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A nitride whisker template for growth of mullite in SiC reticulated porous ceramics
Ceramics International 43 (2017) 11197–11203
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A nitride whisker template for growth of mullite in SiC reticulated porous ceramics
MARK
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Xiong Lianga, Yawei Lia, Qinghu Wanga, , Shaobai Sanga, Yibiao Xua, Yuanyuan Chena, Benwen Lib, Christos G. Anezirisc a b c
The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, China Institute of Thermal Engineering, School of Energy and Power Engineering, Dalian University of Technology, Dalian 116024, China Technical University Bergakademie Freiberg, Institute for Ceramic, Glass- and Construction Materials, Agricolastraße 17, 09596 Freiberg, Germany
A R T I C L E I N F O
A BS T RAC T
Keywords: SiC reticulated porous ceramics Nitride whisker template Column-liked mullite Mechanical properties
Silicon carbide reticulated porous ceramics (SiC RPCs) were fabricated by polymer replica technique. The effects of nitride whisker template on the growth of mullite, the strut structure and mechanical properties of SiC RPCs were investigated. Prepolyurethane (PU) open-cell sponge was first coated by SiC slurry consisting of SiC, reactive Al2O3, microsilica and Si powder, then it was nitridized at 1400 °C in a flowing N2 atmosphere to prepare SiC preforms. Subsequently, these preforms were treated by vacuum infiltration of alumina slurry and fired at 1450 °C in air. The results showed that Si2N2O whiskers grew on the surface and in the matrix of SiC preforms after nitridation. The diameter of struts in SiC RPCs increased after vacuum infiltration process because alumina slurry was easily adhered by the surface nitride whiskers. In addition, such whiskers inside the strut of SiC preforms acted as the template to promote the growth of column-liked mullite in SiC RPCs. The mechanical properties and thermal shock resistance of SiC RPCs were greatly improved due to the special interfacial characteristics of multi-layered struts as well as better interlocked column-liked mullite in SiC skeleton.
1. Introduction Porous media combustion (PMC) has unique characteristics, such as high radiant output, low NOx and CO emissions, high flame speed and power density, etc [1–3]. It has been applied in internal combustion engines, low calorific gas burners, volatile organic compound oxidizers and radiation heaters [4–6]. Reticulated porous ceramics (RPCs) with the open and three-dimensional network structure are considered as the most promising high-temperature components for PMC due to their high permeability, thermal stability and resistance to chemical corrosion [7]. However, during service process, the scour of flue gas and large thermal stress from the sharp temperature gradients made porous burners easily damaged, thus the mechanical properties and thermal shock resistance of RPCs should be taken into account [8,9]. In particular, silicon carbide RPCs are considered as the candidate material for PMC because of their low thermal expansion coefficient, high thermal conductivity and strength [10]. The polymer sponge replica technique was the most common method for producing RPCs [11]. However, the hollow struts with triangular voids and longitudinal cracks would generate when the
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Corresponding author. E-mail address: [email protected] (Q. Wang).
http://dx.doi.org/10.1016/j.ceramint.2017.05.169 Received 27 March 2017; Received in revised form 22 May 2017; Accepted 23 May 2017 Available online 25 May 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
sponge was burnt out during firing process. These defects led to the poor mechanical properties of RPCs and limited their structural application [12–14]. Besides, it was difficult to fabricate SiC RPCs with high strength using traditional sintering methods. In general, SiC ceramics were treated via gas pressure sintering at high temperature. This method was unsuitable for fabrication of SiC RPCs, because the fast heating rate caused lots of cracks during the decomposition of polymer sponge [15]. In order to strengthen SiC RPCs, some approaches were proposed, involving in thickening the ceramic skeleton and optimizing the microstructure of strut. The polymeric template was treated by acid or alkali to roughen the sponge surface [16], and the rheological behavior of SiC slurry was adjusted to increase the weight of coated slurry [17]. Also, recoating technique was employed to eliminate the surface flaws and thicken the struts of SiC RPCs [15,18]. Furthermore, vacuum infiltration process was an effective method to obtain the dense struts via filling ceramic slurry into triangular voids of struts [19]. In addition, the ceramic fiber was introduced into the skeleton to optimize the microstructure, thereby strengthening RPCs [20]. The in-situ formed needle-liked Si3N4 was certified as whiskerreinforced mechanism in Si3N4 RPCs when the whisker integrated with
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the liquid phase as a whole, which needed extremely high sintering temperature [21,22]. For increasing the mechanical properties of SiC RPCs, the in-situ nitridation of Si powder together with vacuum infiltration process were applied to improve the macro- and microstructure of struts in this work. The green SiC RPCs containing Si powder were intendedly nitridized to produce SiC preforms, where the whiskers formed on the strut surface and in the matrix. As vacuum infiltration of alumina slurry was carried out, the hollow strut was densified and alumina slurry would be adhered by the surface whiskers. After the infiltrated SiC preforms were fired in air, the whisker might be oxidized to active silica and promoted the mullitization reaction. In the experiment, the nitridation behavior of Si powder in SiC preforms was first investigated. The effects of nitride whiskers on the growth of mullite, the evolution of microstructure and mechanical properties of SiC RPCs were also discussed in details. 