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SiC ceramics with unidirectionally aligned channels produced by freeze-drying and chemical vapor infiltration
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Porous SiCnw /SiC ceramics with unidirectionally aligned channels produced by freeze-drying and chemical vapor infiltration Daoyang Han, Hui Mei ∗ , Shanshan Xiao, Junchao Xia, Jinlei Gu, Laifei Cheng Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an Shannxi 710072, PR China
a r t i c l e
i n f o
Article history: Received 27 August 2016 Received in revised form 12 October 2016 Accepted 13 October 2016 Available online xxx Keywords: Unidirectional freeze-drying Silicon carbide nanowire (SiCnw ) aerogel Chemical vapor infiltration (CVI) Porous silicon carbide nanowire/silicon carbide (SiCnw /SiC) ceramics
a b s t r a c t The current study introduces a methodology for the fabrication of porous silicon carbide nanowire/silicon carbide (SiCnw /SiC) ceramics with macroscopic unidirectionally aligned channels and reports on their microstructural and mechanical properties. The material was produced by freezing of a water-based slurry of -SiC nanowires (SiCnw ) with control of the ice growth direction. Pores were subsequently generated by sublimation of the columnar ice during freeze-drying. Chemical vapor infiltration (CVI) of SiC into the open pore network of the SiCnw aerogel with unidirectionally aligned channels, resulted in the formation of highly porous SiCnw /SiC ceramics which exhibited a unique microstructure as identified by scanning electron microscopy. The pore size distribution and the mechanical properties of the asfabricated porous ceramics were examined by mercury intrusion porosimetry and three-point bending and compression tests, respectively, while phase composition was investigated through X-ray diffraction. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Porous silicon carbide (SiC) ceramic is the preferable material in a wide range of applications such as catalyst supports [1,2], filters for hot gas or molten metal [3–5], and high efficiency combustion burners [6]. This is due to its unique properties such as low density, good thermal shock resistance [7], high mechanical properties and excellent chemical stability at elevated temperatures [8]. The mechanical performance of a porous SiC ceramic depends primarily on the selected fabrication method which affects the pore size distribution, level of interconnection between pores, pore orientation, as well as on the intrinsic properties of the material used as the matrix. There are four conventional methodologies suggested for manufacturing porous SiC ceramics, namely sacrificial template, reaction forming/bonding, replication and partial sintering [9]. Recently, freeze casting was newly suggested for the preparation of porous ceramics; the technique has attracted increased scientific interest due to its simplicity and environmental friendliness [10–14]. Therein, a ceramic slurry acting as the freezing vehicle is poured into a mold at a temperature below its freezing point. This leads to the formation of a bicontinuous structure, consisting of three dimensionally interconnected frozen vehicle network and networks of concentrated ceramic particles. Porous ceramics are
∗ Corresponding author. E-mail addresses: [email protected], [email protected] (H. Mei).
readily produced by removing the frozen vehicle network via freeze drying, followed by sintering of ceramic walls [15]. In Ref. [15], highly-aligned porous SiC ceramics decorated with SiC nanowires were fabricated by unidirectional freeze casting of SiC/camphene slurry with various contents of polycarbosilane (PCS) preceramic used simultaneously both as binder and precursor for in situ growth of SiC nanowires. In [16], porous SiC ceramics with good high temperature mechanical properties were fabricated by a freeze-drying combined with solid sintering. Yoon [17] et al. fabricated highlyaligned porous SiC ceramics with well-defined pore structures by freezing a PCS/camphene solution, followed by the pyrolysis of the porous PCS objects at 1400 ◦ C for 1 h in a flowing Ar atmosphere. In Refs. [18,19], porous SiC ceramics were also fabricated by freezecasting and solid state sintering. In Ref. [20], SiC whisker-reinforced SiC ceramic matrix composites were prepared by depositing SiC matrix into the SiC whisker preforms using chemical vapor infiltration (CVI). The latter method is renowned as the most effective and mature means of preparing SiC matrix with ultra-pure, minimal residual stresses and controllable grain sizes [21–26]. Driven by previous successful applications of unidirectional freezing [10] and CVI technology [20] in the production of porous ceramics, we report herein a new strategy for the production of porous SiC ceramics. Initially, -SiC nanowire (SiCnw ) aerogels with unidirectionally aligned channels were prepared by successive freezing, storage in the frozen state, and defrosting of SiCnw solutions. The role of SiCnw is to prepare a macroporous preform with an open pore network. The as-fabricated preform was strong
http://dx.doi.org/10.1016/j.jeurceramsoc.2016.10.015 0955-2219/© 2016 Elsevier Ltd. All rights reserved.
Please cite this article in press as: D. Han, et al., Porous SiCnw /SiC ceramics with unidirectionally aligned channels produced by freezedrying and chemical vapor infiltration, J Eur Ceram Soc (2016), http://dx.doi.org/10.1016/j.jeurceramsoc.2016.10.015
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Fig. 1. Schematic of polytetrafluoroethylene mold, unit is mm; 1: sheet metal, 2: polytetrafluoroethylene mold.
