High-performance epoxy composites reinforced with three-dimensional Al2O3 ceramic framework

Journal Pre-proofs High-Performance Epoxy Composites Reinforced with Three-Dimensional Al2O3 Ceramic Framework Liu-Cheng Hao, Zi-Xuan Li, Fan Sun, Ke Ding, Xiao-Nan Zhou, Zhong-Xiao Song, Zhong-Qi Shi, Jian-Feng Yang, Bo Wang PII: DOI: Reference:

S1359-835X(19)30397-5 https://doi.org/10.1016/j.compositesa.2019.105648 JCOMA 105648

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Composites: Part A

Received Date: Revised Date: Accepted Date:

16 May 2019 22 September 2019 25 September 2019

Please cite this article as: Hao, L-C., Li, Z-X., Sun, F., Ding, K., Zhou, X-N., Song, Z-X., Shi, Z-Q., Yang, J-F., Wang, B., High-Performance Epoxy Composites Reinforced with Three-Dimensional Al2O3 Ceramic Framework, Composites: Part A (2019), doi: https://doi.org/10.1016/j.compositesa.2019.105648

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High-Performance Epoxy Composites Reinforced with Three-Dimensional Al2O3 Ceramic Framework Liu-Cheng Hao1, 2, Zi-Xuan Li1, Fan Sun1, Ke Ding1, Xiao-Nan Zhou1, Zhong-Xiao Song1, Zhong-Qi Shi1, Jian-Feng Yang1, Bo Wang1* 1State 2

Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an710049, China

High Voltage Switchgear Insulation Materials Laboratory of State Grid, Pinggao Electic Co.，Ltd， Pingdingshan 467001, China

Abstract High-performance epoxy composites are prepared by a vacuum infiltration method using a functionalized 3D-Al2O3 skeleton as the reinforcement and conducting framework. The preformed, porous thermal conducting framework acts as a highway for phonon transfer and can resist a high external loading to protect the epoxy composites. Consequently, a remarkable thermal conductivity of 4.356 W·m-1·K-1 is achieved for the 3D-43vol%Al2O3/epoxy composite, which is 2.7 times higher than that achieved for epoxy composites with conventional randomly dispersed 40vol% Al2O3 particles. The thermal conductivity of the composites with lower interface resistance was higher than that of 3D-43vol% Al2O3/epoxy composite with non-functionalization (3.796 W·m-1·K-1), because of the enhanced Al2O3/epoxy interface adhesion. Moreover, the 3D skeleton provides a dual advantage in its prominent reinforcement and toughening of the epoxy composites. The inclusion of 80vol% of Al2O3 preform dramatically enhanced the flexural strength (352 MPa) by 318%, compared to that of neat epoxy. Keywords: Polymer-matrix composites (PMCs); 3-Dimensional reinforcement; Thermal properties; Vacuum infiltration;

1

1. Introduction Ceramic/polymer composites could be widespread application in in many high-density electronic devices, including central processing units, radio frequency (RF) units, batteries and displays due to its outstanding properties, such as high thermal conductivity, high strength, good reliability and more importantly, low cost[1,2,3,4]. However, the intrinsic thermal conductivity of polymer is very low, for instance, the thermal conductivity of the widely used epoxy resin in these applications is only about 0.15~0.25 W·m-1·K-1. In order to enhance the thermal conductivity of polymer matrix composite (PMCs), traditional methods usually embedded high thermally conductivity ceramic fillers including AlN, SiO2, SiC, BN, Si3N4, graphite, carbon nanotube, graphene, in the polymer matrix[2, 5-15]. α-Al2O3 particles is often used as filler because of its low cost, a relatively high thermal conductivity, low linear coefficient and high electrical resistivity. Although the thermal conductivity of α-Al2O3 is relative lower, comparing with other inorganic materials (such as AlN, Si3N4, SiC, BN, etal), it is widely used in fabricating PMCs for the gas insulated switchgear (GIS) materials where there is a strict requirement for the dielectric properties under the SF6 atmosphere, due to its excellent corrosion resistance for HF decomposing from SF6 gaseous [16]. Although the thermal conductivity of PMCs can be enhanced by embedding the ceramic fillers into the matrix, a high loading fraction of fillers is generally needed to achieve a satisfied thermal conductivity, because the ceramic particles can intertwine with each other and the resultant particle chains work as thermal conductive pathways. Lee et al.[17] prepared AlN/epoxy composite with 57 vol.% filler content and reported a thermal conductivity of 3.39 W·m-1·K-1. Jang et al.[18] prepared Al2O3/epoxy composite with 80wt.% filler content by modifying the thermal 2