2. Experimental 2.1. Preparation of SiC RPCs A commercial available SiC powder ( < 45 µm, 98 wt%, China), Si (~ 3 µm, 98.5 wt%), α-Al2O3 (~ 1.2 µm, Henan Special Refractories Co., Ltd., China) and microsilica powder (~ 0.5 µm, 951UL, 0.13 wt% Na2O, 0.05 wt% Fe2O3, Elkem, Norway) were used as starting materials for the fabrication of SiC RPCs. Si powder was used as the precursor of whiskers. Microsilica and α-Al2O3 were added as sintering aids, the Al2O3/SiO2 molar ratio was designed for chemical composition of mullite. Ammonium lignosulfonate (Tianjin institute of fines chemicals, China) and sodium carboxymethyl-cellulose (CMC, Sinopharm Chemical Reagent Co., Ltd, China) were added as binder and thickening agent to improve the rheology of SiC slurry. Polycarboxylate (FS, BASF Group, Germany) and contraspum K 1012 (Zschimmer & Schwarz, Lahnstein, Germany) were used as dispersant agent and antifoam agent, respectively. The SiC slurries with initial solid content of 77.4 wt% were prepared by mixing powder mixture (shown in Table 1) with the above organic additives and deionized water. SiC slurry was mechanically stirred for 1 h at a rotate speed of 300 rpm to obtain a high thixotropic SiC slurry. The polyurethane open-cell sponge (10 pores/inch, F. M Co. Ltd., Germany) with a size of 50 mm × 50 mm × 20 mm was immersed into the as-prepared SiC slurry, followed by passing through preset roller to remove excess slurry. Remaining closed pores were eliminated by blowing compressed air through the sponge structure. After dried at room temperature, the coated sponges were treated under N2 atmosphere (99.999%, mass fraction) at 1400 °C with a holding time of 1 h to produce SiC preforms. These SiC preforms were named as S0, S5 and S10 respectively according to the content of Si powder in SiC slurries. In the stage of vacuum infiltration, alumina slurry with the solid content of 77 wt% was prepared without binder. Deionized water was mixed with FS and K 1012 by stirring for 5 min. Subsequently, α-Al2O3 powder were added into the solution and ball-milled for 3 h. Then SiC preforms were totally immersed into the alumina slurry and a vacuum of 5 × 10−6 bar was applied for 20 min. The as-infiltrated SiC preforms were fired in air at 1450 °C for 3 h. SiC preforms of S0, S5 and S10 with vacuum infiltration of alumina slurry were named as AS0, AS5 and AS10, respectively. Table 1 Compositions of SiC slurries. Specimen
SiC (wt%)
Microsilica (wt%)
α-Al2O3 (wt%)
Si (wt%)
S0 S5 S10
80 80 80
5.84 5.84 5.84
14.16 14.16 14.16
– +5 + 10
Furthermore, in order to ascertain the effect of silica source on the morphology of reaction-sintered mullite, the matrix specimens comprising of α-Al2O3 and different silica sources were prepared by uniaxially dry pressing method at 5 MPa. These silica sources were microsilica, silicon powder and silicon carbide, respectively. The matrix specimen with silicon powder addition was first nitridized at 1400 °C for 1 h, which was named as SN. The matrix containing microsilica and silicon carbide were named as SO and SC, respectively. All the prepared matrix specimens were heated in air at 1450 °C for 3 h. 2.2. Characterization The RPCs density ρRPC was the density of the whole specimen and was calculated with Eq. (1):
ρRPC =
m l ∙b∙h
(1)
Where l was the length, b the width, h the height, and m the mass of the RPCs specimen. The macrostructure of SiC RPCs was characterized by digital camera and the diameter of struts were analyzed by Image-Pro plus software (Media Cybernetics, Inc., Netherlands). Mechanical tests on the fired RPCs were performed by determining the cold compressive strength (CCS) with a universal testing machine (ETM, Wance, China), using a load speed of 0.5 mm/min. Specimens of the size 50 mm × 50 mm × 20 mm were positioned between loading plates (Φ100 × 20 mm). A cardboard of 4 mm thickness was placed between loading plate and specimen to obtain a uniform loading. The thermal shock resistance of SiC RPCs was evaluated by the water-quenching technique. The fired SiC RPCs were heated at 1100 °C in air for 20 min in a furnace, then immersed into the flowing water of 25 °C. After 3 thermal shock cycles, the residual strength of specimen was determined at room temperature via compression test. The residual strength ratio of CCS was calculated by the change in CCS before and after quenching test, i.e. the residual strength ratio of CCS = 100CCSTS/CCS, where CCS and CCSTS were the CCS before and after 3 thermal shock cycles, respectively. The statistical results of physical and mechanical properties were obtained based on a dataset of 6 specimens. The Gibbs free energy in the reaction of Si powder nitridation was calculated by FactSage software. X-ray diffraction (XRD, X' pert Pro, Philips, Netherlands) was conducted to examine the phase compositions of SiC preforms and RPCs. The microstructure of SiC preform after nitridation and mullite morphology in SiC RPCs were observed by scanning electron microscope (SEM, Quanta 400, FEI Company, USA), equipped with energy dispersive X-ray spectroscopy (EDS, Noran 623 M-3SUT, Thermo Electron Corporation, Japan). 3. Results and discussion 3.1. SiC preforms 3.1.1. Phase composition Fig. 1 shows the XRD patterns of SiC preforms nitridized at 1400 °C. In specimen S0, SiC, α-Al2O3, quartz phases were detected, indicating that mullitization process had not occurred and the added microsilica transformed into quartz phase at this temperature. For SiC RPCs with Si addition, quartz phase was not detected whereas Si2N2O formed on the contrary. The diffraction peak intensity of Si2N2O phase increased with the content of Si powder. Furthermore, the addition of Si powder led to the weight change of SiC preforms (Table 2). In specimen S5 and S10, Si powder completely transformed into Si2N2O phase, thus resulting in the weight gain of SiC preforms. However, the experimental value of weight gain was much lower than theoretical ones. It plausibly related to the gaseous phases of Si(g) and SiO(g) diffused out of SiC struts during the nitridation process of silicon [23– 25], which escaped from SiC preforms together with the flowing N2.
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Fig. 3. XRD patterns of SiC RPCs after fired at 1450 °C.
Fig. 1. The XRD patterns of SiC preforms nitridized at 1400 °C. Table 2 Phase composition and weigh gain of specimens after nitridation process. Specimen
S0 S5 S10
Phase composition
Weight gain (%)
SiC-4H, SiC-6H, corundum, quartz SiC-4H, SiC-6H, corundum, Si2N2O SiC-4H, SiC-6H, corundum, Si2N2O
Theoretical
Experimental
0 3.74 7.14
0 0.74 1.65
3.1.2. Microstructure SEM micrographs of SiC preforms are shown in Fig. 2. Unlike specimen S0 where the amorphous aluminosilicate embedded in the
matrix (Fig. 2(a-b)), the nitridation of Si powder in-situ formed whiskers in the strut of SiC preforms. In specimen S5, the strut surface was covered by large number of blanket-liked whiskers (Fig. 2(c)). At higher magnification, these bent whiskers with large length-diameter ratio were interlocked with each other and well distributed in the matrix as well as on the strut surface (Fig. 2(d)). According to XRD and EDS analysis, they were confirmed to be Si2N2O whiskers. With increasing the content of Si powder to 10 wt%, the amount of whiskers increased and the strut surface was almost completely covered by whiskers (Fig. 2(e)). Besides, the whiskers with larger length-diameter ratio existed inside the triangular void of specimen S10 (Fig. 2(f)). It was worth noting that the more content of Si powder addition, the looser strut appeared in SiC preforms.
Fig. 2. SEM micrographs of SiC preforms after nitridation at 1400 °C. (a, b) S0, (c, d) S5 and (e, f) S10.
