poly (vinylidene fluoride) nanocomposites

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Available at www.sciencedirect.com

ScienceDirect journal homepage: www.elsevier.com/locate/carbon

Exceptional dielectric properties of chlorine-doped graphene oxide/poly (vinylidene fluoride) nanocomposites Ying Wu, Xiuyi Lin, Xi Shen, Xinying Sun, Xu Liu, Zhenyu Wang, Jang-Kyo Kim

*

Department of Mechanical and Aerospace Engineering, The Hong Kong University of Science and Technology, Clear Water Bay, Kowloon, Hong Kong

A R T I C L E I N F O

A B S T R A C T

Article history:

Herein, we report a facile method to significantly enhance the dielectric performance of

Received 1 December 2014

reduced graphene oxide-based polymer composites. Addition of thionyl chloride into gra-

Accepted 28 February 2015

phene oxide (GO) dispersion induces synergistic modifications of the structure, chemistry,

Available online 5 March 2015

charge carrier density and electrical conductivity of GO, as well as the interfacial interaction and phase of the surrounding matrix in the poly (vinylidene fluoride) (PVDF) composite. The composites reinforced with a very low reduced chlorinated GO (Cl-rGO) content of 0.2 vol% deliver an exceptional dielectric constant of 364 with a moderate dielectric loss of 0.077 at 1 kHz. These values are well contrasted with the corresponding properties of the neat PVDF polymer with a constant of 28 and a loss of 0.0029. Synergistic effects arising from chlorination are identified, including the much enhanced electrical conductivity of Cl-GO sheets by more than 3 orders of magnitude through introducing charge-transfer complexes, the improved interfacial interactions between the fillers and the PVDF matrix through hydrogen bonds, and the transformation of PVDF to b-phase with an inherently high dielectric constant due to dipolar interaction. The comparison with the literature data confirms superior dielectric performance of the present Cl-rGO/PVDF composites. Ó 2015 Elsevier Ltd. All rights reserved.

1.

Introduction

Materials with high dielectric constants are widely applied in electronic and electromechanical systems, such as embedded capacitors in microelectronics, electromagnetic interference shielding and energy-storage devices. Many ceramic-, metaland polymer-based composites have been investigated to obtain high dielectric constants, among which polymer composites are most widely studied due to their easy processing, flexibility, high dielectric strengths and superior environmental resistance [1,2]. Typically, four different strategies have

* Corresponding author. E-mail address: [email protected] (J.-K. Kim). http://dx.doi.org/10.1016/j.carbon.2015.02.074 0008-6223/Ó 2015 Elsevier Ltd. All rights reserved.

been explored to develop polymer-based composites with excellent dielectric performance: namely, (i) selection of proper polymer matrices; (ii) incorporation of ceramic additives with high dielectric constants; (iii) addition of conductive organics like polymers or oligomers; and (iv) addition of inorganic conductive fillers [3–6]. Ferroelectric polymers, such as poly (vinylidene fluorine) (PVDF) and poly [(vinylidene fluoride)-co-trifluoroethylene] [P(VDF-TrFE)], possess outstanding dielectric constants in the polymeric family, but their constant values generally lower than 50 are far from enough for practical applications [7,8]. Ceramic additives, such as

