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Preparation of porous SiC-Al2O3 ceramics via gelcasting utilising a shrinkable pore-forming agent and oxidised coarse-grained SiC
Author’s Accepted Manuscript Preparation of porous SiC-Al2O3 ceramics via gelcasting utilising a shrinkable pore-forming agent and oxidised coarse-grained SiC Li Wan, Qingquan Tian, Yun Xiang, Lin Chen, Jieguang Song, Xianzhong Wang, Xinshuang Guo www.elsevier.com/locate/ceri
PII: DOI: Reference:
S0272-8842(18)33375-3 https://doi.org/10.1016/j.ceramint.2018.12.008 CERI20231
To appear in: Ceramics International Received date: 30 September 2018 Revised date: 30 November 2018 Accepted date: 2 December 2018 Cite this article as: Li Wan, Qingquan Tian, Yun Xiang, Lin Chen, Jieguang Song, Xianzhong Wang and Xinshuang Guo, Preparation of porous SiC-Al 2O3 ceramics via gelcasting utilising a shrinkable pore-forming agent and oxidised coarse-grained SiC, Ceramics International, https://doi.org/10.1016/j.ceramint.2018.12.008 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Preparation of porous SiC-Al2O3 ceramics via gelcasting utilising a shrinkable pore-forming agent and oxidised coarse-grained SiC Li Wana, Qingquan Tianb,c, Yun Xianga, Lin Chena, Jieguang Songa, Xianzhong Wanga, Xinshuang Guoa,c* a
Jiangxi Key Laboratory of Industrial Ceramics, Engineering Technology Research
Centre for Environmental Protection Materials and Equipment of Jiangxi Province, Pingxiang University, Pingxiang 337055, China b
School of Chemistry and Materials, Weinan Normal University, Weinan 714099, China
c
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China *
Corresponding author. [email protected]
Abstract By utilising soaked millet as a shrinkable pore-forming agent, porous silicon carbide-alumina (SiC-Al2O3) ceramics were prepared via gelcasting. The fabrication of SiC-Al2O3 ceramics based on oxidised and unoxidised coarse-grained SiC was also studied. The water swelling, drying shrinkage, and low-temperature carbonisation of the millet were investigated. We found that the shrinkage of the soaked millet was greater than that of gel body during drying, which left large gaps that prevented shrinkage stresses from destroying the gel body. Low-temperature carbonisation of the millet should be performed slowly at 220–240 °C because its expansion rate increases to 45% at 250 °C, resulting in the cracking of samples. At a constant sintering temperature, the flexural strength of the SiC-Al2O3 ceramics prepared with
SiC powders oxidised at 1000 °C was the highest, indicating that oxidised powders can successfully decrease the required sintering temperature and improve the flexural strength of composite ceramics. Based on our optimised process, porous SiC-Al2O3 ceramics were sintered at 1500 °C for 2 h. When their skeletons were fully developed, their pore sizes were in the range of 1.5–2 mm. Their porosity and flexural strength were 60.2–65.1% and 8.3–10.5 MPa, respectively. Keywords: Porous SiC-Al2O3 ceramics; Oxidised coarse-grained SiC; Pore-forming agent; Gelcasting; Flexural strength
1. Introduction Porous silicon carbide (SiC) ceramics have attracted considerable attention for their unique combination of desirable properties, including excellent thermal shock resistance, good chemical resistance, large specific surface area, and a low thermal expansion coefficient [1–5]. Because of its low self-diffusion coefficient and strong covalent bonding, the required sintering temperature for coarse-grained SiC is approximately 2100–2400 K [6, 7]. Lowering the sintering temperature of SiC ceramics is beneficial for reducing production costs and has become a research hotspot. By combining SiC with other inorganic materials, the resulting porous ceramic composites have lower required sintering temperatures and excellent physical properties [8]. There are many different types of porous SiC composite ceramics, including SiC-Al2O3 [9–12], SiC-ZrO2 [13], and SiC-mullite [14, 15] among others [16, 17], where SiC-Al2O3 ceramics are the most widely developed. Low temperature
(≤1600 °C) SiC-Al2O3 ceramics are typically prepared utilising a precursor method [18], reaction bonding [19], and liquid-phase sintering [8, 20]. Liquid-phase sintering has attracted significant attention because it is a relatively simple process. The SiC powders utilised in liquid-phase sintering are typically untreated raw powders. SiC and Al2O3 powders can be bonded utilising mullite and oxidation-derived SiO2 [21]. When the surfaces of SiC powders are oxidised to form SiO2, the benefits in terms of the formation of mullite and reduction of the required sintering temperature for SiC-Al2O3 ceramics is still unclear. The pore-forming agent method is another common method for preparing porous ceramics based on its fabrication simplicity, practicability, and economic viability [22–25]. There are many types of solid pore-forming agents and they are typically combined with traditional slip casting methods to prepare porous ceramics [26–31]. However, solid pore-forming agents typically do not shrink during drying, which hinders the shrinkage of the green body and leads to various defects, such as the breakage of the skeletons of green bodies. Therefore, it is necessary to utilise shrinkable pore-forming agents to prepare porous ceramics with good structural integrity. Additionally, the prepared green bodies have low strength and can be easily destroyed during demoulding and drying. Gelcasting, a near-net technique with high yields, has been widely adopted to prepare high-strength green bodies with homogenous microstructures [32–34]. The combination of gelcasting and the pore-forming agent methods can prevent cracking during demoulding and drying. In this study, porous SiC-Al2O3 ceramics were prepared via gelcasting utilising
millet as a pore-forming agent. Millet was chosen for its water-swelling characteristics and approximately circular shape. We thoroughly investigated the drying shrinkage and low-temperature carbonisation process of the millet. SiC-Al2O3 ceramics were prepared utilising both oxidised and unoxidised coarse-grained SiC particles and their flexural strengths under various sintering conditions were investigated. Finally, the structures and properties of the porous SiC-Al2O3 ceramics prepared utilising an optimised process were studied. 2. Experiments 2.1 Materials and sample preparation Commercial SiC powders (≥99%, D50=37.6 µm) and Al2O3 powders (99.8%, D50=4.3 µm) were utilised as the two starting materials for preparing SiC-Al2O3 ceramics. To investigate the effects of the oxidation of SiC on the sintering properties of composite ceramics, SiC powders were oxidised at 1000 and 1100 °C for 2 h in air. The treated powders were labelled as P1000 and P1100, and the untreated SiC powders were labelled as P0. The ceramic green bodies created by mixing P0, P1000, and P1100 with Al2O3 were labelled as C0, C1000, and C1100, respectively. Fig. 1 presents a flow chart for fabricating porous SiC-Al2O3 ceramics. First, acrylamide
(AM,
monomer),
N,N’-methyleneacrylamide
(cross-linker),
and
polyethylene glycol (dispersant) with a weight ratio of 30:3:2 were added to deionised water. After stirring to obtain premixed solutions, SiC and Al2O3 powders with a weight ratio of 60:40 were added to prepare a slurry with a solid content of 45 vol%. The
slurry
was
ball-milled
for
4
h
and
then
evacuated
for
2
h.