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Table 3 Relative diffraction peak intensity of mullite/SiC in SiC RPCs. Specimen
AS0
AS5
AS10
Mullite/SiC
16.50%
19.83%
23.90%
3.2. SiC RPCs 3.2.1. Phase compositions XRD patterns of SiC RPCs after vacuum infiltration of alumina slurry and treatment at 1450 °C are given in Fig. 3. For specimens with Si powder addition, the phase composition of the fired SiC RPCs was much different to SiC preforms. The nitridation products of Si2N2O phase disappeared whereas mullite phase formed in specimen AS5 and AS10. The relative diffraction peak intensity ratio of mullite and SiC increased with the content of Si powder in SiC RPCs (Table 3), indicating that nitride whiskers promoted the formation of mullite in SiC RPCs after firing in air. 3.2.2. Microstructure The cross section of the etched struts in SiC RPCs are presented in Fig. 4. The formation of nitride whiskers in SiC preforms remarkably affected the final strut structures of SiC RPCs. With the growth of whiskers on the surface of SiC preforms, more alumina slurries adhered on SiC skeleton, thus leading to the formation of thicker coating in SiC RPCs than specimen AS0 (Fig. 4(a-c)). Furthermore, some needle-liked mullite formed in the matrix of SiC RPCs, while the lumpy mullite grew around SiC particles in specimen AS0 (Fig. 4(d)). For SiC RPCs with Si powder addition, besides the mullite exhibited the similar morphology to AS0, some column-liked mullite with large
length-diameter ratio was observed in AS5 (Fig. 4(e)). It was worth noting that the amount of mullite increased with the content of Si powder in SiC preforms, especially for the column-liked mullite. Meanwhile, the column-liked mullite interlocked with each other and showed its network structure in the matrix of SiC RPCs (Fig. 4(f)). In addition, the significant interface formed between multi-layers, and its thickness increased with the amount of nitride whiskers in SiC preforms. At higher magnification of interface in AS10 (Zone 1 in Fig. 4(c)), a dense transition layer located between SiC skeleton and infiltrated alumina area, which contained large amount of column-liked mullite (Fig. 5(a)). For the outer layer of strut (Zone 2 in Fig. 4(c)), the continuous coating with good interfacial characteristics formed, which was connected with SiC skeleton by column-liked mullite (Fig. 5(b)). 3.2.3. Physical and mechanical properties Table 4 lists the physical properties of SiC RPCs fired at 1450 °C. The bulk density of SiC RPCs containing Si powder was larger than that of specimen AS0. Meanwhile, the weight gain of infiltrated alumina in SiC RPCs increased with the content of Si powder, leading to the bulk density of specimens increased from 0.38 g/cm3 to 0.45 g/cm3. Because the addition of Si powder promoted the formation of nitride whiskers in SiC preforms, more alumina slurry was able to coat on the surface of strut, the diameter of struts increased at the same time. Mechanical properties including CCS and its residual ratio of SiC RPCs before and after 3 thermal shock cycles are still presented in Table 4. It is apparent that nitridation of Si powder in SiC preforms produced the positive influence on the mechanical properties of SiC RPCs. For example, CCS of specimen AS0 was 0.60 MPa and its value increased with the content of Si powder in SiC RPCs. The CCS reached the maximum in specimen AS10 of 0.83 MPa. It was worth mentioning that the SiC RPCs with Si powder nitridation were less sensitive to
Fig. 4. Cross section struts of SiC RPCs after etched by hydrofluoric acid. (a, d) AS0, (b, e) AS5 and (c, f) AS10.
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Fig. 5. The interfacial characteristic in AS10. (a) and (b) were the higher magnification micrographs of zone 1 and 2 in Fig. 4(c). Table 4 Physical and mechanical properties of SiC RPCs. Specimens
Weight gain of infiltrated alumina g
Bulk density ρb g/cm3
Diameter of struts mm
CCS MPa
CCSTS MPa
CCSTS/CCS
AS0 AS5 AS10
7.13 ± 0.83 8.18 ± 1.29 8.90 ± 1.03
0.38 ± 0.02 0.42 ± 0.04 0.45 ± 0.03
0.56 ± 0.08 0.76 ± 0.14 0.89 ± 0.13
0.60 ± 0.14 0.73 ± 0.11 0.83 ± 0.08
0.40 ± 0.02 0.55 ± 0.03 0.65 ± 0.13
66.7% 75.1% 78.3%
Fig. 6. SEM micrographs of etched matrix SO (a), SS (b) and SC (c) after fired at 1450 °C.
thermal shock than AS0. The residual strength and its residual ratio persistently increased with the content of Si powder, especially specimen AS10 exhibited the largest values of 0.65 MPa and 78.3%, respectively.
3.3. Discussion The nitridation process of SiC preform was controlled under flowing nitrogen of 99.999 wt% in this experiment, therefore the 0.001 wt% oxygen was taken into account in thermodynamic calculation. In SiC preform, Si(s) preferred to react with O2(g) to form SiO(g) through Reaction (2) because of its negative Gibbs free energy. In the flowing nitrogen, Si3N4(s) was easily formed by Reaction (3) [23]. However, in the case that SiO2(s) existed in SiC preforms, Reaction (4) would occur and resulted in the formation of Si2N2O(s). Besides, The O2(g) was considered as the role of catalytic agent from Reaction (2) and (3), and promoted the formation of SiO(g) in the system. Therefore, the growth of Si2N2O whiskers could also be processed via the V-V reaction between SiO(g) and N2(g) when SiO(g) partial pressure was high enough. Meanwhile, the SiO(g) or Si(g) would escape from SiC preforms with the flowing nitrogen at elevated temperature, thus leading to little weight gain in specimen S5 and S10.