conductivity of composites and reported a thermal conductivity of 6.66 W·m-1·K-1. However, simply increasing the loading fraction of fillers of the composites did not obtain ideal thermal conductivity, mainly due to the high interfacial thermal resistance. This limitation was commonly attributed to the phonon scattering at the filler-filler physical contacts and the interface between the filler and the matrix. Thus, the enhancement in the thermal conductivity of such high loading ceramic/polymer composites is expected to be rather modest by Zhang et al[19] and Bae et al.[20]. Moreover, it should be noted that, excessive high filler loading in polymer matrix will deteriorate the processing ability, mechanical properties and dielectric breakdown strength of the PMCs[21]. Our interest is to fabricate an epoxy based composites with high thermal conductivity for the GIS materials at a moderate loading fraction of alumina fillers. Besides the interfacial thermal resistance between fillers and polymer matrix, the structure and distribution of the inorganic fillers in the polymer matrices has a great influence on the thermal conductive property of PMCs. Construction of 3D thermally conductive networks by the filler in polymer composites was considered as an effective approach to enhance the thermal conductivity of PMCs[22,23]. A continuous thermally conductive pathway in composites formed by thermally conductive fillers is devoted to improve the heat transfer and reduce the interfacial thermal resistance among fillers. Ng et al. [24] fabricated a 3D segregated structure composite by sintering the PP/AlN particles with a core-shell structure, which showed higher thermal conductivity than those of solution mixing and melt mixing methods. Yin et al.[25] synthesized 22vol%Si3N4/epoxy composites with high thermal conductivity of 3.89 W·m-1·K-1 by using sintered Si3N4 foams as fillers, which indicated that the formation of thermally conductive networks could improve phonon transferring and 3

enhance the thermal conductivity of composites. Compared with the particle reinforced polymer composites by mixing method, both the ceramic reinforcement and the polymer matrix exhibited a continuous structure in the MMCs with 3D interpenetrating network structures. Consequently, the characteristics of the reinforcement and the polymer matrix can be preserved simultaneously in the composites, which would contribute to the best overall thermal and mechanical performance. Herein, a facile and environmentally friendly method to construct 3D interpenetrating thermal transport channels in an epoxy matrix composite is reported. As shown in the flow diagram of Fig.1, 3D-Al2O3/epoxy composites was fabricated by a vacuum infiltration method using a surface functionalized 3D porous alumina ceramic skeleton as the 3D reinforcement and conducting framework. These 3D-Al2O3 skeletons with densely interlinked frameworks, excellent mechanical robustness and integrity, and multiple heat transfer paths, enable reliable fabrications of diverse 3D-Al2O3/polymer composites. Moreover, the 3D skeleton provided a dual advantage over conventional Al2O3 particles in its prominent reinforcement and toughening of the epoxy composites. 2. Experimental Procedure

