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High-strength thermal insulating mullite nanofibrous porous ceramics
Journal Pre-proof High-strength thermal insulating mullite nanofibrous porous ceramics Ying Zhang, Yongjun Wu, Xukun Yang, Dinghe Li, Xueying Zhang, Xue Dong, Xinghe Yao, Jiachen Liu, Anran Guo
PII:
S0955-2219(20)30012-1
DOI:
https://doi.org/10.1016/j.jeurceramsoc.2020.01.011
Reference:
JECS 12988
To appear in:
Journal of the European Ceramic Society
Received Date:
15 September 2019
Revised Date:
5 January 2020
Accepted Date:
6 January 2020
Please cite this article as: Zhang Y, Wu Y, Yang X, Li D, Zhang X, Dong X, Yao X, Liu J, Guo A, High-strength thermal insulating mullite nanofibrous porous ceramics, Journal of the European Ceramic Society (2020), doi: https://doi.org/10.1016/j.jeurceramsoc.2020.01.011
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High-strength thermal insulating mullite nanofibrous porous ceramics Ying Zhanga, Yongjun Wub, Xukun Yangb, Dinghe Lic, Xueying Zhanga, Xue Dongc⁎, Xinghe Yaob, Jiachen Liua, Anran Guoa⁎ a
School of Materials Science and Engineering, Key Lab of Advanced Ceramics and
Machining Technology of Ministry of Education, Tianjin University, Tianjin 300072,
b
Aerospace Business Department, China Academy of Launch Vehicle Technology, Beijing 100076, China
College of Aeronautical Engineering, Civil Aviation University of China, Tianjin
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c
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China
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300300, China
Corresponding author:
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Abstract
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Anran Guo, Email: [email protected]; Xue Dong, Email: [email protected]
Mullite fibrous porous ceramics is one of the most commonly used high
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temperature insulation materials. However, how to improve the strength of the mullite fibrous porous ceramics dramatically under the premise of no sacrificing the low
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sample density has always been a difficult scientific problem. In this study, the strategy of using mullite nanofibers to replace the mullite micron-fibers was proposed to fabricate the mullite nanofibrous porous ceramics by the gel-casting method. Results show that mullite nanofibrous porous ceramics present a much higher compressive strength (0.837MPa) than that of mullite micron-fibrous porous ceramics 1
(0.515MPa) even when the density of the mullite nanofibrous porous ceramics (0.202 g/cm3) is only around three quarters of that of the mullite micron-fibrous porous ceramics (0.266 g/cm3). The obtained materials that present the best combination of mechanical and thermal properties can be regarded as potential high-temperature thermal insulators in various thermal protection systems. Key words: Mullite fiber; Porous ceramics; Thermal insulation materials
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1. Introduction Mullite fibrous porous ceramics is one of the most commonly used high
temperature (>1300oC) insulation materials in various thermal protection systems
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owing to their low density, high porosity, excellent thermal shock resistance and low thermal conductivity [1-4]. The mullite fibrous porous ceramics are mainly composed
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of the interlocked micron-sized mullite fibers bound by a high-temperature binder [5].
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This three-dimensional porous skeleton structure not only endows the material with the characteristics of low density and low thermal conductivity, but also ensures that the three-dimensional pores existing in the material don’t collapse at high
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temperatures [6-9].
Density, compressive strength and thermal conductivity are the most important
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properties of the mullite fibrous porous ceramics (MFCs). Generally, the thermal
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conductivity of the MFCs is mainly determined by the solid-phase heat conduction and the gas-phase heat conduction due to its highly porous structure. Owing to the fact that the thermal conductivity of the air is much lower than that of the mullite fiber as the solid phase, the most effective way to reduce the thermal conductivity of the MFCs is to reduce the content of the solid phase, namely to decrease the sample density [10]. However, the decrease of density will inevitably lead to a decrease of the 2
material strength, which in turn limits the further application of the mullite fibrous ceramics. Therefore, how to improve the strength of the mullite fibrous porous ceramics under the premise of no sacrificing the low sample density has always been a difficult scientific problem. Many researchers have tried to change the microstructure of the MFCs, including the distribution of the binder and the arrangement of the mullite fiber, in order to increase the sample strength [11, 12]. However, results show that these methods only increase the sample strength in a
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limited range because they did not change the fundamental unite (mullite fibers) of the mullite fibrous ceramics.
The fracture of mullite fibrous porous ceramics is mainly caused by the fracture of
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the mullite fiber or the binder [13]. Therefore, one effective method to increase the sample strength is increasing the strength of mullite fibers and binders. Based on the
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above-mentioned, it can be assumed that the properties of MFCs can be greatly
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improved if the size of the mullite fiber as the matrix of MFCs reduces to the nanoscale. Firstly, with the reduction of the mullite fiber diameter from the micron scale to the nanometer scale, the strength of the fibers will substantially increase due
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to the decrease of the internal defects existing in mullite fibers [14]. Secondly, the size of the pores lapped by the fibers would reduce dramatically due to the decrease of
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fiber diameter, bringing about the increase of the pore wall content and the bonding
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points, finally resulting in the increase of the sample strength. The forming methods of mullite micron-fibrous porous ceramics include vacuum
filtration forming method and pressure assisted molding method, and both methods will exert a longitudinal pressure on the materials [15-17]. However, as the fiber diameter reduced to the nanometer scale, the longitudinal pressure would destroy of the 3D skeleton structure during the forming process, and consequently lead to the 3
increase of the sample density. Therefore, the best forming method to fabricate the mullite nanofibrous porous ceramics should be an in-situ forming method to prevent the structural collapse. Freeze casting which could obtain samples with small linear shrinkage and high porosity was proposed to prepare nanofibrous porous ceramics by many researchers [18, 19]. However, the mullite nanofibrous porous ceramics prepared by the freeze casting usually showed a multilevel pore structure with small pores formed by the overlapped nanofibers and large pores caused by the sublimation
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of the ice crystal, which was not beneficial to the sample strength [20]. Different from the freeze casting method, the gel casting method as another quite common in-situ forming method is able to fabricate samples with uniform microstructure and high
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strength [21, 22]. However, the water-based gel casting often leads to the break of samples with ultralow solid content during the drying process. Considering that the
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tertbutyl alcohol (TBA) has higher saturated vapor pressure and lower surface tension
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relative to water, and can volatilize rapidly at 50oC, the TBA-based gel-casting is a promising method for the fabrication of mullite nanofibrous porous ceramics [23, 24]. In this work, in order to further increase the strength of the mullite fibrous
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ceramics, the strategy of using mullite nanofibers to replace the mullite micron-fibers was propose to fabricate the mullite nanofibrous porous ceramics. The fabrication
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process of mullite nanofibrous porous ceramics was analyzed. Moreover, the effects
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of the mullite fiber content and the fiber aspect ratio on the sample microstructures and properties were investigated. 2. Experimental 2.1 Materials The
precursor
of
mullite
nanofibers
was
prepared
using
polyhydromethylsiloxane (Alfa Aesar Chemistry Co., Ltd., China) as the silicon 4
source, aluminum-tri-sec-butoxide (Sigma-Aldrich Trading Co., Ltd., China) as the aluminum source, polyvinylpyrrolidone (PVP-K90, Mw=1300000, Sigma-Aldrich Trading Co., Ltd., China) as the spinning aid, respectively. The silica sol which was used as the high temperature binder for 3D skeleton structure was prepared with tetraethyl orthosilicate (TEOS), deionized water and ethanol (ETOH). Polyacrylamide as the low-temperature binder were prepared using tert-butyl alcohol (TBA), acrylamide
(AM),
N,
N-methylenebisacrylamide
(MBAM),
ammonium
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persulfate(APS) and N, N, N, N-tetramethylethylenediamine (TEMED). All the above chemical reagents for binder were analytical grade from Aladdin Reagent Co., Ltd.,
China. Mullite micron-fiber porous ceramics (MMFCs) were purchased from
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Zhejiang Hongda Crystal Fiber Co., Ltd., China. The density of the MMFCs is 0.266
sample open porosity is 88.92%.
