PVDF composites prepared by an electrospinning-hot pressing method

Composites Science and Technology 172 (2019) 58–65

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Composites Science and Technology journal homepage: www.elsevier.com/locate/compscitech

Enhanced dielectric permittivity in surface-modified graphene/PVDF composites prepared by an electrospinning-hot pressing method

T

Bo Lina, Zeng-Tian Lia, Ying Yanga, Ying Lib, Jie-Ci Lina, Xu-Min Zhenga, Fu-An Hea,∗, Kwok-Ho Lamb,∗∗ a b

Institute of Chemical Engineering, Guangdong University of Petrochemical Technology, Maoming 525000, China Department of Electrical Engineering, The Hong Kong Polytechnic University, Hung Hom, Kowloon, Hong Kong

A R T I C LE I N FO

A B S T R A C T

Keywords: Surface-modified graphene/PVDF composite Electrospinning-hot pressing method Dielectric performance Thermal conductivity

In the present work, the surface-modified graphene (SMG)/poly(vinylidene fluoride) (PVDF) fibrous membranes obtained from the electrospinning were treated by the hot pressing in the laminating mode to form the SMG/ PVDF composites. The SMG was prepared by subjecting the graphene oxide to silane modification, NaBH4 reduction, and PVDF grafting in sequence. The successful surface modification of graphene was confirmed by TEM, XPS, Raman spectroscopy, FTIR, WAXD, and TGA. Furthermore, the structures of SMG/PVDF composites fabricated by the electrospinning-hot pressing method were studied by SEM, FTIR, and WAXD, which exhibited the well dispersion of SMG in the PVDF matrix. Finally, the investigation showed that the dielectric permittivities of SMG/PVDF composites increased with the SMG content, which were significantly higher than that of pristine PVDF. The dielectric permittivity of SMG (16 wt%)/PVDF composite (83.8) at 1000 Hz was found to be ten-fold that of the corresponding value of pristine PVDF (8.3) with a relatively low dielectric loss factor (0.34) and a relatively high thermal conductivity (0.679 W/mK).

1. Introduction Flexible electroactive materials play an important role in the electronic applications such as sophisticated sensor, actuator, and capacitor devices [1,2]. Poly(vinylidene fluoride) (PVDF), as a soft polymer, has been received considerable attention owing to its outstanding pyroelectricity, piezoelectricity, ferroelectricity, as well as relatively high dielectric permittivity [3–12]. In particular, many researches are still focusing on the enhancement of dielectric permittivity for PVDF, which gives great potential in the applications of embedded capacitor, highenergy-density storage, and so on [13,14]. The commonly used method to achieve high-dielectric-permittivity polymer-based materials is the incorporation of functional fillers into a polymer matrix. Initially, the high-dielectric-permittivity ceramics, such as BaTiO3 and PbZrTiO3, were added into polymer [15,16]. However, the enhancement of dielectric permittivities for such composites was limited unless the concentration of ceramic was beyond 40 vol% [13,14]. Recently, based on the percolation theory, different conductive filler/polymer composites containing xGNPs, CNTs, silver, and Ni have been developed while their dielectric permittivity could be

∗

enhanced significantly near the percolation threshold [13,14,17–19]. Graphene has been regarded as an ideal filler for the high-dielectricpermittivity composite owing to its excellent electrical and thermal conductivity as well as high aspect ratio [20–24]. On the other hand, as shown in the following equation,

ε ∝ εm (fc − f )−q

(1)

a high dielectric permittivity of conductive filler (ε) can be achieved if the dielectric permittivity of polymer matrix (εm ) is high. PVDF has a relatively high dielectric permittivity of about 10 in comparison with other polymers, which is excellent for being a matrix of high-dielectricpermittivity composites. Besides the properties of filler and matrix, the good dielectric performance of graphene/PVDF composite may also highly depend on (1) the uniform non-contacting distribution of graphene in the polymer matrix because of the poor compatibility of organic PVDF and carboneous graphene, and (2) the high electrically resistive interface between polymer and graphene, resulting in low dielectric loss [21]. It is well known that surface modification can adjust the interfacial interaction between functional fillers and polymer matrixes [21]. On the other hand, electrospinning is a simple but effective technique to

Corresponding author. Corresponding author. E-mail addresses: [email protected] (F.-A. He), [email protected] (K.-H. Lam).