enough to self-assemble and very ideal for the diffusion of gases using a CVI process. After densification using multiple CVI cycles, the mechanical properties of the obtained porous SiCnw /SiC ceramics such as compressive and flexural strength were investigated by compression and three-point bending tests, while scanning electron microscopy and X-ray diffraction were used to examine the microstructure and phase composition of the ceramic. 2. Experimental 2.1. Material preparation protocol Porous SiCnw /SiC ceramics were fabricated following a newly established experimental protocol of the successive steps presented in the following. Initially, 0.75 g of -SiCnw of diameters 0.1–0.6 m, lengths 50–100 m, purity higher than 96% and density of 3.21 g/cm3 (Changsha Sinet Advanced Materials Co., Ltd. Changsha, China) were dispersed in 100 g of aqueous solution of 0.8 wt% sodium carboxymethyl cellulose (CMC), an etherified derivative of cellulose, (Cekal 30000A, Shenzhen Hairenda Technology Co. Ltd. Shenzhen, China). Dispersion was achieved through ultrasonication for 1.0 h in a 250 ml glass beaker using an BILON1500 ultrasonic homogenizer (Xian Bilon Biological Technology Co. Ltd) operating at 20 ± 1 kHz, 300–1500 W with a 10 mm diameter probe and an on/off pulse rate of 1/s. During ultrasonication the solution was retained at room temperature by continuous cooling fluid circulation through the heat exchanger. CMC was added into the suspensions of particles with the aim of dispersing and bonding the particles in the freeze-casting process. The resultant slurry was poured into a polytetrafluoroethylene mold as shown in Fig. 1. The mold was composed of two parts: the metal with high thermal conductivity attached to the mold bottom, and the mold side made of polytetrafluoroethylene. It was subsequently placed on the pre-cooled at Ts = −80 ◦ C stainless-steel circular tray of a freeze-dryer (LGJ-18S, Beijing Song Yuan Hua Xing Science and Technology Develop Co. Ltd). In this manner, ice growth was forced to occur macroscopically in the vertical direction, causing the formation of columnar ice, the phase separation of water and nanowires [27] (Fig. 2), and the piling up of the nanowires between the columnar ice. After the slurry froze completely, vacuum conditions were imposed by a vacuum pump and the stainless steel circular tray was heated and the frozen specimens lyophilized. The thermal cycle used in the lyophilize process is as follows: tray temperature −50 ◦ C for 1 h, −20 ◦ C for 2 h, 10 ◦ C for 2 h, 20 ◦ C for 3 h, 30 ◦ C for 3 h, 40 ◦ C for 4 h, 50 ◦ C for 4 h; gas pressure remained
Fig. 2. Pattern formation and SiCnw segregation during freeze casting of SiCnw slurry.
below 30 Pa. At the end of the sublimation step, a SiCnw aerogel with unidirectionally aligned channels was formed. The aerogel was subsequently inserted into a CVI furnace where it was infiltrated with SiC. During the CVI process, methyltrichlorosilane (CH3 SiCl3 , abbreviated as MTS) was used as precursor gas, hydrogen as carrier and argon as diluting gas. Extensive details of the employed CVI process can be found in [28]. The infiltration time was set to 80 h and the process was repeated 3 consecutive times. After 3 cycles, the as-fabricated ceramic plates were reduced in thickness of 2 mm from each face using 800 # grit emery paper. The resulting 3 mm thick plates were cut into rectangular parallelepipeds of dimensions 4 × 3 × 50 mm3 and 3 × 3 × 50 mm3 using a laser cutting machine. These specimens with smooth surfaces were further infiltrated with SiC matrix for 2 and 4 consecutive CVI runs, respectively. Finally, three sets of specimens with 3, 5 and 7 cycles were subjected to physical characterization. 2.2. Testing and calculations To assess the porosity of the ceramics, mercury intrusion porosimetry (MIP) was conducted on an AutoPore IV 9500 Mercury Porosimeter (Micromeritics Instruments, USA). Three-point bending and compression tests were conducted on an electronic universal testing machine (SANS CMT 4304, Shenzhen, China). The
Please cite this article in press as: D. Han, et al., Porous SiCnw /SiC ceramics with unidirectionally aligned channels produced by freezedrying and chemical vapor infiltration, J Eur Ceram Soc (2016), http://dx.doi.org/10.1016/j.jeurceramsoc.2016.10.015
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Table 1 Porosity and density of the porous SiCnw /SiC ceramic.
Fig. 3. Schematic diagram of the measurement of static compressive strength.
dimensions of compression specimens were 3 × 3 × 3 mm3 and tests was conducted under crosshead displacement control, at a speed of 0.5 mm/min. Compressive tests were performed both perpendicularly (transverse compressive) and parallel (longitudinal compression) to the ice growth orientation as illustrated in Fig. 3. Three-point bending specimens had dimensions of 40 × 4 × 3 mm3 and crosshead speed and load span of 0.5 mm/min and 30 mm, respectively, were used. Compressive strength and flexural strength were calculated as: P a = S
(1)
3PL b = 2bh2
(2)
where P is the load, L is the support span, b is the specimen width, h its height, and S the cross sectional area. Fifteen specimens were tested for statistically accurate calculation of compressive and flexural strengths. The microstructure of the SiCnw aerogel and porous SiCnw /SiC ceramic was observed by scanning electron microscopy (SEM, Hitachi S-4800, Tokyo, Japan). The phase composition of the ceramic was examined by X-ray diffraction (XRD, D8-Advance, Bruker, Germany). 3. Results and discussion 3.1. Microstructure of SiCnw aerogel and porous SiCnw /SiC ceramic During the freezing process, the formation of crystalline ice causes every solute originally dispersed in the aqueous medium to be expelled to the boundaries between adjacent ice crystals [27] (Fig. 2). Due to the higher thermal conductivity of the metal part of the mold compared to that of polytetrafluoroethylene, the slurry in the special mold froze unidirectionally hence causing the formation of columnar ice parallel to the ice growth orientation while the solute was expelled to the boundaries between the columnar ice. Subsequent freeze-drying gave rise to a cryogel, a macroporous structure characterized by walls of matter enclosing empty areas where ice crystals originally resided. The freeze-drying process also allowed for achievement of monoliths which preserved the size and shape of the parent container submitted to freezing. Fig. 4 presents SEM micrographs of the fabricated SiCnw aerogels. Fig. 4a and b shows front and top views of an aerogel, which are the planes parallel and perpendicular to the macroscopic direction of ice formation, respectively. Open micro-pores and micro-pore channels of sizes of several hundred micrometers are observed in the aerogel, aligned along the macroscopic direction of ice formation similarly to channels. The high magnification image in Fig. 4d demonstrates the SiCnw walls in the aerogel, comparable to the parent SiCnw (Fig. 4c), which exhibited mesh-like morphology bonded by CMC. In the walls, there were many smaller micro-pores of dimensions of several micrometers attributed to smaller ice particles existing inside the SiCnw walls in the frozen slurry. These micro-pores and micro-pore channels provide paths for subsequent SiC infiltration by CVI. In summary, the aerogel microstructure was comprised
CVI cycles
3
5
7
Porosity/% Density/g/cm3
32.23 1.99
23.44 2.27
9.23 2.30