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BaTiO3 and lead zirconate titanate (PZT), can significantly enhance the dielectric properties. However, a very large volume fraction is often needed to achieve high constants, which inevitably deteriorates the mechanical properties and dielectric strengths of the composites [9,10]. Addition of conductive organic materials, like polyaniline, showed excellent compatibilities and strong interfacial bonds with the host polymer matrices, resulting in low dielectric losses, but the enhancement of dielectric constant was limited because of their relatively poor electrical conductivities [7,11]. Metals and carbon materials are frequently used as inorganic fillers for the synthesis of composite materials, among which carbon nanotubes (CNTs), graphene and their derivatives have shown great potential. In particular, remarkably high dielectric constants of graphene/polymer composites, as high as 14,000 at 1 kHz, have been reported [12], but these high constants were at the expected expense of very high loss tangents, limiting their practical applications. Therefore, finding a new strategy is required to mitigate the concomitant increase in dielectric loss while maintaining the high constant values. Dielectric properties of composites with conductive fillers can be explained by the Maxwell–Wagner–Sillars (MWS) polarization principle where the nomadic charges are entrapped between the interfaces of different phases having different electrical conductivities [13,14]. Polymer composites possessing high dielectric constants, low dielectric losses and good process compatibility with printed circuit boards have received significant attention as promising candidate dielectric materials for embedded capacitor applications [15]. The dielectric constant of graphene/polymer composites can be maximized by creating a large area of interfaces between graphene and polymer through uniform dispersion of graphene. Reduced graphene oxide (rGO), rather than pristine graphene, is commonly used to achieve uniform dispersion because of its reasonably high electrical conductivity and acceptable hydrophility arising from the attached oxygenated functional groups [12]. Chemical doping further increases the electrical conductivity of rGO sheets by introducing extra charge carriers, either holes or electrons [16,17]. Alkali metals, such as potassium and rubidium, are typical electron donors while chlorine and bromine are popular electron acceptors [18–20]. Among various doping strategies developed for enhanced electrical conductivities, a chemical treatment by thionyl chloride (SOCl2) is known to be one of the most efficient choices [21,22]. The charge–transfer complexes formed on graphene oxide (GO) after chlorination improved the electrical conductivity [23–25]. Chlorinated rGO (Cl-rGO) sheets were found very stable in solvents like N,N-dimethylformamide (DMF) [26], possibly because of the enhanced repulsive forces between the GO sheets arising from the absorbed Cl ions. Thanks to these advantages, chlorination of CNTs and graphene have been widely studied, but their application in the synthesis of composites is rare. It is worth exploring the SOCl2-treated graphene for developing polymer composites possessing excellent dielectric constants and low dielectric losses. PVDF is a very attractive functional semi-crystalline polymer, exhibiting extraordinary piezoelectric and dielectric properties, outstanding physical and chemical stability and acceptable mechanical properties [27–29]. It has been widely

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used to prepare conducting composites for over-current protectors, antistatic shielding, strain sensors and high storage capacitors [30,31]. Five different polymorphs of PVDF have been observed, namely a, b, c, d and e phases. The multiple crystalline structures are attributed to the similar atomic radii of hydrogen and fluorine atoms in repeating units (–CH2–CF2–) which permit chain polarization while avoiding possible steric hindrance effects [32]. To achieve high dielectric constants along with low losses, we report for the first time on chlorine-doped GO (Cl-GO)/ polymer composites by treating GO dispersions directly with SOCl2 (Fig. 1). The significance of chlorine functionalization employed in the work is that the process is sonication-free, thus allowing us to maintain the large-size GO sheets to the benefit of enhanced dielectric constants. The nature of C–Cl bonds is evaluated by X-ray photoelectron spectroscopy (XPS), which is closely related to the charge carrier densities and the corresponding electrical conductivities of Cl-GO. The superior dielectric performance of the b-phase PVDF and the enhanced interfacial interactions in Cl-rGO/PVDF composites are demonstrated.

2.

Experimental

2.1.

Synthesis of GO and chlorinated GO/PVDF composites

GO was prepared using the modified Hummers method which is essentially the same as reported previously [33,34]. The asprepared GO dispersion with polydispersity was sorted to obtain large-size GO dispersion via centrifugation as described elsewhere [35]. The final GO slurry was diluted to a concentration of 1 mg/mL using N,N-dimethylformamide (DMF, Fisher) and ultrasonicated for 15 min using a bath sonicator. To prepare chlorinated rGO/PVDF composites, GO was functionalized by adding SOCl2 (Sigma–Aldrich) into the GO dispersion in the volume ratio of SOCl2:GO = 1:10 and the mixture was heated to 70 °C for different durations up to 6 h. The reacted solution was centrifuged to obtain Cl-GO slurry which was then re-dispersed in DMF. The Cl-GO dispersion was blended with DMF-diluted PVDF (Kynar, 1.78 g/cm3 in density) solution and stirred for 1 h to obtain a homogeneous

Fig. 1 – Simplified flowchart for the preparation of Cl-rGO/ PVDF composites. (A color version of this figure can be viewed online.)

