Enhanced breakdown strength and suppressed dielectric loss of polymer nanocomposites with BaTiO3 fillers modified by fluoropolymer

The introduction of ceramic fillers into a polymer matrix is an effective way to obtain dielectric nanocomposites with high energy storage density. However, the inorganic fillers are difficult to disperse evenly into the polymer matrix because of the poor compatibility, which stems from the large surface energy difference and the mismatch in dielectric constant between the fillers and polymer matrix. Polymer nanocomposites with high dielectric constant while maintaining high breakdown strength have great potential to achieve high energy storage density. In this work, poly(dodecafluoroheptyl methacrylate) terminated with a thiol end group (PDFMA-SH) was synthesized via a two-step process including Reversible Addition-Fragmentation Chain Transfer (RAFT) polymerization and subsequent aminolysis reaction. The polymer was then grafted into the surface of BaTiO3 (BT) nanoparticles by a “thiol–ene” click reaction to reduce the surface energy of BT nanoparticles. A novel nanocomposite consisted of the core–shell structured PDFMA@BT hybrid nanoparticles and poly(vinylidene fluoride–chlorotrifluoroethylene) (P(VDF–CTFE)) matrix was prepared. The influence of the fluoropolymer shell on the dispersion of fillers, the compatibility between the fillers and polymer matrix, dielectric properties and breakdown strength were investigated systematically. The results indicate that the strong interfacial adhesion between the hybrid nanoparticles and P(VDF–CTFE) matrix makes the fillers uniformly dispersed in the polymer matrix. Meanwhile, the excellent compatibility between the two components is favorable for enhancing the breakdown strength and suppressing dielectric loss, providing a condition to prepare dielectric materials with high energy storage density.


Introduction
Dielectric materials with high energy storage density have received great attention due to their wide application in capacitors, pulse power devices, hybrid electric vehicles, and vibration energy collectors. [1][2][3][4][5][6] The energy storage density of the dielectrics is determined by the dielectric constant 7 (k) and breakdown strength (E b ) of the nanocomposites. The energy storage density (U e ) for linear dielectrics can be calculated by U e ¼ 1/23 r 3 0 E b 2 , where 3 r and 3 0 are the relative permittivity of the dielectrics and the vacuum permittivity, 8 respectively. Therefore, U e is determined by the synergy of the breakdown strength and dielectric constant, implying the high energy storage density can be achieved while the dielectric constant and breakdown strength are enhanced simultaneously. In order to simultaneously increase the E b and the dielectric constant, it is proposed to prepare a multi-component dielectric nanocomposite. One approach is to incorporate a conductive ller 9-12 (such as aluminum powder, graphene, carbon nanotubes, etc.) into the polymer matrix. When the llers concentration approaches the percolation threshold, 13 an ultra-high dielectric constant can be achieved. However, the high dielectric constant is oen accompanied by high dielectric loss and low breakdown strength, limited the achievable high energy storage density. 14,15 Another effective method is to mix a high dielectric constant ceramic material [16][17][18][19][20] (such as BaTiO 3 (BT), Pb(Zr,Ti)O 3 (PZT), TiO 2 ) with polymer matrix. The dielectric properties of the polymer nanocomposites can be improved by adjusting the llers contents. However, the nano-size inorganic llers with large specic surface area and high surface energy tend to agglomerating, which limits the increase of energy density of the nanocomposite. 8,18 Thus, the preparation of the polymer nanocomposites with outstanding dispersity and stability remains is urgently required.
Ferroelectric polymers [21][22][23][24][25] poly(vinylidene uoride) (PVDF) and its binary/terpolymers are oen used as a polymer matrix for the preparation of high performance dielectric nanocomposite materials due to their high dielectric constant. Compared to hydrocarbon-based polymers, uoropolymers have lower surface energy, resulting in less compatible with most organic and inorganic materials. Therefore, in order to solve the problem of the agglomeration of ceramic llers and the interfacial compatibility between PVDF matrix and llers, researchers have concentrated a lot of work on the surface modication of ceramic llers. Song et al. 26 have used the selfpolymerization of dopamine to modify BT bers, effectively improving the interfacial compatibility. Zhang et al. 8 have prepared PVDF based nanocomposites with paraffin-coated BT nanoparticles. The nanocomposite possessed a high energy storage density about of 21.1 J cm À3 with 50 vol% llers. These strategies of BT nanoparticle modication with alkane modi-ers still have the issues of llers aggregation in the polymer matrix. That's because the structural difference between the polymer shell layer of the llers and the polymer matrix, creating obstacles to practical application to some extent. Achieving the excellent interfacial adhesion of the nanocomposites, associated with the interfacial coupling effect between the matrix and the modier of llers, is the key to improve both the dielectric properties and breakdown strength of the nanocomposites. 27,28 It has been shown that a large difference in relative permittivity between an inorganic ceramic ller and the polymer matrix leads to an inhomogeneous electric eld distribution, since electric elds tends to concentrate in phases of low permittivity. 29 This can be achieved by introducing a low-permittivity polymer shell layer on the surface of the high-permittivity nanoller to mitigate the permittivity mismatch between the ller particle and polymer matrix. 30 Therefore, we designed and synthesized the core-shell structured hybrid llers using uoropolymer as the organic shells of llers to reduce the surface energy of BT nanoparticles. The polymer shells can decrease the interface difference 31 between the llers and the matrix, which can effectively enhance the interfacial adhesion.
In this work, we rstly synthesized poly(dodecauoroheptyl methacrylate) with a thiol end group (PDFMA-SH) by RAFT polymerization. Meanwhile, the double bonds functional groups were introduced to the BT surface by using a silane coupling agent. Subsequently, the PDFMA was graed onto the surface of the BT nanoparticles by a "thiol-ene" click reaction to obtain PDFMA@BT hybrid llers with a core-shell structure. The results show that the presence of uoropolymer shells have improved the dispersibility of the llers in the polymer matrix, and enhanced the compatibility between the llers and the polymer matrix, which could signicantly improve the breakdown strength of the nanocomposites. The novel hybrid llers modied by uoropolymer provide a strategy for realizing dielectric nanocomposite materials with high energy storage density.

