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High Dielectric and Mechanical Properties Achieved in Crosslinked PVDF/#-SiC Nanocomposites with Elevated Compatibility and Induced Polarization at Interface Yefeng Feng, Bei Miao, Honghong Gong, Yunchuan Xie, Xiaoyong Wei, and Zhicheng Zhang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b04776 • Publication Date (Web): 05 Jul 2016 Downloaded from http://pubs.acs.org on July 9, 2016

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ACS Applied Materials & Interfaces

High Dielectric and Mechanical Properties Achieved in Crosslinked PVDF/α-SiC Nanocomposites with Elevated Compatibility and Induced Polarization at Interface Yefeng Feng†, Bei Miao†, Honghong Gong†, Yunchuan Xie†, Xiaoyong Wei††, Zhicheng Zhang†* †

Department of Applied Chemistry, MOE Key Laboratory for Nonequilibrium Synthesis and

Modulation of Condensed Matter, School of Science, Xi’an Jiaotong University, Xi’an, P. R. China, 710049. ††

Electronic Materials Research Laboratory, Key Laboratory of the Ministry of Education &

International Center for Dielectric Research, Xi’an Jiaotong University, Xi’an, P. R. China, 710049.

ABSTRACT: Remarkably improved dielectric properties including high-k, low loss and high breakdown strength combined with promising mechanical performance such as high flexibility, good heat and chemical resistivity are hardly to be achieved in high-k dielectric composites based on current composite fabrication strategy. In this work, a family of high-k polymer nanocomposites has been fabricated from a facile suspension cast process followed by chemical crosslinking at elevated temperature. Internal double-bonds bearing poly(vinylidene fluoride-chlorotrifluoroethylene) (P(VDFCTFE-DB)) in total amorphous phase is employed as crosslinkable polymer matrix. α-SiC particles with a diameter of 500nm are surface modified with 3-aminpropyltriethoxysilane (KH-550) as fillers for their comparable dielectric performance with PVDF polymer matrix, low conductivity and high breakdown strength. The interface between SiC particles and PVDF matrix has been finely tailored, which leads to the significantly elevated dielectric constant from 10 to over 120 in SiC particles due to the strong induced polarization. As a result, a remarkably improved dielectric constant (c.a. 70) has been observed in c-PVDF/m-SiC composites bearing 36 vol% SiC, which could be perfectly predicted by effective

ACS Paragon Plus Environment

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medium approximation (EMA) model. The optimized interface and enhanced compatibility between two components are also responsible for the depressed conductivity and dielectric loss in the resultant composites. Chemical crosslinking constructed in the composites results into the promising mechanical flexibility, good heat and chemical stability, and elevated tensile performance of the composites. Therefore, excellent dielectric and mechanical properties are finely balanced in the PVDF/α-SiC composites. This work might provide a facile and effective strategy to fabricate high-k dielectric composites with promising comprehensive performances.

KEYWORDS: induced polarization; surface modification; high-k; nanocomposite; interface

1. INTRODUCTION Flexible materials with high dielectric permittivity (high-k) performance have attracted massive attention and research interest in the field of high-performance electronic devices, such as the electroactive actuators, embedded capacitors and organic field-effect transistors (OFETs).1-3 It has been widely acknowledged that polymeric dielectric materials constructed mostly onto covalent bonds possess excellent flexibility and processability but undesirable low dielectric permittivity (εr). While the ionic inorganic materials such as ferroelectric ceramics show ultra-high εr along with undesired brittleness. Therefore, complexing strategy has long been proposed to combine the flexibility of polymer and high permittivity of inorganic materials in the resultant composites. During past decades, polymer/ferroelectric ceramics based 0-3 type high dielectric composites have achieved great success both in theory and practice.4-10 Usually, ceramic powders with large volume content are required to realize desired high dielectric performance in the resultant composites, which could be well predicted with classic series and parallel effective dielectric models11,12. That would result into their dispersion inhomogeneity and aggregation in polymer matrix, thus the significantly depressed flexibility and mechanical properties in the composites.13-19 Especially, when the fine powders in nano scale with extremely high relative surface area are requested to meet the high homogeneity in the ACS Paragon Plus Environment

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resultant composites, the mechanical and dielectric performances of the composites could hardly be balanced. In addition to ferroelectric ceramics, conductive or semi-conductive fillers with low bandgap have also been introduced into polymers to fabricate high-k composites20-23 based on the interface polarization (Maxwell-Wagner-Sillars effect24-29) and electric percolation threshold30-34 theories. Driven by the external electric field, free charges would be accumulated at the interface between different dielectric materials with various relaxation times. Before the conduction network is constructed among conductive particles, namely, the vicinity of the critical percolative volume content reaches, the amount of the accumulated interfacial charges would increase quickly as a function of filler content or the external electric field, which contributes greatly onto the permittivity. Limited by the percolation threshold, usually a low volume content of fillers is allowed to avoid the transition from insulator to conductor, which means the flexible polymeric performance could be fairly well maintained in the composites.35-40 However, these materials could hardly be applied under elevated electric field since the interfacial charges would cause huge conductive losses under high electric field. Even under low electric field, the dielectric loss at low frequency from the interfacial polarization relaxation is usually very high. Despite of the excellent dielectric properties achieved in the reported complexing strategies and the composites, how to balance the desired mechanical properties and the promising comprehensive electrical performances in high-k dielectric polymer composites is still a great challenge to the scientists. More recent results suggest that the mechanical properties of the dielectric materials have even more practical meanings than high εr considering their high field application for static electrical energy storage.41-43 In this work, we would like to offer a practice for fabricating dielectric composites with balanced comprehensive electrical performances and mechanical properties. First of all, α-SiC particles with a relatively high bandgap (~3.09 eV), high breakdown strength (~3200 kV/cm), medium dielectric constant (~10) and low ac conductivity (~5×10-7 S/cm) were employed as fillers44. That would avoid the