N,N,N’,N’-tetramethylethylenediamine (TEMED, catalyst) and ammonium persulfate (APS, initiator) were then added to the degassed slurry. The amounts of TEMED and APS relative to the weight of AM were 0.1 wt% and 3.7 wt%, respectively. The chemical reagents above were procured from Sinopharm Chemical Reagent Co., Ltd. Following mechanical stirring for 2 min, the as-prepared slurry was cast into a glass mould with a millet template (Φ25 mm×50 mm) and consolidated to form green bodies. Next, the prepared green bodies were dried 25 °C for 24 h and 50 °C for 48 h. The dried samples were then carbonised at low temperature and subjected to pressureless sintering in air to obtain porous SiC-Al2O3 ceramics. It is worth noting that an optimised process for fabricating SiC-Al2O3 ceramics was first identified and then the porous SiC-Al2O3 ceramics were prepared according to the optimised process. 2.2 Characterisation The average diameters of the SiC and Al2O3 powder particles were measured utilising a laser particle size analyser (BT-9300H, Dandong Bettersize Instruments Ltd, China). Thermal analysis of the millet (SiC-Al2O3 ceramics) was performed utilising a thermal analyser (STA449, Netzsch Inc., Germany) at a heating rate of 1 (10) °C/min from room temperature to 650 (1550) °C in air. The morphologies and crystal structures of the samples were analysed via scanning electron microscopy (SEM, S8010, Hitachi Inc., Japan) and X-ray diffraction (XRD, D8-Advance, Bruke Inc., Germany), respectively. The porosity and bulk density of the porous SiC-Al2O3 ceramics were measured based on the Archimedes method. The thermal shock
resistance was performed by the water-quenching method, which was referred to the literature [35]. The residual flexural strength of the samples was tested to evaluate the thermal shock resistance of the porous SiC-Al2O3 ceramics. The heat conductivity was measured using the thermal conductivity meter (DRPL-II, Xiangtan Xiangyi instrument Co., Ltd., China). The flexural strength of the ceramics (dimension of 3 mm × 4 mm × 40 mm) was measured utilising a universal testing machine (AGS-X-10KN, Shimadzu Inc., Japan) with a span of 30 mm and loading rate of 0.5 mm/min. 3. Results and discussion 3.1 Optimisation of the process parameters for SiC-Al2O3 ceramics Fig. 2 presents the flexural strengths of the SiC-Al2O3 ceramics prepared with different types of SiC powders. The ceramics were sintered at 1400, 1450, and 1500 °C for 2 h, respectively. When the sintering temperature was 1400 °C, the flexural strengths of the composite ceramics prepared from C1000, C0, and C1100 were 17.6, 17.1 and 14.1 MPa, respectively, meaning the flexural strength of the ceramic prepared from C1100 was the lowest. According to the thermogravimetry (TG)–differential scanning calorimetry (DSC) analysis results presented in Fig. 3(a), the exothermic peak at 1455 °C can be attributed to the formation of mullite through a reaction between the SiO2 obtained from SiC oxidation and Al2O3. Therefore, the different flexural strengths of the three types of composite ceramics are largely related to the connections between SiC particles created by cristobalite at 1400 °C [21]. Regarding the composite ceramics
prepared from C1000, the surfaces of SiC powders treated at 1000 ℃ are relatively uniform (Fig. 3(b). The proper amount of liquid phase was formed at 1400 °C, which is beneficial for close bonding between cristobalite and Al2O3 (Fig. 3(c)). As a result, the composite ceramics prepared from C1000 have the highest flexural strength. Regarding the composite ceramics prepared from C1100, the SiC powders were excessively oxidised at 1100 °C [36], resulting in the generation of partial liquid phase (indicated by white circles) and certain ablation (indicated by black circles) on their surfaces (Fig. 3(d)). Excessive liquid phase consequently generates in the composite ceramics at 1400 °C, which is not beneficial for bonding between Al2O3 and cristobalite (Fig. 3(e)). Therefore, the flexural strength of the composite ceramics prepared from C1100 is lower than that of the ceramics prepared from C1000. When the sintering temperature is 1450 or 1500 °C, the bonding of ceramic particles mainly depends on mullite and cristobalite. Therefore, the flexural strength of the composite ceramics sintered at 1450 and 1500 °C is greater than that of the ceramics sintered at 1400 °C. Additionally, the composite ceramics prepared from C1000 have the highest flexural strength, followed by those prepared from C1100 and C0 (Fig. 2). Some possible reasons for these results are as follows. Regarding the composite ceramics prepared from C1100, with the formation of mullite and oxidation of SiC, various interface defects, such as voids between SiC and Al2O3, can be eliminated, leading to relatively tight bonding (Fig. 3(f)), which results in an increase in flexural strength. Regarding the composite ceramics prepared from C0, because there is less SiO2 on the surface of the unoxidised SiC and less mullite formed