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2.2 Preparation
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g/cm3, the sample thermal conductivity of the MMFCs is 0.1012 Wm-1K-1 and the
Firstly, the diphasic mullite precursor with the alumina/silica molar ratio of 3:1 was prepared by the electrospinning method. The detailed fabrication process was
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described in our previous works [20]. Secondly, the precursor fibers were calcinated at 800oC with a heating rate of 2 oC/min. The average fiber diameter was about 476
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nm as shown in Fig. S1 in the supplementary materials. Then amorphous fibers as
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shown in Fig. 1a were dispersed into individual fibers using a high-speed homogenizer. The dispersion was dried at 120°C and the obtained fibers were shown in Fig. 1b. After that, 14.5 g of the AM and 0.5 g of the MBAM were added into 85 g of TBA to obtain the 100 g of premix, respectively. Then 1.5 wt%, 2.0 wt% and 2.5 wt% of the individual amorphous fibers and 0.25 g of the silica sols (40 wt%) were added into 5ml premix solution, respectively. An aqueous solution of APS (40 wt%) and an 5
aqueous solution of TEMED (4 wt%) were used as the initiator and catalyst, respectively. Then 0.25 ml of the APS and 0.4 ml of the TEMED were dissolved in the mixture sequentially. A gel-casting process would occur after 30 minutes and the green body after demolding was dried at 60oC for 12h. Finally, the samples were sintered and the heating system was 1 oC/min from room temperature to 700oC and 5 o
C/min from 700oC to 1400oC. In order to prevent the cracking in the binder removal
process caused by the rapid decomposition of polymer, it must be noted that the
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samples were kept heating at 390°C and 540°C for one hour. Then the mullite nanofibrous porous ceramics (MNFCs) were obtained after calcination at 1400oC for
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two hours.
Fig.1. SEM images of (a) amorphous fibers after sintering at 800oC, (b) the individual nanofibers
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after mechanical dispersion and (c and d) the green body.
2.3 Characterization The sample microstructure was observed via the scanning electron microscope
(SEM, S-4800, Hitachi Ltd., Japan). The phase transition of samples was investigated by the X-ray Diffraction (XRD, Rigaku D/Max-2500, Japan). The thermogravimetry of the green body was carried out from room temperature to 1400°C in air by a 6
thermal analyzer (TG-DSC, Netzsch Sta 449C, Germany). The linear shrinkage and the bulk density of the sample were calculated based on the length and weight of the samples. The porosity of the sample was measured by the water-immersion technique using the Archimedes method. The thermal conductivity of the ceramics was investigated by the thermal conductivity measuring instrument with a transient hot disk method (Hot Disk TPS 2500S, Sweden). The compressive strength were tested at
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room temperature, 400°C, 600°C, 800°C and 1000°C with a loading speed of 0.02 mm/min by an electronic universal testing machine (CMT4303, Meister Industrial Systems, China). The pore size distribution of the ceramics was investigated using a mercury porosimetry (AutoPore IV 9510, America).
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3. Results and discussions
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3.1 Fabrication of mullite nanofibrous porous ceramics
It is a challenge to improve the strength of the MFCs dramatically without
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sacrificing its low sample density. In this paper, the strategy of using mullite nanofibers to replace the micron-fibers was proposed to fabricate the mullite
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nanofibrous porous ceramics with low density and high strength. The fabrication flow chart of the mullite nanofibrous porous ceramics is shown in Fig. 2a. The matrix in
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MNFCs was mullite nanofibers with the alumina/silica molar ratio of 3:1 prepared by the electrospinning method. In order to ensure the complete decomposition of organic
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components in precursor fibers and get the optimum calcination temperature, the crosslinked fibers were tested by a thermal analyzer from room temperature to 1400°C. As shown in Fig. S2 in the supplementary materials, the weight loss of fibers between 32°C to 800°C was about 73.42%, while the weight loss of fibers in the range of 800-1400°C was only 0.83%. Therefore, the fiber calcination temperature 7
was selected at 800°C, which could both remove the organic component and meanwhile maintain the fiber's flexibility. After calcination, the non-woven fiber mat was dispersed into individual fibers by a high-speed homogenizer. By controlling the mixing rate, fibers with different average lengths (~20 μm and ~7 μm) could be obtained. As an in-situ forming method, the TBA-based gel casting is selected to fabricate the mullite nanofibrous porous ceramics with an ultralow solid loading. The advantages of this method include that the gel-casting time can be controlled by
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changing the addition content of the initiator and catalyst, and the sample density can be easily controlled by changing the solid content. As shown in Fig.1c and Fig. 1d, the green body exhibited an obvious uniform pore structure formed by the overlapped
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mullite fibers. In the green body, fibers are coated by the polyacrylamide as the
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low-temperature binder and the silica gel as the high-temperature binder.
Fig.2. (a) Fabrication flow chart of the MNFCs; (b, c and d) SEM images of MNCFs; Inset of Figure 2b: an optical photograph of MNFCs
With the increase of the sintering temperature, the polyacrylamide decomposed 8
completely and the silica sols transformed into the silica acting as the high temperature binders binding the overlapped mullite fibers. SEM images in Fig. 2b, 2c and 2d demonstrated the microstructure of the MNFCs after sintered at 1400℃. As shown in Fig.2b, the pores were mainly formed by the overlapped mullite nanofibers and the enlarged image of the fiber junction (Fig. 2c and 2d) shows that the high-temperature silica binder firmly bound the connected mullite nanofibers. The inset of Fig. 2b shows the photograph of the mullite nanofibrous porous ceramics after
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sintered at 1400°C.
Fig.3. TG and DSC curves of the green body from room temperature to 1400oC.
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In order to determine the chemical changes of the green body after sintered at
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different temperatures, the thermogravimetric of the green body from room temperature to 1400°C in the air and the X-ray diffraction analyses of samples after sintered at different temperatures were carried out. Fig. 3 shows that the weight loss of the sample mainly consisted of three stages. The first stage is the weight loss about 13.2% between 25°C to 200°C, which was mainly caused by the removal of the free water. In the second weightlessness stage, a dramatical mass loss about 63.4% 9
occurred between 200°C and 656°C due to the oxidative degradation of organic component (acrylamide and N, N-methylenebisacrylamide) at high temperature. Two obvious exothermic peaks appeared at 390°C and 540°C respectively on the DSC curve, and the weight loss rate of green body around the two peaks were quite fast. Therefore, in order to prevent the cracking caused by the rapid oxidative degradation of polymer, the sintering process during this period should be as slow as possible. In the third stage (656°C~1400°C), the weightlessness was about 4.5% and the
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exothermic peak at 998°C was caused by the formation of mullite phase. As the temperature continued to rise, the weight of MNFCs remained stable which indicates
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that the MNFCs exhibited an excellent thermal stability.
