Novel aluminum borate foams with controllable structures as exquisite high-temperature thermal insulators

Journal of the European Ceramic Society 40 (2020) 173–180

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Original Article

Novel aluminum borate foams with controllable structures as exquisite hightemperature thermal insulators ⁎

T

⁎

Han Luoa, Yuanbing Lia,b, , Ruofei Xianga,b, , Shujing Lia,b, Jun Luoc, Hailu Wanga, Xuesong Lia a

The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, 430081 Wuhan, PR China National-provincial Joint Engineering Research Center of High Temperature Materials and Lining Technology, Wuhan University of Science and Technology, 430081 Wuhan, PR China c School of optical information and energy engineering, Wuhan Institute of Technology, 430081 Wuhan, PR China b

A R T I C LE I N FO

A B S T R A C T

Keywords: Aluminum borate foam Foam-casting Process parameter Pore size distribution Thermal conductivity

In this contribution, a manageable foam-casting technique for the preparation of novel aluminum borate foams (ABFs) as thermal insulators with highly controllable performances is presented. ABFs were fabricated from αAl2O3 and 2Al2O3·B2O3 with the addition of various amounts of foaming agents, thickening agents, and slurry solid contents. The dispersions and rheological properties of the slurries were then examined, followed by exploration of microstructural evolution and testing of mechanical/thermal properties. It should be noted that the generated micro-pores generated and interlocking rod-like 9Al2O3·2B2O3 crystals may lead to superior mechanical tolerances and lower thermal conductivities for the ABFs. In general, the as-prepared ABFs with porosities ranging from 73.8 to 96.3 vol%, compressive strengths of 8.20–0.15 MPa, and thermal conductivities of 0.228–0.046 W/(mK) (200–800 °C) could render them suitable for application as high-temperature thermal insulating materials.

1. Introduction In the past few decades, ceramic foams have been extensively studied owing to their excellent properties, such as a low density, high melting point, low thermal conductivity, large specific surface area, good thermal shock resistance, and perfect chemical durability, which render these ceramic foams widely applicable to industrial fields, especially under high-temperature condition. In terms of the pore architecture, ceramic foams with open and interconnected pore structures have been employed as carriers for catalysts [1], for the capture of particulate matter [2], and in molten metal filters [3], gas permeable bricks [4] and bone substitutes [5]. In addition, ceramic foams with partial or complete closed pore structures have been applied as soundabsorbing materials [6] and thermal insulators [7]. To date, ceramic foams have been successfully fabricated by a range of methods including packed arrays of particles [8], in situ decomposition [9], foaming methods [10], template methods [11], and 3D printing [12]. Among these techniques, the foaming method is a versatile fabrication technology for ceramic foams owing to its simplicity, environmental friendliness, low cost, and versatility. In addition, the process parameters for this method can be varied to regulate the product porosity, with example parameters including the foaming [13] or sintering [14]

⁎

temperature, the solid content of the slurry [15], and the addition of different types or quantities of each additives [16]. This in turn can improve the mechanical properties of the samples and allow investigations to be carried out into the relationship between the pore architecture and the physical properties of these ceramic foams. However, an excess porosity lowers the ceramic foam strength, thereby limiting their further application. As such, improving the mechanical properties of ceramic foams whilst optimizing the pore structure [17] is a priority. The introduction of a reinforcing phase such as fibers [18], whiskers [19], and plate crystals [20] has also been examined. In contrast to regulation of the complex pore architecture or the introduction of an enhanced phase that is expensive and poorly dispersible, the selection of a material with superior properties is a potential solution. In this context, aluminum borate (9Al2O3·2B2O3) is a binary compound with the highest melting point of the Al2O3-B2O3 phase diagram, in addition to possessing outstanding mechanical properties, an excellent corrosion resistance, and a good chemical stability [21]. Compared with the properties of materials commonly used as thermal barrier layers (e.g., mullite, spinel, tieillite, cordierite, and calcium hexaluminate), aluminum borate exhibits unparalleled comprehensive properties with a theoretical strength up to 8000 MPa [22]. More

Corresponding authors at: The State Key Laboratory of Refractories and Metallurgy, Wuhan University of Science and Technology, 430081 Wuhan, PR China. E-mail addresses: [email protected] (Y. Li), [email protected] (R. Xiang).

https://doi.org/10.1016/j.jeurceramsoc.2019.08.018 Received 15 June 2019; Received in revised form 15 August 2019; Accepted 15 August 2019 Available online 10 September 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.