by networks of unidirectionally-aligned channels achieved by the employed freeze-drying process. SEM micrographs of the porous SiCnw /SiC ceramics are shown in Fig. 5. As seen in Fig. 5a and the inset micrograph therein, all SiCnw walls in the aerogel were completely filled with SiC matrix. The top face of the ceramic showed a morphology similar to the one developed during SiCnw aerogel formation. Fig. 5b shows a typical microstructure of the cross-section of the SiCnw /SiC walls; the red line in the picture indicates the SiCnw walls in the aerogel, on both sides of which dense SiC matrix has deposited. 3.2. Density, porosity and mechanical properties Fig. 6a and b demonstrates the unique microstructure encountered on the surfaces and the fractured sections of the porous ceramic specimens after 3 (termed “a” specimen) and 5 (termed “b” specimen) CVI cycles, respectively. As seen in Fig. 6a, some aligned grooves were observed on the surface after polishing, which were parallel to the macroscopic direction of ice formation. After 2 additional CVI runs the aligned grooves were filled with SiC matrix as seen in Fig. 6b. The inset micrographs of Fig. 6a and b are top views of the fractured sections of specimens “a” and “b”; the views are similar to the top view of the SiCnw aerogel. Open pores with a flat shape and unidirectionally aligned channels encircled by SiCnw /SiC walls are seen isolated in the ceramic. However, by examination of the two inset micrographs, it is found that the pore size distribution, level of interconnection between pores and pore orientation of the two specimens are similar. This finding is believed to be linked to the CVI process and pore structure of the aerogel as explained in the following. As the CVI process is strongly dependent on diffusion mechanisms, the morphological conditions of the aerogel are conducive to gas flow throughout the material volume where a chemical reaction takes place between the gaseous species diffusing through the pores. However, diffusion conditions are superior at the surfaces as compared to the inside, hence deposition of SiC matrix inside the aerogel is anticipated to be much suppressed compared to its outer surfaces resulting in ceramic with unfilled channels existing between SiCnw /SiC walls and closed surfaces. With time, a pure SiC coating was formed on the outer surface which obstructed further SiC infiltration inside the material. Therefore, it is believed that the effect of the additional 2 and 4 consecutive CVI runs on the pore structure of the ceramics is small and limited to filling of surface grooves and increasing of SiC coating thickness. The outer wall thickness increased with increasing CVI cycles and the inter wall thickness changed a little. The micro-pore channels were closed at the ends/surfaces due to SiC deposition but they were open inside the body. This effect is also linked to the pore size distributions discussed in the following section. An enlarged micrograph of the fractured section of the SiCnw /SiC walls is shown in Fig. 7 and manifests the existence of dense SiC matrix on both sides. Porosity and density of the ceramics after 3, 5, and 7 CVI cycles were assessed by MIP; the results are summarized in Table 1 while pore size distributions of three as-fabricated porous ceramics are shown in Fig. 8. Even though absolute porosity values varied from specimen to specimen, all samples exhibited an almost identical, non-uniform, unimodal pore size distribution shape varying in the range of 20 to 350 m, with a maximum peak at a pore size of 30–80 m. The relative volume content of pore diameters rang-
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Fig. 4. SEM micrographs of the SiCnw aerogel; (a) front view, which is the plane parallel to the macroscopic ice growth, (b) top view, which is the plane perpendicular to the macroscopic ice growth, (c) starting SiCnw , and (d) the enlarged images of the SiCnw “walls”.
Fig. 5. SEM micrographs of the porous SiCnw /SiC ceramic; (a) top view, which is the plane of the ceramic after one CVI cycle perpendicular to the macroscopic ice growth, and (b) cross-section surface of the SiCnw /SiC “walls”.
Fig. 6. SEM micrographs of the surface and fractured section of the specimens; (a) “a” specimen, and (b) “b” specimen.
ing from 30 to 80 m, and 120 to 350 m were 4–16 and 0–3%, respectively. The biggest pore sizes of the three specimens were
36 m with relative volume contents of 15.45, 11.25 and 14.21%, the highest observed. The particular pore size distributions can
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Fig. 9. Variation in flexural and compressive strength of the porous ceramics as a function of porosity. Fig. 7. The enlarged SEM micrograph of the fractured section of the SiCnw /SiC “walls”.
be attributed to the microstructure of the specimens and the specific CVI process. The effect of 2 and 4 additional CVI runs on the pore structure of the ceramics is small as discussed in the previous section. Fig. 8b shows the differential pore volume of the three specimens. It was found that the differential pore volume of specimens with different CVI cycles exhibited almost identical trends and decreased significantly with increasing CVI cycles, a finding which can be attributed to the thicker SiC coating and the relative reduction of the pore volume. Fig. 9 shows the variation in flexural strength of the porous ceramics as a function of porosity. Flexural strengths of 122 ± 10, 180 ± 16 and 270 ± 25 MPa were established for porosity of 32.23, 23.44 and 9.23%, respectively, hence proving a monotonically decreasing linear relation between the two properties. The increase in strength is mainly due to the increase in SiC coating thickness. The typical flexural stress-displacement responses of the specimen are shown in Fig. 10. It is observed that all specimens exhibited typical brittle fractures, and reflected some smaller waved path, which are associated with undulations of low and high-density regions (according to the channels and SiCnw /SiC “walls” in the porous ceramic). Due to the anisotropic macrostructure of the fabricated ceramics, their strength and failure mode under compressive loading was anticipated to be highly dependent on loading direction. Eventually, transverse compression strengths of 54 ± 3, 65 ± 5, 110 ± 15 MPa, and longitudinal compression strengths of 249 ± 40, 401 ± 20, 496 ± 29 MPa were established for 3, 5 and 7 consecutive CVI runs, respectively; the longitudinal compressive strength
Fig. 10. Typical flexural stress-displacement curves of the porous SiCnw /SiC ceramic with different porosity.
was much higher than its transverse counterparts. The typical compressive stress-displacement curves of the porous SiCnw /SiC ceramics in the directions parallel and perpendicular to the freezing direction are shown in Fig. 11. It is observed that the stressstrain curve in the freezing direction features a different shape from that perpendicular to the freezing direction. Recalling that the compressive strengths in the freezing direction were much higher than that perpendicular to the freezing direction, the nature of the observed mechanical response can be attributed to the anisotropic microstructure produced by the unidirectional freezedrying process. When the compressive load and freezing directions are parallel, i.e., the load is parallel to the SiCnw /SiC walls in the specimens, load distributes equally on all SiCnw /SiC walls until sud-
Fig. 8. Pore size distribution curve of the porous ceramic specimens; (a) relative volume content, and (b) differential pore volume.