Fig. 4. XRD patterns of the samples after sintering at different temperatures.
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Fig. 4 shows the X-ray diffraction patterns of samples after sintering at different
temperatures. The broad backgrounds at around 23° indicate the existence of amorphous silica in the green body and the sample at 800℃. When the sintering temperature was above 1000°C, the characteristic peaks of γ-alumina (PDF#75-0921) [20] and mullite phase (PDF#15-0776) could be clearly observed, indicating the 10
aluminum source firstly transformed into γ-alumina phase and then some metastable alumina reacted with amorphous silica forming the mullite phase. As the sintering temperature rose to 1200°C, the SiO2 peak at 23° (PDF#29-0085) was observed, and the diffraction peaks of γ-alumina phase in the fibers disappeared, suggesting the mullitization process basically completed. When the sintering temperature continued to rise to 1400°C, the peak intensity is further strengthened, indicating the enhanced crystallinity of mullite [25]. Moreover, the excessive silicon oxide attaches to the fiber
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joints as the binder to increase the sample strength. Ultimately, only the mullite phase and the high-temperature binder (SiO2 phase) exist in the MNFCs.
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3.2 Effect of fiber content
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Fig. 5. SEM images of MNFCs fabricated with the addition of different fiber contents. (a and d) 1.5 wt% fiber content, (b and e) 2.0 wt% fiber content, and (c and f) 2.5 wt% fiber content.
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Solid loading is a basic parameter to control the sample density in the TBA-based gel-casting system. In this experiment, the pore structure and the densities of MNFCs were adjusted by changing the fiber content (1.5 wt%, 2 wt% and 2.5 wt%), further leading to the variation on the mechanical strength and thermal conductivity of the MNFCs. All the MNFCs with different fiber contents were prepared with the long nanofibers as the starting materials. Fig 5 shows the SEM images of the MNFCs with 11
the addition of different fiber contents. For the MNFCs with 1.5 wt% fiber content (Fig. 5a and 5d), mullite nanofibers in the sample overlapped with each other forming a 3D network structure and the pore size was about 5~20 μm. Due to the low fiber content, the crossing point between the fibers mainly existed at the end of each fiber. Therefore, the sample showed a loose 3D structure and the pore size was close to the length of the fiber. As the fiber content increased, the 3D skeleton structure of the MNFCs became compact and the pore size of the sample with 2.5 wt% fiber content
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reduced to 1~8μm as shown in Fig. 5c and 5f.
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Fig. 6. (a) The compressive strength, linear shrinkage density, (b) open porosity and thermal conductivity of MNFCs fabricated with the addition of different fiber contents.
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As shown in Fig. 6a, the bulk density of MNFCs increased from 0.061 g/cm3 to 0.108 g/cm3 as the fiber content increased from 1.5 wt% to 2.5 wt%, resulting in the
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increase of mechanical strength from 0.114 MPa to 0.153 MPa. As shown in Fig.6b, the porosities decreased from 97.9% to 96.4% as the fiber content increased, and all
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the products exhibited a relatively high porosity due to the abundant pores formed by the unique nanofibrous skeleton. The increase of fiber content resulted in the increase of the material density, further leading to the increase of thermal conductivity (from 0.0597 Wm-1K-1 to 0.0715 Wm-1K-1). 3.3 Effect of the fiber aspect ratio 12
As shown in Fig. 5, the MNFCs mainly consisted of the randomly arranged mullite nanofibers, and therefore the aspect ratio of the fibers should be one of the most important factors influencing the microstructure and properties of the MNFCs. In this experiment, as shown in Fig.7a and Fig.7c, two kinds of mullite nanofibers with different fiber lengths (7 μm and 20 μm) were used as the raw materials to prepare the mullite nanofibrous porous ceramics with the fiber content of 1.5wt%. The samples using the long fibers and the short fibers as the starting materials were
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expressed as L-MNFCs and S-MNFCs, respectively. In addition, the properties of the purchased mullite micron-fibrous (10 μm in diameter) porous ceramics (MMFCs)
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were also investigated to show the effect of fiber diameter.
Fig. 7. SEM images of (a) short mullite nanofibers, (b) S-MNFCs, (c) long mullite nanofibers, (d) L-MNFCs and (e and f) MMFCs. 13
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Fig. 8. Pore size distribution of different mullite fibrous porous ceramics.
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The morphologies and the pore size distribution of different mullite fibrous porous
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ceramics (S-MNFCs, L-MNFCs and MMFCs) are shown in Fig. 7 and Fig. 8. The pore size distribution curve of MNFCs shows an obvious bimodal pore size
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distribution (pore I in 1-5 μm and pore II in 5-15 μm). As shown in Fig. 7b, the pore I was the fundamental porous structure of the nanofibrous porous ceramics, and was
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mainly formed by the overlapped neighboring nanofibers. These relatively densely distributed fibers compose the pore walls of pore II. Different from the composition of
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the pore I, the contact points of the overlapped mullite fibers in the pore II tend to be at the end of nanofibers. Therefore, as the length of nanofiber increases, the size of
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pore II in Fig.7d is larger than that in Fig. 7b. Meanwhile, this information could also be confirmed in the curve of pore size distribution of different mullite fibrous porous ceramics. Compared with the pore size distribution of S-MNFCs (red curve in Fig.8), the content of pore I in L-MNFCs (blue curve in Fig. 8) decrease and that of pore II increase with the increase of nanofiber length. Furthermore, the average pore size of two samples increased from 3.185 μm to 7.991 μm. 14
Fig. 9. (a) The compressive strength, the thermal conductivity and bulk density of samples
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fabricated with different fibers. (b) Compressive strength of the S-MNFCs at different test temperature.
When the fiber diameter increased from nanometer scale to micron scale, the
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porosity of MMFCs decreased greatly (88.92%), and the crossed mullite fibers
formed 45-120 μm macropore (pore III). Compared the properties of S-MNFCs and
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MMFCs (Fig.9a), it can be found that although the density of S-MNFCs (0.202 g/cm3)
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was still a little lower than that of MMFCs (0.266 g/cm3), the compressive strength of S-MNFCs (0.837MPa) was much higher than that of MMFCs (0.515MPa), which was mainly attributed to the difference of the fiber diameter. As the fiber diameter reduced
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from the micron scale to the nanometer scale, the strength of the individual nanofiber substantially increased due to the decrease of the internal defects in the fibers, and the
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amount of pore wall and bonding points also increased obviously because of the
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decrease of the pore size, finally resulting in the increase of the material strength. Taking into account the potential application of the mullite nanofibrous porous ceramics, the compressive strength at different temperature are shown in Fig.9b. It can be seen that the compressive strength of the S-MNFCs decreased with the increase of the test temperature. Moreover, when the test temperature increased to 1000°C, the compressive strength of the material was still quite high, about 0.726 15
MPa. In addition, the thermal conductivity of the S-MNFCs (0.1001 Wm-1K-1) is a little lower than that of the MMFCs (0.1012 Wm-1K-1) due to the low density and small pore structure of nanofibrous ceramics. 4. Conclusions In summary, mullite nanofibrous porous ceramics were successfully prepared by the TBA-based gel casting, and the effects of fiber content and fiber aspect ratio on the microstructure and properties of the samples were studied. The strength of the
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MNFCs increased from 0.114 MPa to 0.158 MPa with the solid content of long nanofibers increasing from 1.5wt% to 2.5wt%. In addition, compared with the
S-MNFCs, the L-MNFCs showed an ultralow density (0.061 g/cm3) and low thermal
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conductivity (0.0597 Wm-1K-1). The S-MNFCs showed a much higher compressive strength (0.837MPa) than that of MMFCs (0.515MPa) although the density of
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MMFCs was still a little higher than that of S-MNFCs, which was mainly attributed to
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the difference of the fiber diameter. All the results prove that using mullite nanofibers to replace the mullite micron-fibers to prepare the mullite fibrous porous ceramics could effectively improve the sample strength under the premise of no sacrificing the
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low sample density. Based on the obtained results, the mullite nanofibrous porous ceramics provides a new insight into the development of high efficient thermal
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insulation materials.
Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgement This work is supported by the National Natural Science Foundation of China (Project No. 51502196 and 51872194). References [1] P. Soltani, M.S. Johari, M. Zarrebini, Eff ect of 3D fiber orientation on permeability of realistic fibrous porous networks, Powder Technol. 254 (2014)
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44-56. [2] A.K. Pradhan, D. Das, R. Chattopadhyay, S.N. Singh, Eff ect of 3D fiber orientation distribution on transverse air permeability of fibrous porous media,
-p
Powder Technol. 221 (2012) 101-104.
[3] T. Stylianopoulos, A. Yeckel, J.J. Derby, X.J. Luo, M.S. Shephard, E.A. Sander,
re
V.H. Barocas, Permeability calculations in three-dimensional isotropic and
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oriented fiber networks, Phys. Fluids. 20 (2008) 123601. [4] R.B. Zhang, X.B. Hou, C.S. Ye, B.L. Wang, D.N. Fang, Fabrication and properties
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of fibrous porous mullite–zirconia fiber networks with a quasi-layered structure, J. Eur. Ceram. Soc. 36 (2016) 3539-3544.
ur
[5] J. Zhang, X. Dong, F. Hou, H.Y. Du, J.C. Liu, A.R. Guo, Eff ect of mullite fiber
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content on the microstructure and properties of porous mullite fiber/silica composite, Ceram. Int. 42 (2016) 6520-6524.
[6] D. Rasky, F.S. Milos, T.H. Squire, Thermal protection system materials and costs for future reusable launch vehicles. J Spacecr Rockets. 38 (2001) 294-296. [7] S. Liu, J.C. Liu, H.Y. Du, F. Hou, A.R. Guo, Microstructure of mullite fiber-based 17
hierarchical structures adjusted by Al/Si mole ratio of the raw material powders, Ceram. Int. 40 (2014) 11405-11410. [8] H. Schneider, J. Schreuer, B. Hildmann, Structure and properties of mullite—A review, J. Eur. Ceram. Soc. 28 (2008) 329-344. [9] R. Baetens, B.P. Jelle, A. Gustavsen, Aerogel insulation for building applications: A state-of-the-art review, Energ. Buildings. 43 (2011) 761-769.
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[10] J. König, V. Nemanič, M. Žumer, R.R. Petersen, M.B. Østergaard, Y.Z. Yue, D. Suvorov, Evaluation of the contributions to the effective thermal conductivity of an open-porous-type foamed glass, Constr. Build. Mater. 214 (2019) 337-343.
-p
[11] W.J. Zang, T. Jia, X. Dong, J.C. Liu, H.Y. Du, F. Hou, A.R. Guo, Preparation of
re
homogeneous mullite-based fibrous ceramics by starch consolidation, J. Am. Ceram. Soc. 101 (2018) 3138–3147.
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[12] T. Jia, H. Chen, X. Dong, W.J. Zang, A.R. Guo, J.C. Liu, Preparation of
na
homogeneous mullite fibrous porous ceramics consolidated by propylene oxide, Ceram. Int. 45 (2019) 2474-2482
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[13] X. Dong, G.F. Sui, Z.Q. Yun, M.C. Wang, A.R. Guo, J. Zhang, J.C. Liu, Effect of temperature on the mechanical behavior of mullite fibrous ceramics with a 3D
Jo
skeleton structure prepared by molding method, Mater. Des. 90 (2016) 942–948.
[14] Z.X. Chen, Z. Zhang, C.-C. Tsai, K. Kornev, I. Luzinov, M.H. Fang, F. Peng, Electrospun mullite fibers from the sol–gel precursor, J. Sol-Gel Sci. Technol. 74 (2015) 208-219. [15] C. Soares, N. Padoin, D. Muller, D. Hotza, C.R. Rambo, Evaluation of resistances 18
to fluid flow in fibrous ceramic medium, Appl. Math. Model. 39 (2015) 7197-7210. [16] B. Zhang, et al., Ultra-low cost porous mullite ceramics with excellent dielectric properties and low thermal conductivity fabricated from kaolin for radome applications, Ceram. Int. 45 (2019) 18865-18870. [17] R.B. Zhang, C.S. Ye, X.B. Hou, S.H. Li, B.L. Wang, Microstructure and
ro of
properties of lightweight fibrous porous mullite ceramics prepared by vacuum squeeze moulding technique, Ceram. Int. 42 (2016) 14843-14848.
[18] Y. Si, J.Y. Yu, X.M. Tang, J.L. Ge, B. Ding. Ultralight nanofibre-assembled
-p
cellular aerogels with superelasticity and multifunctionality, Nat. Commun. 5
re
(2014) 5802.
[19] Y. Si, X.Q. Wang, L.Y. Dou, J.Y. Yu, B. Ding. Ultralight and fire-resistant
na
4 (2018) 8925.
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ceramic nanofibrous aerogels with temperature-invariant superelasticity, Sci. Adv.
[20] R.L. Liu, X. Dong, S.T. Xie, T. Jia, Y.J. Xue, J.C. Liu, W. Jing, A.R. Guo,
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Ultralight, thermal insulating, and high-temperature-resistant mullite-based nanofibrous aerogels, Chem. Eng. J. 360 (2019) 464-472.
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[21] L. Yuan, B.Y. Ma, Q. Zhu, X.D. Zhang, H. Zhang, J.K. Yu, Preparation and properties of mullite-bonded porous fibrous mullite ceramics by an epoxy resin gel-casting process, Ceram. Int. 43 (2017) 5478-5483. [22] L. Yuan, Z.Y. Liu, X.H. Hou, Z.Q. Liu, Q. Zhu, S.G. Wang, B.Y. Ma, J.K. Yu, Fibrous ZrO2-mullite porous ceramics fabricated by a hydratable alumina based 19
aqueous gel-casting process, Ceram. Int. 45 (2019) 8824-8831. [23] H. Xu, H.Y. Du, J.C. Liu, A.R. Guo, Preparation of sub-micron porous yttria-stabilized ceramics with ultra-low density by a TBA-based gel-casting method, Chem. Eng. J. 173 (2011) 251– 257. [24] Z.G. Hou, H.Y. Du, J.C. Liu, R.H. Hao, X. Dong, M.X. Liu, Fabrication and properties of mullite fiber matrix porous ceramics by a TBA-based gel-casting
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process, Chem. Eng. J. 33 (2013) 717-725.
[25] W. Zhang, Q. MA, Zeng K, S. Liang, W. Mao. Mechanical properties and thermal
stability of carbon fiber cloth reinforced sol-derived mullite composites, J. Adv.