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furnace to 1200 °C for 3 h to produce the desired ABFs.

importantly, its unique rod-like crystal structure may lead to self-reinforcing as the basis for preparing lightweight, heat-insulating, and high-strength porous ceramics. However, to the best of our knowledge, very few studies have focused on the utilization of aluminum borate foams (ABFs) as both thermal insulators and lightweight structured components, or on the relationship between the pore parameters, physical strength, and high temperature thermal conductivity of ABFs. Thus, we herein report the use of a water-based slurry foam-casting method for the preparation of a porous green skeleton followed by the fabrication of ABFs via thermal treatment. In addition, various process parameters (e.g., the addition of dispersing, foaming, and thickening agents, and the solid content of the slurries) will be optimized to prepare ABFs with a wide porosity distribution. Furthermore, the effects of pore structure on the compressive strength and thermal conductivity of the ABFs will be investigated in detail based on the mechanism of the foaming process.

2.4. Evaluation methods The particle size distribution of the MP was determined using a laser particle analyzer (Mastersizer 2000, Malvern Instruments, UK). The Zeta potential of the MP was measured by a Zeta Potential Analyzer (ZetaProbe, Colloidal Dynamics Co., Ltd., Ponte Vedra Beach, FL, USA) using hydrochloric acid and sodium hydroxide to adjust the pH value. The rheological properties of the slurry were measured using a stresscontrolled rheometer (Physica MCR301, Anton Paar, Graz, Austria) with a concentric cylinder measurement geometry CC27 with shear rates of 0.01–500 s−1 at 25 °C. Differential thermal analysis and thermogravimetry (DTA/TG) measurements were then carried out for gelatin between 30 °C and 900 °C at a heating rate of 10 °C/min under a flow of air using a thermal analyzer (Netzsch Sta 449C, Germany). The phase compositions of the MP and ABFs were analyzed by X-ray diffraction (Philips, X'Pert PRO, Cu Kα), while the ABF microstructures were observed by scanning electron microscopy (SEM, JSM-6610, JEOL, Japan). Prior to examination in backscattering mode, the ABFs were impregnated with epoxy resin and triethanolamine under vacuum, then cured, sliced, and polished to improve the grayscale contrast and edge definition between the pore and the solid, and to ensure that the so-called “shine-through” effect could be avoided. Furthermore, the pore size distribution of the ABFs was statistically analyzed using “Micro-image Analysis and Process System” (MIAPS) software according to the binarization images of the SEM images shown in Fig. 1. A full-automatic true density analyzer (AccuPyc1330, Micromeritics Instrument Corp, USA) was used to measure the true density of the sintered samples. The bulk density of each ABF was calculated from its mass-to-volume ratio, and the total porosity was calculated according to the following equation:

2. Materials and methods 2.1. Preparation of primary raw materials In order to prepare mixed powder of α-Al2O3 and 2Al2O3·B2O3, αAl2O3 (Kaifeng Special Refractories Co., Ltd., China, D50 = 2.4 μm, ≥99wt% pure), boric acid (Shanghai Macklin Biochemical Co., Ltd., China, D50 = 238.5 μm, ≥99.5 wt% pure) and polyvinyl alcohol (PVA) solution (concentration of 5 wt%) were first mixed using an Erich mixer at 1800 rpm for 30 min. The weight ratio of α-Al2O3: boric acid: PVA solution was 14:6:1. Then, rectangular bars (25 mm × 25 mm × 140 mm) were prepared using a die-pressing technique at 5 MPa. After dried at 110 °C for 24 h, all the samples were heated to 800 °C in a furnace at a heating rate of 5 °C/min and a holding time of 3 h. After cooling, the sintered body was crushed and then sieved using a 1 mm sieve. Next obtained particles were wet-milled in a planetary ball mill containing ZrO2 balls and alcohol at 300 rpm. The weight ratio of ball: particles: alcohol was 1:2:1. The milled mixtures were oven-dried at 110 ◦C for 24 h, and then sieved using a 100 mesh sieve to obtain mixed powder of α-Al2O3 and 2Al2O3·B2O3.

ρ ρt = ⎜⎛1 − b × 100%⎞⎟ ρtr ⎝ ⎠

(1)

where ρt represents the total porosity of the porous ceramic, and ρtr and ρb represent the true density and bulk density of the ABF, respectively. The compressive strength was measured on a hydraulic universal testing machine (ETM, Wance, China) using a load speed of 0.5 mm/ min using samples of 24 mm × 24 mm × 12 mm size (ASTM standard C365/C365M-05). Five samples were used to calculate the average compressive strength, total porosity and bulk density. Thermal conductivity measurements were performed on disk-shaped samples (20 mm thickness, 180 mm diameter), at 200, 400, 600, and 800 °C using a water flow plate thermal conductivity apparatus (PBDR-02, Precondar, PR China).

2.2. Other raw materials Other admixtures are listed in Table 1. It should be noted that BA is a gelatin solution prepared in-house that contains a 20 wt% solid content at 60 °C. 2.3. Sample preparation The batch composition used in the experiment is outlined in Table 2. Initially, a ceramic slurry containing pre-ball-milling-treated MP, deionized water, DA, and TA were prepared by planetary ball milling using zirconia balls as the milling media for 3 h to obtain a uniform slurry. Subsequently, FA was added and the mixture was subjected to high-speed stirring and foaming using a beater for 2–3 min. Subsequently, BA was added to the foamed slurry and mixed by mechanical agitation to prepare a homogeneous mixture. The foams were then casted into the mold, and after drying, porous ceramic-gelatin composites were obtained. Finally, the obtained products were heated in a

3. Results and discussion The results obtained for the variation in particle size with deagglomeration time are shown in Fig. 2. As indicated, the MP milled for 1 h exhibited a wide particle size distribution (1.1–31.8 μm) and a large D50 value of 9.2 μm, which confirms the presence of particle agglomerates. The size distribution gradually narrows and the D50 value noticeably decreases from 9.2 to 2.9 μm upon increasing the ball milling time from 1 to 4 h. Although a higher degree of deagglomeration can

Table 1 Introduction of additives. Additive Dispersing agent (DA) Foaming agent (FA) Thickening agent (TA) Binding agent (BA)

WSM-R SH-1261 CMC Gelatin solution

Main ingredient

Purchased from

Poly-carboxylic acid composite water reducer Polymer composite animal protein sodium carboxymethylcellulose Gelatin

Smile New Material Co., Ltd., China Dongguan Shenhai Energy-saving Building materials technology Co., Ltd., China Shanghai Macklin Biochemical Co., Ltd., China

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Table 2 The batch composition of the samples. Sample no.

SC50

SC55

SC60

SC65

SC70 FA1

FA0.75

FA0.5

FA0.25

FA0.1 TA0

TA0.05

TA0.1

TA0.15

TA0.2

MP Distilled water DA FA TA BA

50 50 +0.25 +1 0 +10

55 45 +0.25 +1 0 +10

60 40 +0.25 +1 0 +10

65 35 +0.25 +1 0 +10

70 30 +0.25 +1 0 +10

70 30 +0.25 +0.75 0 +10

70 30 +0.25 +0.5 0 +10

70 30 +0.25 +0.25 0 +10

70 30 +0.25 +0.1 0 +10

70 30 +0.25 +0.1 +0.05 +10

70 30 +0.25 +0.1 +0.1 +10

70 30 +0.25 +0.1 +0.15 +10

70 30 +0.25 +0.1 +0.2 +10

#:“+”denote the amount of additives added is based on the overall quality of water and mixed powders.