∗∗

https://doi.org/10.1016/j.compscitech.2019.01.003 Received 4 November 2018; Received in revised form 3 January 2019; Accepted 5 January 2019 Available online 11 January 2019 0266-3538/ © 2019 Elsevier Ltd. All rights reserved.

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resulting mixture was treated with ultrasonication for 0.5 hour and stirring for 3 hours at 70 °C to form the homogeneous SMG/PVDF solution for electrospinning. (2) The above SMG/PVDF solution was put into a syringe for electrospinning by adjusting the following parameters: the positive and negative loaded voltages (+15 kV and −2 kV, respectively), the distance of collection (10 cm), and the rotation speed of collecting drum (2800 rpm). The resulting electrospun SMG/PVDF membranes were dried at 85 °C under vacuum for 24 hours. (3) The SMG/PVDF composites were obtained from the hot pressing method as follows. The electrospun SMG/PVDF membranes were cut into small pieces and then several pieces of SMG/PVDF membrane were laminated in the warp-warp arrangement [20–22]. Finally, the laminated electrospun SMG/PVDF membranes were compressed (200 °C, 12–14 MPa, 4 minutes) to prepare the SMG/PVDF composites.

fabricate membranes containing well-dispersed non-contacting functional fillers with good alignment [25–29]. Recently, Gu et al. have successfully prepared boron nitride/polyimide composites and silicon carbide particle/polystyrene composites with the enhanced thermal conductivity by the electrospinning-hot press technique, which provides a significant guidance for the fabrication of high-performance functional filler/polymer composites [25–29]. Currently, there are few reports on the study of the high-dielectric-permittivity graphene/PVDF composite prepared by the combined method of surface modification and electrospinning. We have previously reported the fabrication of dielectric composite basing on PVDF and PVDF-grafted graphene by the electrospinning-hot pressing method [30]. However, the dielectric permittivity of such a composite was relatively low of 32.7 at 1000 Hz near percolation threshold probably because the PVDF-grafted graphene obtained from the harsh Friedel-Crafts reaction hasn't been fully modified. Moreover, in that work, the removal of AlCl3 after FriedelCrafts reaction was complicated. For the purpose of avoiding these problems, in the present work, surface modified graphene (SMG) was prepared by a convenient radical-initiated reaction using a nonionic benzoyl peroxide (BPO) to graft the PVDF chains on the graphene surface. Next, SMG/PVDF fibrous membranes with different amounts of SMG were obtained from electrospinning and then treated by the hot pressing in laminating arrangement to form the SMG/PVDF composites, and the dielectric performance of the resulting SMG/PVDF composites were investigated.

2.4. Characterization The JEOL JEM-2010 transmission electron microscopy (TEM) was utilized to investigate the microstructures of GO, amino-GO, amino-G, and SMG with an acceleration voltage of 200 kV. The Zeiss Merlin scanning electron microscopy (SEM) was utilized to investigate the morphologies of the electrospun SMG/PVDF membranes and the SMG/ PVDF composites with an acceleration voltage of 10 kV. The samples were coated with a thin gold layer before the SEM measurement. The Nicolet/Nexus 670 spectrometer were utilized to obtain the Fouriertransform infrared (FTIR) spectra of the samples from 700 to 1500 cm−1. The Bruker D8 diffractometer was utilized to obtain the Wide-angle X-ray diffraction (WAXD) patterns of the samples with the 2 theta ranging from 2 theta = 10–35 degree and the scan rate of 1 degree/minute. The Shimadzu DTG-60A Thermal Gravimetric Analyzer (TGA) was utilized to investigate the thermal stabilities of the samples (N2 atmosphere, temperature range = 30–800 °C, heating rate = 10 °C/ minute). The Renishaw 2000 equipment was utilized to obtain the Raman spectra of the samples (laser wavelength = 514 nm, grating = 1800 lines/mm, exposure time = 10 s, slit = 50 μm). The Thermal Scientific Escalab 250Xi equipment was utilized to conduct the X-ray photoelectron spectroscopy (XPS) measurements (hυ = 1486.6 eV, power = 150 W, beam = 500 μm). The Agilent 4294A impedance analyzer was utilized to investigate the dielectric properties of the pristine PVDF and the SMG/PVDF composites from 103 Hz to 107 Hz. The TPS2500 thermal conductivity equipment (Hot Disk, Germany) was utilized to measure the thermal conductivities of the pristine PVDF and the SMG/PVDF composites.