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re
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Ceram. 8 (2019) 218-227.
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PII:
S0955-2219(20)30012-1
DOI:
https://doi.org/10.1016/j.jeurceramsoc.2020.01.011
Reference:
JECS 12988
To appear in:
Journal of the European Ceramic Society
Received Date:
15 September 2019
Revised Date:
5 January 2020
Accepted Date:
6 January 2020
Please cite this article as: Zhang Y, Wu Y, Yang X, Li D, Zhang X, Dong X, Yao X, Liu J, Guo A, High-strength thermal insulating mullite nanofibrous porous ceramics, Journal of the European Ceramic Society (2020), doi: https://doi.org/10.1016/j.jeurceramsoc.2020.01.011
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High-strength thermal insulating mullite nanofibrous porous ceramics Ying Zhanga, Yongjun Wub, Xukun Yangb, Dinghe Lic, Xueying Zhanga, Xue Dongc⁎, Xinghe Yaob, Jiachen Liua, Anran Guoa⁎ a
School of Materials Science and Engineering, Key Lab of Advanced Ceramics and
Machining Technology of Ministry of Education, Tianjin University, Tianjin 300072,
b
Aerospace Business Department, China Academy of Launch Vehicle Technology, Beijing 100076, China
College of Aeronautical Engineering, Civil Aviation University of China, Tianjin
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c
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China
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300300, China
Corresponding author:
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Abstract
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Anran Guo, Email: [email protected]; Xue Dong, Email: [email protected]
Mullite fibrous porous ceramics is one of the most commonly used high
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temperature insulation materials. However, how to improve the strength of the mullite fibrous porous ceramics dramatically under the premise of no sacrificing the low
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sample density has always been a difficult scientific problem. In this study, the strategy of using mullite nanofibers to replace the mullite micron-fibers was proposed to fabricate the mullite nanofibrous porous ceramics by the gel-casting method. Results show that mullite nanofibrous porous ceramics present a much higher compressive strength (0.837MPa) than that of mullite micron-fibrous porous ceramics 1
(0.515MPa) even when the density of the mullite nanofibrous porous ceramics (0.202 g/cm3) is only around three quarters of that of the mullite micron-fibrous porous ceramics (0.266 g/cm3). The obtained materials that present the best combination of mechanical and thermal properties can be regarded as potential high-temperature thermal insulators in various thermal protection systems. Key words: Mullite fiber; Porous ceramics; Thermal insulation materials
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1. Introduction Mullite fibrous porous ceramics is one of the most commonly used high
temperature (>1300oC) insulation materials in various thermal protection systems
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owing to their low density, high porosity, excellent thermal shock resistance and low thermal conductivity [1-4]. The mullite fibrous porous ceramics are mainly composed
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of the interlocked micron-sized mullite fibers bound by a high-temperature binder [5].
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This three-dimensional porous skeleton structure not only endows the material with the characteristics of low density and low thermal conductivity, but also ensures that the three-dimensional pores existing in the material don’t collapse at high
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temperatures [6-9].
Density, compressive strength and thermal conductivity are the most important
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properties of the mullite fibrous porous ceramics (MFCs). Generally, the thermal
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conductivity of the MFCs is mainly determined by the solid-phase heat conduction and the gas-phase heat conduction due to its highly porous structure. Owing to the fact that the thermal conductivity of the air is much lower than that of the mullite fiber as the solid phase, the most effective way to reduce the thermal conductivity of the MFCs is to reduce the content of the solid phase, namely to decrease the sample density [10]. However, the decrease of density will inevitably lead to a decrease of the 2
material strength, which in turn limits the further application of the mullite fibrous ceramics. Therefore, how to improve the strength of the mullite fibrous porous ceramics under the premise of no sacrificing the low sample density has always been a difficult scientific problem. Many researchers have tried to change the microstructure of the MFCs, including the distribution of the binder and the arrangement of the mullite fiber, in order to increase the sample strength [11, 12]. However, results show that these methods only increase the sample strength in a
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limited range because they did not change the fundamental unite (mullite fibers) of the mullite fibrous ceramics.
The fracture of mullite fibrous porous ceramics is mainly caused by the fracture of
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the mullite fiber or the binder [13]. Therefore, one effective method to increase the sample strength is increasing the strength of mullite fibers and binders. Based on the
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above-mentioned, it can be assumed that the properties of MFCs can be greatly
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improved if the size of the mullite fiber as the matrix of MFCs reduces to the nanoscale. Firstly, with the reduction of the mullite fiber diameter from the micron scale to the nanometer scale, the strength of the fibers will substantially increase due
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to the decrease of the internal defects existing in mullite fibers [14]. Secondly, the size of the pores lapped by the fibers would reduce dramatically due to the decrease of
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fiber diameter, bringing about the increase of the pore wall content and the bonding
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points, finally resulting in the increase of the sample strength. The forming methods of mullite micron-fibrous porous ceramics include vacuum
filtration forming method and pressure assisted molding method, and both methods will exert a longitudinal pressure on the materials [15-17]. However, as the fiber diameter reduced to the nanometer scale, the longitudinal pressure would destroy of the 3D skeleton structure during the forming process, and consequently lead to the 3
increase of the sample density. Therefore, the best forming method to fabricate the mullite nanofibrous porous ceramics should be an in-situ forming method to prevent the structural collapse. Freeze casting which could obtain samples with small linear shrinkage and high porosity was proposed to prepare nanofibrous porous ceramics by many researchers [18, 19]. However, the mullite nanofibrous porous ceramics prepared by the freeze casting usually showed a multilevel pore structure with small pores formed by the overlapped nanofibers and large pores caused by the sublimation
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of the ice crystal, which was not beneficial to the sample strength [20]. Different from the freeze casting method, the gel casting method as another quite common in-situ forming method is able to fabricate samples with uniform microstructure and high
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strength [21, 22]. However, the water-based gel casting often leads to the break of samples with ultralow solid content during the drying process. Considering that the
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tertbutyl alcohol (TBA) has higher saturated vapor pressure and lower surface tension
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relative to water, and can volatilize rapidly at 50oC, the TBA-based gel-casting is a promising method for the fabrication of mullite nanofibrous porous ceramics [23, 24]. In this work, in order to further increase the strength of the mullite fibrous
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ceramics, the strategy of using mullite nanofibers to replace the mullite micron-fibers was propose to fabricate the mullite nanofibrous porous ceramics. The fabrication
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process of mullite nanofibrous porous ceramics was analyzed. Moreover, the effects
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of the mullite fiber content and the fiber aspect ratio on the sample microstructures and properties were investigated. 2. Experimental 2.1 Materials The
precursor
of
mullite
nanofibers
was
prepared
using
polyhydromethylsiloxane (Alfa Aesar Chemistry Co., Ltd., China) as the silicon 4
source, aluminum-tri-sec-butoxide (Sigma-Aldrich Trading Co., Ltd., China) as the aluminum source, polyvinylpyrrolidone (PVP-K90, Mw=1300000, Sigma-Aldrich Trading Co., Ltd., China) as the spinning aid, respectively. The silica sol which was used as the high temperature binder for 3D skeleton structure was prepared with tetraethyl orthosilicate (TEOS), deionized water and ethanol (ETOH). Polyacrylamide as the low-temperature binder were prepared using tert-butyl alcohol (TBA), acrylamide
(AM),
N,
N-methylenebisacrylamide
(MBAM),
ammonium
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persulfate(APS) and N, N, N, N-tetramethylethylenediamine (TEMED). All the above chemical reagents for binder were analytical grade from Aladdin Reagent Co., Ltd.,
China. Mullite micron-fiber porous ceramics (MMFCs) were purchased from
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Zhejiang Hongda Crystal Fiber Co., Ltd., China. The density of the MMFCs is 0.266
sample open porosity is 88.92%.