4

Fig.1. Flow diagram for the preparation of the Al2O3 skeletons and 3D-Al2O3/EP composites. 2.1 Preparation of porous Al2O3 skeletons α-Al2O3 powder (0.5 μm, Shanghai Chaowei Nano Co. Ltd.) was used as starting powder. The green compacts with dimension of 50 mm×5 mm ×5 mm and Φ15mm×5 mm were formed by biaxial compression of the powder under pressure of 28~240 MPa for 1 min. The green bodies were placed in an electrical furnace and sintered at 1250°C ~1450°C for 2 h in air atmosphere with a heating rate of 300 K/h. Via using this simple method, porous Al2O3 with the open porosity in the range of 16~57% had been fabricated and the porosity can be controlled by varying the uniaxial pressure and the sintering temperature. 2.2 Surface functionalization of porous Al2O3 skeletons A silane coupling agent, 3-glycidoxypropyltrimethoxy silane (KH-560, Sinopharm Chemical Reagent Co. Ltd.) was used to treat the surfaces of alumina. Prior to surface functionalization, porous Al2O3 skeletons were dried in an oven for 4 h at 110°C. After cooling down to room temperature, the porous Al2O3 skeletons were soaked in the mixture of ethanol and silane coupling agent stirred at 80°C for 3 h. The weight ratio of porous Al2O3/silane coupling agent was 100/3. The value of PH was adjusted to ~4 by adding ethylic acid. The functionalized Al2O3 skeletons were isolated through centrifugation and quickly washed by fresh ethanol. Finally, the product was dried in a vacuum oven for 12 h to yield fried body. 2.3 Preparation of Al2O3/epoxy composites Bisphenol-A epoxy resin (E51, DGEBA, Shanghai Resin Factory Co. Ltd.,) was used as polymer matrix. Methyltetrahydrophthalic anhydride (MeTHPA, Sinopharm Chemical Reagent Co. Ltd) and 2,4,6-three-two methylamino methyl phenol 5

(DMP-30, Sinopharm Chemical Reagent Co. Ltd.) were used as curing agent and cross-linker catalyst, respectively. The mixture of epoxy resin and curing agent MeTHPA in ethanol solution with a weight ratio of 1/0.86 weight ratio was stirring under vacuum at 50 °C for 60 min. And then, stoichiometric quantity of the DMP-30 accelerator (epoxy/DMP-30=1:0.01 weight ratio) was added into the mixture and string under vacuum at 60 °C for 30 minutes. The resulted mixture was infiltrated into the preheated functionalized porous Al2O3 skeletons under vacuum condition for 4h. Finally, the composites cured as follows: 95°C, 2 h; 140°C, 3 h and 170°C, 4 h. 2.4 Characterization The relative density and porosity of the composite was determined by the Archimedes method. Infrared absorption spectras of the functionalized porous alumina ceramics were measured using a Fourier transform infrared (FTIR) spectrometer (Bruker Vertex 70). Morphological observation was determined by the scanning electron microscope (SEM, GeminiSEM500). Thermo gravimetric analysis (TGA) of Al2O3/epoxy composites was conducted by using TA Q500 Instruments. Samples (10 mg) were placed in aluminum crucibles and dynamic tests were performed by heating the samples from 25 °C to 1000 °C with a heating ramp of 10 °C/min, under nitrogen (thermolysis conditions) and air flow (thermo-oxidation conditions) mixture. The samples with a dimension of Φ12.7 mm×2 mm were using for thermal conductivity measurement. The thermal conductivity of the composites were measured using a light flash system (Netzsch Model LFA44) at 25°C and 100 °C, respectively. The flexural strength was measured by three-point bending method with a 20 mm span at a cross-head speed of 0.5 mm/min at room temperature. Each final value was averaged over five measurements. 6

3. Results and Discussion To demonstrate the actual occurrence of chemical reactions between modifiers and Al2O3, FTIR spectra of raw porous Al2O3 ceramics and functionalized Al2O3 ceramics were presented in Fig. 2. Compared with raw Al2O3, functionalized Al2O3 exhibits new absorption at 2950 and 2840 cm-1, which represent the valence stretching vibration of aliphatic C-H. The peak at 1050 cm-1 was assigned to the Si-O-Si stretch. The result of FTIR indicated that the epoxy groups had been successfully introduced to the surface of porous Al2O3 skeletons via chemical grafting. The presence of the hydroxide group coverage on its surface is one of the features of aluminum oxide. After surface functionalization, hydroxide group interacted with silane coupling agent in ethanol solution. And Si-O-Si stretch was introduced to the surface of porous Al2O3 ceramics. After treatment of epoxy groups, the interfacial adhesion strength could be improved by forming covalent bonds with the polymer matrix.