between the SiC and Al2O3 particles, the flexural strength of this composite is lower than that of the ceramics prepared from C1100. Fig. 4 presents the XRD patterns of the composite ceramics prepared from C0 and C1000, which were sintered at 1500 °C for 2 h. The main crystalline phases include SiC, Al2O3, cristobalite, and mullite for both types of composite ceramics. Compared to the composite ceramics prepared from C0, the peak intensity of the cristobalite and mullite peaks was significantly higher in the composite ceramics prepared from C1000, indicating that the levels of cristobalite and mullite are higher. As a result, the mullite and cristobalite produce a better binding effect on the ceramic particles, resulting in improved flexural strength for the composite ceramics prepared from C1000. Therefore, to obtain the SiC-Al2O3 ceramics with high flexural strength at a sintering temperature of 1500 °C, the raw SiC powders should be oxidised at 1000 °C. Fig. 5 presents a schematic explaining the effects of the unoxidised and oxidised SiC powders on the microstructures of the prepared composite ceramics. When the SiC powders are oxidised at 1000 °C, a layer of SiO2 is formed on the SiC surface (Fig. 5(a)). Fig. 5(b) presents the microstructure of a green body prepared via gelcasting utilising oxidised SiC and Al2O3 powders. The SiC and Al2O3 particles are bonded together by polymer strands, which facilitates close contact between the SiO2 and Al2O3. When the green body is sintered, the closely contacting SiO2 and Al2O3 react easily to form mullite. This mullite can tightly bond the SiC and Al2O3 particles together to improve the flexural strength of the composite ceramics. Regarding the unoxidised SiC, the polymer strands link the SiC with Al2O3 in the green body
prepared via gelcasting (Fig. 5(d)). During sintering, the SiC surfaces are gradually oxidised to form cristobalite, but the cristobalite is thinner than that in the oxidised SiC (Fig. 5(e)). Additionally, because of the direct contact between Al2O3 and SiC, the Al2O3 hinders the surface oxidation of SiC, which leads to incomplete oxidation of the SiC surfaces (Fig. 5(e)). As a result, the amount of mullite formed is small and its distribution is inhomogeneous, resulting in low flexural strength for the ceramics prepared from unoxidised powders. Fig. 6 presents the flexural strength versus holding time for the composite ceramics prepared from C1000 and C0. As the holding time increases from 1 to 4 h, the flexural strength of the composite ceramics prepared from C1000 and C0 both increases. In comparison, the flexural strength of the composite ceramics prepared from C0 increases much more significantly. This is because as the holding time increases, the SiO2 content from the oxidation of SiC powder increases and more mullite is generated, resulting in a greater increase in flexural strength. However, when the sintering temperature and holding time are the same, the flexural strength of the composite ceramics prepared from C0 is still lower than that of the ceramics prepared from C1000. Therefore, by considering the flexural strength and production cost, the holding time for SiC-Al2O3 ceramics sintered at 1500 °C was set to 2 h. In summary, SiC-Al2O3 ceramics can be fabricated utilising the oxidised coarse-grained SiC and fine-grained Al2O3 particles. The optimised process parameters for the fabrication of SiC-Al2O3 ceramics with superior flexural strength are as follows. SiC powders are oxidised at 1000 °C for 2 h. The sintering temperature
is 1500 °C and the holding time is 2 h. Hereafter, porous SiC-Al2O3 ceramics will be prepared according to these parameters. 3.2 Water swelling, drying shrinkage, and low-temperature carbonisation of millet Figs. 7(a) and (b) present the diameter changes of the millet before and after water absorption. Before soaking, the millet is approximately circular and with diameters ranging from 1.5 to 1.8 mm (Fig. 7(a)). After soaking, the volume of the millet increases significantly and its diameters range from 1.6 to 2.1 mm (Fig. 7(b)). To obtain a larger volume shrinkage for the millet during the drying process, soaked millet is utilised as a pore-forming agent. Fig. 7(c) presents an image of the millet shrinkage in the gel-cast green body after drying. The gel-cast green body and millet both shrink, but the shrinkage of the millet is greater than that of the gel-cast green body. There are many gaps left between the millet and gel-cast green body, which is favourable for the removal of the millet and reduction of shrinkage stress during carbonisation. During the conversion from millet to carbon, volatile species, such as H2O and carbonyl groups, will be released [37]. If the carbonisation process is poorly optimised, the dried body will be destroyed by these materials. Fig. 8 presents the TG curve of the millet. The millet begins to lose weight at 50 °C and the weight loss approaches 100% at 600 °C. The weight loss of the millet drops sharply in the range of 200–300 °C with a weight loss value of approximately 46% in this range, which implies that the millet is mostly oxidised and decomposed in this temperature range. To ensure full oxidation of the millet, the heating rate in this interval was set to