Si(s) + 1/2O2(g) = SiO(g) ΔGθ= −109,391 −78.71Τ (298 Κ < Τ < 1685 K) (2) 3SiO(g) + 2N2(g) = Si3N4(s) + 3/2O2(g) ΔGθ= −408,343 + 560.89Τ (298 Κ < Τ < 1685 K) (3) Si3N4(s) + SiO2(s) = 2Si2N2O(s) ΔGθ= −73,137 − 4.76Τ (298 Κ < Τ < 1685 K) (4) 2SiO(g) + N2(g) = Si2N2O(s) + 1/2O2(g) ΔGθ= −688,096 + 429.05Τ (298 Κ < Τ < 1685 K) (5) As the vacuum infiltration of alumina slurry was carried out, the porous and loose skeleton of SiC RPCs was filled up by alumina (Fig. 4). After the infiltrated SiC RPCs was fired at 1450 °C in air, the mullite with three kinds of morphologies formed in the matrix of SiC skeleton. In order to reveal this phenomenon, the mullitization reaction between active alumina and different silica source was analyzed, which was shown in Fig. 6. It was clearly seen that mullite with needle-shape and large grain size formed in the matrix specimen comprising of microsilica and alumina, while the lumpy mullite grew around SiC particles when the silica source was silicon carbide (Fig. 6(a-b)). In addition, nitridation-derived whiskers were beneficial for the growth of mullite
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Fig. 7. Schematic diagram of structure evolution of SiC RPC with Si powder addition. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article).
grain and facilitated the formation of column-liked mullite (Fig. 6(c)). For the growth mechanism of mullite, it was greatly accepted that mullite formation was controlled by “dissolution-precipitation” in reaction sintering containing SiO2 and Al2O3 [26,27]. The SiO2-rich liquid phase was first generated at elevated temperature, then the Al2O3 powders were dissolved in it. When the critical concentration of Al2O3 in SiO2-rich liquid phase was achieved, mullite began to nucleate and grow [28]. The dissolution velocity of Al2O3 into the SiO2-rich liquid phase as well as the template of silica source determined the crystal nucleation and mullite morphology. In the matrix consisting of microsilica and alumina, the SiO2-rich liquid phase with low viscosity easily formed due to the existence of impurity in microsilica [29]. This accelerated atomic mobility and increased the dissolution velocity of α-Al2O3, thereby a large number of mullite crystals started to nucleate when the Al2O3 concentration reached a critical value in SiO2-rich liquid. Furthermore, the liquid reaction medium with its low viscosity contributed to rapid chemical reaction and the growth of mullite, thus the large needle-liked mullite formed though the silica source was spherical microsilica. As the silica source was silicon carbide, the SiO2 film oxidized from silicon carbide easily formed and wrapped on the surface of SiC particles. Meanwhile, α-Al2O3 dissolved into SiO2 film, and mullite nucleation as well as crystal growth occurred. After the mullite crystal formed around the surface of SiC particle layer by layer, the growth of mullite was determined by the diffusion of Al2O3 in mullite layer. The low diffusion velocity of Al2O3 slowed down the mullitization process, only a small amount of initially formed mullite crystals grew to become large crystals, thereby the lumpy mullite formed around SiC particles. Compared to silicon carbide, the nitride whisker with its large specific surface area exhibited high activity and would be oxidized to SiO2 liquid at lower temperature. Besides, the highly reactive and fine SiO2 derived by the oxidation of nitride whiskers would accelerate mullitization reaction and benefited to the formation of mullite crystal [30]. In addition, the nitride whisker played the role of template and induced the mullite to grow along the whisker direction during the mullitization process, resulting in formation of mullite with column-liked morphology [31]. From XRD and SEM results of SiC RPCs with Si powder addition, a schematic diagram could be drawn to reveal the phase composition and structure evolution of struts at elevated temperature (Fig. 7). The green SiC RPCs were prepared after the PU template was coated by SiC slurry via polymer sponge replica technique (Fig. 7(a)). After being nitridized at 1400 °C, the nitridation of Si powder was performed and amounts of nitride whiskers formed both in the matrix of SiC skeleton and on the surface of strut, resulting in a loose skeleton and rough surface (Fig. 7(b)). As the vacuum infiltration process was carried out, the nitride whiskers were wrapped by alumina slurry and the hollow skeleton was densified. Furthermore, the strut with rough surface was beneficial to the adhesion of alumina slurry (Fig. 7(c)), thus thickening the strut and strengthening the SiC RPCs [32]. After the
treatment of infiltrated SiC RPCs at elevated temperature, these nitride whiskers were oxidized to active silica and acted as the template in mullitization process, which promoted the growth of column-liked mullite (Fig. 7(d)). The larger amount and well interlocked columnliked mullite endowed SiC RPCs with better mechanical properties and thermal shock resistance [33,34]. Besides, the multi-layered strut with good interface characteristics also strengthened SiC RPCs. 4. Conclusions The following conclusions can be drawn by studying the influence of Si powder nitridation on the structure and mechanical properties of SiC RPCs. (1) Si powder nitridation resulted in the formation of Si2N2O whiskers both in the matrix of SiC skeleton and on the surface of SiC preforms. (2) The strut of infiltrated SiC RPCs could be thickened after vacuum infiltration process because alumina slurry was easily adhered by the surface nitride whiskers. In addition, in-situ formed whiskers in SiC preforms acted as the template to promote the growth of column-liked mullite in SiC RPCs. (3) The mechanical properties and thermal shock resistance of SiC RPCs with Si powder nitridation were greatly improved compared to non-nitridation ones, which was attributed to the special interface characteristics of multi-layered struts and better interlocked column-liked mullite in SiC skeleton. Acknowledgment The authors thank the subprojects A01 and C03 of German Research Foundation (DFG) for supporting these investigations in terms of the Collaborative Research Centre 920, Multi-Functional Filter s for Metal Melt Filtration-A Contributions towards Zero Defect Materials. References
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[22] X.M. Yao, Y. Yang, X.J. Liu, et al., Effect of recoating slurry compositions on the microstructure and properties of SiC reticulated porous ceramics, J. Eur. Ceram. Soc. 33 (2013) 2909–2914. [23] M.L. Long, Y. Li, H.X. Qin, et al., Formation mechanism of Si3N4 in reactionbonded Si3N4-SiC composites, Ceram. Int. 42 (2016) 16448–16452. [24] J. Yong, L. Wu, P. Wang, et al., Pretreatment and sintering of Si3N4, powder synthesized by the high-temperature self-propagation method, Mater. Res. Bull. 44 (2009) 21–24. [25] Y.H. Zhang, Microstructures and mechanical properties of silicon nitride bonded silicon carbide ceramic foams, Mater. Res. Bull. 39 (2004) 755–761. [26] L.B. Kong, Y.B. Gan, J. Ma, et al., Mullite phase formation and reaction sequences with the presence of pentoxides, J. Alloy. Compd. 351 (2003) 264–272. [27] S.H. Hong, W. Cermignani, G.L. Messing, Anisotropic grain growth in seeded and B2O3-doped diphasic mullite gels, J. Eur. Ceram. Soc. 16 (1996) 133–141. [28] S. Sankaran, Mullitization of diphasic aluminosilicate gels, J. Am. Ceram. Soc. 74 (1991) 2388–2392. [29] Y.Q. Niu, Z.Z. Wang, Y.M. Zhu, et al., Experimental evaluation of additives and K2O –SiO2 –Al2O3, diagrams on high-temperature silicate melt-induced slagging during biomass combustion, Fuel 179 (2016) 52–59. [30] J.C. Huling, G.L. Messing, Hybrid gels for homoepitactic nucleation of mullite, J. Am. Ceram. Soc. 72 (1989) 1725–1729. [31] C.W. Cui, J.L. Huang, L.H. Gao, et al., Fabrication of textured SrBi2Nb2O9 ceramics by the template grain growth method, J. Chin. Ceram. Soc. 35 (2007) 1298–1301. [32] L. Gibson, M. Ashby, Cellular Solids: Structure and Properties, Cambridge University Press, New York, 1997. [33] A. Mukhopadhyay, B.T.T. Chu, M.L.H. Green, et al., Understanding the mechanical reinforcement of uniformly dispersed multiwalled carbon nanotubes in aluminoborosilicate glass ceramic, Acta Mater. 58 (2010) 2685–2697. [34] M. Mazaheri, D. Mari, R. Schaller, G. Bonnefont, et al., Processing of yttria stabilized zirconia reinforced with multi-walled carbon nanotubes with attractive mechanical properties, J. Eur. Ceram. Soc. 31 (2011) 2691–2698.