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solution which was casted into a glass mold and cured at 80 °C for 24 h. The composite films were post-cured at 170 °C for 6 h after a 2 h reduction by hydrazine vapor at 95 °C, after which mildly reduced fillers (labeled as Cl-rGO) were obtained. 10 films were stacked together for compression at 200 °C for 30 min at a pressure of 100 MPa to form composite plates of 15 mm in thickness. The above process is schematically illustrated in Fig. 1. Neat PVDF and untreated rGO/PVDF composites were also prepared using the same procedure. Filler volume fraction in composites were estimated using the density of both rGO and Cl-rGO as 2.2 g/cm3 [36].

2.2.

Characterization

The images of GO sheets deposited on a quartz substrate were taken on a scanning electron microscope (SEM, JEOL JSM 6390) and were analyzed using the software ImageJ to measure the size distribution. X-ray photoelectron spectroscopy (XPS, Axis Ultra DLD) was used to characterize the elemental compositions of GO before and after chlorination. The fractions of chlorine in ionic and covalent states were specifically analyzed to identify an optimized degree of chlorination that could produce the most balanced properties of Cl-GO. Raman spectroscopy (Reinshaw MicroRaman/ Photoluminescence System) was used to characterize the changes in GO structure due to chlorination. The electrical conductivities and charge carrier densities of GO and Cl-GO were measured using the four-point probe method on a resistivity/Hall measurement system (Bio-Rad HL5500PC). Fourier transform infrared spectroscopy (FT-IR, Bruker Vertex 70 Hyperion 1000) was used to study the functional groups of

GO before and after chlorination, and thus to evaluate the interfacial interactions between Cl-rGO and PVDF matrix. The phase compositions of the rGO/PVDF and Cl-rGO/PVDF composites were examined by X-ray diffraction (XRD, PANalytical X’pert Pro) using Cu Ka1 (k = 0.154 nm) radiation. The dielectric properties of the composites with different filler contents were measured using an impedance/gain-phase analyzer (Hewlett Packard 4149A) in the range of frequencies from 100 Hz to 40 MHz. The crystallizing behaviors of the neat PVDF and composites with different filler contents were characterized by differential scanning calorimetry (DSC, Setaram 90/ 39324), and the crystallinities were calculated using the following equation [37]: XC ¼

DHm  100% DH0m  x

ð1Þ

where DHm is the melting enthalpy obtained during heating, DH0m is the melting enthalpy of 100% crystallized pure PVDF (104.7 J/g) [37], and x is the mass percentage of PVDF in the composites.

3.

Results and discussion

3.1. Nature of C–Cl bonds and electrical conductivities of chlorinated GO Large-size monolayer GO sheets with an average area of 477 lm2 (Fig. 2a) were synthesized from natural graphite flakes using the modified Hummers method. Raman spectroscopy was used to study the changes in GO structure after chlorination, as shown in Fig. 2b. The Raman spectra showed

Fig. 2 – (a) Size distribution of GO sheets and a typical SEM image in inset (scale bar: 10 lm); (b) Raman spectra of Cl-GO with different degrees of chlorination and the corresponding G-band shifts in inset; and (c) electrical conductivity and charge carrier density of Cl-GO as a function of chlorination duration. (A color version of this figure can be viewed online.)

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two major peaks, D- and G-bands, representing the disordered sp3 structure and the doubly degenerate zone center E2g mode, respectively. The G-band peaks of all Cl-GO underwent significant shifts compared to that of the untreated GO, showing a maximum (of 10 cm1, as shown in the inset) after 2 h of chlorination. The G-band shift is a consequence of charge transfer between the dopant and the GO sheets. No quantitative relationship between the G-band shift and doping concentration has been reported previously, and the trends for different doping methods were not identical [17,38]. It is worth noting, however, that the G-band position was qualitatively related to the increase in charge carrier density due to doping: a higher charge carrier density resulted in a larger shift [17,38]. Thus, it can be said that the treatment by SOCl2 for 2 h resulted in the most significant doping effect. The charge carrier densities and the corresponding electrical conductivities of GOs with and without chlorination are plotted in Fig. 2c. Both properties varied with chlorination duration in a functionally very similar manner, indicating a significant correlation between them. Two different types of C–Cl bonds existed in the chlorinated carbon, namely charge-transfer complexes and covalent bonds [20,39]. The formation of charge-transfer complexes generally increases the electrical conductivities by adding electrons to the conduction band or holes to the valence bond [40]. Here, the ionic Cl acted as an electron acceptor, i.e. p-type dopant, to enhance the charge carrier density. The maximum value of charge carrier density corresponded to 2 h chlorination, which coincided with the maximum G-band shift (Fig. 2b) as well as the maximum fraction of ionic Cl analyzed from the XPS results as discussed below. It is worth mentioning that