2. Experimental 2.1. Materials Graphene oxide (GO) was prepared according to our previous works [18,19]. PVDF with the trademark of Solef 6008 was provided by the Shanghai branch of the Solvay Company, China. 3-aminopropyltriethoxysilane (APTES), sodium tetrahydroborate (NaBH4), N,N-dimethylformamide (DMF), absolute ethanol, BPO, acetone were provided by the Guangzhou Chemical Company, China. 2.2. Preparation of amino-GO, amino-G, and SMG The preparation of the amino-GO was as follow. The mixed solution containing 2 g GO, 0.2 g APTES, 10 mL distilled water, and 190 mL absolute ethanol was ultrasonicated for 0.5 hour and then stirred for 24 hours at 60 °C. The resulting mixture was filtered, washed with distilled water for several times, and dried at 60 °C under vacuum for 24 hours to obtain the amino-GO. The preparation of the amino-G was as follow. 2 g as-prepared amino-GO was dispersed in 200 mL distilled water by ultrasonication for 0.5 hour followed by the addition of 10 g NaBH4 slowly at room temperature. Next, the reduction of the amino-GO was carried out by stirring for 24 hours at 90 °C. The resulting mixture was filtered, washed with distilled water for several times, and dried at 60 °C under vacuum for 24 hours to obtain the amino-G. The preparation of the SMG was as follow. 2 g as-prepared amino-G was dispersed in 200 mL DMF by ultrasonication followed by the addition of 2 g PVDF and 0.1 g BPO. Next, the grafting of PVDF on aminoG was carried out by stirring for 6 hours at 85 °C in N2 atmosphere. The resulting mixture was added into 200 mL distilled water, filtered, washed with distilled water for several times, and dried at 60 °C under vacuum for 24 hours to obtain the SMG.

3. Results and discussion 3.1. Characterization of amino-GO, amino-G, and SMG The morphologies of the GO, the amino-GO, and the amino-G were investigated by TEM. GO has the high-transparence and wrinkle-surface morphology, as shown in Fig. 1a. After surface modification by APTES and reduction by NaBH4 in sequence, no obvious morphological changes for resulting amino-GO (see Fig. 1b) and amino-G (see Fig. 1c) can be observed. Hence, to confirm the structures of GO, amino-GO, and amino-G, they were further studied by FTIR, TGA (see Fig. S1), Raman spectroscopy, WAXD, and XPS. Fig. 2 gives the FTIR spectra of the GO, the amino-GO, and the amino-G. It can be found that the GO exhibits characteristic FTIR peaks of oxygenated groups at 3428 cm−1 for O–H, at 1726 cm−1 for C]O, at 1386 cm−1 for C–OH, at 1195 cm−1 for C–O–C, as well as at 1041 cm−1 for C–O, respectively [31]. After modification by APTES, the broadened FTIR peak of the resulting amino-GO in comparison with GO at 1080 cm−1 should be ascribed to Si–O–Si group and Si–O–C group. After reduction by NaBH4, the characteristic peaks at 2930 cm−1 and 2850 cm−1 for –CH2– as well as at 1465 cm−1 for N–H can be identified from the FTIR spectrum of the resulting amino-G due to the existence of

2.3. Preparation of SMG/PVDF composites The preparation process of SMG/PVDF composites by an electrospinning-hot pressing method is illustrated in Scheme 1. (1) The desired amounts of SMG and PVDF were blended together and then added into the mixed solvent of DMF (70 wt%) and acetone (30 wt%). Next, the 59

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Scheme 1. Preparation process of SMG/PVDF composites by an electrospinning-hot pressing method.