1 °C/min. Our study indicated that temperature has a significant influence on the diameter change of the millet during carbonisation. Fig. 9(a) presents the morphologies of the millet carbonised at 210, 220, 230, 240, and 250 °C for 2 h. As the temperature increases from 210 to 240 °C, the colour of the millet becomes darker and its diameter decreases. When the temperature reaches 240 °C, the colour of the millet becomes completely black and it diameter shrinks by approximately 11.5% (Fig. 9(b)). However, when the temperature increases to 250 °C, the expansion ratio of the millet reaches 45 %. Fig. 9(c) presents the morphology of the sample carbonised at 260 °C. It can be seen that the millet expands too much, which leads to cracking of the sample. Figs. 9(d) and (e) present the morphologies of the samples carbonised at 240 °C. The magnified image in Fig. 9(e) corresponds to the rectangular region in Fig. 9(d). This sample is relatively complete. Therefore, to ensure the integrity of the sample, the millet should be slowly carbonised at 220-240 °C. 3.3 Preparation of porous SiC-Al2O3 ceramics The porous SiC-Al2O3 ceramics were prepared by casting the gel-cast slurry into the millet template and then removing the template via sintering. Fig. 10(a) presents the morphology of the prepared porous SiC-Al2O3 ceramic with optimised process parameters. Its skeletons are relatively complete and the number of interconnected pores is large. The pore sizes range from 1.5-2 mm. Fig. 10(b) illustrates the microstructure of the skeleton for the porous SiC-Al2O3 ceramic. Al2O3 particles are nearly uniformly dispersed in the SiC particles, and the Al2O3 and SiC particles are
tightly bonded together, which suggests that the optimised sintering process ensures good bonding between SiC particles based on the oxidation-derived SiO2 and mullite formation. Fig. 11 shows the residual flexural strength as a function of the quenching temperature for the prepared porous SiC-Al2O3 ceramics with optimised process parameters. When the quenching temperature is below 400 ℃, the residual flexural strength of the quenched ceramics decreases obviously from the initial 9.6 to 5.65 MPa. With increasing the quenching temperature from 400 to 1000 ℃, the residual flexural strength decreases slowly from 5.65 to 3.85 MPa. Because the fracture of porous ceramics does not happen at a high quenching temperature of 1000 ℃, indicating their good thermal shock resistance to fracture. Generally, the prepared SiC-Al2O3 porous ceramics have the total porosity of 60.2-65.1%, bulk density of 1.06-1.31 g/cm3, thermal conductivity of 0.63-0.81 W/m K and flexural strength of 8.3-10.5 MPa at room temperature. The prepared porous SiC-Al2O3 ceramics have good thermal shock resistance to fracture. 4. Conclusions Porous SiC-Al2O3 ceramics were prepared via gelcasting utilising soaked millet as a pore-forming agent. Oxidised coarse-grained SiC and fine-grained Al2O3 powders were utilised to successfully fabricate SiC-40wt% Al2O3 ceramics with a high flexural strength of 39.8 MPa. Compared to untreated SiC powders, the oxidised powders were beneficial for improving the flexural strength of composite ceramics and lowering their required sintering temperature. The shrinkage of the soaked millet was
greater than that of the green body during drying. To prevent cracking of the green body, the millet must be carbonised slowly at 220-240 °C. The pore sizes of the prepared porous SiC-Al2O3 ceramics utilising the optimised process were 1.5-2 mm. Additionally, their skeletons were relatively complete and the number of interconnected pores was large. Their total porosity and flexural strength were 60.2-65.1% and 8.3-10.5 MPa, respectively. Our work provides a facile and effective method for lowering the required sintering temperature for SiC ceramics based on the pre-oxidation of SiC and developing porous ceramics with superior pore integrity by utilising shrinkable pore-forming agents. The prepared porous ceramics have good thermal shock resistance to fracture, and expected to have potential applications in the high temperature catalytic substrates and molten-metal filters. Acknowledgements Financial support from the Key Technologies R&D Program of China (No. SQ2016ZYC01003836), and Science and Technology Project of the Jiangxi Provincial Department of Education (No. GJJ171143) is gratefully acknowledged. Appendix A. Supplementary data The oxidized SiC powders slightly affected the rheological properties of ceramic slurries, which could be ignored (supplementary Fig. S1). References [1] H.B. Wu, Y.S. Li, S. Yun, X.J. Liu, Z.R. Huang, D.J. Jiang, Effects of particle grading on porous gelcasted and solid-state-sintered SiC ceramics with improved connectivity, J Alloy Compd 732 (2018) 547-554.
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Fig. 1 Flow chart for preparation of porous SiC-Al2O3 ceramics via gelcasting utilising the soaked millet as a pore-forming agent. Fig. 2 Flexural strengths of the different types of SiC-Al2O3 ceramics at different sintering temperatures. Fig. 3 (a) TG-DSC curve of the SiC-Al2O3 ceramics; (b) morphologies of SiC oxidised at 1000 °C; (c) microstructure of the composite ceramic prepared from C1000 at 1400 °C; (d) morphologies of SiC oxidised at 1100 °C; (e) microstructure of the composite ceramic prepared from C1100 at 1400 °C; and (f) microstructure of the composite ceramic prepared from C1100 at 1500 °C. Fig. 4 XRD patterns of composite ceramics prepared from C0 and C1000, sintered at 1500 °C for 2 h (A is alumina, C is cristobalite, M is mullite, and S is silicon carbide). Fig. 5 Schematic showing the effects of unoxidised and oxidised SiC powders on the microstructures of the prepared composite ceramics. Fig. 6 Flexural strength of composite ceramics prepared from C1000 and C0, sintered at 1500 °C with different holding times. Fig. 7 Diameter changes of the millet (a) before and (b) after soaking; and (c) shrinkage of the millet in the green body after drying. Fig. 8 TG curve of the millet. Fig. 9 (a) Morphologies of the original millet and millet carbonised at 210, 220, 230, 240, and 250 °C for 2 h; (b) shrinkage ratios of the millet carbonised at different temperatures; and morphologies of the samples carbonised at (c) 260 °C and (d and e) 240 °C.