Fig. 2. FT-IR spectra of raw Al2O3 and modified porous Al2O3. The SEM morphologies of the fractured surfaces of the porous Al2O3 ceramics sintered at different sintering temperature are shown in Fig. 3. The sintered Al2O3 ceramics generally exhibit high porosity. The morphology of the open pores between 7

Al2O3 particles was irregular and homogeneous dispersed in the Al2O3. The porosity of porous Al2O3 ceramic decreased and the Al2O3 grain size increased with increasing the sintering temperature. It was due to the enhanced surface diffusion at high temperature. These as-sintered Al2O3 skeletons exhibited good mechanical properties, which enabled them to maintain their open pore structure during the infiltration of epoxy. The relative density of the porous Al2O3 ceramic skeletons sintered at 1250°C, 1350°C, and 1450°C were 43%, 60%, and 80%, respectively. For convenience, the relative density of the porous Al2O3 ceramic skeletons was also called “Al2O3 content” in the following context on the Al2O3/EP composites.

Fig. 3. SEM morphologies of the fractured surfaces of the porous Al2O3 ceramics sintered at different temperature for 2 h with a heating rate of 5 °C/min. Fig. 4 shows the SEM images of the fractured surfaces of the 3D interpenetrating Al2O3/epoxy composites with different volume content of Al2O3. As shown in it, after the epoxy infiltration under vacuum condition, all the pores in the porous Al2O3 8

ceramics are uniformly filled with epoxy matrix. Thus, the Al2O3/EP composites are composed

of

two

interpenetrating

networks.

For

the

composites

with

non-functionalization Al2O3 skeleton (Fig. 4(a)), the Al2O3/epoxy interface with the weak interfacial bonding was obviously observed. As shown in Fig .4(a), large gap between alumina and epoxy could be obviously observed, reflecting weak interfacial adhesion between inorganic filler and polymer. It was due to the adhesion force was lower the shrinkage force during curing processing. It means that the controllable tuning of the surface properties of the oxide particles is an important factor while fabricating the polymer-based composites. After treatment of epoxy groups-Al2O3, the interfacial adhesion strength was improved by forming covalent bonds with the polymer matrix[26]. There were mainly two characteristics of the Al2O3/epoxy composites fracture surface: (i) pull out of the Al2O3 grains; when the orientation of the Al2O3 grains was approximately perpendicular to the crack propagation way, few Al2O3 grains were pulled out from the epoxy matrix and some equiaxed shaped holes appeared; (ii) debonding of the Al2O3/epoxy interface; if the angle between the Al2O3 grains orientation and the fracture surface was small, the primary crack propagated against along the Al2O3/epoxy interface with the weak interfacial bonding, deflected from initial direction, grew along the weak interfacial. As a result, the intergranular fracture mode appeared for the Al2O3 grains. The crack deflection and the pullout of the Al2O3 grains would absorb the fracture energy, led to high fracture toughness. The well adhesion of the Al2O3/epoxy interface after the flexural trials suggested that the bonding between the Al2O3 grains and the epoxy matrix was strong enough to support the load transfer from the matrix to the reinforcements. It was indicated that the 3D Al2O3 skeleton provided a dual advantage over conventional Al2O3 particles in its prominent reinforcement and toughening of the epoxy composites. 9

Fig. 4. SEM photographs of the fractured surfaces of the 3D interpenetrating Al2O3/epoxy composites with (a) 43vol% raw Al2O3, (b) 43vol%, (c) 60%vol%, and (d) 80vol% functionalized Al2O3.

Fig. 5. Thermal conductivity of the 3D interpenetrating Al2O3/epoxy composites as a function of the filler content. Fig. 5 shows the thermal conductivity of the 3D interpenetrating Al2O3/epoxy composites as a function of the volume fraction of Al2O3 fillers. The thermal conductivity considerably increased from 4.36 to 30.84 W·m-1·K-1 when the Al2O3 10