APTES. On the other hand, the FTIR peak intensities of the amino-G for oxygenated groups are clearly lower than those of the GO as a result of the successful reduction [23,24]. All of the Raman spectra for the GO, the amino-GO, and the aminoG (see Fig. 3) exhibited obvious peaks of D band at about 1350 and G band at about 1590 cm−1. The intensity ratio of D band to G band (ID/ IG) has been previously used to study the graphene structure [32,33]. In this work, the ID/IG values of the GO, the amino-GO, and the amino-G were calculated to be 1.48, 1.62, and 1.71, respectively. The GO had an obvious lower ID/IG value than those of the amino-G and the amino-GO, which could be ascribed to the formation of smaller in-plane sp2 domains during the modification process of GO by the APTES and the NaBH4, suggesting that many oxygenated groups have been removed in the amino-G and the amino-GO [32,33]. The WAXD patterns of the GO, the amino-GO, and the amino-G are displayed in Fig. 4. Due to the existence of lots of oxygenated functional group on its surface, GO had a WAXD diffraction peak of (001) crystal plane at 2 theta = 10.8°, which was lower than that of the as-reported

Fig. 2. FTIR results of GO, amino-GO, amino-G, and SMG.

Fig. 1. TEM images of (a) GO, (b) amino-GO, (c) amino-G, and (d) SMG. 60

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which the BPO attacked both amino-G and PVDF to produce radicals, resulting in the grafting of PVDF on the amino-G. A possible mechanism for the grafting of PVDF on the skeleton of the amino-G is shown in Fig. S2. According to the SEM image (see Fig. 1d), the transparency of the SMG was reduced in comparison with the amino-G, which is probably due to the coverage of PVDF on the graphene surface. The newly occurred peak of Raman spectrum for SMG compared to amino-GO should be ascribed to the existence of PVDF (see Fig. 3). Similarly, the FTIR spectrum of SMG exhibits the characteristic peaks of PVDF at 1276 cm−1, 880 cm−1, and 839 cm−1 [5], respectively, which confirmed the successful attachment of PVDF on graphene (see Fig. 2). Moreover, the WAXD pattern of the SMG also simultaneously possessed the characteristic diffractive peaks of graphene at 2 theta = 24.6° and PVDF at 2 theta = 20.5° (see Fig. 4).

Fig. 3. Raman results of GO, amino-GO, amino-G, and SMG.

3.2. Morphology and structure of SMG/PVDF composites The solutions containing acetone, DMF, PVDF, and SMG with different mass ratios were electrospun to form fibrous membranes. Since the rotation speed of rotating disk collector reached as high as 2800 rpm, most of the fibers in the electrospun SMG/PVDF membrane could be effectively aligned in the rolling direction even after a long time of electrospinning collection, as shown in the SEM images at low and high magnifications (see Fig. 6). However, many randomly distributed fibers in the electrospun SMG/PVDF membrane could still be found, which was similar to the work of the aligned electrospun pristine PVDF membrane reported by Yee et al. [35]. They suggested that, as the electrospinning time increased, a number of the residual charges accumulated on the surface of electrospun fibers, which could repel the fibers with one another by electrostatic repelling, resulting in the random distribution of some fibers from the rotating direction. Moreover, it seems that the amino-G graphene should have been embedded in the electrospun SMG/PVDF membrane with well dispersion, which would be helpful in the formation of microcapacitors after the hotpressing treatment in the laminating arrangement particularly for the improvement of dielectric permittivity of the resulting solid SMG/PVDF film. The SEM images of cross-sectional surface for SMG/PVDF film at low and high magnifications are shown in Fig. 7. It can be seen that the porous electrospun fibrous SMG/PVDF membrane has been pressed into a solid SMG/PVDF film after hot-pressing in the warp-warp laminating arrangement. Since most of SMGs were embedded deeply in the PVDF matrix, some SMGs could only be observed in the SEM image of SMG/ PVDF film at high magnification (see Fig. 7b). The FTIR spectra of the pristine PVDF film obtained from direct hot pressing, and the as-prepared SMG/PVDF composite films are displayed in Fig. 8. The characteristic peaks of nonpolar α phase crystalline can be found in the FTIR spectrum of the pristine PVDF film at 764 cm−1, 798 cm−1, 854 cm−1, 974 cm−1, 1212 cm−1, and 1380 cm−1 [5], respectively. As well known, the most useful crystal form of PVDF with excellent piezoelectricity, pyroelectricity, and ferroeletricity is its β phase crystalline. For all SMG/PVDF composite films, the intensities of their FTIR peaks at 1276 cm−1 and 839 cm−1 originating from the polar β phase crystalline increased with the SMG content, suggesting that the incorporation of SMG into PVDF matrix could partly transfer the α phase of PVDF to the useful β one. It was possible that the β phase crystalline with TTTT sequence could be formed from the interfacial interaction between PVDF and graphene including the π-dipole interaction and hydrogen bonding (see Fig. S3) [35]. In particular, when the SMG content was 16 wt%, the polar β phase crystalline of PVDF has already become the dominated crystal form in the SMG/PVDF composite film because of the existence of a great deal of interfacial interaction. Furthermore, the content of β phase for different samples were calculated as follow [36]:

Fig. 4. WAXD patterns of GO, amino-GO, amino-G, and SMG.

natural graphite at 2 theta = 26.6°. After the treatment of surface modification, the WAXD diffraction peak of (001) crystal plane for the resulting amino-GO slightly decreased to 2 theta = 10.4°, which is probably due to the intercalation of APTES. After the reduction by NaBH4, the WAXD diffraction peak of (001) crystal plane for the resulting amino-G significantly increased to 2 theta = 24.6°. It means that the crystal structure of graphene for the amino-G had been mostly recovered [24]. The XPS C1s spectra of the GO, the amino-GO, and the amino-G along with their peak-fitting curves are shown in Fig. 5, and the corresponding contents of different carbon components according to the peak-fitting curves are listed in Table 1. It can be found that, in addition to the C]C(sp2)/C–C (sp3) non-oxygenated carbon component, GO had a lot of oxygenated carbon components in relation to different functional groups such as C–OH, C–O–C, C]O, and O–C]O [22]. After the modification by the APTES, the higher atomic percentage of C]C/C–C and the lower content of oxygenated carbon components of amino-GO compared to GO can be attributed to the reduction effect of the APTES, as reported by Yang et al. [34]. For the amino-G, the oxygenated carbon components decreased significantly and an obvious increment of C]C/ C–C component from 15.3% to 71.1% could also be found as a result of the reduction by NaBH4, implying most of the oxygenated groups had been removed. The resulting amino-G was further modified by PVDF in DMF solvent using BPO to obtain the SMG. Since the amino-G contained NH2 groups, it could disperse well in DMF during the reaction process in 61

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Fig. 5. XPS results and XPS fitting curves of GO, amino-GO, amino-G, and SMG. Table 1 The fitting XPS peak positions of C 1s and the relative atomic percentage of various functional groups in GO, amino-GO, and amino-G. Sample

C-C/C]C (eV) (%)

C-OH (eV) (%)

C-O-C (eV) (%)

C=O (eV) (%)

O-C]O (eV) (%)

GO

284.7 15.3 284.8 31.5 284.6 71.1

285.4 20.9 285.5 11.4 285.5 9.5

286.5 19.5 286.6 13.3 286.4 9.3

287.4 36.0 287.5 35.4. 287.5 5.8

288.8 8.3 289.0 8.4 288.8 4.3

amino-GO amino-G

F (β ) =

Aβ (K β / K α ) Aα + Aβ

× 100%

(2)

where F (β ) : the content of β phase crystalline; Aα : the absorbance at 764 cm−1; Aβ : the absorbance at 839 cm−1; K α (6.1 × 104 cm2/mol) and K β (7.7 × 104 cm2/mol) are the absorption coefficients at the corresponding wavenumber. Accordingly, the F (β ) of pristine PVDF, SMG(4 wt%)/PVDF composite, SMG(8 wt%)/PVDF composite, SMG(12 wt%)/PVDF composite, and SMG(16 wt%)/PVDF composite were calculated to be 0.22.8%, 49.5%, 52.0%, 75.7%, and 83.2%. Moreover, the F (β ) of SMG (16 wt%)/PVDF composite was higher than that (78.5%) of the electrospun SMG(16 wt%)/PVDF membrane (see Fig. S4). It was possible that the SMG had better ability to facilitate the formation of β phase crystalline in SMG(16 wt%)/PVDF composite

Fig. 6. SEM images of electrospun SMG (16 wt%)/PVDF membrane at (a) low and (b) high magnification. 62

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Fig. 7. SEM images of cross-section surface for SMG (16 wt%)/PVDF film at (a) low and (b) high magnification.