Synthesis of the thiol-terminated PDMFA
PDFMA-CDB 32 is synthesized typically as follows. DFMA (5.0 g, 1.25 Â 10 À2 mol), a mixture of CDB (0.226 g, 8.3 Â 10 À4 mol), AIBN (0.028 g, 1.66 Â 10 À4 mol) and THF (5 mL) was charged sequentially in a single-necked ask, and remove the air from the bottle. The reactor was purged with nitrogen and then placed in a water bath at 65 C. Aer 12 h, the product was puried by precipitating into excess methanol (three times), and suction ltration, then drying in a vacuum for 24 h to obtain a pink powder PDFMA-CDB (4.021 g, 80.7 wt% yield).

Modication of BT nanoparticles by uoropolymer
PDFMA@BT with the core-shell structure was prepared as follows: 10.0 g BT nanoparticles were dispersed in 100 mL H 2 O 2 (30 wt%) solution with ultrasonic treatment for 30 min, and then reuxed at 105 C in an oil bath for 6 h. The BT nanoparticles were separated by centrifugation at 9000 rpm and washed with deionized water thoroughly, and then the product was dried at 60 C for 24 h to obtain hydroxylated BT (BT-OH). Subsequently, 10.0 g of BT-OH nanoparticles were dispersed in 100 mL anhydrous toluene solution with ultrasonic treatment for 30 min. Then 2.0 g KH-570 and 1 mL acetic acid was added. The reaction system was heated to 80 C for 12 h under the atmosphere of nitrogen. The nanoparticles were separated by centrifugation at 9000 rpm and washed with ethanol thoroughly, then the product was dried at 60 C for 24 h to obtain vinyl-functionalized BT nanoparticles (BT-ene). Fluoropolymerfunctionalized BT nanoparticles were obtained in the last step: 10.0 g of BT-ene nanoparticles and 100 mL of THF were added to a ask, and the mixture was sonicated for 30 min followed by the addition of PDFMA-SH (1.0 g) and DMPA (0.1 g). Under the protection of nitrogen, the dispersion was stirred at room temperature for 12 h with ultraviolet light irradiation (l ¼ 365 nm). Aer centrifugation and washed with THF, the nanoparticles were dried at 60 C for 24 h to obtain PDFMA@BT.

Preparation of nanocomposite
The core-shell structured PDFMA@BT nanoparticles were rstly dispersed into DMF and sonicated for 30 min. P(VDF-CTFE) was then added with vigorous stirring, and the temperature was raised to 60 C to form a stable suspension. The suspension was cast onto a glass substrate and the thickness of the lm was controlled using a handy scraper. The cast lm was dried under vacuum at 80 C for 24 h to remove DMF. The resulting lm has a thickness of about 10 mm.

Characterization
The 1 H nuclear magnetic resonance ( 1 H NMR) was characterized by an Advance III 400 MHz spectrometer at room temperature, and the deuterated acetone was used as a solvent. The molecular weight distribution (Đ) of the polymer was determined by gel permeation chromatography (GPC) with THF as the eluent (a ow rate of 1.0 mL min À1 ). UV-vis spectra were recorded on TU-1901 Spectrophotometer with a sampling interval of 1.00 nm. The analysis of the functional groups of BT nanoparticles was carried out by Fourier transform infrared spectroscopy (FTIR, Nicolet iS10). X-ray diffraction (XRD, PANalytical) was adopted to identify the phase composition of the PDFMA@BT nanocomposite. Transmission electron microscope (TEM) images of PDFMA@BT nanoparticles were obtained via a JEOL-1400 electron microscope. Scanning electron microscopy (SEM, S-2500, Hitachi Seiki) was used for the characterization of the dispersion of llers in the nanocomposites. The dielectric properties were measured using a broadband dielectric spectrometer (Agilent E4990A) at the frequency range from 10 2 Hz to 10 7 Hz. The breakdown strength (E b ) was characterized using a Programmable Withstanding Voltage Tester (CS9916AX) under a DC 500 V s À1 until sample failure.