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A nitride whisker template for growth of mullite in SiC reticulated porous ceramics
MARK
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Xiong Lianga, Yawei Lia, Qinghu Wanga, , Shaobai Sanga, Yibiao Xua, Yuanyuan Chena, Benwen Lib, Christos G. Anezirisc a b c
The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, Wuhan 430081, China Institute of Thermal Engineering, School of Energy and Power Engineering, Dalian University of Technology, Dalian 116024, China Technical University Bergakademie Freiberg, Institute for Ceramic, Glass- and Construction Materials, Agricolastraße 17, 09596 Freiberg, Germany
A R T I C L E I N F O
A BS T RAC T
Keywords: SiC reticulated porous ceramics Nitride whisker template Column-liked mullite Mechanical properties
Silicon carbide reticulated porous ceramics (SiC RPCs) were fabricated by polymer replica technique. The effects of nitride whisker template on the growth of mullite, the strut structure and mechanical properties of SiC RPCs were investigated. Prepolyurethane (PU) open-cell sponge was first coated by SiC slurry consisting of SiC, reactive Al2O3, microsilica and Si powder, then it was nitridized at 1400 °C in a flowing N2 atmosphere to prepare SiC preforms. Subsequently, these preforms were treated by vacuum infiltration of alumina slurry and fired at 1450 °C in air. The results showed that Si2N2O whiskers grew on the surface and in the matrix of SiC preforms after nitridation. The diameter of struts in SiC RPCs increased after vacuum infiltration process because alumina slurry was easily adhered by the surface nitride whiskers. In addition, such whiskers inside the strut of SiC preforms acted as the template to promote the growth of column-liked mullite in SiC RPCs. The mechanical properties and thermal shock resistance of SiC RPCs were greatly improved due to the special interfacial characteristics of multi-layered struts as well as better interlocked column-liked mullite in SiC skeleton.
1. Introduction Porous media combustion (PMC) has unique characteristics, such as high radiant output, low NOx and CO emissions, high flame speed and power density, etc [1–3]. It has been applied in internal combustion engines, low calorific gas burners, volatile organic compound oxidizers and radiation heaters [4–6]. Reticulated porous ceramics (RPCs) with the open and three-dimensional network structure are considered as the most promising high-temperature components for PMC due to their high permeability, thermal stability and resistance to chemical corrosion [7]. However, during service process, the scour of flue gas and large thermal stress from the sharp temperature gradients made porous burners easily damaged, thus the mechanical properties and thermal shock resistance of RPCs should be taken into account [8,9]. In particular, silicon carbide RPCs are considered as the candidate material for PMC because of their low thermal expansion coefficient, high thermal conductivity and strength [10]. The polymer sponge replica technique was the most common method for producing RPCs [11]. However, the hollow struts with triangular voids and longitudinal cracks would generate when the
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Corresponding author. E-mail address: [email protected] (Q. Wang).
http://dx.doi.org/10.1016/j.ceramint.2017.05.169 Received 27 March 2017; Received in revised form 22 May 2017; Accepted 23 May 2017 Available online 25 May 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
sponge was burnt out during firing process. These defects led to the poor mechanical properties of RPCs and limited their structural application [12–14]. Besides, it was difficult to fabricate SiC RPCs with high strength using traditional sintering methods. In general, SiC ceramics were treated via gas pressure sintering at high temperature. This method was unsuitable for fabrication of SiC RPCs, because the fast heating rate caused lots of cracks during the decomposition of polymer sponge [15]. In order to strengthen SiC RPCs, some approaches were proposed, involving in thickening the ceramic skeleton and optimizing the microstructure of strut. The polymeric template was treated by acid or alkali to roughen the sponge surface [16], and the rheological behavior of SiC slurry was adjusted to increase the weight of coated slurry [17]. Also, recoating technique was employed to eliminate the surface flaws and thicken the struts of SiC RPCs [15,18]. Furthermore, vacuum infiltration process was an effective method to obtain the dense struts via filling ceramic slurry into triangular voids of struts [19]. In addition, the ceramic fiber was introduced into the skeleton to optimize the microstructure, thereby strengthening RPCs [20]. The in-situ formed needle-liked Si3N4 was certified as whiskerreinforced mechanism in Si3N4 RPCs when the whisker integrated with
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the liquid phase as a whole, which needed extremely high sintering temperature [21,22]. For increasing the mechanical properties of SiC RPCs, the in-situ nitridation of Si powder together with vacuum infiltration process were applied to improve the macro- and microstructure of struts in this work. The green SiC RPCs containing Si powder were intendedly nitridized to produce SiC preforms, where the whiskers formed on the strut surface and in the matrix. As vacuum infiltration of alumina slurry was carried out, the hollow strut was densified and alumina slurry would be adhered by the surface whiskers. After the infiltrated SiC preforms were fired in air, the whisker might be oxidized to active silica and promoted the mullitization reaction. In the experiment, the nitridation behavior of Si powder in SiC preforms was first investigated. The effects of nitride whiskers on the growth of mullite, the evolution of microstructure and mechanical properties of SiC RPCs were also discussed in details. 