Fig. 10 (a) Porous SiC-Al2O3 ceramic prepared via gelcasting utilising soaked millet as a pore-forming agent and (b) SEM image of the microstructure of the skeleton of the porous SiC-Al2O3 ceramic. Fig. 11 Residual flexural strength as a function of the quenching temperature for the prepared porous SiC-Al2O3 ceramics.
PII: DOI: Reference:
S0272-8842(18)33375-3 https://doi.org/10.1016/j.ceramint.2018.12.008 CERI20231
To appear in: Ceramics International Received date: 30 September 2018 Revised date: 30 November 2018 Accepted date: 2 December 2018 Cite this article as: Li Wan, Qingquan Tian, Yun Xiang, Lin Chen, Jieguang Song, Xianzhong Wang and Xinshuang Guo, Preparation of porous SiC-Al 2O3 ceramics via gelcasting utilising a shrinkable pore-forming agent and oxidised coarse-grained SiC, Ceramics International, https://doi.org/10.1016/j.ceramint.2018.12.008 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Preparation of porous SiC-Al2O3 ceramics via gelcasting utilising a shrinkable pore-forming agent and oxidised coarse-grained SiC Li Wana, Qingquan Tianb,c, Yun Xianga, Lin Chena, Jieguang Songa, Xianzhong Wanga, Xinshuang Guoa,c* a
Jiangxi Key Laboratory of Industrial Ceramics, Engineering Technology Research
Centre for Environmental Protection Materials and Equipment of Jiangxi Province, Pingxiang University, Pingxiang 337055, China b
School of Chemistry and Materials, Weinan Normal University, Weinan 714099, China
c
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China *
Corresponding author. [email protected]
Abstract By utilising soaked millet as a shrinkable pore-forming agent, porous silicon carbide-alumina (SiC-Al2O3) ceramics were prepared via gelcasting. The fabrication of SiC-Al2O3 ceramics based on oxidised and unoxidised coarse-grained SiC was also studied. The water swelling, drying shrinkage, and low-temperature carbonisation of the millet were investigated. We found that the shrinkage of the soaked millet was greater than that of gel body during drying, which left large gaps that prevented shrinkage stresses from destroying the gel body. Low-temperature carbonisation of the millet should be performed slowly at 220–240 °C because its expansion rate increases to 45% at 250 °C, resulting in the cracking of samples. At a constant sintering temperature, the flexural strength of the SiC-Al2O3 ceramics prepared with
SiC powders oxidised at 1000 °C was the highest, indicating that oxidised powders can successfully decrease the required sintering temperature and improve the flexural strength of composite ceramics. Based on our optimised process, porous SiC-Al2O3 ceramics were sintered at 1500 °C for 2 h. When their skeletons were fully developed, their pore sizes were in the range of 1.5–2 mm. Their porosity and flexural strength were 60.2–65.1% and 8.3–10.5 MPa, respectively. Keywords: Porous SiC-Al2O3 ceramics; Oxidised coarse-grained SiC; Pore-forming agent; Gelcasting; Flexural strength
1. Introduction Porous silicon carbide (SiC) ceramics have attracted considerable attention for their unique combination of desirable properties, including excellent thermal shock resistance, good chemical resistance, large specific surface area, and a low thermal expansion coefficient [1–5]. Because of its low self-diffusion coefficient and strong covalent bonding, the required sintering temperature for coarse-grained SiC is approximately 2100–2400 K [6, 7]. Lowering the sintering temperature of SiC ceramics is beneficial for reducing production costs and has become a research hotspot. By combining SiC with other inorganic materials, the resulting porous ceramic composites have lower required sintering temperatures and excellent physical properties [8]. There are many different types of porous SiC composite ceramics, including SiC-Al2O3 [9–12], SiC-ZrO2 [13], and SiC-mullite [14, 15] among others [16, 17], where SiC-Al2O3 ceramics are the most widely developed. Low temperature
(≤1600 °C) SiC-Al2O3 ceramics are typically prepared utilising a precursor method [18], reaction bonding [19], and liquid-phase sintering [8, 20]. Liquid-phase sintering has attracted significant attention because it is a relatively simple process. The SiC powders utilised in liquid-phase sintering are typically untreated raw powders. SiC and Al2O3 powders can be bonded utilising mullite and oxidation-derived SiO2 [21]. When the surfaces of SiC powders are oxidised to form SiO2, the benefits in terms of the formation of mullite and reduction of the required sintering temperature for SiC-Al2O3 ceramics is still unclear. The pore-forming agent method is another common method for preparing porous ceramics based on its fabrication simplicity, practicability, and economic viability [22–25]. There are many types of solid pore-forming agents and they are typically combined with traditional slip casting methods to prepare porous ceramics [26–31]. However, solid pore-forming agents typically do not shrink during drying, which hinders the shrinkage of the green body and leads to various defects, such as the breakage of the skeletons of green bodies. Therefore, it is necessary to utilise shrinkable pore-forming agents to prepare porous ceramics with good structural integrity. Additionally, the prepared green bodies have low strength and can be easily destroyed during demoulding and drying. Gelcasting, a near-net technique with high yields, has been widely adopted to prepare high-strength green bodies with homogenous microstructures [32–34]. The combination of gelcasting and the pore-forming agent methods can prevent cracking during demoulding and drying. In this study, porous SiC-Al2O3 ceramics were prepared via gelcasting utilising
millet as a pore-forming agent. Millet was chosen for its water-swelling characteristics and approximately circular shape. We thoroughly investigated the drying shrinkage and low-temperature carbonisation process of the millet. SiC-Al2O3 ceramics were prepared utilising both oxidised and unoxidised coarse-grained SiC particles and their flexural strengths under various sintering conditions were investigated. Finally, the structures and properties of the porous SiC-Al2O3 ceramics prepared utilising an optimised process were studied. 2. Experiments 2.1 Materials and sample preparation Commercial SiC powders (≥99%, D50=37.6 µm) and Al2O3 powders (99.8%, D50=4.3 µm) were utilised as the two starting materials for preparing SiC-Al2O3 ceramics. To investigate the effects of the oxidation of SiC on the sintering properties of composite ceramics, SiC powders were oxidised at 1000 and 1100 °C for 2 h in air. The treated powders were labelled as P1000 and P1100, and the untreated SiC powders were labelled as P0. The ceramic green bodies created by mixing P0, P1000, and P1100 with Al2O3 were labelled as C0, C1000, and C1100, respectively. Fig. 1 presents a flow chart for fabricating porous SiC-Al2O3 ceramics. First, acrylamide
(AM,
monomer),
N,N’-methyleneacrylamide
(cross-linker),
and
polyethylene glycol (dispersant) with a weight ratio of 30:3:2 were added to deionised water. After stirring to obtain premixed solutions, SiC and Al2O3 powders with a weight ratio of 60:40 were added to prepare a slurry with a solid content of 45 vol%. The
slurry
was
ball-milled
for
4
h
and
then
evacuated
for
2
h.