of PVDF molecular chains into SMG and un-intercalated SMG, correspondingly. 3.3. Dielectric properties and thermal conductivity of SMG/PVDF composites The dependency of the dielectric permittivity (ε) and the dielectric loss factor (tanδ) at frequency from 50 Hz to 107 Hz for the pristine PVDF and its composites at room temperature were investigated, as displayed in Fig. 10a and Fig. 10b, respectively. Owing to the declined moving abilities of the dipoles and the charge carriers, ε of the pristine PVDF and its composites decreased with the increasing frequency (see Fig. 10a). It seems that ε of SMG/PVDF composite containing 16 wt% amino-G exhibited a very strong frequency-dependence behavior particularly in the low-frequency range because the existing of a lot of interfaces between graphene and PVDF would result in a strong Maxwell-Wagner-Sillars interfacial polarization [3]. More importantly, the ε values of SMG/PVDF composites increased with the amino-G content, which were much greater than that of the pristine PVDF. For example, ε of the pristine PVDF at 1000 Hz was only 8.3 (see Fig. 10c), while those of the SMG/PVDF composites containing 4 wt%, 8 wt%, 12 wt%, and 16 wt% SMG were 12.4, 22.2, 44.4, and 83.8, respectively, which were ∼1.5-fold, 2.6-fold, 5-fold, and 10-fold, respectively, that of the pristine PVDF. Such a significant enhancement of ε for the SMG/PVDF composites compared to the pristine PVDF can be attributed to (1) the effective formation of a large number of conductive graphene/insulative PVDF/conductive graphene microcapacitor structures, as shown in Scheme 1, and (2) the outstanding Maxwell-Wagner-Sillars interfacial polarization caused by the existence of a lot of interface [3]. It has been reported that the necessary loading amounts of commonly used ceramics for high-dielectric-permittivity polymeric composites should be usually over 40 vol%. For instance, ε increased from ∼10 for pure poly(vinylidene fluoride-trifluoroethylene) [P(VDF-TrFE)] to ∼32 and ∼50 for P(VDF-TrFE)-based composite filled with 30 vol% (Bi0.5Na0.5)0.94Ba0.06TiO3 and 40 vol% PMN-PT ceramics at 1000 Hz [15,16], respectively. While for the SMG/PVDF composite, the actual graphene content was only 8 wt% even with high SMG filling content of 16 wt%, which is beneficial to remain the flexibility of the resulting SMG/PVDF composite. On the other hand, in addition to ε, tanδ is another parameter with important meaning for the utilization of the high-dielectric-permittivity materials. It can be seen that the tanδ values of the SMG/PVDF composites were higher than that of the pristine PVDF particularly in the low-frequency range (see Fig. 10b), but still could be accepted for the practical application [37]. For instance, the dielectric loss factors of the SMG/PVDF composites containing 12 wt% and 16 wt% SMG at 1000 Hz were only 0.25 and 0.34, respectively. Based on the percolation theory, when the content of conductive filler is beyond the percolation threshold, the tanδ value of conductive filler/ PVDF composite will be seriously high. For comparison, the dielectric

Fig. 8. FTIR results of PVDF and its composites.

subjected to hot-pressing than the electrospinning process because the PVDF molecular chains could move more easily at high temperature to crystallize. The WAXD results of the pristine PVDF film obtained from direct hot pressing and the as-prepared SMG/PVDF composite films are shown in Fig. 9. The WAXD pattern of PVDF exhibited four diffractive peaks at 2 theta = 17.9°, 18.5°, 20.1°, and 26.8° which were in relation to the α phase crystalline. While for all SMG/PVDF films, their WAXD peak intensities of α phase crystalline reduced and the WAXD diffractive peak at 2 theta = 20.4° corresponding to β phase crystalline of PVDF increased gradually with the SMG content [5]. In addition, the WAXD peaks at about 20.9° and 24.3° should be attributed to the intercalation

Fig. 9. WAXD patterns of PVDF and its composites. 63

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Fig. 10. Dependencies of (a) dielectric permittivity and (b) dielectric loss factor on frequency from 50 Hz to 107 Hz for pristine PVDF and its composites, (c) dielectric constants and dielectric loss factors at 1000 Hz for pristine PVDF and its composites, (d) thermal conductivities for pristine PVDF and its composites.