2. Experimental 2.1. Preparation of SiC RPCs A commercial available SiC powder ( < 45 µm, 98 wt%, China), Si (~ 3 µm, 98.5 wt%), α-Al2O3 (~ 1.2 µm, Henan Special Refractories Co., Ltd., China) and microsilica powder (~ 0.5 µm, 951UL, 0.13 wt% Na2O, 0.05 wt% Fe2O3, Elkem, Norway) were used as starting materials for the fabrication of SiC RPCs. Si powder was used as the precursor of whiskers. Microsilica and α-Al2O3 were added as sintering aids, the Al2O3/SiO2 molar ratio was designed for chemical composition of mullite. Ammonium lignosulfonate (Tianjin institute of fines chemicals, China) and sodium carboxymethyl-cellulose (CMC, Sinopharm Chemical Reagent Co., Ltd, China) were added as binder and thickening agent to improve the rheology of SiC slurry. Polycarboxylate (FS, BASF Group, Germany) and contraspum K 1012 (Zschimmer & Schwarz, Lahnstein, Germany) were used as dispersant agent and antifoam agent, respectively. The SiC slurries with initial solid content of 77.4 wt% were prepared by mixing powder mixture (shown in Table 1) with the above organic additives and deionized water. SiC slurry was mechanically stirred for 1 h at a rotate speed of 300 rpm to obtain a high thixotropic SiC slurry. The polyurethane open-cell sponge (10 pores/inch, F. M Co. Ltd., Germany) with a size of 50 mm × 50 mm × 20 mm was immersed into the as-prepared SiC slurry, followed by passing through preset roller to remove excess slurry. Remaining closed pores were eliminated by blowing compressed air through the sponge structure. After dried at room temperature, the coated sponges were treated under N2 atmosphere (99.999%, mass fraction) at 1400 °C with a holding time of 1 h to produce SiC preforms. These SiC preforms were named as S0, S5 and S10 respectively according to the content of Si powder in SiC slurries. In the stage of vacuum infiltration, alumina slurry with the solid content of 77 wt% was prepared without binder. Deionized water was mixed with FS and K 1012 by stirring for 5 min. Subsequently, α-Al2O3 powder were added into the solution and ball-milled for 3 h. Then SiC preforms were totally immersed into the alumina slurry and a vacuum of 5 × 10−6 bar was applied for 20 min. The as-infiltrated SiC preforms were fired in air at 1450 °C for 3 h. SiC preforms of S0, S5 and S10 with vacuum infiltration of alumina slurry were named as AS0, AS5 and AS10, respectively. Table 1 Compositions of SiC slurries. Specimen
SiC (wt%)
Microsilica (wt%)
α-Al2O3 (wt%)
Si (wt%)
S0 S5 S10
80 80 80
5.84 5.84 5.84
14.16 14.16 14.16
– +5 + 10
Furthermore, in order to ascertain the effect of silica source on the morphology of reaction-sintered mullite, the matrix specimens comprising of α-Al2O3 and different silica sources were prepared by uniaxially dry pressing method at 5 MPa. These silica sources were microsilica, silicon powder and silicon carbide, respectively. The matrix specimen with silicon powder addition was first nitridized at 1400 °C for 1 h, which was named as SN. The matrix containing microsilica and silicon carbide were named as SO and SC, respectively. All the prepared matrix specimens were heated in air at 1450 °C for 3 h. 2.2. Characterization The RPCs density ρRPC was the density of the whole specimen and was calculated with Eq. (1):
ρRPC =
m l ∙b∙h
(1)
Where l was the length, b the width, h the height, and m the mass of the RPCs specimen. The macrostructure of SiC RPCs was characterized by digital camera and the diameter of struts were analyzed by Image-Pro plus software (Media Cybernetics, Inc., Netherlands). Mechanical tests on the fired RPCs were performed by determining the cold compressive strength (CCS) with a universal testing machine (ETM, Wance, China), using a load speed of 0.5 mm/min. Specimens of the size 50 mm × 50 mm × 20 mm were positioned between loading plates (Φ100 × 20 mm). A cardboard of 4 mm thickness was placed between loading plate and specimen to obtain a uniform loading. The thermal shock resistance of SiC RPCs was evaluated by the water-quenching technique. The fired SiC RPCs were heated at 1100 °C in air for 20 min in a furnace, then immersed into the flowing water of 25 °C. After 3 thermal shock cycles, the residual strength of specimen was determined at room temperature via compression test. The residual strength ratio of CCS was calculated by the change in CCS before and after quenching test, i.e. the residual strength ratio of CCS = 100CCSTS/CCS, where CCS and CCSTS were the CCS before and after 3 thermal shock cycles, respectively. The statistical results of physical and mechanical properties were obtained based on a dataset of 6 specimens. The Gibbs free energy in the reaction of Si powder nitridation was calculated by FactSage software. X-ray diffraction (XRD, X' pert Pro, Philips, Netherlands) was conducted to examine the phase compositions of SiC preforms and RPCs. The microstructure of SiC preform after nitridation and mullite morphology in SiC RPCs were observed by scanning electron microscope (SEM, Quanta 400, FEI Company, USA), equipped with energy dispersive X-ray spectroscopy (EDS, Noran 623 M-3SUT, Thermo Electron Corporation, Japan). 3. Results and discussion 3.1. SiC preforms 3.1.1. Phase composition Fig. 1 shows the XRD patterns of SiC preforms nitridized at 1400 °C. In specimen S0, SiC, α-Al2O3, quartz phases were detected, indicating that mullitization process had not occurred and the added microsilica transformed into quartz phase at this temperature. For SiC RPCs with Si addition, quartz phase was not detected whereas Si2N2O formed on the contrary. The diffraction peak intensity of Si2N2O phase increased with the content of Si powder. Furthermore, the addition of Si powder led to the weight change of SiC preforms (Table 2). In specimen S5 and S10, Si powder completely transformed into Si2N2O phase, thus resulting in the weight gain of SiC preforms. However, the experimental value of weight gain was much lower than theoretical ones. It plausibly related to the gaseous phases of Si(g) and SiO(g) diffused out of SiC struts during the nitridation process of silicon [23– 25], which escaped from SiC preforms together with the flowing N2.
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Fig. 3. XRD patterns of SiC RPCs after fired at 1450 °C.