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above disadvantages of the high conductive loss, low breakdown strength, and narrow percolation threshold when conductive or low bandgap semi-conductive fillers are utilized. Secondly, a totally amorphous poly(vinylidene fluoride) based fluoropolymer bearing certain content of internal double bonds for crosslinking purpose, P(VDF-CTFE-DB), is applied as polymer matrix. The dielectric difference and interfaces between the crystal and amorphous phase of PVDF based semi-crystallized fluoropolymers could be neglected. Its dielectric constant (~12) is rather close to α-SiC (~10) to avoid the uneven electric field formed among the components with different dielectric constant, which is recognized as one of the reasons for the reduced breakdown strength in the dielectric composites. Thirdly, in an effort to achieve the good compatibility of the two components and flexibility in the resultant composites, α-SiC particles with nanoscale were modified with 3-aminpropyltriethoxysilane (KH-550). The significantly elevated permittivity of SiC induced by the strong induced polarization has been proved by the classic effective medium approximation (EMA) model45,46 for the first time. The optimized interface between high polar PVDF matrix and SiC fillers is responsible for the excellent dielectric properties including high-k, low dielectric loss, low conductivity and high breakdown strength observed in the composites filled with modified SiC particles. The chemical crosslinking of PVDF matrix provides fine flexibility and high chemical and thermal stability for the resultant films, which could overcome the disadvantages of brittleness and poor stability in many high-k dielectric polymer composites. This work might provide a facile and effective strategy to fabricate dielectric composites with promising comprehensive performances.

2. EXPERIMENTAL SECTION 2.1. Materials. α-SiC filler with nanoscale was kindly provided by Xi’an Tongxin Semi-conductor Accessory Co. Ltd. and it was washed with anhydrous alcohol for three times to remove impurities and then dried thoroughly at 200 °C for 10 h before use. P(VDF-CTFE) (80/20) with 20 mol% CTFE and a molecular weight (Mw) of 120,000 was purchased from Zhonghao Chenguang Research Institute of Chemical Industry. 3-aminpropyltriethoxysilane (KH-550, Alfa Aesar, AR grade), absolute ethanol


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(Tianjin Reagents Co. Ltd., AR grade), glacial acetic acid (Tianjin Reagents Co. Ltd., AR grade), aqueous hydrogen peroxide solution (H2O2, Tianjin Reagents Co. Ltd., AR grade, 30 wt%), acetone (Tianjin Reagents Co. Ltd., AR grade), and benzoyl peroxidewere (BPO, Tianjin Reagents Co. Ltd., AR grade) were used as received. P(VDF-CTFE-DB) (VDF/CTFE/DB=80/10/10 in molar ratio) was synthesized from P(VDF-CTFE) (VDF/CTFE=80/20 in molar ratio) through the dehydrochlorination process.47 P(VDF-CTFE) (5.0 g) was completely dissolved in 100 mL NMP into a 250 mL double-necked flask equipped with a condenser and a magnetic stirrer. Triethylamine (30 mL, 214 mmol) was introduced into the flask and the reaction was carried out at 70 °C under vigorous stirring for 1 h. The resultant hybrid was slowly precipitated in 1 L deionized water (DI water) and then the precipitant was thoroughly washed with DI water for 4 times and with methanol for 2 times. The resultant polymer was dried at 45 °C for 48 h under reduced pressure. 2.2. Surface modification onto nano SiC particles via KH-550. Firstly, surface activation of nano α-SiC particles was carried out by dispersing SiC particles (4.0 g) into H2O2 aqueous solution (30 wt%, 40.0 mL) with stirring for 24 h at ambient temperature followed by washing the particles thoroughly with deionized (DI) water and drying them for 24 h at 45 °C. Then, surface modification of SiC particles with a kind of silane coupling agent KH-550 was facilely conducted by scattering neat SiC particles (4.0 g) into a mixture of absolute ethanol (60.0 mL), DI water (15.0 mL), KH-550 (4.0 g) and glacial acetic acid (2.0 mL) with stirring for 24 h at ambient temperature followed by washing the resultant particles drastically with absolute ethanol for 5 times and drying them for 12 h at 45 °C under reduced pressure. 2.3. Fabrication of polymer based composites containing nano SiC semi-conductive particles. All P(VDF-CTFE-DB) based crosslinked composites with modified nano α-SiC and uncrosslinked composites bearing neat SiC were fabricated through the solution casting method48 from a suspension containing designed volume fraction of filler in high viscosity polymer solution with easily volatile


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acetone as the solvent (avoiding the sedimentation of high density filler during solvent evaporation) at ambient temperature on glass slides. After kept at 160 °C under reduced pressure for 4 hours, the films with a thickness of ~60 µm were peeled off followed by sputtering with Au on both surfaces as electrodes for electric properties measurements. For the thermal crosslinking among the polymers containing unsaturated internal double bonds, BPO initiator (3 wt% of polymer mass) was introduced into composite. The fabricating process of the composite with modified nano SiC filler and crosslinked polymer matrix was clearly shown by the schematic diagram in Fig. 1. For a facile expression, the materials and composites in present work were indicated by the following abbreviations as shown in Tab. 1.