high-content fillers, such as flaky Al (40 vol%)/PVDF composite with 0.8 W/mK, β-SiC (20 vol%)/PVDF with 0.75 W/mK, and ZnO (40 wt %)/PVDF with 0.55 W/mK [38–40]. The relatively high thermal conductivity of the SMG/PVDF composite can be ascribed to the high thermal conductivity of the SMG filler and the low interfacial thermal resistance between SMG and PVDF owing to their good compatibility [28,41]. This good compatibility originated from the interfacial interaction between the PVDF matrix and the grafted PVDF on the SMG surface. Accordingly, the existing of the SMG filler in the PVDF matrix can not only enhance the dielectric permittivity of the resulting SMG/ PVDF composite with the relatively low dielectric loss factor but also effectively dissipate the heat generated from the dielectric loss.

properties of reduced graphene oxide (RGO)/PVDF composite near the percolation threshold with RGO content of 1.8 wt% was studied. It can be observed that the resulting RGO(1.8 wt%)/PVDF composite exhibited lower ε value of 37.0 and larger tanδ value of 0.73 at 1000 Hz when compared to the SMG/PVDF composites containing 12 wt% and 16 wt% SMG (see Fig. 10c). The explanation of this phenomenon was as follows: with the help of the surface modification and electrospinning, a high filling content of well-dispersed graphene in the SMG/PVDF composites could be achieved, resulting in not only the formation of plenty of microcapacitors but also the enhanced MWS interfacial polarization as a result of more interface area. Moreover, the insulating PVDF grafted on graphene surface could impede the direct touch of the conductive SMG with one another by the shielding effect of the grafted insulative PVDF and then reduced the dielectric loss in the SMG/PVDF composites significantly. High thermal dissipation also plays an important role in the application of high-dielectric-permittivity electrically-conductive filler/ polymer composite due to their relatively high dielectric loss factor. It is fortune that the addition of electrically conductive filler into the polymer matrix usually can also improve the thermal conductivity of the resulting composite. The thermal conductivities of the pristine PVDF and RGO (1.8 wt%)/PVDF composite are 0.196 W/mK and 0.245 W/mK (see Fig. 10d), respectively, indicating that the electrically conductive RGO filler was also thermally conductive with the ability to transfer the heat. While for the SMG/PVDF composites, their thermal conductivities increased with the SMG content. Especially, the thermal conductivities of the SMG/PVDF composites containing 12 wt% and 16 wt% SMG reached as high as 0.466 W/mK and 0.679 W/mK, respectively. This result is comparable to or even better than those reported thermal conductivities of PVDF-based composites containing

4. Conclusions Novel SMG/PVDF composites have been obtained from the incorporation of the PVDF-surface-modified graphene into the PVDF matrix by the combination of electrospinning and hot pressing. Owing to the shielding effect of the PVDF coating layer on the graphene filler and the well distribution of the PVDF-surface-modified graphene in the PVDF matrix, it could be found that the as-prepared SMG/PVDF composites displayed much better dielectric performance in comparison with those of pristine PVDF and RGO/PVDF composite near the percolation threshold. Especially, when the SMG content was 16 wt%, the dielectric permittivity of the SMG/PVDF composite reached as high as 83.8 with a relatively low dielectric loss factor of 0.34 at 1000 Hz and a relatively high thermal conductivity of 0.679 W/mK, which could be regarded as a potential material for the applications in high-charge storage and embedded capacitor. Moreover, the strategy demonstrated in this work provides an effective way for the construction of other 64

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advanced graphene/polymer composites. Acknowledgments

[19]

The authors would like to acknowledge the support of the Guangdong province natural science foundation, China (2014A030307037, 2016A030310308, 2017A030313268, and 2017A030313080), the college student training program of Guangdong University of Petrochemical Technology (2017pyA012), national innovation and entrepreneurship training program for college students, China (201711656003), and the Hong Kong Polytechnic University (1ZVGH, G-YBLM, and G-YBPN).

[20]

[21]

[22]

Appendix A. Supplementary data

[23]

Supplementary data to this article can be found online at https:// doi.org/10.1016/j.compscitech.2019.01.003.

[24]

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