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2.2 Preparation
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g/cm3, the sample thermal conductivity of the MMFCs is 0.1012 Wm-1K-1 and the
Firstly, the diphasic mullite precursor with the alumina/silica molar ratio of 3:1 was prepared by the electrospinning method. The detailed fabrication process was
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described in our previous works [20]. Secondly, the precursor fibers were calcinated at 800oC with a heating rate of 2 oC/min. The average fiber diameter was about 476
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nm as shown in Fig. S1 in the supplementary materials. Then amorphous fibers as
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shown in Fig. 1a were dispersed into individual fibers using a high-speed homogenizer. The dispersion was dried at 120°C and the obtained fibers were shown in Fig. 1b. After that, 14.5 g of the AM and 0.5 g of the MBAM were added into 85 g of TBA to obtain the 100 g of premix, respectively. Then 1.5 wt%, 2.0 wt% and 2.5 wt% of the individual amorphous fibers and 0.25 g of the silica sols (40 wt%) were added into 5ml premix solution, respectively. An aqueous solution of APS (40 wt%) and an 5
aqueous solution of TEMED (4 wt%) were used as the initiator and catalyst, respectively. Then 0.25 ml of the APS and 0.4 ml of the TEMED were dissolved in the mixture sequentially. A gel-casting process would occur after 30 minutes and the green body after demolding was dried at 60oC for 12h. Finally, the samples were sintered and the heating system was 1 oC/min from room temperature to 700oC and 5 o
C/min from 700oC to 1400oC. In order to prevent the cracking in the binder removal
process caused by the rapid decomposition of polymer, it must be noted that the
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samples were kept heating at 390°C and 540°C for one hour. Then the mullite nanofibrous porous ceramics (MNFCs) were obtained after calcination at 1400oC for
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two hours.
Fig.1. SEM images of (a) amorphous fibers after sintering at 800oC, (b) the individual nanofibers
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after mechanical dispersion and (c and d) the green body.
2.3 Characterization The sample microstructure was observed via the scanning electron microscope
(SEM, S-4800, Hitachi Ltd., Japan). The phase transition of samples was investigated by the X-ray Diffraction (XRD, Rigaku D/Max-2500, Japan). The thermogravimetry of the green body was carried out from room temperature to 1400°C in air by a 6
thermal analyzer (TG-DSC, Netzsch Sta 449C, Germany). The linear shrinkage and the bulk density of the sample were calculated based on the length and weight of the samples. The porosity of the sample was measured by the water-immersion technique using the Archimedes method. The thermal conductivity of the ceramics was investigated by the thermal conductivity measuring instrument with a transient hot disk method (Hot Disk TPS 2500S, Sweden). The compressive strength were tested at
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room temperature, 400°C, 600°C, 800°C and 1000°C with a loading speed of 0.02 mm/min by an electronic universal testing machine (CMT4303, Meister Industrial Systems, China). The pore size distribution of the ceramics was investigated using a mercury porosimetry (AutoPore IV 9510, America).
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3. Results and discussions
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3.1 Fabrication of mullite nanofibrous porous ceramics
It is a challenge to improve the strength of the MFCs dramatically without
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sacrificing its low sample density. In this paper, the strategy of using mullite nanofibers to replace the micron-fibers was proposed to fabricate the mullite
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nanofibrous porous ceramics with low density and high strength. The fabrication flow chart of the mullite nanofibrous porous ceramics is shown in Fig. 2a. The matrix in
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MNFCs was mullite nanofibers with the alumina/silica molar ratio of 3:1 prepared by the electrospinning method. In order to ensure the complete decomposition of organic
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components in precursor fibers and get the optimum calcination temperature, the crosslinked fibers were tested by a thermal analyzer from room temperature to 1400°C. As shown in Fig. S2 in the supplementary materials, the weight loss of fibers between 32°C to 800°C was about 73.42%, while the weight loss of fibers in the range of 800-1400°C was only 0.83%. Therefore, the fiber calcination temperature 7
was selected at 800°C, which could both remove the organic component and meanwhile maintain the fiber's flexibility. After calcination, the non-woven fiber mat was dispersed into individual fibers by a high-speed homogenizer. By controlling the mixing rate, fibers with different average lengths (~20 μm and ~7 μm) could be obtained. As an in-situ forming method, the TBA-based gel casting is selected to fabricate the mullite nanofibrous porous ceramics with an ultralow solid loading. The advantages of this method include that the gel-casting time can be controlled by
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changing the addition content of the initiator and catalyst, and the sample density can be easily controlled by changing the solid content. As shown in Fig.1c and Fig. 1d, the green body exhibited an obvious uniform pore structure formed by the overlapped
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mullite fibers. In the green body, fibers are coated by the polyacrylamide as the
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low-temperature binder and the silica gel as the high-temperature binder.
Fig.2. (a) Fabrication flow chart of the MNFCs; (b, c and d) SEM images of MNCFs; Inset of Figure 2b: an optical photograph of MNFCs
With the increase of the sintering temperature, the polyacrylamide decomposed 8
completely and the silica sols transformed into the silica acting as the high temperature binders binding the overlapped mullite fibers. SEM images in Fig. 2b, 2c and 2d demonstrated the microstructure of the MNFCs after sintered at 1400℃. As shown in Fig.2b, the pores were mainly formed by the overlapped mullite nanofibers and the enlarged image of the fiber junction (Fig. 2c and 2d) shows that the high-temperature silica binder firmly bound the connected mullite nanofibers. The inset of Fig. 2b shows the photograph of the mullite nanofibrous porous ceramics after
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sintered at 1400°C.
Fig.3. TG and DSC curves of the green body from room temperature to 1400oC.
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In order to determine the chemical changes of the green body after sintered at
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different temperatures, the thermogravimetric of the green body from room temperature to 1400°C in the air and the X-ray diffraction analyses of samples after sintered at different temperatures were carried out. Fig. 3 shows that the weight loss of the sample mainly consisted of three stages. The first stage is the weight loss about 13.2% between 25°C to 200°C, which was mainly caused by the removal of the free water. In the second weightlessness stage, a dramatical mass loss about 63.4% 9
occurred between 200°C and 656°C due to the oxidative degradation of organic component (acrylamide and N, N-methylenebisacrylamide) at high temperature. Two obvious exothermic peaks appeared at 390°C and 540°C respectively on the DSC curve, and the weight loss rate of green body around the two peaks were quite fast. Therefore, in order to prevent the cracking caused by the rapid oxidative degradation of polymer, the sintering process during this period should be as slow as possible. In the third stage (656°C~1400°C), the weightlessness was about 4.5% and the
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exothermic peak at 998°C was caused by the formation of mullite phase. As the temperature continued to rise, the weight of MNFCs remained stable which indicates
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that the MNFCs exhibited an excellent thermal stability.
Fig. 4. XRD patterns of the samples after sintering at different temperatures.