filler content increased from 43 to 80vol%. For the 3D interpenetrating composites with 60vol% fillers, the thermal conductivity was 12 W·m-1·K-1. It is 1.8~3.5 times higher than that of the polymer composites with the same content of Al2O3 filler [27, 28, 29]. Remarkably, the measured thermal conductivity of the 3D interpenetrating composites was much higher than the prediction of Byrggeman equation. This demonstrated that 3D-Al2O3 thermally conductive pathways enhanced the thermal conductivity of the polymer matrix significantly. The intrinsic reason can be explored by Agari model [30], which considers the effect of dispersion state by introducing factor Cf and CP: lg   VCf lg  f  (1  V ) lg(Cm m) , where λ, λf, λm, and V represent thermal conductivities of the composite, filler, matrix, and volume fraction of the filler in the composite, respectively. Cm is a factor relating to the effect of the filler on the secondary structure of the polymer, and Cf is a factor relating to the ease in forming conductive chains of the filler. The closer Cf values are to 1, the more easily conductive chains are formed in composite. In the present study, the value of Cf was ~1.1, which indicated that the formation of thermal conductivity pathways in the composites strongly enhanced by the 3D-Al2O3 ceramic skeleton. The thermal conduction was primarily carried out by the Al2O3 ceramic skeleton, of which the thermal conductivity was largely determined by the grain boundaries. On the other hand, heat in Al2O3 ceramics is usually dominated by phonons, which can be scattered by defects such as grain boundaries and closed pores, etc. Higher sintering temperature led to larger particle size and fewer clean grain boundaries (Fig. 4). It would contribute to higher thermal conductivity in this study.

11

Fig. 6. a) Thermal conductivity of the Al2O3/epoxy composites fabricated by different methods, and (b) Comparison of thermal conductivity results in this work with Al2O3/polymer systems in other work. In order to better understand this effective 3D-Al2O3 skeletons thermal transport highway, the thermal conductivity of composites prepared by the powder mixing method is compared with that of the composites processed by vacuum assisted infiltrating method. Fig. 6(a) shows the thermal conductivity of Al2O3/epoxy composites fabricated by different methods. A remarkable thermal conductivity of 4.356 W·m-1·K-1 is achieved for the 3D-43vol%Al2O3/epoxy composite, which is 2.7 times higher than that achieved for epoxy composites with conventional randomly dispersed 40vol% Al2O3 particles (1.598 W·m-1·K-1).

And it was higher than that of

3D-43vol% Al2O3/epoxy composites by using raw alumina (3.796 W·m-1·K-1), which was due to the interface resistance caused by the weaker interfacial bonded Al2O3/epoxy interface (as shown in Fig. 4(a)). The comparison of our results with the previous literature values using in polymer-based composites is shown in Fig. 6 (b). Clearly, our composite shows thermal transport properties superior to those of most Al2O3/epoxy systems, accomplishing higher thermal conductivity in polymer-based Al2O3 composites but for lower filler loading. Our 3D-Al2O3/EP composite (~50 vol%) exhibits by far the 12

highest thermal conductivity for comparable filler loading, exceeding even that for higher volume fractions up to 55 vol% [31, 33], therefore indicating the existence of a unique thermal transport highway structure in our composites. Specifically, after surface functionalization, the thermal conductivity of our composites is higher than that of 3D-Al2O3/EP composites with non-functionalization fabricated by Yu etal [30] for comparable filler loading, because of the enhanced interface adhesion between alumina and polymer matrix. Although the filler particles contacting each other form continuous chains of particles over the whole composite samples and the resultant particle chains work as thermal conductive pathways. The efficacy of the interfacial modification became less remarkable. The interfacial thermal resistance partly comes from a barrier to the heat flow induced at the interface lead to deteriorative thermal conductivity, because of the phonon spectra mismatch between the inorganic and polymer phase. The surface chemical functionalization of fillers can form covalent bridge bonds, leading to the improved interfacial adhesion, which can minimize the interfacial phonon scattering and decrease the interfacial heat resistance by interpenetrating particle-resin and particle-particle interfaces.

Figure 7. Schematic illustration of three kinds of ceramic/polymer composites: (a) dispersed, (b) physically contacted, and (c) chemically bonded. As shown in Fig. 7(a), few Al2O3 particles are randomly dispersed in the epoxy matrix in case of lower loading of fillers, resulting lower thermal conductivity. From 13