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Fig. 11. Typical compressive stress-displacement curves of the porous SiCnw /SiC ceramic with different porosity; (a) longitudinal, and (b) transverse.
Fig. 12. Schematic diagram of the compressive load in the direction perpendicular to the freezing direction.
den fracture, hence the evidenced brittle fracture behavior with elastic response to failure. On the opposite scenario, when load is perpendicular to the freezing direction, (the loading plane is assumed to be the white line marked in Fig. 12), the compressive load direction is perpendicular to the SiCnw /SiC walls (a), and possesses certain discrete angles (ranging from 0 to 90◦ ) between loading direction and SiCnw /SiC walls (b). Due to the irregular crosslinking among SiCnw /SiC walls, load-transmitting path becomes zig-zag which causes stress concentration at curvature and branching positions and rapid propagation of the dominant macro-crack leading to ultimate material failure. The inflection points in the curves are believed to be correlated to load transfer from one SiCnw /SiC wall to another due to the aligned channels encircled by the walls. However, after load reaches its peak value, the mechanical response exhibits relatively flat and low curves in the failure region. This behavior can be attributed to the sequential gradual failure of SiCnw /SiC walls prolonging the failure stage. By examination of the curves’ shape, it was also observed that the height of curves in the elastic region increases with CVI run count; the finding can be attributed to the increased SiC coating thickness resulting in the improvement of the load-bearing capacity of the porous ceramics.
Fig. 13. SEM micrographs of the nanowire pull-out on the cross-section surface of the porous SiCnw /SiC ceramic, in which red arrows indicate nanowire pulled out from the matrix, and the white arrows indicate holes due to nanowire pull-out. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.).
due to the low preparation temperature [20]. In Fig. 13, locations of nanowire pull-out from the matrix are indicated by red arrows while white arrows are used to highlight locations of residual holes due to complete nanowire disengagement from the matrix. Nanowire pull-out due to debonding and friction due to sliding of the nanowire along the debonded interface are toughening mechanisms associated with the consumption of fracture energy attributed to the rough surface and large aspect ratio of nanowires which are effective in improving the resistance of nanowire pulling out and crack propagating [20]. Typical X-ray diffraction patterns of porous SiCnw /SiC ceramics are shown in Fig. 14. Findings suggest that the as-fabricated porous ceramics include -SiC, stacking fault (SF), and SiO2 . The diffraction peaks at 35.60◦ , 59.98◦ , 71.78◦ and 75.49◦ were aligned with the (111), (220), (311) and (222) planes of -SiC, respectively. There was a low-density peak (marked with SF) [29,30] at a diffraction angle of 33.7◦ near the strong (111) peak, which suggested that SFs were contained in the SiC. In addition, diffraction peak at 26.6◦ corresponds to SiO2 characteristic peak, which is attributed to the reaction of oxygen contained in CMC with SiC.
3.3. Mechanisms and phase composition
4. Conclusions
Fig. 13 presents SEM micrographs of the cross-section of failed porous ceramics. It is observed that the nanowires retain their original appearance due to their high tensile strength; their physical and mechanical properties could be maintained after the CVI process
Porous SiCnw /SiC ceramics with unidirectionally aligned channels were synthesized by unidirectional freeze-drying and subsequent infiltration of SiC matrix by the CVI method. The ceramics had high porosities with unique macroscopically aligned
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Fig. 14. XRD patterns of the porous SiC ceramic.
channel structure attributed to the freeze-drying process. After 3, 5, and 7 CVI runs, the porosity and density of the ceramics were 32.23, 23.44, 9.23% and 1.99, 2.27 and 2.30 g/cm3 , respectively. Specimens exhibited typical brittle fracture behaviors with flexural strengths of 122 ± 10, 180 ± 16 and 270 ± 25 MPa for porosity of 32.23, 23.44 and 9.23%, respectively, while transverse compression strengths of 54 ± 3, 65 ± 5, 110 ± 15 MPa, and much higher longitudinal compression strengths of 249 ± 40, 401 ± 20, 496 ± 29 MPa were established for 3, 5 and 7 CVI runs, respectively. Indications of toughening mechanisms were identified by SEM as extensive nanowire pull-out on the cross-section of failed specimens while the phase composition of the porous ceramic included -SiC, stacking fault (SF), and SiO2 . Acknowledgements This work had been financially supported by Natural Science Foundation of China (51272210 and 50902112), Program for New Century Excellent Talents in University (NCET-13-0474), Foreign Talents Introduction and Academic Exchange Program of China (B08040), and Scientific Research Fund of Jilin Provincial Education Department (2015437). References [1] C. Pham-Huu, C. Bouchy, T. Dintzer, G. Ehret, C. Estournes, M.J. Ledoux, High surface area silicon carbide doped with zirconium for use as catalyst support. Preparation, characterization and catalytic application, Appl. Catal. A: Gen. 180 (1–2) (1999) 385–397. [2] N. Keller, C. Pham-Huu, S. Roy, M.J. Ledoux, C. Estournes, J. Guille, Influence of the preparation conditions on the synthesis of high surface area SiC for use as a heterogeneous catalyst support, J. Mater. Sci. 34 (13) (1999) 3189–3202. [3] P.H. Pastila, V. Helanti, A.P. Nikkilä, T.A. Mäntylä, Effect of crystallization on creep of clay bonded SiC- filters, in: D.E. Bray (Ed.), 22nd Annual Conference on Composites, Advanced Ceramics, Materials, and Structures: B: Ceramic Engineering and Science Proceedings (1998) 37–44. [4] I. Nettleship, Applications of porous ceramics, Key Eng. Mater. 122–124 (122) (1996) 305–324. [5] P. Pastila, V. Helanti, A.P. Nikkilä, T. Mäntylä, Environmental effects on microstructure and strength of SiC-based hot gas filters, J. Eur. Ceram. Soc. 21 (9) (2001) 1261–1268. [6] U.F. Vogt, L. Györfy, A. Herzog, T. Graule, G. Plesch, Macroporous silicon carbide foams for porous burner applications and catalyst supports, J. Phys. Chem. Solids 68 (5–6) (2007) 1234–1238.