(Figure 1) (Table 1) Full names

Abbreviations

Uncrosslinked P(VDF-CTFE-DB)

PVDF

Crosslinked P(VDF-CTFE-DB)

c-PVDF

Neat SiC

SiC

Oxidized SiC

Si-OH

Modified SiC

m-SiC

Uncrosslinked P(VDF-CTFE-DB)/neat SiC composite

PVDF/SiC

Uncrosslinked P(VDF-CTFE-DB)/x vol% neat SiC composite

PVDF/SiC-x

Crosslinked P(VDF-CTFE-DB)/modified SiC composite

c-PVDF/m-SiC

Crosslinked P(VDF-CTFE-DB)/x vol% modified SiC composite

c-PVDF/m-SiC-x

2.4. Characterization. X-ray diffraction (XRD) measurement was conducted on a Rigaku D/max


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2400 diffractometer (Rigaku Industrial Corp, Japan) with the X-ray wavelength of 1.542 Å (Cu Kα radiation, 40 kV and 100 mA), the 2θ diffraction angle from 30° to 75°, the rate of 15°/min and the step of 0.02°. Field-emission scanning electron microscopy (FE-SEM) results were obtained on JEOL JSM6700F. Fourier transform infrared (FTIR) spectroscopy was made with an IR spectrophotometer (Bruker-Tensor 27). Thermogravimetric analysis (TGA) was performed on a NJKHTG-1. Proton nuclear magnetic resonance (1H NMR) spectra were carried out on a Bruker (Advance III) 400 MHz spectrometer with acetone-d6 as the solvent and tetramethylsilane as an internal standard. Relative surface area of nano SiC particles was obtained by a 3H-2000PS2 BET tester. The dielectric and alternative current (ac) conductive properties of all the composites were obtained via a HP4284A LCR meter at frequencies ranging from 100 Hz to 1 MHz with 1 V voltage. The dielectric performances of composites under five different frequencies were obtained by the E4980A LCR meter at temperatures from 30 to 150 °C with the voltage of 1 V. Electric breakdown strength results were obtained through an auto voltage withstanding tester (RK2674B). Electric conductivity results under direct current (dc) electric filed (1 kV/cm) were achieved by a digit megger (PC68). Stress-strain curves were obtained at a tensile speed of 2 mm/min through a material testing machine CMT 6503 (Shenzhen Suns Technology Stock Co. Ltd.) and the composite films were cut into dumbbell shapes (the testing part with 4 mm width and about 0.14 mm thickness). Gold electrodes were sputtered on both surfaces of the composite films by JEOL JFC-1600 auto fine coater (Japan) for all electric properties characterizations. Different from the measurement of composite films, nano modified/neat SiC particles (the samples containing 100 vol% filler) were compacted in a cylinder-shaped mold with a diameter of 20 mm and height of 5 mm equipped with two copper plates at bottom and top of them as electrodes for electric properties measurement.

3. RESULTS AND DISCUSSION 3.1. Characterization of nano-particles and polymer. The composition and morphology of the used neat nano SiC particles were characterized via XRD and surface SEM results as presented in Fig. 2


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and 3, respectively. The structure of neat nano SiC particles were confirmed by XRD result in Fig. 2 to be in α-form crystal phase with high purity belonging to hexagonal crystal system via MDI Jade 5.0 analysis software. The diffraction angles 2θ at 34°, 36°, 38°, 41°, 60°, 66° and 72° are assigned to the crystal indices of (1 0 1), (1 0 2), (1 0 3), (1 0 4), (1 1 0), (1 0 9) and (2 0 2) for α-SiC, respectively49. Neither the surface modification of nano SiC with KH-550 nor the fabrication process of composites leads to crystal form alternation of SiC as shown in Fig. 2. A rather broad grain size distribution of neat nano SiC particles ranging from 250 nm to 2 µm was observed as shown in Fig. 3a. The geometry shape of neat nano α-SiC is very irregular, namely the filler has rather high geometry asymmetry as exhibited in Fig. 3b. The average diameter of SiC particles was statistically estimated as 400-500 nm. The relative surface area was measured with the 3H-2000PS2 BET tester to be 6.2 m2/g. The composition of surface modified nano SiC was characterized via FTIR and TGA as shown in Fig. 4. As shown in Fig. 4a, the peaks at 2921 and 2851 cm-1 on the FTIR spectra of surface modified SiC particles corresponding to the asymmetric and symmetric stretching vibrations of methylene groups from KH-550 strongly suggests the successful attachment of KH-550 molecules onto the surface of SiC.50 TGA results comparison of neat and modified nano SiC particles as exhibited in Fig. 4b could further confirm that. According to the inset of Fig. 4b, the weight of modified nano SiC particles is decreased more quickly than neat nano SiC particles with the elevation of testing temperature from 50 °C to 500 °C, which might be ascribed to the decomposition of the organic KH-550 grafted onto surface of inorganic nano SiC. The grafting of organic KH-550 onto the surface of SiC particles could be further confirmed by the dispersion behavior of the particles into acetone solvent. Both modified and neat nano SiC particles were finely dispersed into acetone respectively. The suspensions obtained were put into two vials and stood still simultaneously after a vigorous shaking. Five hours later, two particles exhibited very different dispersion behaviors in acetone as shown in Fig. 4c. Pristine SiC particles were found to be mostly settled onto the bottom, while almost all the modified nano SiC particles are still well suspended in acetone. That means the grafted organic molecules from KH-550 favor the stabilizing