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2 h chlorination gave rise to remarkable 3 orders of magnitude improvements in both charge carrier density and electrical conductivity. Moreover, chlorine doping had a negligible influence on electron mobility while it enhanced the hole mobility of rGO sheets by several times [41]. The above observations were further confirmed by the XPS analysis, as shown in Fig. 3. The general spectra (Fig. 3a) contained obvious peaks of Cl 2p, Cl 2s and S 2p after chlorination, which were absent in GO. The atomic ratios of Cl to S as measured by the intensities of the Cl and S peaks were larger than 2 (Table 1) for all chlorination durations, suggesting chemical reactions between the dopant SOCl2 and GO, rather than simple intercalation of SOCl2. The N 1s peak located at 401.2 eV is assigned to the C–N bonding and attributed to DMF which was not completely evaporated after drying at 80 °C because of its high boiling point of 153 °C. Any remaining DMF was removed after the composites were post-cured at 170 °C for 6 h. The deconvoluted C 1s peaks for GO (Fig. 3b) consisted of C@C/C–C (284.5 eV), C–O (epoxy/hydroxyl, 286.5 eV), C@O (carbonyl, 287.8 eV) and COOH (carboxyl, 288.9 eV) groups [42–44]. A new peak for the C–S (286.1 eV) bonds [45] occurred in ClGO, while the C–Cl bonds at 286.5 eV [46] were not clearly seen because of the overlapping with the C–O groups. The binding energies of 2p spectra appeared as doublets associated with 2p3/2 and 2p1/2 levels due to the spin–orbit coupling separated by 1.6 and 1.95 eV for Cl and S, respectively [23] and the further details are discussed in the following. The two different C–Cl bonds (Fig. 3c) were quantitatively analyzed according to the previous studies [47]. The Cl 2p peaks were split into two pairs where the doublet located at the lower binding

Fig. 3 – XPS spectra of GO obtained before and after chlorination: (a) general XPS spectra of GOs with different chlorination durations; (b) deconvoluted spectra of C 1s of GO and Cl-GO after chlorination for 2 h; deconvoluted XPS curve fittings of (c) Cl 2p and (d) S 2p. (A color version of this figure can be viewed online.)

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Table 1 – Summary of elemental compositions of GO with different chlorination durations. Chlorination duration

0 1h 2h 3h 4h 5h 6h

Elements (at.%) O

C

Cl

S

Cl/S ratio

29.9 21.3 21.9 22.5 22.2 20.3 21.9

70.1 72.3 70.3 73.4 73.9 75.4 74.1

– 4.5 5.6 3.1 3.1 3.4 3.1

– 1.9 2.2 1.0 0.8 0.9 0.9

– 2.4 2.5 3.1 4.0 3.8 3.5

energies corresponds to the ionic state of Cl and that at the higher binding energies is covalently-bonded Cl. The peaks appeared at 201.4 and 199.8 eV represented the Cl–C@O bonds and covalent C–Cl bonds, respectively, while the doublet whose 2p1/2 peak located at 198.6 eV was a reflection of ionically bonded chlorine on the GO basal plane [26,39,48]. Similar rules were applied to the deconvoluted S 2p spectra (Fig. 3d) to reveal that the 2p3/2 peak appeared at 164.4 eV was assigned to the C–S bonds formed between the GO sheets whereas the 2p3/2 peak located at 168.9 eV was attributed to the SO2 4 groups [45]. The atomic concentration of Cl is plotted as a function of chlorination duration, as shown in Fig. 4. With increasing chlorination duration, the Cl atomic concentration initially increased to a maximum of 5.5 at.% after 2 h of chlorination, followed by a decrease to 3.3 at.% before saturation for further chlorination. This phenomenon can be explained by the equilibrium between the chemical reaction and the physical adsorption. It is well-known that SOCl2 molecules need to be in close contact with the GO sheets for interactions to occur between them [45], making both chemical reaction and physical adsorption with the GO sheets possible. When the chlorination process began, enough doping/reactive sites were available for both chemical reaction and physical