Fig. 1. The XRD patterns of SiC preforms nitridized at 1400 °C. Table 2 Phase composition and weigh gain of specimens after nitridation process. Specimen
S0 S5 S10
Phase composition
Weight gain (%)
SiC-4H, SiC-6H, corundum, quartz SiC-4H, SiC-6H, corundum, Si2N2O SiC-4H, SiC-6H, corundum, Si2N2O
Theoretical
Experimental
0 3.74 7.14
0 0.74 1.65
3.1.2. Microstructure SEM micrographs of SiC preforms are shown in Fig. 2. Unlike specimen S0 where the amorphous aluminosilicate embedded in the
matrix (Fig. 2(a-b)), the nitridation of Si powder in-situ formed whiskers in the strut of SiC preforms. In specimen S5, the strut surface was covered by large number of blanket-liked whiskers (Fig. 2(c)). At higher magnification, these bent whiskers with large length-diameter ratio were interlocked with each other and well distributed in the matrix as well as on the strut surface (Fig. 2(d)). According to XRD and EDS analysis, they were confirmed to be Si2N2O whiskers. With increasing the content of Si powder to 10 wt%, the amount of whiskers increased and the strut surface was almost completely covered by whiskers (Fig. 2(e)). Besides, the whiskers with larger length-diameter ratio existed inside the triangular void of specimen S10 (Fig. 2(f)). It was worth noting that the more content of Si powder addition, the looser strut appeared in SiC preforms.
Fig. 2. SEM micrographs of SiC preforms after nitridation at 1400 °C. (a, b) S0, (c, d) S5 and (e, f) S10.
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Table 3 Relative diffraction peak intensity of mullite/SiC in SiC RPCs. Specimen
AS0
AS5
AS10
Mullite/SiC
16.50%
19.83%
23.90%
3.2. SiC RPCs 3.2.1. Phase compositions XRD patterns of SiC RPCs after vacuum infiltration of alumina slurry and treatment at 1450 °C are given in Fig. 3. For specimens with Si powder addition, the phase composition of the fired SiC RPCs was much different to SiC preforms. The nitridation products of Si2N2O phase disappeared whereas mullite phase formed in specimen AS5 and AS10. The relative diffraction peak intensity ratio of mullite and SiC increased with the content of Si powder in SiC RPCs (Table 3), indicating that nitride whiskers promoted the formation of mullite in SiC RPCs after firing in air. 3.2.2. Microstructure The cross section of the etched struts in SiC RPCs are presented in Fig. 4. The formation of nitride whiskers in SiC preforms remarkably affected the final strut structures of SiC RPCs. With the growth of whiskers on the surface of SiC preforms, more alumina slurries adhered on SiC skeleton, thus leading to the formation of thicker coating in SiC RPCs than specimen AS0 (Fig. 4(a-c)). Furthermore, some needle-liked mullite formed in the matrix of SiC RPCs, while the lumpy mullite grew around SiC particles in specimen AS0 (Fig. 4(d)). For SiC RPCs with Si powder addition, besides the mullite exhibited the similar morphology to AS0, some column-liked mullite with large
length-diameter ratio was observed in AS5 (Fig. 4(e)). It was worth noting that the amount of mullite increased with the content of Si powder in SiC preforms, especially for the column-liked mullite. Meanwhile, the column-liked mullite interlocked with each other and showed its network structure in the matrix of SiC RPCs (Fig. 4(f)). In addition, the significant interface formed between multi-layers, and its thickness increased with the amount of nitride whiskers in SiC preforms. At higher magnification of interface in AS10 (Zone 1 in Fig. 4(c)), a dense transition layer located between SiC skeleton and infiltrated alumina area, which contained large amount of column-liked mullite (Fig. 5(a)). For the outer layer of strut (Zone 2 in Fig. 4(c)), the continuous coating with good interfacial characteristics formed, which was connected with SiC skeleton by column-liked mullite (Fig. 5(b)). 3.2.3. Physical and mechanical properties Table 4 lists the physical properties of SiC RPCs fired at 1450 °C. The bulk density of SiC RPCs containing Si powder was larger than that of specimen AS0. Meanwhile, the weight gain of infiltrated alumina in SiC RPCs increased with the content of Si powder, leading to the bulk density of specimens increased from 0.38 g/cm3 to 0.45 g/cm3. Because the addition of Si powder promoted the formation of nitride whiskers in SiC preforms, more alumina slurry was able to coat on the surface of strut, the diameter of struts increased at the same time. Mechanical properties including CCS and its residual ratio of SiC RPCs before and after 3 thermal shock cycles are still presented in Table 4. It is apparent that nitridation of Si powder in SiC preforms produced the positive influence on the mechanical properties of SiC RPCs. For example, CCS of specimen AS0 was 0.60 MPa and its value increased with the content of Si powder in SiC RPCs. The CCS reached the maximum in specimen AS10 of 0.83 MPa. It was worth mentioning that the SiC RPCs with Si powder nitridation were less sensitive to
Fig. 4. Cross section struts of SiC RPCs after etched by hydrofluoric acid. (a, d) AS0, (b, e) AS5 and (c, f) AS10.
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Fig. 5. The interfacial characteristic in AS10. (a) and (b) were the higher magnification micrographs of zone 1 and 2 in Fig. 4(c). Table 4 Physical and mechanical properties of SiC RPCs. Specimens
Weight gain of infiltrated alumina g
Bulk density ρb g/cm3
Diameter of struts mm
CCS MPa
CCSTS MPa
CCSTS/CCS
AS0 AS5 AS10
7.13 ± 0.83 8.18 ± 1.29 8.90 ± 1.03
0.38 ± 0.02 0.42 ± 0.04 0.45 ± 0.03
0.56 ± 0.08 0.76 ± 0.14 0.89 ± 0.13
0.60 ± 0.14 0.73 ± 0.11 0.83 ± 0.08
0.40 ± 0.02 0.55 ± 0.03 0.65 ± 0.13
66.7% 75.1% 78.3%
Fig. 6. SEM micrographs of etched matrix SO (a), SS (b) and SC (c) after fired at 1450 °C.
thermal shock than AS0. The residual strength and its residual ratio persistently increased with the content of Si powder, especially specimen AS10 exhibited the largest values of 0.65 MPa and 78.3%, respectively.