cause wet foams to exhibit instability due to large (heavy) particles tending to descend under gravity [23], upon increasing the ball milling time from 4 to 5 h, the D50 value decreased only slightly from 2.9 to 2.4 μm. This is due to the fact that as the ball milling time increases, the ball milling efficiency inevitably decreases [24]. As such, a deagglomeration time of 4 h was selected to obtain the pre-ball-milling treated MP for use in the subsequent dispersions/rheological properties tests and foaming stage. Fig. 3 shows the Zeta potential of the MP (D50 = 2.9 μm) suspended in distilled water as a function of pH both in the presence and absence of DA. For the preparation of high performance foam ceramics, a ceramic slurry with a good dispersibility is required. According to the DLVO theory, the repulsion energy produced by the electric double layer on the particle surface is proportional to the square of the surface potential, and so improving the surface potential of a powder using appropriate dispersants is an effective means to improve the powder dispersion. It should be noted that adsorption was attributed to hydrogen bonding between the carboxylate groups on the long chain of DA and the ceramic surface groups. Upon the addition of DA, the absolute value of the slurry Zeta potential reached a high level in the initial time state compared with the case where no DA was added, and a large value of ˜40 was achieved over a pH range of 6–11.5. As such, the pH value of the suspension should not be adjusted to any great extent to avoid the introduction of impurities in subsequent processes. Fig. 4 shows the effect of the DA, TA, and solid contents on the viscosity of the ceramic slurry, in addition to an increase in the shear rate from 0.01 to 500 s−1 at 25 °C. Indeed, we wished to carry out a detailed study into the rheological properties of the slurry as a premise to the foaming process. It was clearly observed that all slurries exhibited pseudo-plastic behavior, and so the viscosity decreased rapidly upon increasing the shear rate, and interestingly, in the low shear rate region, the slurry viscosity decreased rapidly upon increasing the shear rate, then remained relatively constant in the high shear rate region. This was attributed to the fact that during the initial shear action, the internal network structure of the slurry is destroyed, and a new network is subsequently generated between the particles. Thus, Fig. 4 (a) shows the effect of the DA content on the viscosity and the flow properties of the suspensions with a 70 wt% solid phase content and in the absence of

Fig. 2. Evolution of particle size distributions of MP with different ball milling time.

Fig. 3. Zeta potential as a function of pH and DA addition of MP suspended in distilled water.

Fig. 1. (a) SEM micrographs of polished samples and (b) images of binarized photographs. 175

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Fig. 4. Effect of (a) DA content, (b) solid content and (c) TA content on rheological properties of ceramic slurry.

involved heating to 250 °C at a rate of 5 °C/min, followed by heating to 650 °C at a rate of 1 °C/min, then heating to 1200 °C at 5 °C/min and holding for 3 h. The XRD patterns of the MP and the ABFs are shown in Fig. 6, where almost pure phase 9Al2O3·2B2O3 was obtained from a mixture of 2Al2O3·B2O3 and α-alumina. This indicates that in situ reactive sintering is a convenient and effective approach for the syntheses of singlephase products (see Eqs. 3 and 4).