Fig. 4 – Atomic concentration of Cl in GO and fractions of Cl in ionic and covalent states as a function of chlorination duration. (Solid line is only a guide to the eyes for atomic concentration of Cl). (A color version of this figure can be viewed online.)

adsorption. The quick interaction between GO and SOCl2 resulted in a rapid increase in Cl atomic concentration on GO sheets, reaching a maximum after 2 h. When all the doping/reactive sites were occupied, a further treatment did not increase the Cl atomic concentration. Instead, de-doping may have taken place probably because of the transformation between the physical adsorption and chemical reaction [19], resulting in changes in the fraction of covalent- and ionicCl. With further chlorination, the atomic concentration of Cl in GO was reduced to a low constant value. A similar phenomenon was reported that the fraction of Br atomic concentration in graphene nanoplatelets reached a peak after 10 h of Br exposure and then reduced with further Br-treatment [19]. The fractions of the ionic Cl (two blue peaks in Fig. 3c) and the covalent Cl (two green peaks in Fig. 3c) were measured from the intensities of the corresponding peaks and are superimposed in Fig. 4. The fraction of ionic Cl showed a trend functionally analogous to the atomic concentration of Cl, and a maximum ionic fraction was obtained after 2 h of chlorination, very much consistent with the observations from the Raman spectroscopy (Fig. 2b) and the charge carrier density (Fig. 2c). As a result, the fraction of covalently-bonded Cl showed an opposite trend: it initially decreased and then gradually increased with treatment duration.

3.2. Interfacial interaction and polymorphs of Cl-rGO/ PVDF nanocomposites FT-IR analysis was conducted to study the interfacial reactions taking place between GO and SOCl2 as well as between Cl-rGO and PVDF in the composites, as shown in Fig. 5. There were a few prominent changes in the spectrum obtained after chlorination of GO (Fig. 5a): a new peak appeared at 710 cm1, a reflection of C–Cl bonds [49], confirming the doping of chlorine. The peak located at 1740 cm1 attenuated as a consequence of the reduction in C@O bonds in the carboxyl groups, while the area under the C–OH bonds also decreased after chlorination. Apparently, SOCl2 reacted with C–OH and COOH oxygenated groups in GO to form C–Cl and Cl–C@O groups. The broad banks at 3400 cm1, a characteristic of stretching vibration of O–H bonds, downshifted from 3415 to 3402 cm1 when chlorinated GO sheets were used (inset 2 in Fig. 5b). This indicates enhanced hydrogen bonds in the ClrGO/PVDF composite compared to the rGO/PVDF composite [50]. Strong hydrogen bonds were formed between the hydroxyl (–OH) groups in Cl-rGO and the F atoms in PVDF. The peak located at 870 cm1 assigned to >CF2 and the skeletal bending vibration underwent an up-shift from 871 to 874 cm1 (inset 1 in Fig. 5b), indicating the strong dipolar interaction between >CF2 groups in PVDF chains and the newly formed >C@O groups in chlorinated rGO sheets [32,51]. The improved interfacial interactions between the adsorbed F atoms and the functional groups on Cl-rGO resulted in the modification of chain conformation of the PVDF matrix at the interfacial region, as discussed below. Fig. 6a shows the XRD spectra of pristine GO, rGO/PVDF and Cl-rGO/PVDF composites. The pristine GO showed a peak ˚ between adjaat 2h = 10.33°, indicating a d-spacing of 8.55 A cent GO sheets. This means that the EG was fully exfoliated

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Fig. 5 – FT-IR spectra of (a) GO before and after chlorination; and (b) rGO/PVDF and Cl-rGO/PVDF composites with 0.4 vol% filler content. (A color version of this figure can be viewed online.)

