [38,47,81], Ref. [58], and Ref. [64] were measured by a single-edge notched beam, a chevron-notched beam, and an indentation-strength method, respectively. Thus, it cannot be ruled out that the difference in the toughness values was partially originated from the difference in the measuring method.