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Porous SiCnw /SiC ceramics with unidirectionally aligned channels produced by freeze-drying and chemical vapor infiltration Daoyang Han, Hui Mei ∗ , Shanshan Xiao, Junchao Xia, Jinlei Gu, Laifei Cheng Science and Technology on Thermostructural Composite Materials Laboratory, Northwestern Polytechnical University, Xi’an Shannxi 710072, PR China
a r t i c l e
i n f o
Article history: Received 27 August 2016 Received in revised form 12 October 2016 Accepted 13 October 2016 Available online xxx Keywords: Unidirectional freeze-drying Silicon carbide nanowire (SiCnw ) aerogel Chemical vapor infiltration (CVI) Porous silicon carbide nanowire/silicon carbide (SiCnw /SiC) ceramics
a b s t r a c t The current study introduces a methodology for the fabrication of porous silicon carbide nanowire/silicon carbide (SiCnw /SiC) ceramics with macroscopic unidirectionally aligned channels and reports on their microstructural and mechanical properties. The material was produced by freezing of a water-based slurry of -SiC nanowires (SiCnw ) with control of the ice growth direction. Pores were subsequently generated by sublimation of the columnar ice during freeze-drying. Chemical vapor infiltration (CVI) of SiC into the open pore network of the SiCnw aerogel with unidirectionally aligned channels, resulted in the formation of highly porous SiCnw /SiC ceramics which exhibited a unique microstructure as identified by scanning electron microscopy. The pore size distribution and the mechanical properties of the asfabricated porous ceramics were examined by mercury intrusion porosimetry and three-point bending and compression tests, respectively, while phase composition was investigated through X-ray diffraction. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Porous silicon carbide (SiC) ceramic is the preferable material in a wide range of applications such as catalyst supports [1,2], filters for hot gas or molten metal [3–5], and high efficiency combustion burners [6]. This is due to its unique properties such as low density, good thermal shock resistance [7], high mechanical properties and excellent chemical stability at elevated temperatures [8]. The mechanical performance of a porous SiC ceramic depends primarily on the selected fabrication method which affects the pore size distribution, level of interconnection between pores, pore orientation, as well as on the intrinsic properties of the material used as the matrix. There are four conventional methodologies suggested for manufacturing porous SiC ceramics, namely sacrificial template, reaction forming/bonding, replication and partial sintering [9]. Recently, freeze casting was newly suggested for the preparation of porous ceramics; the technique has attracted increased scientific interest due to its simplicity and environmental friendliness [10–14]. Therein, a ceramic slurry acting as the freezing vehicle is poured into a mold at a temperature below its freezing point. This leads to the formation of a bicontinuous structure, consisting of three dimensionally interconnected frozen vehicle network and networks of concentrated ceramic particles. Porous ceramics are
∗ Corresponding author. E-mail addresses: [email protected], [email protected] (H. Mei).
readily produced by removing the frozen vehicle network via freeze drying, followed by sintering of ceramic walls [15]. In Ref. [15], highly-aligned porous SiC ceramics decorated with SiC nanowires were fabricated by unidirectional freeze casting of SiC/camphene slurry with various contents of polycarbosilane (PCS) preceramic used simultaneously both as binder and precursor for in situ growth of SiC nanowires. In [16], porous SiC ceramics with good high temperature mechanical properties were fabricated by a freeze-drying combined with solid sintering. Yoon [17] et al. fabricated highlyaligned porous SiC ceramics with well-defined pore structures by freezing a PCS/camphene solution, followed by the pyrolysis of the porous PCS objects at 1400 ◦ C for 1 h in a flowing Ar atmosphere. In Refs. [18,19], porous SiC ceramics were also fabricated by freezecasting and solid state sintering. In Ref. [20], SiC whisker-reinforced SiC ceramic matrix composites were prepared by depositing SiC matrix into the SiC whisker preforms using chemical vapor infiltration (CVI). The latter method is renowned as the most effective and mature means of preparing SiC matrix with ultra-pure, minimal residual stresses and controllable grain sizes [21–26]. Driven by previous successful applications of unidirectional freezing [10] and CVI technology [20] in the production of porous ceramics, we report herein a new strategy for the production of porous SiC ceramics. Initially, -SiC nanowire (SiCnw ) aerogels with unidirectionally aligned channels were prepared by successive freezing, storage in the frozen state, and defrosting of SiCnw solutions. The role of SiCnw is to prepare a macroporous preform with an open pore network. The as-fabricated preform was strong
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Fig. 1. Schematic of polytetrafluoroethylene mold, unit is mm; 1: sheet metal, 2: polytetrafluoroethylene mold.
enough to self-assemble and very ideal for the diffusion of gases using a CVI process. After densification using multiple CVI cycles, the mechanical properties of the obtained porous SiCnw /SiC ceramics such as compressive and flexural strength were investigated by compression and three-point bending tests, while scanning electron microscopy and X-ray diffraction were used to examine the microstructure and phase composition of the ceramic. 2. Experimental 2.1. Material preparation protocol Porous SiCnw /SiC ceramics were fabricated following a newly established experimental protocol of the successive steps presented in the following. Initially, 0.75 g of -SiCnw of diameters 0.1–0.6 m, lengths 50–100 m, purity higher than 96% and density of 3.21 g/cm3 (Changsha Sinet Advanced Materials Co., Ltd. Changsha, China) were dispersed in 100 g of aqueous solution of 0.8 wt% sodium carboxymethyl cellulose (CMC), an etherified derivative of cellulose, (Cekal 30000A, Shenzhen Hairenda Technology Co. Ltd. Shenzhen, China). Dispersion was achieved through ultrasonication for 1.0 h in a 250 ml glass beaker using an BILON1500 ultrasonic homogenizer (Xian Bilon Biological Technology Co. Ltd) operating at 20 ± 1 kHz, 300–1500 W with a 10 mm diameter probe and an on/off pulse rate of 1/s. During ultrasonication the solution was retained at room temperature by continuous cooling fluid circulation through the heat exchanger. CMC was added into the suspensions of particles with the aim of dispersing and bonding the particles in the freeze-casting process. The resultant slurry was poured into a polytetrafluoroethylene mold as shown in Fig. 1. The mold was composed of two parts: the metal with high thermal conductivity attached to the mold bottom, and the mold side made of polytetrafluoroethylene. It was subsequently placed on the pre-cooled at Ts = −80 ◦ C stainless-steel circular tray of a freeze-dryer (LGJ-18S, Beijing Song Yuan Hua Xing Science and Technology Develop Co. Ltd). In this manner, ice growth was forced to occur macroscopically in the vertical direction, causing the formation of columnar ice, the phase separation of water and nanowires [27] (Fig. 2), and the piling up of the nanowires between the columnar ice. After the slurry froze completely, vacuum conditions were imposed by a vacuum pump and the stainless steel circular tray was heated and the frozen specimens lyophilized. The thermal cycle used in the lyophilize process is as follows: tray temperature −50 ◦ C for 1 h, −20 ◦ C for 2 h, 10 ◦ C for 2 h, 20 ◦ C for 3 h, 30 ◦ C for 3 h, 40 ◦ C for 4 h, 50 ◦ C for 4 h; gas pressure remained
Fig. 2. Pattern formation and SiCnw segregation during freeze casting of SiCnw slurry.