into single-layer GO sheets. The characteristic peak of pristine GO disappeared either in rGO/PVDF or Cl-rGO/PVDF composites because of excellent dispersion of rGO or Cl-rGO sheets in the PVDF matrix. The XRD spectra in Fig. 6a show that the rGO/PVDF composite mainly consisted of a- and cphases, and a large amount of b-phase emerged in the ClrGO/PVDF composite. The inset shows that the crystallinity of the composites was improved by chlorination and with increasing filler content. It is thought that the rGO sheets served as nuclei for PVDF crystallization while chlorination enhanced their nucleation effect to a certain extent. The rGO/PVDF composite exhibited characteristic peaks a/c (0 2 0), a/c (1 1 0) and a (0 2 1)/c (0 2 2) at diffraction angles 2h = 18.30°, 20.04° and 26.56°, respectively, as well as a small peak of b (3 1 0) at 2h = 36.3° [52–54], suggesting predominant a/c phases together with a small fraction of b-phase. The ClrGO/PVDF composite had an additional small peak at 2h = 17.66° assigned to a (1 0 0), and two peaks at 2h = 20.7° and 36.3° corresponding to b (2 0 0)/(1 1 0) and b (3 1 0), respectively, both of which were intensified as a result of chlorination of rGO. Among the five crystals of PVDF (a, b, c, d and e), the aphase is the most commonly obtained polymorph, possessing a monoclinic unit cell with non-polar TGTG 0 chain conformation where T, G and G 0 represent trans, gauche+ and gauche conformation, respectively [32]. The b-phase has an orthorhombic unit cell in polar all trans (TTTT) conformation

and shows outstanding piezoelectric, pyroelectric and dielectric performances [55]. The c-phase also has a polar orthorhombic unit cell with a TTTGTTTG 0 chain conformation and can be produced during crystallization from the melt at high isothermal crystallization temperatures and solution-evaporation process at temperatures higher than 70 °C [56]. The schematic chain conformations of the aand b-phases are shown in Fig. 6b. The a-phase with a TGTG 0 conformation is thermodynamically more stable than the b-phase with a TTTT chain conformation, and there is a big energy barrier between them [57]. The direct formation of b-phase is rather difficult due to the high energy of the all-trans conformations, regardless of its slightly lower crystal lattice energy than that of the a-phase [58]. Therefore, special techniques have been developed to obtain the polar b-PVDF, like mechanical drawing at certain temperatures, electrospinning or the application of high electric field polarization [43]. Although the b-phase was difficult to form, a very small amount of b (3 1 0) appeared in rGO/PVDF even without special treatments because of the rGO-induced crystallinity and the chain confinement of PVDF itself [59]. It grew more prominent after the chlorination at the expense of a (0 2 1)/c (0 2 2). With increasing the fraction of b-phase, the piezoelectric, pyroelectric and dielectric properties of the polymer can be greatly enhanced. The strong b-phase peak observed in the Cl-rGO/ PVDF composite can be attributed to the improved hydrogen bond between the Cl-rGO sheets and the PVDF chains, as well

Fig. 6 – (a) XRD spectra of pristine GO, rGO/PVDF, and Cl-rGO/PVDF composites, as well as crystallinity of composites with different filler contents in inset; and (b) schematic views of TGTG 0 (a-phase) and TTTT type chains (b-phase) of PVDF. (A color version of this figure can be viewed online.)

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as the dipolar interaction between the C@O bonds and the F atoms. The enhanced interfacial interactions allowed further adsorption of F atoms onto the basal plane of Cl-rGO sheets, leading to TTTT conformation (b-phase) at certain parts of PVDF chains.

3.3. Electrical conductivity and dielectric performance of Cl-rGO/PVDF Dielectric properties of polymer-based nanocomposites are closely related to the intrinsic electrical conductivities of fillers. Cl-GO sheets produced after chlorination for 2 h having the superior doping effect were chosen for the fabrication of PVDF-based composites and their electrical conductivities and dielectric properties were studied in comparison with the corresponding properties of rGO/PVDF composites, as shown in Fig. 7. The AC conductivities of the rGO/PVDF composites were very low and similar to that of the neat polymer regardless of filler content and typically proportional to frequency because of the inherently low conductivities of mildly reduced GO due to the presence of many defects in the sp3 carbon structure. In contrast, the Cl-rGO/PVDF composites presented much higher conductivities, varying with filler content especially at low frequencies below 10 kHz. At filler contents of 0.32 vol% or above, the conductivities became much higher than those of the corresponding rGO/PVDF composites, especially at the low frequency domain where they remained almost constant. Thus, it can be assumed that conducting networks were formed near the percolation threshold between 0.2 and 0.32 vol% Cl-rGO [60]. As a result of chlorination, the dielectric constants of the Cl-rGO/PVDF composites were greatly improved with moderate increases in loss tangent (Fig. 7b and c). The Cl-rGO/