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the particles in organic solvent and thus the suspension in polymer solution for their good solubility into organic solvents. The composition of P(VDF-CTFE-DB) was determined from 1H NMR as shown in Fig. 5a. The peaks at 2.2-2.7 ppm and 2.7-3.2 ppm were assigned to head-head (-CF2-CH2-CH2-CF2-) and head-tail (-CF2-CH2-CF2-CH2-) connections of VDF units. The peak at 3.2-3.6 ppm was identified as the protons on VDF units adjacent to CTFE units (-CF2-CH2-CFCl-CF2-). In contrast with pristine P(VDF-CTFE), the new multiple peaks at 6.2-6.7 ppm in the resultant P(VDF-CTFE-DB) were assigned to the protons on the double bonds (-CF2-CF=CH-CF2-) after the removal of HCl from (-CF2-CFCl-CH2-CF2-).47 The chemical composition was determined to be containing 80 mol% PVDF, 10 mol% CTFE and 10 mol% DB from the integrals of 1H NMR result. Moreover, the successful introduction of -CF=CH- bonds could be further confirmed by the characteristic stretching vibration absorption of C=C bonds at 1720 cm-1 and the out-plane bending vibration signal of C-H bonds in -CF=CH- at 704 cm-1 through FTIR result of P(VDF-CTFE-DB) in Fig. 5b.47

(Figure 2)


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(Figure 3)

(Figure 4)

(Figure 5) 3.2. Dispersion of nano particles in composites and its effect on the mechanical properties. High loading content of fillers with high surface area and poor compatibility between filler and organic matrix usually result into the formation of air voids at the interface zone and loose morphology in the resultant composites51. Surface modification of inorganic filler with organic surfactants is regarded as ACS Paragon Plus Environment

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one of the most effective methods to improve interface compatibility between inorganic filler and organic polymer matrix52-54. As indicated in the cross section morphology of composites filled with pristine SiC in Fig. 6e-6h, a large quantity of voids were observed in the uncrosslinked composites filled with pristine SiC at even low SiC content of 20 vol%. As filler content increases, the amount of voids was increased quickly and the polymer matrix among adjacent particles was reduced accordingly. That means SiC and P(VDF-CTFE-DB) show poor compatibility as well although the relative surface area of SiC is not as high as the other fine powders with even smaller size. Besides of the air voids, the homogeneity of neat SiC in composites is rather poor as well. As discussed above, neat SiC particles would quickly settle down in organic solvents due to the larger density of SiC together with the poor compatibility between particles and organic solvents. That means most of the SiC particles have already reached the glass substrate before the solvent evaporated completely and the formation of films, which is quite common in polymer composites filled with high density inorganic particles. After surface modified with KH-550, the content of air voids in the interface zone was significantly reduced in all crosslinked composites filled with modified nano SiC particles as shown in Fig. 6a-6d, comparing with corresponding neat SiC filled uncrosslinked composites with the consistent composition. No obvious voids were detected in the composites filled with up to 47 vol% SiC particles. The surface of the SiC particles was finely wrapped with polymer matrix and no naked SiC particles were observed. Once the filler content increases to 60 vol%, a lot of voids between particles were detected as well. But the surface of the particles was still partially coated with polymer different from the smooth surface of pristine SiC particles, which strongly suggests that the voids are mostly originated from the over loaded filler particles instead of poor compatibility between fillers and polymer matrix. That indicates SiC particles are kindly immersed in P(VDF-CTFE-DB) matrix with the bonding assistance of KH-550 bearing NH2 high polar groups. Meanwhile, the modified SiC particles were much more evenly distributed in matrix comparing to the rather heterogeneous dispersion in pristine SiC particles filled composites. That could be attributed to the significantly elevated suspension stability due to the


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compatibility improvement of surface modified SiC particles within organic solvents as discussed above. The slowed sedimentation of SiC particles during the solvent evaporation and formation of the film allows the particles evenly embedded in the crosslinked polymer matrix network.