below 30 Pa. At the end of the sublimation step, a SiCnw aerogel with unidirectionally aligned channels was formed. The aerogel was subsequently inserted into a CVI furnace where it was infiltrated with SiC. During the CVI process, methyltrichlorosilane (CH3 SiCl3 , abbreviated as MTS) was used as precursor gas, hydrogen as carrier and argon as diluting gas. Extensive details of the employed CVI process can be found in [28]. The infiltration time was set to 80 h and the process was repeated 3 consecutive times. After 3 cycles, the as-fabricated ceramic plates were reduced in thickness of 2 mm from each face using 800 # grit emery paper. The resulting 3 mm thick plates were cut into rectangular parallelepipeds of dimensions 4 × 3 × 50 mm3 and 3 × 3 × 50 mm3 using a laser cutting machine. These specimens with smooth surfaces were further infiltrated with SiC matrix for 2 and 4 consecutive CVI runs, respectively. Finally, three sets of specimens with 3, 5 and 7 cycles were subjected to physical characterization. 2.2. Testing and calculations To assess the porosity of the ceramics, mercury intrusion porosimetry (MIP) was conducted on an AutoPore IV 9500 Mercury Porosimeter (Micromeritics Instruments, USA). Three-point bending and compression tests were conducted on an electronic universal testing machine (SANS CMT 4304, Shenzhen, China). The
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Table 1 Porosity and density of the porous SiCnw /SiC ceramic.
Fig. 3. Schematic diagram of the measurement of static compressive strength.
dimensions of compression specimens were 3 × 3 × 3 mm3 and tests was conducted under crosshead displacement control, at a speed of 0.5 mm/min. Compressive tests were performed both perpendicularly (transverse compressive) and parallel (longitudinal compression) to the ice growth orientation as illustrated in Fig. 3. Three-point bending specimens had dimensions of 40 × 4 × 3 mm3 and crosshead speed and load span of 0.5 mm/min and 30 mm, respectively, were used. Compressive strength and flexural strength were calculated as: P a = S
(1)
3PL b = 2bh2
(2)
where P is the load, L is the support span, b is the specimen width, h its height, and S the cross sectional area. Fifteen specimens were tested for statistically accurate calculation of compressive and flexural strengths. The microstructure of the SiCnw aerogel and porous SiCnw /SiC ceramic was observed by scanning electron microscopy (SEM, Hitachi S-4800, Tokyo, Japan). The phase composition of the ceramic was examined by X-ray diffraction (XRD, D8-Advance, Bruker, Germany). 3. Results and discussion 3.1. Microstructure of SiCnw aerogel and porous SiCnw /SiC ceramic During the freezing process, the formation of crystalline ice causes every solute originally dispersed in the aqueous medium to be expelled to the boundaries between adjacent ice crystals [27] (Fig. 2). Due to the higher thermal conductivity of the metal part of the mold compared to that of polytetrafluoroethylene, the slurry in the special mold froze unidirectionally hence causing the formation of columnar ice parallel to the ice growth orientation while the solute was expelled to the boundaries between the columnar ice. Subsequent freeze-drying gave rise to a cryogel, a macroporous structure characterized by walls of matter enclosing empty areas where ice crystals originally resided. The freeze-drying process also allowed for achievement of monoliths which preserved the size and shape of the parent container submitted to freezing. Fig. 4 presents SEM micrographs of the fabricated SiCnw aerogels. Fig. 4a and b shows front and top views of an aerogel, which are the planes parallel and perpendicular to the macroscopic direction of ice formation, respectively. Open micro-pores and micro-pore channels of sizes of several hundred micrometers are observed in the aerogel, aligned along the macroscopic direction of ice formation similarly to channels. The high magnification image in Fig. 4d demonstrates the SiCnw walls in the aerogel, comparable to the parent SiCnw (Fig. 4c), which exhibited mesh-like morphology bonded by CMC. In the walls, there were many smaller micro-pores of dimensions of several micrometers attributed to smaller ice particles existing inside the SiCnw walls in the frozen slurry. These micro-pores and micro-pore channels provide paths for subsequent SiC infiltration by CVI. In summary, the aerogel microstructure was comprised
CVI cycles
3
5
7
Porosity/% Density/g/cm3
32.23 1.99
23.44 2.27
9.23 2.30
by networks of unidirectionally-aligned channels achieved by the employed freeze-drying process. SEM micrographs of the porous SiCnw /SiC ceramics are shown in Fig. 5. As seen in Fig. 5a and the inset micrograph therein, all SiCnw walls in the aerogel were completely filled with SiC matrix. The top face of the ceramic showed a morphology similar to the one developed during SiCnw aerogel formation. Fig. 5b shows a typical microstructure of the cross-section of the SiCnw /SiC walls; the red line in the picture indicates the SiCnw walls in the aerogel, on both sides of which dense SiC matrix has deposited. 3.2. Density, porosity and mechanical properties Fig. 6a and b demonstrates the unique microstructure encountered on the surfaces and the fractured sections of the porous ceramic specimens after 3 (termed “a” specimen) and 5 (termed “b” specimen) CVI cycles, respectively. As seen in Fig. 6a, some aligned grooves were observed on the surface after polishing, which were parallel to the macroscopic direction of ice formation. After 2 additional CVI runs the aligned grooves were filled with SiC matrix as seen in Fig. 6b. The inset micrographs of Fig. 6a and b are top views of the fractured sections of specimens “a” and “b”; the views are similar to the top view of the SiCnw aerogel. Open pores with a flat shape and unidirectionally aligned channels encircled by SiCnw /SiC walls are seen isolated in the ceramic. However, by examination of the two inset micrographs, it is found that the pore size distribution, level of interconnection between pores and pore orientation of the two specimens are similar. This finding is believed to be linked to