TA. It is notable that the viscosity of the slurries initially decreased but then increased upon increasing the DA content from 0.05 to 0.3 wt%; the minimum viscosity was observed upon the addition of 0.25 wt% DA. This was attributed to the fact that when the DA content is low, it cannot completely cover the particle surface, while in the presence of excess DA, micelles were easily formed in solution between the functional groups of the DA itself, resulting in a reduction in the slurry stability and an increased viscosity. In addition, the use of lower admixtures is desirable to reduce costs, and so 0.25 wt% DA was selected for the following experiment. It can be also seen in Fig. 4 (b) that the slurry viscosity decreased upon reducing the solid content at the same shear rate, in the presence of 0.25 wt% DA, and in the absence of TA. Upon decreasing the solid content from 70 to 50 wt%, the viscosity of the slurry decreased markedly and sharply. It should be noted that the Woodcock formula was developed to investigate the effect of solid loading on the viscosity of ceramic slurry, as outlined in Eq. 2 [25]: 1

h 1 5 = ⎜⎛ + ⎟⎞ d 3 πφ 6⎠ ⎝

(3)

9Al2O3 (s)+ 2B2O3 (l)→Al18B4O33 (s)

(4)

Fig. 7 shows the microstructures of the ABFs obtained by polishing, in addition to the pore size distributions with different FA, TA, and solid contents. Pore formation was predominantly determined by the introduction of highly spherical bubbles during mechanical stirring, although a small fraction of pores may originate from the burning loss of organics (i.e., BA, TA, and FA). It should also be noted that the pore size distribution shifted to larger diameters and changed from unimodal to bimodal. This unique hierarchical aperture distribution was determined by the essence of the foaming process, and has been studied as the classical “cell-window” model. In general, the slurry viscosity is known to affect both the stability of the bubble and the foaming ability. In addition, the foaming ability is closely related to the swelling properties of the slurry after mechanical stirring, while the stability of the bubble is indicated by its shape retention after foaming, which involves resistance to spontaneous accumulation and floating to the top of the slurry. These factors would result in the formation of large interconnected bubbles or even the collapse of the wet foams. A schematic model representing the effect of the slurry viscosity on pore formation in the ABFs is shown in Fig. 8. Under low viscosity conditions, the slurry is readily and absolutely foamed, but the bubbles are unstable, coalescing with each other to grow and form an interconnected pore structure. In contrast, although a high-viscosity slurry cannot be fully foamed, the bubbles are particularly stable, rarely coalescing with each other, thereby resulting in the formation of closed pore structures with a narrow pore size distribution. Fig. 9 shows the fracture morphology of the ABFs (TA0.2), where it can be observed that 3D non-directional interlocking rod-like crystals are formed. According to previous studies, ceramics with interlocked anisotropic grains such as rods or flakes possessed a significantly higher strength than those with isotropous grains. In our study, elongated aluminum borate grains grew in the struts between the pores and became interlocked with one another, thereby improving the strength of the material. Fig. 10 shows the porosity, bulk density, and associated compressive strength of the various ABFs prepared using different FA, TA, and solid contents. More specifically, upon increasing the FA content from 0.1 to 1.0 wt%, the total porosity increased from 85.2 to 92.6 vol%, while the bulk density and compressive strength decreased from 0.44 to 0.22 g/ cm3 and 3.56 to 1.00 MPa, respectively. In addition, upon decreasing the solid content from 70 to 50 wt%, the total ABF porosity increased from 92.6 to 96.3 vol%, while the bulk density and compressive

2

−1

9(Al4B2O9) (s)→2(Al18B4O33) (s)+5B2O3 (l)

(2)