the schematic of interfacial heat dissipation in the particles reinforced composite (as shown in Fig. 7b), it is clearly comprehended that the dispersed Al2O3 particles cannot form the thermal transport network required for efficient heat transfer. Instead, heat would be dissipated at the Al2O3/Epoxy interfaces, which obstruct the thermal transport. Once a high loading fraction of ceramic fillers is reach to the threshold value, in which the filler physically intertwine with each other, which facilitate thermal transport. However, the interfacial thermal resistance between fillers is higher than that of chemically boned Al2O3 skeletons after high temperature sintering. The bridge structures of 3D ceramic/polymer composites are like highways through the epoxy matrix (as shown in Fig. 7c) for heat flow to propagate on the highway, therefore decreasing the interfacial thermal loss with the polymer phase. Furthermore, the 3D thermal transfer paths are made of homogeneous Al2O3-bonded interconnections. It is different from the hetero-contact forms relying on weaker hydrogen bonds and/or Van der Waals*s interactions between Al2O3 and/or Al2O3 and other structures, which lead to a low-efficient phonon transfer and high interface-layer thermal resistance once heat flows along heterogeneous Al2O3 passageways. The Al2O3-self-interconnexted paths within 3D Al2O3/epoxy composites maintain a Al2O3 intrinsic phonon transfer performance, these are much more effective than those for multiple hetero-linked Al2O3 cases. The homogeneous interconnections within 3D-Al2O3 greatly reduce heter-interface resistance during heat transfer. Therefore for the same Al2O3 filler loading (~40 vol%), the thermal transport performance of 3D-Al2O3/EP composites is superior to that of particles reinforced composites. Consequently, 3D ceramic/polymer composites have a great advantage over particles reinforced composites in terms of thermal transport performance. There are several reasons that result in the unprecedented TC values. First, the 14

3D thermal transfer paths are made of homogeneous Al2O3-bondedinterconnections. It is different from the hetero-contact forms relying on weaker hydrogen bonds and/or Van der Waals’s interactions between Al2O3 and/or Al2O3 and other structures, which lead to a low-efficient phonon transfer and high interface-layer thermal resistance once

heat

flows

along

heterogeneous

Al2O3

passageways.

The

Al2O3-self-interconnexted paths within 3D Al2O3/epoxy composites maintain Al2O3 intrinsic phonon transfer performance, these are much more effective than those for multiple hetero-linked Al2O3 cases. The homogeneous interconnections within 3D-Al2O3 greatly reduce heter-interface resistance during heat transfer. Second, the densely interconnected Al2O3 porous skeletons provide high-densely Al2O3-made pathways, these are preferential to transfer heat and greatly enhance the flow efficiency. What is more, the isotropically constructed 3D-Al2O3 networks embedded into the composites render a uniform thermal disspation towards the entire 3D space, make 3D-Al2O3/Epoxy TC particularly isotropic.

Fig. 8. TGA thermograms of Al2O3/epoxy composites. The thermal stability of the Al2O3/epoxy composites was measured by TGA at a heating rate of 5 °C/min under an air atmosphere. The results are presented in Fig. 8. The degradation of the composites exhibited a two-stage process. The first stage occurs at the temperature between 300 and 460 °C (the temperature of 5% weight loss, T d5), corresponding to degradation of the cured epoxy network [34, 35]. It 15

should be noted that the thermal stabilization (T d5) of the composites is increased with the increase of Al2O3 loading (as shown in Fig. 6a). The fastest degradation temperatures (the peak of DTG profiles) of the Al2O3/epoxy composites exhibit a similar trend in Fig. 6b and increase with increasing the Al2O3 loading. The second stage takes place from 550 to 650°C, which was due to the pyrolysis of the degradation products of the epoxy. The char at 750 °C of the composites increased gradually with increasing Al2O3 content. This was attributed to the formation of char by the Al2O3 particles. The thermal stability of composites increased with increasing Al2O3 filler content in the composites. This suggests that the Al2O3 particles can retard the epoxy degradation and improve the thermal stability of the neat epoxy resin. Similar observations was concluded the resin/filler hybrid composites [36, 37, 38]. The presence of Al2O3 particles in the composites would inhibit the mobility of epoxy polymer and thus cause a ‘‘tortuous path’’ effect to delay the escape of volatile degradation products and the permeation of oxygen and char formation, which is likely responsible for the improvement in thermal stability [39, 40].