the CVI process and pore structure of the aerogel as explained in the following. As the CVI process is strongly dependent on diffusion mechanisms, the morphological conditions of the aerogel are conducive to gas flow throughout the material volume where a chemical reaction takes place between the gaseous species diffusing through the pores. However, diffusion conditions are superior at the surfaces as compared to the inside, hence deposition of SiC matrix inside the aerogel is anticipated to be much suppressed compared to its outer surfaces resulting in ceramic with unfilled channels existing between SiCnw /SiC walls and closed surfaces. With time, a pure SiC coating was formed on the outer surface which obstructed further SiC infiltration inside the material. Therefore, it is believed that the effect of the additional 2 and 4 consecutive CVI runs on the pore structure of the ceramics is small and limited to filling of surface grooves and increasing of SiC coating thickness. The outer wall thickness increased with increasing CVI cycles and the inter wall thickness changed a little. The micro-pore channels were closed at the ends/surfaces due to SiC deposition but they were open inside the body. This effect is also linked to the pore size distributions discussed in the following section. An enlarged micrograph of the fractured section of the SiCnw /SiC walls is shown in Fig. 7 and manifests the existence of dense SiC matrix on both sides. Porosity and density of the ceramics after 3, 5, and 7 CVI cycles were assessed by MIP; the results are summarized in Table 1 while pore size distributions of three as-fabricated porous ceramics are shown in Fig. 8. Even though absolute porosity values varied from specimen to specimen, all samples exhibited an almost identical, non-uniform, unimodal pore size distribution shape varying in the range of 20 to 350 m, with a maximum peak at a pore size of 30–80 m. The relative volume content of pore diameters rang-
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Fig. 4. SEM micrographs of the SiCnw aerogel; (a) front view, which is the plane parallel to the macroscopic ice growth, (b) top view, which is the plane perpendicular to the macroscopic ice growth, (c) starting SiCnw , and (d) the enlarged images of the SiCnw “walls”.
Fig. 5. SEM micrographs of the porous SiCnw /SiC ceramic; (a) top view, which is the plane of the ceramic after one CVI cycle perpendicular to the macroscopic ice growth, and (b) cross-section surface of the SiCnw /SiC “walls”.
Fig. 6. SEM micrographs of the surface and fractured section of the specimens; (a) “a” specimen, and (b) “b” specimen.
ing from 30 to 80 m, and 120 to 350 m were 4–16 and 0–3%, respectively. The biggest pore sizes of the three specimens were
36 m with relative volume contents of 15.45, 11.25 and 14.21%, the highest observed. The particular pore size distributions can
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Fig. 9. Variation in flexural and compressive strength of the porous ceramics as a function of porosity. Fig. 7. The enlarged SEM micrograph of the fractured section of the SiCnw /SiC “walls”.
be attributed to the microstructure of the specimens and the specific CVI process. The effect of 2 and 4 additional CVI runs on the pore structure of the ceramics is small as discussed in the previous section. Fig. 8b shows the differential pore volume of the three specimens. It was found that the differential pore volume of specimens with different CVI cycles exhibited almost identical trends and decreased significantly with increasing CVI cycles, a finding which can be attributed to the thicker SiC coating and the relative reduction of the pore volume. Fig. 9 shows the variation in flexural strength of the porous ceramics as a function of porosity. Flexural strengths of 122 ± 10, 180 ± 16 and 270 ± 25 MPa were established for porosity of 32.23, 23.44 and 9.23%, respectively, hence proving a monotonically decreasing linear relation between the two properties. The increase in strength is mainly due to the increase in SiC coating thickness. The typical flexural stress-displacement responses of the specimen are shown in Fig. 10. It is observed that all specimens exhibited typical brittle fractures, and reflected some smaller waved path, which are associated with undulations of low and high-density regions (according to the channels and SiCnw /SiC “walls” in the porous ceramic). Due to the anisotropic macrostructure of the fabricated ceramics, their strength and failure mode under compressive loading was anticipated to be highly dependent on loading direction. Eventually, transverse compression strengths of 54 ± 3, 65 ± 5, 110 ± 15 MPa, and longitudinal compression strengths of 249 ± 40, 401 ± 20, 496 ± 29 MPa were established for 3, 5 and 7 consecutive CVI runs, respectively; the longitudinal compressive strength
Fig. 10. Typical flexural stress-displacement curves of the porous SiCnw /SiC ceramic with different porosity.
was much higher than its transverse counterparts. The typical compressive stress-displacement curves of the porous SiCnw /SiC ceramics in the directions parallel and perpendicular to the freezing direction are shown in Fig. 11. It is observed that the stressstrain curve in the freezing direction features a different shape from that perpendicular to the freezing direction. Recalling that the compressive strengths in the freezing direction were much higher than that perpendicular to the freezing direction, the nature of the observed mechanical response can be attributed to the anisotropic microstructure produced by the unidirectional freezedrying process. When the compressive load and freezing directions are parallel, i.e., the load is parallel to the SiCnw /SiC walls in the specimens, load distributes equally on all SiCnw /SiC walls until sud-
Fig. 8. Pore size distribution curve of the porous ceramic specimens; (a) relative volume content, and (b) differential pore volume.
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Fig. 11. Typical compressive stress-displacement curves of the porous SiCnw /SiC ceramic with different porosity; (a) longitudinal, and (b) transverse.