N,N,N’,N’-tetramethylethylenediamine (TEMED, catalyst) and ammonium persulfate (APS, initiator) were then added to the degassed slurry. The amounts of TEMED and APS relative to the weight of AM were 0.1 wt% and 3.7 wt%, respectively. The chemical reagents above were procured from Sinopharm Chemical Reagent Co., Ltd. Following mechanical stirring for 2 min, the as-prepared slurry was cast into a glass mould with a millet template (Φ25 mm×50 mm) and consolidated to form green bodies. Next, the prepared green bodies were dried 25 °C for 24 h and 50 °C for 48 h. The dried samples were then carbonised at low temperature and subjected to pressureless sintering in air to obtain porous SiC-Al2O3 ceramics. It is worth noting that an optimised process for fabricating SiC-Al2O3 ceramics was first identified and then the porous SiC-Al2O3 ceramics were prepared according to the optimised process. 2.2 Characterisation The average diameters of the SiC and Al2O3 powder particles were measured utilising a laser particle size analyser (BT-9300H, Dandong Bettersize Instruments Ltd, China). Thermal analysis of the millet (SiC-Al2O3 ceramics) was performed utilising a thermal analyser (STA449, Netzsch Inc., Germany) at a heating rate of 1 (10) °C/min from room temperature to 650 (1550) °C in air. The morphologies and crystal structures of the samples were analysed via scanning electron microscopy (SEM, S8010, Hitachi Inc., Japan) and X-ray diffraction (XRD, D8-Advance, Bruke Inc., Germany), respectively. The porosity and bulk density of the porous SiC-Al2O3 ceramics were measured based on the Archimedes method. The thermal shock
resistance was performed by the water-quenching method, which was referred to the literature [35]. The residual flexural strength of the samples was tested to evaluate the thermal shock resistance of the porous SiC-Al2O3 ceramics. The heat conductivity was measured using the thermal conductivity meter (DRPL-II, Xiangtan Xiangyi instrument Co., Ltd., China). The flexural strength of the ceramics (dimension of 3 mm × 4 mm × 40 mm) was measured utilising a universal testing machine (AGS-X-10KN, Shimadzu Inc., Japan) with a span of 30 mm and loading rate of 0.5 mm/min. 3. Results and discussion 3.1 Optimisation of the process parameters for SiC-Al2O3 ceramics Fig. 2 presents the flexural strengths of the SiC-Al2O3 ceramics prepared with different types of SiC powders. The ceramics were sintered at 1400, 1450, and 1500 °C for 2 h, respectively. When the sintering temperature was 1400 °C, the flexural strengths of the composite ceramics prepared from C1000, C0, and C1100 were 17.6, 17.1 and 14.1 MPa, respectively, meaning the flexural strength of the ceramic prepared from C1100 was the lowest. According to the thermogravimetry (TG)–differential scanning calorimetry (DSC) analysis results presented in Fig. 3(a), the exothermic peak at 1455 °C can be attributed to the formation of mullite through a reaction between the SiO2 obtained from SiC oxidation and Al2O3. Therefore, the different flexural strengths of the three types of composite ceramics are largely related to the connections between SiC particles created by cristobalite at 1400 °C [21]. Regarding the composite ceramics
prepared from C1000, the surfaces of SiC powders treated at 1000 ℃ are relatively uniform (Fig. 3(b). The proper amount of liquid phase was formed at 1400 °C, which is beneficial for close bonding between cristobalite and Al2O3 (Fig. 3(c)). As a result, the composite ceramics prepared from C1000 have the highest flexural strength. Regarding the composite ceramics prepared from C1100, the SiC powders were excessively oxidised at 1100 °C [36], resulting in the generation of partial liquid phase (indicated by white circles) and certain ablation (indicated by black circles) on their surfaces (Fig. 3(d)). Excessive liquid phase consequently generates in the composite ceramics at 1400 °C, which is not beneficial for bonding between Al2O3 and cristobalite (Fig. 3(e)). Therefore, the flexural strength of the composite ceramics prepared from C1100 is lower than that of the ceramics prepared from C1000. When the sintering temperature is 1450 or 1500 °C, the bonding of ceramic particles mainly depends on mullite and cristobalite. Therefore, the flexural strength of the composite ceramics sintered at 1450 and 1500 °C is greater than that of the ceramics sintered at 1400 °C. Additionally, the composite ceramics prepared from C1000 have the highest flexural strength, followed by those prepared from C1100 and C0 (Fig. 2). Some possible reasons for these results are as follows. Regarding the composite ceramics prepared from C1100, with the formation of mullite and oxidation of SiC, various interface defects, such as voids between SiC and Al2O3, can be eliminated, leading to relatively tight bonding (Fig. 3(f)), which results in an increase in flexural strength. Regarding the composite ceramics prepared from C0, because there is less SiO2 on the surface of the unoxidised SiC and less mullite formed