Fig. 9. Flexural strengths and specific strength of the 3D interpenetrating Al2O3/epoxy composites with different Al2O3 content. In addition to its excellent thermal conductivity and thermal stability, the 16

3D-Al2O3 porous ceramics also provide a remarkable reinforcement to the epoxy composites, which is necessary for shielding materials that are used in areas where mechanical properties are important. Fig. 7 shows the flexural strengths of the 3D interpenetrating Al2O3/epoxy composites with different Al2O3 contents. For comparison, the flexural strength data of the degreased green body, the sintered porous Al2O3 ceramic skeletons and the pure epoxy are also included in Fig. 7. The flexural strengths of both the Al2O3/epoxy composites and the sintered porous Al2O3 ceramic preforms rapidly increased with the Al2O3 content. The inclusion of 80vol% of Al2O3 preform dramatically enhanced the flexural strength (352 MPa) by 318%, compared to that of neat epoxy. Whereas, the addition of separated inorganic particles did not show markedly reinforcement as opposed to that of neat epoxy (167 MPa for AlN/Epoxy composites [41], and 160 MPa for 40vol%Al2O3 particles reinfroced Epoxy composites [42]). The strongest Al2O3 preform led to the highest strength (352 MPa) because the load was transferred from the epoxy matrix to Al2O3 skeletons. As shown in Fig. 9(b), the specific flexural strength of porous Al2O3 ceramics decreases with the volume content of alumina. It was mainly due to the weaker neck between Al2O3 grains of sintered body at lower temperature. The bonding strength between the Al2O3 grains increases with the sintering temperature. Whereas, the specific flexural strength of composites was almost constant with increasing the alumina content from 43vol% to 70vol%, and slightly increase as the alumina content increasing to 80%. For the 3D interpenetrating Al2O3/EP composites, both the ceramic reinforcement and the polymer matrix exhibited a continuous structure. Consequently, the characteristics of the Al2O3 reinforcement and the polymer matrix can be preserved simultaneously in the composites, which would contribute to the best overall performance. 17

Usually, the neat epoxy had a smooth and featureless fracture surface, which is typical for brittle epoxy. As shown in Fig.3, using a preformed 3D-Al2O3 skeleton, the rougher fracture surface provided a larger surface area for fracture energy absorption, which would contribute to high fracture toughness. Several pores left by the pulled-out Al2O3 particles suggested a moderate interfacial adhesion between the Al2O3 preform and the matrix. On the other hand, the preformed strong and robust 3D-Al2O3 framework throughout the epoxy matrix ensured a uniform dispersion of the Al2O3 particles, thereby providing a large contact area with the polymer chains, which greatly increased the energy required to pull out the numerous Al2O3 particles, comparing to the aggregation of mechanically dispersed Al2O3 particles. Furthermore, the interconnected 3D-Al2O3 ceramic frameworks can improve the mechanical properties of the epoxy by isotropic and effective stress transfer. Therefore, the excellent thermal and mechanical properties of 3D Al2O3/epoxy composites can be attributed to the more continuous and stronger conductive network for phonon transport, and load transfer. 4. Conclusions 3D interpenetrating Al2O3/epoxy composites with high thermal conductivity were prepared by impregnating epoxy into a preformed, porous, surface functionalized 3D alumina ceramic skeleton. The highly conductive framework resolved the dispersion problem of separated Al2O3 fillers and acted as an efficient bunch of channels for phonon transfer and struts to resist external loading after compounding with the epoxy. A remarkable thermal conductivity of 4.356 W·m-1·K-1 was achieved for the epoxy composites with 43vol% of 3D-Al2O3 framework, which is 2.7 times higher than that achieved for epoxy composites with conventional mechanically dispersed 40vol% Al2O3 particles. The thermal conductivity of the 18