Fig. 12. Schematic diagram of the compressive load in the direction perpendicular to the freezing direction.
den fracture, hence the evidenced brittle fracture behavior with elastic response to failure. On the opposite scenario, when load is perpendicular to the freezing direction, (the loading plane is assumed to be the white line marked in Fig. 12), the compressive load direction is perpendicular to the SiCnw /SiC walls (a), and possesses certain discrete angles (ranging from 0 to 90◦ ) between loading direction and SiCnw /SiC walls (b). Due to the irregular crosslinking among SiCnw /SiC walls, load-transmitting path becomes zig-zag which causes stress concentration at curvature and branching positions and rapid propagation of the dominant macro-crack leading to ultimate material failure. The inflection points in the curves are believed to be correlated to load transfer from one SiCnw /SiC wall to another due to the aligned channels encircled by the walls. However, after load reaches its peak value, the mechanical response exhibits relatively flat and low curves in the failure region. This behavior can be attributed to the sequential gradual failure of SiCnw /SiC walls prolonging the failure stage. By examination of the curves’ shape, it was also observed that the height of curves in the elastic region increases with CVI run count; the finding can be attributed to the increased SiC coating thickness resulting in the improvement of the load-bearing capacity of the porous ceramics.
Fig. 13. SEM micrographs of the nanowire pull-out on the cross-section surface of the porous SiCnw /SiC ceramic, in which red arrows indicate nanowire pulled out from the matrix, and the white arrows indicate holes due to nanowire pull-out. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.).
due to the low preparation temperature [20]. In Fig. 13, locations of nanowire pull-out from the matrix are indicated by red arrows while white arrows are used to highlight locations of residual holes due to complete nanowire disengagement from the matrix. Nanowire pull-out due to debonding and friction due to sliding of the nanowire along the debonded interface are toughening mechanisms associated with the consumption of fracture energy attributed to the rough surface and large aspect ratio of nanowires which are effective in improving the resistance of nanowire pulling out and crack propagating [20]. Typical X-ray diffraction patterns of porous SiCnw /SiC ceramics are shown in Fig. 14. Findings suggest that the as-fabricated porous ceramics include -SiC, stacking fault (SF), and SiO2 . The diffraction peaks at 35.60◦ , 59.98◦ , 71.78◦ and 75.49◦ were aligned with the (111), (220), (311) and (222) planes of -SiC, respectively. There was a low-density peak (marked with SF) [29,30] at a diffraction angle of 33.7◦ near the strong (111) peak, which suggested that SFs were contained in the SiC. In addition, diffraction peak at 26.6◦ corresponds to SiO2 characteristic peak, which is attributed to the reaction of oxygen contained in CMC with SiC.
3.3. Mechanisms and phase composition
4. Conclusions
Fig. 13 presents SEM micrographs of the cross-section of failed porous ceramics. It is observed that the nanowires retain their original appearance due to their high tensile strength; their physical and mechanical properties could be maintained after the CVI process
Porous SiCnw /SiC ceramics with unidirectionally aligned channels were synthesized by unidirectional freeze-drying and subsequent infiltration of SiC matrix by the CVI method. The ceramics had high porosities with unique macroscopically aligned
Please cite this article in press as: D. Han, et al., Porous SiCnw /SiC ceramics with unidirectionally aligned channels produced by freezedrying and chemical vapor infiltration, J Eur Ceram Soc (2016), http://dx.doi.org/10.1016/j.jeurceramsoc.2016.10.015
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ARTICLE IN PRESS D. Han et al. / Journal of the European Ceramic Society xxx (2016) xxx–xxx
Fig. 14. XRD patterns of the porous SiC ceramic.
channel structure attributed to the freeze-drying process. After 3, 5, and 7 CVI runs, the porosity and density of the ceramics were 32.23, 23.44, 9.23% and 1.99, 2.27 and 2.30 g/cm3 , respectively. Specimens exhibited typical brittle fracture behaviors with flexural strengths of 122 ± 10, 180 ± 16 and 270 ± 25 MPa for porosity of 32.23, 23.44 and 9.23%, respectively, while transverse compression strengths of 54 ± 3, 65 ± 5, 110 ± 15 MPa, and much higher longitudinal compression strengths of 249 ± 40, 401 ± 20, 496 ± 29 MPa were established for 3, 5 and 7 CVI runs, respectively. Indications of toughening mechanisms were identified by SEM as extensive nanowire pull-out on the cross-section of failed specimens while the phase composition of the porous ceramic included -SiC, stacking fault (SF), and SiO2 . Acknowledgements This work had been financially supported by Natural Science Foundation of China (51272210 and 50902112), Program for New Century Excellent Talents in University (NCET-13-0474), Foreign Talents Introduction and Academic Exchange Program of China (B08040), and Scientific Research Fund of Jilin Provincial Education Department (2015437). References [1] C. Pham-Huu, C. Bouchy, T. Dintzer, G. Ehret, C. Estournes, M.J. Ledoux, High surface area silicon carbide doped with zirconium for use as catalyst support. Preparation, characterization and catalytic application, Appl. Catal. A: Gen. 180 (1–2) (1999) 385–397. [2] N. Keller, C. Pham-Huu, S. Roy, M.J. Ledoux, C. Estournes, J. Guille, Influence of the preparation conditions on the synthesis of high surface area SiC for use as a heterogeneous catalyst support, J. Mater. Sci. 34 (13) (1999) 3189–3202. [3] P.H. Pastila, V. Helanti, A.P. Nikkilä, T.A. Mäntylä, Effect of crystallization on creep of clay bonded SiC- filters, in: D.E. Bray (Ed.), 22nd Annual Conference on Composites, Advanced Ceramics, Materials, and Structures: B: Ceramic Engineering and Science Proceedings (1998) 37–44. [4] I. Nettleship, Applications of porous ceramics, Key Eng. Mater. 122–124 (122) (1996) 305–324. [5] P. Pastila, V. Helanti, A.P. Nikkilä, T. Mäntylä, Environmental effects on microstructure and strength of SiC-based hot gas filters, J. Eur. Ceram. Soc. 21 (9) (2001) 1261–1268. [6] U.F. Vogt, L. Györfy, A. Herzog, T. Graule, G. Plesch, Macroporous silicon carbide foams for porous burner applications and catalyst supports, J. Phys. Chem. Solids 68 (5–6) (2007) 1234–1238.
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Please cite this article in press as: D. Han, et al., Porous SiCnw /SiC ceramics with unidirectionally aligned channels produced by freezedrying and chemical vapor infiltration, J Eur Ceram Soc (2016), http://dx.doi.org/10.1016/j.jeurceramsoc.2016.10.015