PVDF composite with 0.2 vol% Cl-rGO had a remarkable dielectric constant of 364 at 1 kHz, which was 78% and 13 times higher than those of the rGO/PVDF composite and the neat PVDF, respectively. The corresponding loss tangent values of the three materials were 0.077, 0.088 and 0.0029, respectively. A 0.4 vol% Cl-rGO gave rise to an extremely high dielectric constant of 7488, equivalent to nearly 270 times that of the neat PVDF, with a moderately increased loss tangent of 2.88 at 1 kHz. Both the dielectric constants and losses underwent rapid increases towards the low-frequency region, as a result of relaxation arising from interfacial polarization. The interfacial polarization refers to the charge accumulation at the interfaces of heterogeneous media of different electrical conductivities when an external electric field is applied, which is known as the MWS effect [61]. For the Cl-rGO/PVDF composites with filler contents of 0.2 vol% or lower and all rGO/PVDF composites involved in this study, the dielectric losses were very low in the low-frequency region and increased gradually with increasing frequency. This was attributed to the weak interfacial polarization because of the low filler content or the low filler electrical conductivities in the composites, under which condition the dielectric relaxation of PVDF dominates the loss behavior of the composites. According to the percolation theory [62], both the electrical conductivity and permittivity increase sharply in the vicinity of percolation threshold, which is clearly seen from Fig. 7a and b. Both the Cole–Cole equation given below and the dielectric loss peak can be used to estimate the relaxation time as a reflection of the rate of polarization [63,64]. e ¼ e0 þ ie00 ¼ e1 þ

e0  e1 1a

1 þ ði2pf kÞ

ð2Þ

Fig. 7 – Electrical conductivities and dielectric properties of rGO/PVDF and Cl-rGO/PVDF nanocomposites containing different filler contents: (a) AC conductivity; (b) dielectric constant; (c) dielectric loss with the inset plotted on a linear scale; and (d) dielectric properties at 1 kHz of composites with different filler contents. (A color version of this figure can be viewed online.)

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The equation relates the relaxation time, k, with the real and imaginary parts of the complex dielectric constant, e0 and e00 , at a frequency, f ; e0 and e1 are the DC dielectric constant and that measured at the high frequency limit, respectively; and a is a constant [64]. The relaxation time can be calculated by: k ¼ 1=2pf max

ð3Þ

where f max is the frequency at the dielectric loss peak [63,64]. According to the inset of Fig. 7c, the relaxation peaks of the Cl-rGO/PVDF composite with filler contents 0.2 vol% or lower and the rGO/PVDF composite appeared at 1:2  107 Hz, resulting in a similar relaxation time k 13 ns due to the predominant dipolar relaxation of the PVDF matrix [65]. For the Cl-rGO/PVDF composites with filler contents 0.32 and 0.4 vol%, there was no obvious peak over the whole frequency range studied, probably because of the more prominent interfacial polarization induced by the larger amount of conductive fillers. These dielectric properties obtained at 1 kHz are summarized in Fig. 7d where the performance between the two composites with and without chlorination can be clearly distinguished. It is seen that both the dielectric constant and the loss maintained very low in the rGO/PVDF composites regardless of filler content. The drastic increase in dielectric constant of the Cl-rGO/PVDF composite with increasing filler content was achieved at the expense of initially moderate and later large increases of dielectric loss. These two properties were affected by the electrical conductivity of fillers: a high conductivity had an ameliorating effect on dielectric constant whereas it also increased the dielectric loss. This means that these two dielectric parameters cannot be simultaneously ameliorated, but rather need to be balanced. The above exceptional dielectric properties of the Cl-rGO/ PVDF composites are attributed to synergistic effects of several important modifications brought by chlorination of GO sheets: namely (i) increased electrical conductivity due to ptype doping induced by the formation of charge–transfer complexes on Cl-rGO basal planes, (ii) the transformation from the dominating a-phase into the b-phase of PVDF matrix which favors higher dielectric constants, and (iii) the enhanced interfacial adhesion between the Cl-rGO and PVDF matrix through the hydrogen bonds between O–H in Cl-rGO and F in PVDF matrix [50] and strong dipolar interaction between C@O in Cl-rGO and F in PVDF. The high permanent dipole moment and high polarizability of the formed C– Cl bonds also favorably gave rise to an improved dielectric performance, as demonstrated previously [26]. Apart from the above ameliorating influences of GO chlorination, the large-size GO sheets employed in this study also positively affected the dielectric performance by forming a large area of interfaces with the polymer matrix, especially when these large graphene sheets are preferentially aligned in one direction [15]. The dielectric properties of the current Cl-rGO/PVDF composites are compared with those reported in the literature for composites with CNTs or rGO as reinforcement, as shown in Fig. 8 [66–71]. In general, the dielectric constants of all these composites increased with increasing filler content, which was always accompanied by increases in dielectric loss.