(Figure 6) Besides of great influence onto the dispersion behavior of fillers in matrix, the surface compatibility between two components is also responsible for the significantly altered mechanical properties. It has been well accepted that the introduction of inorganic fillers in fine powder into the polymer materials always results into the quickly reduced mechanical properties including the broken strength, elongation, and the flexibility in the resultant composites, especially the large quantity is inevitable for improving the dielectric performances in present case. As expected, the broken strength and elongation of composites containing pristine SiC particles filled P(VDF-CTFE-DB) with no BPO as crosslinking initiator is rather low. As shown in Fig. 7, about 3 % and 1 % elongation was observed in uncrosslinked composites filled with 36 vol% and 60 vol% pristine SiC particles, and their broken strength was detected to be 2.2 MPa and 0.9MPa, respectively. Such low mechanical strength indicates the composite films are easily to be broken and difficult to handle, which is consistent with the morphology observed with SEM. However, if the modified SiC particles instead were utilized and the polymer matrix was chemically crosslinked with BPO, both the broken strength and the elongation were significantly improved comparing to the pristine SiC particles filled uncrosslinked P(VDF-CTFE-DB). As shown in Fig. 7, the highest strength of 3.5 MPa and an elongation of 10% were obtained in 36 vol% modified ACS Paragon Plus Environment

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SiC filled crosslinked film. Apparently, the increased mechanical properties should be ascribed to not only the significantly improved compatibility between two components but also the strengthened polymer matrix by crosslinking. The number of parallel samples measured in present tensile test was 5 and the data deviation was calculated to be ±12% as shown in Tab. 2. Different from the brittle and easy broken performance of pristine SiC particles filled uncrosslinked composite films, the crosslinked composite films bearing modified SiC particles exhibit great organic solvent (acetone) resistivity and flexibility as indicated in Fig. 8. That allows the films to be readily handled even under the conditions with organic solvent pollution.

(Figure 7) (Table 2) Samples

Broken strength (MPa)

Broken elongation (%)

a

2.2±0.26

3.0±0.36

b

3.5±0.42

10.0±1.20

c

0.9±0.11

1.0±0.12

d

1.2±0.14

2.3±0.28


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(Figure 8) 3.3. Electric breakdown strength. Electric breakdown strength (Eb) of the polymer composites has been found to reduce very quickly with the increasing of filler content in most of dielectric polymer composites. With respect to its remarkable influence onto the energy storage capability of the composites as indicated in equation Ue=1/2ε0εrEb2, rendering an Eb as high as possible is crucial to achieve high Ue in dielectric composites as well. As shown in Fig. 9, Eb results of PVDF/SiC composites as a function of filler volume content from 0 to 60 vol% were presented. With the increasing of filler content, Eb of composites filled with neat SiC was reduced linearly. That might be attributed to the uneven electric field formed in the composites induced by the components with varied dielectric performance, the high conduction and the air voids in large quantity as discussed above. Once the surface of SiC particles was modified, the composites exhibited much larger Eb (over 100 kV/cm) than that filled with neat SiC. That could be mostly ascribed to the reduced conductivity as well as the air voids content caused by the improvement of filler/polymer interface compatibility and filler dispersion in matrix. Meanwhile, chemical crosslinking could improve the Young’s modulus of the composites thus the elevated Eb according to the equation Eb=0.6(Y/ε0εr)1/2. Interestingly, a relatively high Eb of 320 kV/cm, which is over 64% of the pristine polymer matrix (~500 kV/cm), has been preserved in the composites even filled with 36 vol% SiC.


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(Figure 9) 3.4. Dielectric and conductive properties of modified nano SiC filled crosslinked composites. Measured on a HP4284A LCR meter at frequencies from 100 Hz to 1 MHz with 1V voltage bias, the dielectric permittivity, dielectric loss and ac conductivity of all crosslinked composites bearing nano modified SiC were obtained as a function of testing frequency in Fig. 10. As exhibited in Fig. 10a, in the testing frequency range from 100Hz to 1MHz, εr of polymer matrix was reduced from about 12 to 9 while εr of SiC particles was slightly changed from 11 to 10. After complexed, the permittivity of these composites was slightly decreased as frequency increases, which is similar as that of polymer or filler. As SiC content increases, the permittivity of the composites was firstly improved until the φSiC is up to 36 vol%, and further increasing of φSiC led to gradually depressed εr. Meanwhile, as εr increases, the reducing speed of εr against testing frequency was increasing as well. All the composites possess larger εr than that of both pristine polymer and SiC particles, which would be discussed more detailed later on. As shown in Fig. 10b, polymer matrix exhibited a “U” shape tanδ dependence onto the testing frequency from 100Hz to 1MHz. It has been well recognized the high dielectric loss at frequency below 100 Hz is due to the ion conduction for the low modulus of P(VDF-CTFE-DB) in totally amorphous phase. The elevated loss at frequency from 100 kHz to 1 MHz was well assigned to the dipole orientation relaxation mostly originating from the high polar VDF units. The introduction of modified SiC with a low loading content (c.a. 9 vol%) into P(VDF-CTFE-DB) would result into significantly ACS Paragon Plus Environment

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reduced tanδ in the whole range. At low frequency the reduced ion conduction by the improved modulus might address for the decreased dielectric loss. As more SiC particles were added, the increased interface between polymer matrix and fillers should be responsible for the improved tanδ at low frequency. At high frequency (100 kHz to 1MHz), the significantly reduced tanδ in the composites could be attributed to the hindered polarization relaxation of high polar VDF units, which is caused by the strong interaction between SiC particles and VDF units. Besides, ac conductivity by logarithm of composites was found to be increasing linearly against testing frequency by logarithm from 100 Hz to 1 MHz in Fig. 10c, which confirms their dielectric character under low electric filed. The rather close ac conductivity of these composites to the pristine polymer and fillers strongly indicates the increased permittivity of composites is not caused by elevated conductivity or leakage current. As SiC filling ratio increases from 9 to 77 vol%, the slightly improved ac conductivity should be responsible for the continuous increase of dielectric loss in the testing frequency range from 100 Hz to 1 MHz.