3.3. Discussion The nitridation process of SiC preform was controlled under flowing nitrogen of 99.999 wt% in this experiment, therefore the 0.001 wt% oxygen was taken into account in thermodynamic calculation. In SiC preform, Si(s) preferred to react with O2(g) to form SiO(g) through Reaction (2) because of its negative Gibbs free energy. In the flowing nitrogen, Si3N4(s) was easily formed by Reaction (3) [23]. However, in the case that SiO2(s) existed in SiC preforms, Reaction (4) would occur and resulted in the formation of Si2N2O(s). Besides, The O2(g) was considered as the role of catalytic agent from Reaction (2) and (3), and promoted the formation of SiO(g) in the system. Therefore, the growth of Si2N2O whiskers could also be processed via the V-V reaction between SiO(g) and N2(g) when SiO(g) partial pressure was high enough. Meanwhile, the SiO(g) or Si(g) would escape from SiC preforms with the flowing nitrogen at elevated temperature, thus leading to little weight gain in specimen S5 and S10.
Si(s) + 1/2O2(g) = SiO(g) ΔGθ= −109,391 −78.71Τ (298 Κ < Τ < 1685 K) (2) 3SiO(g) + 2N2(g) = Si3N4(s) + 3/2O2(g) ΔGθ= −408,343 + 560.89Τ (298 Κ < Τ < 1685 K) (3) Si3N4(s) + SiO2(s) = 2Si2N2O(s) ΔGθ= −73,137 − 4.76Τ (298 Κ < Τ < 1685 K) (4) 2SiO(g) + N2(g) = Si2N2O(s) + 1/2O2(g) ΔGθ= −688,096 + 429.05Τ (298 Κ < Τ < 1685 K) (5) As the vacuum infiltration of alumina slurry was carried out, the porous and loose skeleton of SiC RPCs was filled up by alumina (Fig. 4). After the infiltrated SiC RPCs was fired at 1450 °C in air, the mullite with three kinds of morphologies formed in the matrix of SiC skeleton. In order to reveal this phenomenon, the mullitization reaction between active alumina and different silica source was analyzed, which was shown in Fig. 6. It was clearly seen that mullite with needle-shape and large grain size formed in the matrix specimen comprising of microsilica and alumina, while the lumpy mullite grew around SiC particles when the silica source was silicon carbide (Fig. 6(a-b)). In addition, nitridation-derived whiskers were beneficial for the growth of mullite
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Fig. 7. Schematic diagram of structure evolution of SiC RPC with Si powder addition. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article).
grain and facilitated the formation of column-liked mullite (Fig. 6(c)). For the growth mechanism of mullite, it was greatly accepted that mullite formation was controlled by “dissolution-precipitation” in reaction sintering containing SiO2 and Al2O3 [26,27]. The SiO2-rich liquid phase was first generated at elevated temperature, then the Al2O3 powders were dissolved in it. When the critical concentration of Al2O3 in SiO2-rich liquid phase was achieved, mullite began to nucleate and grow [28]. The dissolution velocity of Al2O3 into the SiO2-rich liquid phase as well as the template of silica source determined the crystal nucleation and mullite morphology. In the matrix consisting of microsilica and alumina, the SiO2-rich liquid phase with low viscosity easily formed due to the existence of impurity in microsilica [29]. This accelerated atomic mobility and increased the dissolution velocity of α-Al2O3, thereby a large number of mullite crystals started to nucleate when the Al2O3 concentration reached a critical value in SiO2-rich liquid. Furthermore, the liquid reaction medium with its low viscosity contributed to rapid chemical reaction and the growth of mullite, thus the large needle-liked mullite formed though the silica source was spherical microsilica. As the silica source was silicon carbide, the SiO2 film oxidized from silicon carbide easily formed and wrapped on the surface of SiC particles. Meanwhile, α-Al2O3 dissolved into SiO2 film, and mullite nucleation as well as crystal growth occurred. After the mullite crystal formed around the surface of SiC particle layer by layer, the growth of mullite was determined by the diffusion of Al2O3 in mullite layer. The low diffusion velocity of Al2O3 slowed down the mullitization process, only a small amount of initially formed mullite crystals grew to become large crystals, thereby the lumpy mullite formed around SiC particles. Compared to silicon carbide, the nitride whisker with its large specific surface area exhibited high activity and would be oxidized to SiO2 liquid at lower temperature. Besides, the highly reactive and fine SiO2 derived by the oxidation of nitride whiskers would accelerate mullitization reaction and benefited to the formation of mullite crystal [30]. In addition, the nitride whisker played the role of template and induced the mullite to grow along the whisker direction during the mullitization process, resulting in formation of mullite with column-liked morphology [31]. From XRD and SEM results of SiC RPCs with Si powder addition, a schematic diagram could be drawn to reveal the phase composition and structure evolution of struts at elevated temperature (Fig. 7). The green SiC RPCs were prepared after the PU template was coated by SiC slurry via polymer sponge replica technique (Fig. 7(a)). After being nitridized at 1400 °C, the nitridation of Si powder was performed and amounts of nitride whiskers formed both in the matrix of SiC skeleton and on the surface of strut, resulting in a loose skeleton and rough surface (Fig. 7(b)). As the vacuum infiltration process was carried out, the nitride whiskers were wrapped by alumina slurry and the hollow skeleton was densified. Furthermore, the strut with rough surface was beneficial to the adhesion of alumina slurry (Fig. 7(c)), thus thickening the strut and strengthening the SiC RPCs [32]. After the
treatment of infiltrated SiC RPCs at elevated temperature, these nitride whiskers were oxidized to active silica and acted as the template in mullitization process, which promoted the growth of column-liked mullite (Fig. 7(d)). The larger amount and well interlocked columnliked mullite endowed SiC RPCs with better mechanical properties and thermal shock resistance [33,34]. Besides, the multi-layered strut with good interface characteristics also strengthened SiC RPCs. 4. Conclusions The following conclusions can be drawn by studying the influence of Si powder nitridation on the structure and mechanical properties of SiC RPCs. (1) Si powder nitridation resulted in the formation of Si2N2O whiskers both in the matrix of SiC skeleton and on the surface of SiC preforms. (2) The strut of infiltrated SiC RPCs could be thickened after vacuum infiltration process because alumina slurry was easily adhered by the surface nitride whiskers. In addition, in-situ formed whiskers in SiC preforms acted as the template to promote the growth of column-liked mullite in SiC RPCs. (3) The mechanical properties and thermal shock resistance of SiC RPCs with Si powder nitridation were greatly improved compared to non-nitridation ones, which was attributed to the special interface characteristics of multi-layered struts and better interlocked column-liked mullite in SiC skeleton. Acknowledgment The authors thank the subprojects A01 and C03 of German Research Foundation (DFG) for supporting these investigations in terms of the Collaborative Research Centre 920, Multi-Functional Filter s for Metal Melt Filtration-A Contributions towards Zero Defect Materials. References
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