where h is the space between the ceramic particles in the slurry; d is the particle diameter; and ϕ is the solid loading. As deduced from Eq. 2, the space between the ceramic particles in the slurry increased upon decreasing the solid loading, which leads to an increase in the repulsive interactions and reduced flow of the slurry due to the network structure between the ceramic particles reducing the viscosity. In addition, Fig. 4 (c) shows the effect of TA addition on the slurry viscosity with fixed DA and solid loading of 0.25 and 70 wt%. As indicated, upon increasing the amount of added TA from 0.05 to 0.2 wt%, there was a significant increase in viscosity at low shear rates, but no significant difference at high shear rates. It should be noted that a slurry viscosity < 1 Pa·s at a shear rate of 100 s−1 has been widely accepted as a standard for its compatibility for the foam-casting process [26]. In terms of the rheological behavior, it can be clearly seen that the viscosities of all slurries meet this criteria, and so it could be inferred that the addition of an optimal quantity of TA to the slurry will aid in controlling the properties of the product without unfavorably affecting the foaming performance. For the preparation of high performance ceramic materials, a suitable heat treatment system must be employed. For example, burnout conditions must be controlled to avoid an over pressure of the burnout gases resulting from additives in the green bodies, as this may lead to failure of the fragile ceramics. In this case, during heat treatment, the process most likely to cause collapse of the materials is the burning out of gelatin. Thus, the thermal analysis of gelatin was carried out as presented in Fig. 5 (a), where it is apparent that a weight loss of almost 100 wt% took place between 200 and 660 °C, while an exothermic peak was observed at 570 °C. The initial stage of heat treatment was therefore adjusted based on the thermal analysis of gelatin to avoid excessive or rapid weight loss, in addition to any risks of disruption through the evolution of large volumes of gases due to the burnout of gelatin. This stage is shown as thick lines in Fig. 5 (b) where the firing cycle again 176

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Fig. 5. (a): Thermal analyse curves of gelatin; (b): Thermal cycle for the green body. n σ = Cρ r σ0

(5)

where σ and σ0 are the mechanical strength of the porous sample and the sample whose porosity is close to 0, respectively. In addition, ρr is the relative density of the porous sample. As outlined in Fig. 10 (d), the results obtained herein correlated well with this model (R2 = 0.983). Fig. 11 shows the thermal conductivity values of the ABFs with various porosities and under different testing temperatures, and it is apparent that at a given porosity level, the thermal conductivity increased upon increasing the testing temperature, which was attributed to enhanced heat radiation [28]. For example, with a porosity of 73.8%, an increase in the testing temperature from 200 to 800 °C resulted in an increase in the thermal conductivity from 0.129 to 0.228 W/mK. However, at a given lower testing temperature, the thermal conductivity increased significantly with increasing porosity, i.e., at 200 °C, an increase in the porosity from 73.8 to 96.3% resulted in a decrease in the thermal conductivity from 0.129 to 0.046 W/mK. In contrast, at a given higher testing temperature, such as 800 °C, a different behavior was observed, as an increase in porosity from 73.8 to 91.0% gave a decrease in the thermal conductivity from 0.228 to 0.087 W/mK; upon increasing the porosity further to 96.3%, the thermal conductivity increased to 0.130 W/mK. This was attributed to the ABFs being composed of a solid skeleton of aluminum borate rods with air present in the pores; since the thermal conductivity of air is significantly lower than that of aluminum borate at relatively low testing temperatures, the air trapped in the pores plays the role of a thermal insulator. However, at higher porosities and testing temperatures, the radiative conductivity

Fig. 6. XRD patterns of raw material (MP) and ABFs.

strength decreased from 0.22 to 0.11 g/cm3 and 1.00 to 0.15 MPa, respectively. Upon increasing the TA from 0 to 0.2 wt%, the total porosity decreased from 85.2 to 73.8 vol%, while the bulk density and compressive strength increased from 0.44 to 0.77 g/cm3 and 3.56 to 8.20 MPa, respectively. According to the Gibson and Ashby models [27], the mechanical strength of a brittle cellular solid is related to its relative density:

Fig. 7. The microstructure and pore size distribution of ABFs. 177

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Fig. 8. A schematic model representing the effect of viscosity of the slurry on pore formation.

Fig. 9. Fracture morphology of ABFs.

Fig. 10. Effect of (a): FA addition; (b) Solid content; (c) TA addition on properties of ABFs; (d) Curves of Inρr versus Inσ.