(Figure 10)


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Increasing interface compatibility through filler surface modification is one of the most favorable way to improve dielectric properties of 0-3 high-k ceramic/polymer composites55. As discussed above, the surface modification onto nano α-SiC could elevate the polymer/filler interface compatibility and the dispersion of the filler in polymer matrix. How the improved compatibility affecting the dielectric properties of P(VDF-CTFE-DB)/SiC composites has been investigated by side-by-side comparing the dielectric performances of composites filled with neat SiC and modified SiC particles. εr at 1 kHz of the composites filled with neat SiC and modified SiC as a function of filler volume content was shown in Fig. 11, respectively. For explaining these measured dielectric permittivity results, three models45,46 based on the effective medium approximation (EMA), namely the Bruggeman, Hashin-Shtrikman and Lichtenecker models, were employed to fit the present permittivity data. Especially, according to the core-shell structure shown in Fig. 12, the Hashin-Shtrikman model should be very suitable. According to the previous work56, the permittivity of SiC would be remarkably improved after complexed with high polar polymers in 2-2 model for the strong induced polarization between fillers and polymer, whereas, the permittivity of polymer matrix was rarely varied in the 2-2 polymer/SiC composites. Therefore, the permittivity real part of polymer matrix labeled as εp′ in present these 0-3 composites could be regarded as 12@1kHz based on Fig. 10a, regardless of the permittivity imaginary part of polymer matrix. However, the permittivity real part of m-SiC filler marked by εf′ should be obtained through following calculation (equations (1)-(3)),

ε ∗f = ε 'f − iε ''f = ε 'f − i

σf σf = ε 'f − i ε 0ω 2π f ε 0

ε ''f tan δ = ' εf ε = ' f

ε ''f tan δ

(1)

(2)

=

σf 2π f ε 0 tan δ

(3)

, where εf*, εf′,εf″,σf,ω,ε0, f and tanδ are referring to the complex dielectric permittivity of m-SiC, the


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permittivity real part of m-SiC, the permittivity imaginary part of m-SiC, the dc conductivity of m-SiC (1.0×10-6 S/cm), the angular frequency of electromagnetic wave, the vacuum permittivity (c.a. 8.85× 10-12 F/m), the frequency of applied electric field (1 kHz) and the dielectric loss of m-SiC (0.14@1 kHz), respectively. Substituting these data into equation (3), the permittivity real part of m-SiC filler was obtained as 128.5@1kHz employed to make model fitting. This result agreed very well with that measured in 2-2 composites and could confirm the significantly improved permittivity real part of SiC particles in composites from 10 to 128.5 at 1 kHz. The Lichtenecker, Bruggeman and Hashin-Shtrikman models models were shown as equations (4)-(6) respectively,

ε cα = ϕ f ε αf + ϕ pε αp

(4)

ε f − ε c (1 − ϕ f )( ε f − ε p ) = ε c1/3 ε 1/3 p

(5)

ε fεp  ϕ p − n   1−ϕp  +   1 − n   1 − n  (1 − n ) ε p + nε f

εc = ε p 

(6)

, where εc, εf, εp, ϕf and ϕp are referring to the permittivity of composites, the permittivity of filler, the permittivity of polymer, the volume content of filler and the volume content of polymer, respectively. Moreover, the variable α in equation (4) can change from -1 to 1, and α indicates the anisotropy and isotropy distribution of filler in matrix at -1 and 1 respectively. In present case, α should be equal to 1 attributed to isotropy distribution of filler. In equation (6), n=gϕp and g (geometrical factor) is closely dependent on the 3D shape and topology of filler (0≤g≤1). For ideal sphere filler, g=1/3. However, the 3D shape of SiC filler was confirmed to be very irregular as shown in Fig. 3. Therefore, g should be 1/100 rather than 1/3 based on the effective data fitting, namely n=0.01ϕp in present case. According to the various volume compositions of composite system and the permittivity real parts@1kHz of both components (εp′=12 and εf′=128.5), the fitted permittivity results were obtained as exhibited in Fig. 11 using Lichtenecker (line 3), Bruggeman (line 4) and Hashin-Shtrikman (line 5) models respectively. All