between the SiC and Al2O3 particles, the flexural strength of this composite is lower than that of the ceramics prepared from C1100. Fig. 4 presents the XRD patterns of the composite ceramics prepared from C0 and C1000, which were sintered at 1500 °C for 2 h. The main crystalline phases include SiC, Al2O3, cristobalite, and mullite for both types of composite ceramics. Compared to the composite ceramics prepared from C0, the peak intensity of the cristobalite and mullite peaks was significantly higher in the composite ceramics prepared from C1000, indicating that the levels of cristobalite and mullite are higher. As a result, the mullite and cristobalite produce a better binding effect on the ceramic particles, resulting in improved flexural strength for the composite ceramics prepared from C1000. Therefore, to obtain the SiC-Al2O3 ceramics with high flexural strength at a sintering temperature of 1500 °C, the raw SiC powders should be oxidised at 1000 °C. Fig. 5 presents a schematic explaining the effects of the unoxidised and oxidised SiC powders on the microstructures of the prepared composite ceramics. When the SiC powders are oxidised at 1000 °C, a layer of SiO2 is formed on the SiC surface (Fig. 5(a)). Fig. 5(b) presents the microstructure of a green body prepared via gelcasting utilising oxidised SiC and Al2O3 powders. The SiC and Al2O3 particles are bonded together by polymer strands, which facilitates close contact between the SiO2 and Al2O3. When the green body is sintered, the closely contacting SiO2 and Al2O3 react easily to form mullite. This mullite can tightly bond the SiC and Al2O3 particles together to improve the flexural strength of the composite ceramics. Regarding the unoxidised SiC, the polymer strands link the SiC with Al2O3 in the green body
prepared via gelcasting (Fig. 5(d)). During sintering, the SiC surfaces are gradually oxidised to form cristobalite, but the cristobalite is thinner than that in the oxidised SiC (Fig. 5(e)). Additionally, because of the direct contact between Al2O3 and SiC, the Al2O3 hinders the surface oxidation of SiC, which leads to incomplete oxidation of the SiC surfaces (Fig. 5(e)). As a result, the amount of mullite formed is small and its distribution is inhomogeneous, resulting in low flexural strength for the ceramics prepared from unoxidised powders. Fig. 6 presents the flexural strength versus holding time for the composite ceramics prepared from C1000 and C0. As the holding time increases from 1 to 4 h, the flexural strength of the composite ceramics prepared from C1000 and C0 both increases. In comparison, the flexural strength of the composite ceramics prepared from C0 increases much more significantly. This is because as the holding time increases, the SiO2 content from the oxidation of SiC powder increases and more mullite is generated, resulting in a greater increase in flexural strength. However, when the sintering temperature and holding time are the same, the flexural strength of the composite ceramics prepared from C0 is still lower than that of the ceramics prepared from C1000. Therefore, by considering the flexural strength and production cost, the holding time for SiC-Al2O3 ceramics sintered at 1500 °C was set to 2 h. In summary, SiC-Al2O3 ceramics can be fabricated utilising the oxidised coarse-grained SiC and fine-grained Al2O3 particles. The optimised process parameters for the fabrication of SiC-Al2O3 ceramics with superior flexural strength are as follows. SiC powders are oxidised at 1000 °C for 2 h. The sintering temperature
is 1500 °C and the holding time is 2 h. Hereafter, porous SiC-Al2O3 ceramics will be prepared according to these parameters. 3.2 Water swelling, drying shrinkage, and low-temperature carbonisation of millet Figs. 7(a) and (b) present the diameter changes of the millet before and after water absorption. Before soaking, the millet is approximately circular and with diameters ranging from 1.5 to 1.8 mm (Fig. 7(a)). After soaking, the volume of the millet increases significantly and its diameters range from 1.6 to 2.1 mm (Fig. 7(b)). To obtain a larger volume shrinkage for the millet during the drying process, soaked millet is utilised as a pore-forming agent. Fig. 7(c) presents an image of the millet shrinkage in the gel-cast green body after drying. The gel-cast green body and millet both shrink, but the shrinkage of the millet is greater than that of the gel-cast green body. There are many gaps left between the millet and gel-cast green body, which is favourable for the removal of the millet and reduction of shrinkage stress during carbonisation. During the conversion from millet to carbon, volatile species, such as H2O and carbonyl groups, will be released [37]. If the carbonisation process is poorly optimised, the dried body will be destroyed by these materials. Fig. 8 presents the TG curve of the millet. The millet begins to lose weight at 50 °C and the weight loss approaches 100% at 600 °C. The weight loss of the millet drops sharply in the range of 200–300 °C with a weight loss value of approximately 46% in this range, which implies that the millet is mostly oxidised and decomposed in this temperature range. To ensure full oxidation of the millet, the heating rate in this interval was set to
1 °C/min. Our study indicated that temperature has a significant influence on the diameter change of the millet during carbonisation. Fig. 9(a) presents the morphologies of the millet carbonised at 210, 220, 230, 240, and 250 °C for 2 h. As the temperature increases from 210 to 240 °C, the colour of the millet becomes darker and its diameter decreases. When the temperature reaches 240 °C, the colour of the millet becomes completely black and it diameter shrinks by approximately 11.5% (Fig. 9(b)). However, when the temperature increases to 250 °C, the expansion ratio of the millet reaches 45 %. Fig. 9(c) presents the morphology of the sample carbonised at 260 °C. It can be seen that the millet expands too much, which leads to cracking of the sample. Figs. 9(d) and (e) present the morphologies of the samples carbonised at 240 °C. The magnified image in Fig. 9(e) corresponds to the rectangular region in Fig. 9(d). This sample is relatively complete. Therefore, to ensure the integrity of the sample, the millet should be slowly carbonised at 220-240 °C. 3.3 Preparation of porous SiC-Al2O3 ceramics The porous SiC-Al2O3 ceramics were prepared by casting the gel-cast slurry into the millet template and then removing the template via sintering. Fig. 10(a) presents the morphology of the prepared porous SiC-Al2O3 ceramic with optimised process parameters. Its skeletons are relatively complete and the number of interconnected pores is large. The pore sizes range from 1.5-2 mm. Fig. 10(b) illustrates the microstructure of the skeleton for the porous SiC-Al2O3 ceramic. Al2O3 particles are nearly uniformly dispersed in the SiC particles, and the Al2O3 and SiC particles are