composites with lower interface resistance was higher than that of 3D-43vol% Al2O3/epoxy composite with non-functionalization (3.796 W·m-1·K-1), because of the enhanced Al2O3/epoxy interface adhesion. Moreover, the 3D skeleton provides a dual advantage over conventional Al2O3 particles in its prominent reinforcement and toughening of the epoxy composites. The inclusion of 80vol% of Al2O3 preform dramatically enhanced the flexural strength (352 MPa) by 318%, compared to that of neat epoxy, which was due to the interconnected isotropic 3D framework that effectively transferred the external stress. These results indicate that the potential of 3D Al2O3 framework is an ideal candidate for high-performance polymer composites. Acknowledgments This work was supported by the National Key R&D Program of China (Grant Nos. 2017YFB0903803, 2017YFB0903800). Reference [1] J.S. Park, Y.J. An, K. Shin, J.H. Han, C.S. Lee. Enhanced thermal conductivity of epoxy/three-dimensional carbon hybrid filler composites for effective heat dissipation. RSC Adv 5 (2015) 46989-46996. [2] B. Tang, G.X. Hu, H.Y. Gao, L.Y. Hai. Application of graphene as filler to improve thermal transport property of epoxy resin for thermal interface materials. Int. J. Heat Mass. Transf. 85 (2015):420-429. [3] X. Huang, P. Jiang, T. Tanaka, A review of dielectric polymer composites with high thermal conductivity, IEEE Electr. Insul. Mag. 27 (4) (2011) 8-16. [4] L. Yin, X. Zhou, J. Yu, H. Wang, C. Ran. Fabrication of a polymer composite with high thermal conductivity based on sintered silicon nitride foam. Compos. Part A: Appl. Sci. Manufac. 90 (2016) 626-632. [5] C.Y. Hsieh, S.L. Chung, High thermal conductivity epoxy molding compound 19

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Flexural Properties of Functionally

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Figure caption:

Figure 1. Flow diagram for the preparation of the Al2O3 skeleton and 3D-Al2O3/EP composites. Figure 2. FT-IR spectra of raw Al2O3 and functionalized porous Al2O3, and the schematic diagram of surface functionalization of porous Al2O3 ceramics. Figure 3. SEM morphologies of the fractured surfaces of the porous Al2O3 ceramics sintered at different temperature for 2 h with a heating rate of 5 °C/min. Figure 4. SEM photographs of the fractured surfaces of the 3D interpenetrating Al2O3/epoxy composites with (a) 43vol% raw Al2O3, (b) 43vol%, (c) 60%vol%, and (d) 80vol% functionalized Al2O3 Fig. 5. Thermal conductivity of the 3D interpenetrating Al2O3/epoxy composites as a function of the filler content.

Figure 6. (a) Thermal conductivity of the Al2O3/epoxy composites fabricated by different methods, and (b) Comparison of thermal conductivity results in this work with Al2O3/polymer systems in other work. Figure 7. Schematic illustration of three kinds of ceramic/polymer composites: (a) dispersed, (b) mechanically contacted, and (c) chemically bonded. Figure 8. TGA thermograms of Al2O3/epoxy composites. Figure 9. Flexural strengths and specific strength of the 3D interpenetrating Al2O3 /epoxy composites with different Al2O3 content. Table 1. Thermal conductivity, linear coefficient and dielectric constant of thermal expansion for various materials at room temperature.

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Table 1. Thermal Conductivity, Linear Coefficient and Dielectric constant of Thermal Expansion for Various Materials at Room Temperature Materials

Thermal

Coefficient of

Dielectric

Application of

conductivity

thermal

constant

ceramic/polymer

(W·m-1·K-1)

expansion

composites

(ppm/°C) α-Al2O3

38~42

7

9.3~11.5

Excellent corrosion resistance of SF6 atmosphere in HV insulation

Fused SiO2

1.5~1.6

0.4~0.5

~3.9

Additive in insulation and

Crystalline

3

10

~3.9

materials

SiO2 BeO

packaging

300

5.5

high

High toxicity and the high cost

ZnO

60

2.0~3.0

7.8~8.8

Additive in rubber industry

Si3N4

86-120

2.7~3.1

4.2

26

AUTHOR DECLARATION We wish to draw the attention of the Editor to the following facts which may be considered as potential conflicts of interest and to significant financial contributions to this work. We confirm that there are no known conflicts of interest associated with this publication and there has been no significant financial support for this work that could have influenced its outcome. We confirm that the manuscript has been read and approved by all named authors and that there are no other persons who satisfied the criteria for authorship but are not listed. We further confirm that the order of authors listed in the manuscript has been approved by all of us. We confirm that we have given due consideration to the protection of intellectual property associated with this work and that there are no impediments to publication, including the timing of publication, with respect to intellectual property. In so doing we confirm that we have followed the regulations of our institutions concerning intellectual property. We further confirm that any aspect of the work covered in this manuscript that has involved either experimental animals or human patients has been conducted with the ethical approval of all relevant bodies and that such approvals are acknowledged within the manuscript.

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