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these models take into account the depolarization fields on the interfaces between both components, but interfacial air voids are not regarded. With the increase of ϕf, εc of composites filled with modified SiC and neat SiC particles showed the similar “∧” shape changing tendency, namely, increasing quickly and reaching maximum value at ϕf = 36 vol% followed by quickly reducing to the lowest permittivity value of SiC filler. Meanwhile, εc of composites filled with modified SiC is much higher than the composites filled with neat SiC bearing the consistent ϕf, especially, at the larger ϕf, (c.a. over 20 vol%). The permittivity real part of SiC filler should be closely dependent on the interaction between polymer matrix and filler, which is dominated by many factors including the compatibility between the two components, the polar nature of polymer matrix, the size of particles, surface area of polymer covered, the filler content, and even the content of air voids. When ϕf is below the critical value (36 vol% in this case), the fitted results from Lichtenecker and Hashin-Shtrikman models agreed perfectly with the measured results of the composites filled with modified SiC particles. Further increasing SiC loading content led to deviated εc increasing tendency and even lowered εc. That could be attributed to the quickly reduced εf′ due to the increased air voids as indicated in Fig. 6 and illustrated in Fig. 12. As more SiC particles introduced, the polymer matrix is not sufficient to cover all the surface of modified SiC particles and the air voids at interface zone would be formed. The existence of low polar air voids may form two more types of interfaces, one is between SiC and air voids, and the other is between air voids and polymer matrix. That makes the system more complicated and even more difficult to be predicted. But both interfaces are unfavorable for the elevation of εf′, which may address for the lowered εc observed in the composites with increased ϕf. Meanwhile, the negative effect of air voids onto εc could be further confirmed by the lowered εc value obtained in composites filled with neat SiC than that of modified SiC particles. As compared in Fig. 6, many more air voids were introduced into the composites filled with neat SiC than that of modified SiC even the filler content is below the critical 36


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vol%, which may significantly reduce εf′ in composites. From this point of view, we could demonstrate that the strong interfacial interaction between PVDF and SiC may remarkably change the permittivity real part of SiC. The εf′ of SiC shows a tight dependence onto the interface content, namely, the surface area of SiC particles coated with PVDF. Surface modification of SiC particles may improve εc of composites by means of reducing the formation of air voids. Besides of the remarkable influence onto the permittivity of the composites, the improved compatibility between two components is also favorable for the lowered dielectric loss and dc conductivity of the composites. Dielectric loss@1kHz and dc conductivity@1kV/cm results of crosslinked P(VDF-CTFE-DB)/SiC nano composites were shown in Fig. 13a and 13b respectively as a function of ϕSiC. Tanδ of all the composites was found to be lower than pristine P(VDF-CTFE-DB) matrix, which might be ascribed to two opposite effects including lowered ion conduction caused by the improved Young’s modulus and the increased dielectric loss from interfacial relaxation. First of all, the introduction of hard SiC particles would lead to the increasing of Young’s modulus of composites. Therefore, the ion conduction induced dielectric loss would be depressed. As more SiC added, the interphase area were increased, and the interfacial polarization induced dielectric loss would be increased. When ϕSiC is over the critical value (36 vol%), both the Young’s modulus and interphase exhibited little influence onto the dielectric loss. The dielectric loss of composites filled with neat SiC showed the similar changing trend as that filled with modified SiC. Differently, their dielectric loss is 40 to 100% larger as shown in Fig. 13a. That might be attributed to the higher conduction loss caused by the aggregation of neat SiC particles with higher conductivity than polymer matrix. As indicated in Fig. 13b, where the dc conductivity of two sets of composites measured under electric field of 1 kV/cm was compared, the conductivity of composites filled with modified SiC particles was much lower than that of neat SiC particles. In addition, the dc conductivity@1kV/cm of m-SiC and SiC was measured to be 1.0×10-6 S/cm and 0.98×10-6 S/cm, respectively. Apparently, the improved compatibility between two components is responsible for the lowered dielectric loss as well by avoiding the aggregation of ACS Paragon Plus Environment

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particles induced conduction loss. To clarify the surface modification of SiC particles onto the dielectric performance of composites at elevated temperature, the dielectric properties of 36 vol% modified SiC filled crosslinked composite, 36 vol% neat SiC filled uncrosslinked composite and crosslinked P(VDF-CTFE-DB) were obtained with the testing temperature ranging from 30 to 150 °C and under five various frequencies as shown in Fig. 14. Similar temperature dependence has been observed in the dielectric constant and loss of the composites filled with neat SiC and modified SiC particles. When the temperature was elevated from 30 to 60 °C or the testing frequency was higher than 1 kHz, the permittivity in both composites was independent on testing temperature. At 100 Hz, their permittivity was significantly improved as the temperature was increased beyond 60 °C due to the improved mobility of dipoles in PVDF polymer matrix, which could be confirmed by the similar increasing trend observed in crosslinked P(VDFCTFE-DB) (Fig. 14(e) and (f)). Differently, larger εc and smaller tanδ of c-PVDF/m-SiC composite than that of PVDF/SiC sample were found under the same testing conditions. Meanwhile, the crosslinking of P(VDF-CTFE-DB) is also responsible for the elevated temperature where tanδ at low frequency (100Hz) starts to increase quickly. The results suggest that the dielectric performance dependence of the composites onto the temperature is mostly originating from the P(VDF-CTFE-DB) matrix and chemical crosslinking could improve the thermal stability of the dielectric performance. Based on the temperature dependence of dielectric properties in PVDF/SiC-36, the elevated permittivity in present composite system could be presumably interpreted by the weak nonzero conductivity of both components.