Fig. 8 – Comparison of dielectric properties of polymerbased composites with CNTs or rGO as reinforcement, showing the relationship between dielectric constant and loss. In the brackets are the ranges of volume fractions of fillers reported in the references. The arrow in the top indicates the direction of increasing filler content. (A color version of this figure can be viewed online.) Materials with the data points lying in the top-left region are considered ideal for high-k and low-loss applications, such as energy storage and embedded capacitors. A much smaller Cl-rGO content was needed in this study to achieve a dielectric constant comparable to the other composites, probably because of the large aspect ratio and better dispersion of the Cl-rGO sheets used. It can be said that the present Cl-rGO/PVDF composites delivered dielectric properties among the best with a combination of exceptionally high dielectric constants and reasonably low dielectric losses.

4.

Conclusion

This paper reports the effect of chlorination of GO on dielectric properties of Cl-rGO/PVDF composites in comparison with those of the neat PVDF and rGO/PVDF composites without chlorination. The observed results are explained in terms of the modifications in structure, elemental composition, electrical conductivity, Cl-rGO/PVDF interfacial interaction and PVDF phase transformation. The following can be highlighted from the experimental study: (i) The doping effect of SOCl2 on GO sheets was p-type, and 2 h chlorination gave rise to the maximum ionic Cl of 4.3 at.% formed on GO sheets, which coincided with the highest charge carrier density and electrical conductivity as well as the maximum G-band shift. Remarkable 3 orders of magnitude improvements in both charge carrier density and electrical conductivity were achieved compared to the untreated GO. (ii) The enhanced interfacial interactions between the ClrGO sheets and PVDF chains due to chlorination allowed the adsorption of F atoms onto the basal plane of GO sheets. This led to the transformation of PVDF into b-phase with TTTT chain conformation which had an inherently higher dielectric constant than the other phases of PVDF. (iii) The Cl-rGO/PVDF composites delivered exceptionally high dielectric constants and moderately increased loss tangents with 0.2–0.4 vol% graphene: e.g. a dielectric

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constant of 7488 – equivalent to almost 270 times that of the neat PVDF – with a moderately increased loss tangent of 2.88 at 1 kHz with 0.4 vol% graphene. These values are among the best dielectric performance ever reported to date for composites containing graphene and CNTs. These excellent dielectric properties of the Cl-rGO/PVDF composites are attributed to synergistic effects of several important modifications arising from chlorination of GO sheets: namely the much enhanced electrical conductivities of Cl-rGO sheets and their composites due to the p-type Cl-doping; the induced b-phase of PVDF matrix; the enhanced interfacial adhesion between Cl-rGO and PVDF matrix; and the Cl–C bonds with high polarity and polarizability. The large area of interfaces created by the large GO sheets used in this study also contributed to a certain extent.

Acknowledgments This project was financially supported by the Research Grants Council of Hong Kong SAR. Technical assistance from the Materials Characterization and Preparation Facilities (MCPF), Advanced Engineering Material Facility (AEMF) and Department of Electronic and Computer Engineering (Prof. Philip K.T. Mok) of HKUST is appreciated. R E F E R E N C E S

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