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Fig. 4 shows the X-ray diffraction patterns of samples after sintering at different
temperatures. The broad backgrounds at around 23° indicate the existence of amorphous silica in the green body and the sample at 800℃. When the sintering temperature was above 1000°C, the characteristic peaks of γ-alumina (PDF#75-0921) [20] and mullite phase (PDF#15-0776) could be clearly observed, indicating the 10
aluminum source firstly transformed into γ-alumina phase and then some metastable alumina reacted with amorphous silica forming the mullite phase. As the sintering temperature rose to 1200°C, the SiO2 peak at 23° (PDF#29-0085) was observed, and the diffraction peaks of γ-alumina phase in the fibers disappeared, suggesting the mullitization process basically completed. When the sintering temperature continued to rise to 1400°C, the peak intensity is further strengthened, indicating the enhanced crystallinity of mullite [25]. Moreover, the excessive silicon oxide attaches to the fiber
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joints as the binder to increase the sample strength. Ultimately, only the mullite phase and the high-temperature binder (SiO2 phase) exist in the MNFCs.
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3.2 Effect of fiber content
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Fig. 5. SEM images of MNFCs fabricated with the addition of different fiber contents. (a and d) 1.5 wt% fiber content, (b and e) 2.0 wt% fiber content, and (c and f) 2.5 wt% fiber content.
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Solid loading is a basic parameter to control the sample density in the TBA-based gel-casting system. In this experiment, the pore structure and the densities of MNFCs were adjusted by changing the fiber content (1.5 wt%, 2 wt% and 2.5 wt%), further leading to the variation on the mechanical strength and thermal conductivity of the MNFCs. All the MNFCs with different fiber contents were prepared with the long nanofibers as the starting materials. Fig 5 shows the SEM images of the MNFCs with 11
the addition of different fiber contents. For the MNFCs with 1.5 wt% fiber content (Fig. 5a and 5d), mullite nanofibers in the sample overlapped with each other forming a 3D network structure and the pore size was about 5~20 μm. Due to the low fiber content, the crossing point between the fibers mainly existed at the end of each fiber. Therefore, the sample showed a loose 3D structure and the pore size was close to the length of the fiber. As the fiber content increased, the 3D skeleton structure of the MNFCs became compact and the pore size of the sample with 2.5 wt% fiber content
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reduced to 1~8μm as shown in Fig. 5c and 5f.
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Fig. 6. (a) The compressive strength, linear shrinkage density, (b) open porosity and thermal conductivity of MNFCs fabricated with the addition of different fiber contents.
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As shown in Fig. 6a, the bulk density of MNFCs increased from 0.061 g/cm3 to 0.108 g/cm3 as the fiber content increased from 1.5 wt% to 2.5 wt%, resulting in the
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increase of mechanical strength from 0.114 MPa to 0.153 MPa. As shown in Fig.6b, the porosities decreased from 97.9% to 96.4% as the fiber content increased, and all
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the products exhibited a relatively high porosity due to the abundant pores formed by the unique nanofibrous skeleton. The increase of fiber content resulted in the increase of the material density, further leading to the increase of thermal conductivity (from 0.0597 Wm-1K-1 to 0.0715 Wm-1K-1). 3.3 Effect of the fiber aspect ratio 12
As shown in Fig. 5, the MNFCs mainly consisted of the randomly arranged mullite nanofibers, and therefore the aspect ratio of the fibers should be one of the most important factors influencing the microstructure and properties of the MNFCs. In this experiment, as shown in Fig.7a and Fig.7c, two kinds of mullite nanofibers with different fiber lengths (7 μm and 20 μm) were used as the raw materials to prepare the mullite nanofibrous porous ceramics with the fiber content of 1.5wt%. The samples using the long fibers and the short fibers as the starting materials were
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expressed as L-MNFCs and S-MNFCs, respectively. In addition, the properties of the purchased mullite micron-fibrous (10 μm in diameter) porous ceramics (MMFCs)
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were also investigated to show the effect of fiber diameter.
Fig. 7. SEM images of (a) short mullite nanofibers, (b) S-MNFCs, (c) long mullite nanofibers, (d) L-MNFCs and (e and f) MMFCs. 13
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Fig. 8. Pore size distribution of different mullite fibrous porous ceramics.
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The morphologies and the pore size distribution of different mullite fibrous porous
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ceramics (S-MNFCs, L-MNFCs and MMFCs) are shown in Fig. 7 and Fig. 8. The pore size distribution curve of MNFCs shows an obvious bimodal pore size
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distribution (pore I in 1-5 μm and pore II in 5-15 μm). As shown in Fig. 7b, the pore I was the fundamental porous structure of the nanofibrous porous ceramics, and was
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mainly formed by the overlapped neighboring nanofibers. These relatively densely distributed fibers compose the pore walls of pore II. Different from the composition of
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the pore I, the contact points of the overlapped mullite fibers in the pore II tend to be at the end of nanofibers. Therefore, as the length of nanofiber increases, the size of
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pore II in Fig.7d is larger than that in Fig. 7b. Meanwhile, this information could also be confirmed in the curve of pore size distribution of different mullite fibrous porous ceramics. Compared with the pore size distribution of S-MNFCs (red curve in Fig.8), the content of pore I in L-MNFCs (blue curve in Fig. 8) decrease and that of pore II increase with the increase of nanofiber length. Furthermore, the average pore size of two samples increased from 3.185 μm to 7.991 μm. 14
Fig. 9. (a) The compressive strength, the thermal conductivity and bulk density of samples
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fabricated with different fibers. (b) Compressive strength of the S-MNFCs at different test temperature.
When the fiber diameter increased from nanometer scale to micron scale, the
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porosity of MMFCs decreased greatly (88.92%), and the crossed mullite fibers
formed 45-120 μm macropore (pore III). Compared the properties of S-MNFCs and
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MMFCs (Fig.9a), it can be found that although the density of S-MNFCs (0.202 g/cm3)
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was still a little lower than that of MMFCs (0.266 g/cm3), the compressive strength of S-MNFCs (0.837MPa) was much higher than that of MMFCs (0.515MPa), which was mainly attributed to the difference of the fiber diameter. As the fiber diameter reduced
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from the micron scale to the nanometer scale, the strength of the individual nanofiber substantially increased due to the decrease of the internal defects in the fibers, and the
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amount of pore wall and bonding points also increased obviously because of the
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decrease of the pore size, finally resulting in the increase of the material strength. Taking into account the potential application of the mullite nanofibrous porous ceramics, the compressive strength at different temperature are shown in Fig.9b. It can be seen that the compressive strength of the S-MNFCs decreased with the increase of the test temperature. Moreover, when the test temperature increased to 1000°C, the compressive strength of the material was still quite high, about 0.726 15
MPa. In addition, the thermal conductivity of the S-MNFCs (0.1001 Wm-1K-1) is a little lower than that of the MMFCs (0.1012 Wm-1K-1) due to the low density and small pore structure of nanofibrous ceramics. 4. Conclusions In summary, mullite nanofibrous porous ceramics were successfully prepared by the TBA-based gel casting, and the effects of fiber content and fiber aspect ratio on the microstructure and properties of the samples were studied. The strength of the
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MNFCs increased from 0.114 MPa to 0.158 MPa with the solid content of long nanofibers increasing from 1.5wt% to 2.5wt%. In addition, compared with the
S-MNFCs, the L-MNFCs showed an ultralow density (0.061 g/cm3) and low thermal
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conductivity (0.0597 Wm-1K-1). The S-MNFCs showed a much higher compressive strength (0.837MPa) than that of MMFCs (0.515MPa) although the density of
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MMFCs was still a little higher than that of S-MNFCs, which was mainly attributed to
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the difference of the fiber diameter. All the results prove that using mullite nanofibers to replace the mullite micron-fibers to prepare the mullite fibrous porous ceramics could effectively improve the sample strength under the premise of no sacrificing the
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low sample density. Based on the obtained results, the mullite nanofibrous porous ceramics provides a new insight into the development of high efficient thermal
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insulation materials.
Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgement This work is supported by the National Natural Science Foundation of China (Project No. 51502196 and 51872194). References [1] P. Soltani, M.S. Johari, M. Zarrebini, Eff ect of 3D fiber orientation on permeability of realistic fibrous porous networks, Powder Technol. 254 (2014)
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44-56. [2] A.K. Pradhan, D. Das, R. Chattopadhyay, S.N. Singh, Eff ect of 3D fiber orientation distribution on transverse air permeability of fibrous porous media,
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Powder Technol. 221 (2012) 101-104.
[3] T. Stylianopoulos, A. Yeckel, J.J. Derby, X.J. Luo, M.S. Shephard, E.A. Sander,
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V.H. Barocas, Permeability calculations in three-dimensional isotropic and
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oriented fiber networks, Phys. Fluids. 20 (2008) 123601. [4] R.B. Zhang, X.B. Hou, C.S. Ye, B.L. Wang, D.N. Fang, Fabrication and properties
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of fibrous porous mullite–zirconia fiber networks with a quasi-layered structure, J. Eur. Ceram. Soc. 36 (2016) 3539-3544.
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[5] J. Zhang, X. Dong, F. Hou, H.Y. Du, J.C. Liu, A.R. Guo, Eff ect of mullite fiber
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content on the microstructure and properties of porous mullite fiber/silica composite, Ceram. Int. 42 (2016) 6520-6524.
[6] D. Rasky, F.S. Milos, T.H. Squire, Thermal protection system materials and costs for future reusable launch vehicles. J Spacecr Rockets. 38 (2001) 294-296. [7] S. Liu, J.C. Liu, H.Y. Du, F. Hou, A.R. Guo, Microstructure of mullite fiber-based 17
hierarchical structures adjusted by Al/Si mole ratio of the raw material powders, Ceram. Int. 40 (2014) 11405-11410. [8] H. Schneider, J. Schreuer, B. Hildmann, Structure and properties of mullite—A review, J. Eur. Ceram. Soc. 28 (2008) 329-344. [9] R. Baetens, B.P. Jelle, A. Gustavsen, Aerogel insulation for building applications: A state-of-the-art review, Energ. Buildings. 43 (2011) 761-769.
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[10] J. König, V. Nemanič, M. Žumer, R.R. Petersen, M.B. Østergaard, Y.Z. Yue, D. Suvorov, Evaluation of the contributions to the effective thermal conductivity of an open-porous-type foamed glass, Constr. Build. Mater. 214 (2019) 337-343.
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[11] W.J. Zang, T. Jia, X. Dong, J.C. Liu, H.Y. Du, F. Hou, A.R. Guo, Preparation of
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homogeneous mullite-based fibrous ceramics by starch consolidation, J. Am. Ceram. Soc. 101 (2018) 3138–3147.
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[12] T. Jia, H. Chen, X. Dong, W.J. Zang, A.R. Guo, J.C. Liu, Preparation of
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homogeneous mullite fibrous porous ceramics consolidated by propylene oxide, Ceram. Int. 45 (2019) 2474-2482
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[13] X. Dong, G.F. Sui, Z.Q. Yun, M.C. Wang, A.R. Guo, J. Zhang, J.C. Liu, Effect of temperature on the mechanical behavior of mullite fibrous ceramics with a 3D
Jo
skeleton structure prepared by molding method, Mater. Des. 90 (2016) 942–948.
[14] Z.X. Chen, Z. Zhang, C.-C. Tsai, K. Kornev, I. Luzinov, M.H. Fang, F. Peng, Electrospun mullite fibers from the sol–gel precursor, J. Sol-Gel Sci. Technol. 74 (2015) 208-219. [15] C. Soares, N. Padoin, D. Muller, D. Hotza, C.R. Rambo, Evaluation of resistances 18
to fluid flow in fibrous ceramic medium, Appl. Math. Model. 39 (2015) 7197-7210. [16] B. Zhang, et al., Ultra-low cost porous mullite ceramics with excellent dielectric properties and low thermal conductivity fabricated from kaolin for radome applications, Ceram. Int. 45 (2019) 18865-18870. [17] R.B. Zhang, C.S. Ye, X.B. Hou, S.H. Li, B.L. Wang, Microstructure and
ro of
properties of lightweight fibrous porous mullite ceramics prepared by vacuum squeeze moulding technique, Ceram. Int. 42 (2016) 14843-14848.
[18] Y. Si, J.Y. Yu, X.M. Tang, J.L. Ge, B. Ding. Ultralight nanofibre-assembled
-p
cellular aerogels with superelasticity and multifunctionality, Nat. Commun. 5
re
(2014) 5802.
[19] Y. Si, X.Q. Wang, L.Y. Dou, J.Y. Yu, B. Ding. Ultralight and fire-resistant
na
4 (2018) 8925.
lP
ceramic nanofibrous aerogels with temperature-invariant superelasticity, Sci. Adv.
[20] R.L. Liu, X. Dong, S.T. Xie, T. Jia, Y.J. Xue, J.C. Liu, W. Jing, A.R. Guo,
ur
Ultralight, thermal insulating, and high-temperature-resistant mullite-based nanofibrous aerogels, Chem. Eng. J. 360 (2019) 464-472.
Jo
[21] L. Yuan, B.Y. Ma, Q. Zhu, X.D. Zhang, H. Zhang, J.K. Yu, Preparation and properties of mullite-bonded porous fibrous mullite ceramics by an epoxy resin gel-casting process, Ceram. Int. 43 (2017) 5478-5483. [22] L. Yuan, Z.Y. Liu, X.H. Hou, Z.Q. Liu, Q. Zhu, S.G. Wang, B.Y. Ma, J.K. Yu, Fibrous ZrO2-mullite porous ceramics fabricated by a hydratable alumina based 19
aqueous gel-casting process, Ceram. Int. 45 (2019) 8824-8831. [23] H. Xu, H.Y. Du, J.C. Liu, A.R. Guo, Preparation of sub-micron porous yttria-stabilized ceramics with ultra-low density by a TBA-based gel-casting method, Chem. Eng. J. 173 (2011) 251– 257. [24] Z.G. Hou, H.Y. Du, J.C. Liu, R.H. Hao, X. Dong, M.X. Liu, Fabrication and properties of mullite fiber matrix porous ceramics by a TBA-based gel-casting
ro of
process, Chem. Eng. J. 33 (2013) 717-725.
[25] W. Zhang, Q. MA, Zeng K, S. Liang, W. Mao. Mechanical properties and thermal
stability of carbon fiber cloth reinforced sol-derived mullite composites, J. Adv.
Jo
ur
na
lP
re
-p
Ceram. 8 (2019) 218-227.
20