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Fig. 11. Thermal conductivity values of ABFs with various porosity and testing temperature. Table 3 Comparison of properties of as-prepared ABFs and other ceramic foams reported in the literature fabricated via foam-casting. Material of ceramic foams

porosity

compressive strength (MPa)

thermal conductivity (W/mK)

Anorthite [30] Mullite-corundum [14] (Y1-xHox)2Si2O7 [31] Zirconia [32] Our work

82-94 vol% 65.9-82.7 vol% 78.3-81.2 vol% 96-98% vol% 73.8-96.3 vol%

– 1.61-9.75 11-13.9 0.4-2 0.15-8.2

0.04-0.08 (room temperature) 0.13-0.375 (200-1000 °C) 0.18-0.31 (30-300 °C) 0.027 (room temperature) 0.046-0.228 (200-800 °C)

(200–800 °C). We therefore propose that these materials exhibit great potential for application as high-temperature insulators to replace traditional ceramic materials especially in nuclear energy systems.

likely played a key role; this value could be calculated using Eq. 6 as a function of the pore characteristics [29], namely the emissivity ε of the pore walls, the Stefan-Boltzmann constant σ, the temperature T, the diameter dmax,pore of the pore, and a shape factor γ. λrad,pore=4*ε*σ*dmax,pore*γ*T3

Acknowledgments

(6)

This work was supported by the National Natural Science Foundation of China (No. 51772221), the National Key R＆D Program of China (No. 2017YFB0310701) and the Key Program of Natural Science Foundation of Hubei Province, China (No. 2017CFA004).

Finally, Table 3 compares the porosities, compressive strengths, and thermal conductivities of the ABFs described herein with those of previously reported ceramic foams fabricated via the foam-casting method. As outlined in the table, the prepared ABFs exhibited a highly-tunable porosity range, which is beneficial in the context of industrialization, future applications, and practical value.

References [1] R.Y. Chen, W.B. Jia, Q. Shan, Y.S. Ling, D. Lao, Y.F. Wang, S.J. Li, D.Q. Hei, A novel design of Al2O3-ZrO2 reticulated porous ceramics with hierarchical pore structures and excellent properties, J. Eur. Ceram. Soc. 39 (2019) 1877–1886. [2] J.J. Liu, B. Ren, Y.L. Wang, Y.J. Lu, L. Wang, Y.G. Chen, J.L. Yang, Y. Huang, Hierarchical porous ceramics with 3D reticular architecture and efficient flowthrough filtration towards high-temperature particulate matter capture, Chem. Eng. J. 362 (2019) 504–512. [3] P. Gehre, A. Schmidt, S. Dudczig, J. Hubalkova, C.G. Aneziris, N. Child, I. Delaney, G. Rancoule, D. DeBastiani, Interaction of slip- and flame-spray coated carbonbonded alumina filters with steel melts, J. Am. Ceram. Soc. 101 (2018) 3222–3233. [4] Q.H. Wang, Y.B. Li, S.J. Li, N.N. Xu, R.F. Xiang, O. Yang, Y.E. Ni, Effects of nanoalumina content on the formation of interconnected pores in porous purging plug materials, Ceram. Int. 43 (2017) 16722–16726. [5] A.D.R. Silva, W.R. Rigoli, D.C.R. Mello, L.M.R. Vasconcellos, E.M.J.A. Pallone, A.O. Lobo, Porous alumina scaffolds chemically modified by calcium phosphate minerals and their application in bone grafts, Int. J. Appl. Ceram. Technol. 16 (2019) 562–573.

4. Conclusions We herein reported the successful preparation of aluminum borate foams (ABFs) exhibiting a high mechanical strength and low thermal conductivity. This was achieved using 0.25 wt% dispersing agent, 0.1–1 wt% foaming agent, and 0–0.2 wt% thickening agent with firing at 1200 °C for 3 h. According to microstructural evaluations, it was speculated that the needle-shaped Al18B4O33 crystals played an important role in improving the strength of the prepared ABFs owing to the 3D interlocking microstructure, which was similar to a whiskerreinforced mechanism. In addition, the as-prepared ABFs exhibited a porosity ranging from 73.8 to 96.3 vol%, a compressive strength of 8.20–0.15 MPa, and a thermal conductivity of 0.228–0.046 W/(mK) 179

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