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(Figure 11)

(Figure 12)

(Figure 13)


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(Figure 14)

4. CONCLUSIONS In general, a facile strategy for high dielectric and flexible polymer composites fabrication has been presented. Instead of high conductive particles or high dielectric ferroelectric ceramics, α-SiC particles with comparable dielectric constant with polymer matrix, high breakdown strength and low conductivity have been employed as fillers to overcome the disadvantages due to the introduction of conductive and


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ferroelectric fillers. After complexed with high polar PVDF based polymer, the dielectric constant of αSiC particles has been significantly elevated thanks to the strong induced polarization constructed at the interface, which could be confirmed by the calculated permittivity real part of modified SiC filler. Improving the interface compatibility by surface modification of SiC particles leads to even higher dielectric constant, reduced dielectric loss and ac conduction in the resultant composites, which could be finely addressed by both of classic EMA models with the permittivity real part of modified SiC filler. The improved interface compatibility and the crosslinking of polymer matrix are responsible for the enhanced mechanical flexibility, fine thermal and chemical stability, and elevated tensile feature of the composites. Therefore, excellent dielectric and mechanical properties are finely balanced in the resultant c-PVDF/m-SiC composites.

AUTHOR INFORMATION Corresponding Author. Tel: 86-29-82663937. E-mail: [email protected] Coauthors.E-mails: [email protected]; [email protected]; [email protected]; [email protected]; [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGEMENT This work was financially supported by the National Nature Science Foundation of China (Grant Nos. 51573146, 51103115, 50903065), Fundamental Research Funds for the Central Universities (Grant Nos. XJJ2013075, cxtd2015003), International Science & Technology Cooperation Program of China (Grant No. 2013DFR50470), Natural Science Basic Research Plan in Shanxi Province of China (Grant Nos. 2013JZ003, 2015JZ009), and National Basic Research Program of China 973 (Grant No. 6132620101).


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Supporting Information Available: Monopolar D-E hysteresis loops of c-PVDF, c-PVDF/m-SiC-9 and c-PVDF/m-SiC-20). This material is available free of charge via the Internet at http://pubs.acs.org.

Captions of all figures and tables: Figure 1. Schemed fabrication process of c-PVDF/m-SiC composite films. Figure 2. XRD curves of (a) m-SiC particles, (b) SiC particles and (c) PVDF/SiC-47 composite. Figure 3. Surface SEM of SiC particles with a magnification of (a) 5000 and (b) 10000 times. Figure 4. (a) FTIR comparison of SiC, SiC-OH, m-SiC particles with KH-550, (b) TGA of m-SiC and SiC particles, and (c) dispersion behavior comparison of m-SiC and SiC particles in acetone.

Figure 5. (a) 1H NMR and (b) FTIR of P(VDF-CTFE-DB) and P(VDF-CTFE). Figure 6. Cross-section SEM of c-PVDF/m-SiC and PVDF/SiC composites with various filler contents. (a) c-PVDF/m-SiC-20, (b) c-PVDF/m-SiC-36, (c) c-PVDF/m-SiC-47, (d) c-PVDF/m-SiC-60, (e) PVDF/SiC-20, (f) PVDF/SiC-36, (g) PVDF/SiC-47 and (h) PVDF/SiC-60.

Figure 7. Stress-strain curves of (a) PVDF/SiC-36, (b) c-PVDF/m-SiC-36, (c) PVDF/SiC-60, and (d) cPVDF/m-SiC-60 composites.

Figure 8. Digital figures of (a) PVDF/SiC-36 and c-PVDF/m-SiC-36 composites in acetone and (b) cPVDF/m-SiC-60 composite film rolled onto a NMR tube.

Figure 9. Measured Eb results of (1) PVDF/SiC and (2) c-PVDF/m-SiC composites as a function of filler volume fraction.

Figure 10. (a) dielectric properties of both fluoropolymers, (b) permittivity, (c) dielectric loss and (d) ac conductivity of c-PVDF/m-SiC composites with various filler contents as a function of frequency.

Figure 11. Permittivity of c-PVDF/m-SiC (bar 1) and PVDF/SiC (bar 2) measured at 1kHz as a function of SiC volume content, and the fitted permittivity of c-PVDF/m-SiC from Lichtenecker (line 3), Bruggeman (line 4) and Hashin-Shtrikman (line 5) models using both permittivity real parts of m-SiC (εf′=128.5) and polymer matrix (εp′=12).

Figure 12. Schematic diagrams of SiC particles surrounded by PVDF matrix in (a) PVDF/SiC and (b)


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c-PVDF/m-SiC composites. The grey arrows are referring to the induced dipoles pairs in SiC particles.

Figure 13. (a) Dielectric loss measured at 1kHz and (b) dc conductivity measured under 1kV/cm electric field of c-PVDF/m-SiC composites with varied SiC volume content.

Figure 14. Dielectric temperature spectra of (a, b) PVDF/SiC-36 composite, (c, d) c-PVDF/m-SiC-36 composite and (e, f) c-PVDF. The insets were partially enlarged results.

Table 1. Abbreviations for materials and composites. Table 2. Data deviation in tensile test.

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SYNOPSIS TOC


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High Dielectric and Mechanical Properties ... - ACS Publications

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