tightly bonded together, which suggests that the optimised sintering process ensures good bonding between SiC particles based on the oxidation-derived SiO2 and mullite formation. Fig. 11 shows the residual flexural strength as a function of the quenching temperature for the prepared porous SiC-Al2O3 ceramics with optimised process parameters. When the quenching temperature is below 400 ℃, the residual flexural strength of the quenched ceramics decreases obviously from the initial 9.6 to 5.65 MPa. With increasing the quenching temperature from 400 to 1000 ℃, the residual flexural strength decreases slowly from 5.65 to 3.85 MPa. Because the fracture of porous ceramics does not happen at a high quenching temperature of 1000 ℃, indicating their good thermal shock resistance to fracture. Generally, the prepared SiC-Al2O3 porous ceramics have the total porosity of 60.2-65.1%, bulk density of 1.06-1.31 g/cm3, thermal conductivity of 0.63-0.81 W/m K and flexural strength of 8.3-10.5 MPa at room temperature. The prepared porous SiC-Al2O3 ceramics have good thermal shock resistance to fracture. 4. Conclusions Porous SiC-Al2O3 ceramics were prepared via gelcasting utilising soaked millet as a pore-forming agent. Oxidised coarse-grained SiC and fine-grained Al2O3 powders were utilised to successfully fabricate SiC-40wt% Al2O3 ceramics with a high flexural strength of 39.8 MPa. Compared to untreated SiC powders, the oxidised powders were beneficial for improving the flexural strength of composite ceramics and lowering their required sintering temperature. The shrinkage of the soaked millet was
greater than that of the green body during drying. To prevent cracking of the green body, the millet must be carbonised slowly at 220-240 °C. The pore sizes of the prepared porous SiC-Al2O3 ceramics utilising the optimised process were 1.5-2 mm. Additionally, their skeletons were relatively complete and the number of interconnected pores was large. Their total porosity and flexural strength were 60.2-65.1% and 8.3-10.5 MPa, respectively. Our work provides a facile and effective method for lowering the required sintering temperature for SiC ceramics based on the pre-oxidation of SiC and developing porous ceramics with superior pore integrity by utilising shrinkable pore-forming agents. The prepared porous ceramics have good thermal shock resistance to fracture, and expected to have potential applications in the high temperature catalytic substrates and molten-metal filters. Acknowledgements Financial support from the Key Technologies R&D Program of China (No. SQ2016ZYC01003836), and Science and Technology Project of the Jiangxi Provincial Department of Education (No. GJJ171143) is gratefully acknowledged. Appendix A. Supplementary data The oxidized SiC powders slightly affected the rheological properties of ceramic slurries, which could be ignored (supplementary Fig. S1). References [1] H.B. Wu, Y.S. Li, S. Yun, X.J. Liu, Z.R. Huang, D.J. Jiang, Effects of particle grading on porous gelcasted and solid-state-sintered SiC ceramics with improved connectivity, J Alloy Compd 732 (2018) 547-554.
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Fig. 1 Flow chart for preparation of porous SiC-Al2O3 ceramics via gelcasting utilising the soaked millet as a pore-forming agent. Fig. 2 Flexural strengths of the different types of SiC-Al2O3 ceramics at different sintering temperatures. Fig. 3 (a) TG-DSC curve of the SiC-Al2O3 ceramics; (b) morphologies of SiC oxidised at 1000 °C; (c) microstructure of the composite ceramic prepared from C1000 at 1400 °C; (d) morphologies of SiC oxidised at 1100 °C; (e) microstructure of the composite ceramic prepared from C1100 at 1400 °C; and (f) microstructure of the composite ceramic prepared from C1100 at 1500 °C. Fig. 4 XRD patterns of composite ceramics prepared from C0 and C1000, sintered at 1500 °C for 2 h (A is alumina, C is cristobalite, M is mullite, and S is silicon carbide). Fig. 5 Schematic showing the effects of unoxidised and oxidised SiC powders on the microstructures of the prepared composite ceramics. Fig. 6 Flexural strength of composite ceramics prepared from C1000 and C0, sintered at 1500 °C with different holding times. Fig. 7 Diameter changes of the millet (a) before and (b) after soaking; and (c) shrinkage of the millet in the green body after drying. Fig. 8 TG curve of the millet. Fig. 9 (a) Morphologies of the original millet and millet carbonised at 210, 220, 230, 240, and 250 °C for 2 h; (b) shrinkage ratios of the millet carbonised at different temperatures; and morphologies of the samples carbonised at (c) 260 °C and (d and e) 240 °C.
Fig. 10 (a) Porous SiC-Al2O3 ceramic prepared via gelcasting utilising soaked millet as a pore-forming agent and (b) SEM image of the microstructure of the skeleton of the porous SiC-Al2O3 ceramic. Fig. 11 Residual flexural strength as a function of the quenching temperature for the prepared porous SiC-Al2O3 ceramics.