Preparation and characterization of thiol-terminated PDFMA
The RAFT reagent CDB can effectively regulate the molecular weight and molecular weight distribution 32 of PDFMA according to the previous reports. Aer polymerization, the polymer chains contain a thioester end group (Scheme 1), which can be converted into a thiol group by a reduction reaction. The thiol groups are used as an active site for graing the polymer chain onto the surface of BT nanoparticles. 33 As shown in Fig. 1a, the GPC curves of the homopolymers before and aer the aminolysis are all single peaks, and the molecular weight distribution ($1.1) is narrow. The GPC results certify that the disuldes are not formed during the aminolysis. Meanwhile, the ultraviolet absorption peak at 310 nm of the dithiocarbonate completely disappeared (Fig. 1b) aer aminolysis, and the characteristic absorption of the phenyl (d ¼ 7.2-8.0 ppm) also disappeared (Fig. 1c). As a result, the color of the products is changed from pink to pale yellow as shown in Fig. 1d. According to the above results, the thioester groups have been successfully converted into thiol groups.

Preparation and characterization of PDFMA@BT nanoparticles
Scheme 2 shows the preparation process of the core-shell structured PDFMA@BT hybrid nanoparticles. From the FTIR spectra as shown in Fig. 2a, a small amount of -OH groups (3400 cm À1 ) are presented in the raw BT, and the content of the -OH group is signicantly increased aer the hydroxylation with hydrogen peroxide. Aer that, the double bonds are introduced on the surface of BT-OH nanoparticles by the reaction of BT-OH with KH-570. New absorptions at 1040 cm À1 , 1140 cm À1 , 1720 cm À1 and 2800-3000 cm À1 corresponding to the stretching vibration of Si-O-Si, Si-O-Ba, C]O and -CH 2respectively appear compared with the raw BT nanoparticles, indicating that the silane coupling agent KH-570 is successfully graed onto the surface of the BT nanoparticles. In order to obtain the core-shell structured nanoparticles PDFMA@BT, the "thiol-ene" click reaction between BT-ene and PMDFA-SH is initiated by the DMPA photoinitiator at room temperature under ultraviolet light (365 nm). The FTIR characterization is used to examine the functional groups of the obtained hybrid nanoparticles repeatedly washed by THF. The new absorption peaks at 1600 cm À1 and 1420 cm À1 originated from the vibration of the C-F bonds, conrming that the PDFMA was successfully graed onto the surface of the BT nanoparticles.
Normally, the polymer will decompose at high temperature, but the BT nanoparticles will not show mass loss even above 1000 C. Therefore, the TGA measurement can be used to analyze the amount of uoropolymer graed on the surface of BT nanoparticles. As shown in Fig. 2b, the weight loss of unmodied BT nanoparticles at 750 C is only 2.55%, while the weight loss of BT-OH, BT-ene and PDFMA@BT gradually increase to 10.1%, 15.99%, and 21.5%, respectively. The weight loss of the modied BT nanoparticles under 300 C may be caused by hydroxyl, solvent or other impurities in the sample, while the weight loss above 300 C can be attributed to the decomposition of modiers and uoropolymers graed onto the surface of the modied BT nanoparticles. Overall, it was conrmed by the TGA analysis that the content of the Scheme 1 Preparation route of PDFMA-SH.
This journal is © The Royal Society of Chemistry 2020 RSC Adv., 2020, 10, 7065-7072 | 7067 uoropolymer graed onto the surface of the BT nanoparticles was about 6 wt%.
The TEM images as shown in Fig. 2c and d further demonstrate the core-shell structured PDFMA@BT. It can be seen that the BT nanoparticles have an average size of about 100 nm. A rough polymer layer can be observed on the surface of the BT nanoparticles, and the thickness of the PDFMA shells was about $5 nm, which indicates that the core-shell structured PDFMA@BT nanoparticles are successfully synthesized.

Preparation and characterization of the nanocomposite
Nanocomposites with different volume fraction (0-20 vol%) have been prepared to investigate the effects of the interfacial adhesion on the micromorphology and electrical properties. Fig. 3 shows the cross-sectional images of the two kinds of nanocomposites doping with different volume load with BT and PDFMA@BT llers, respectively. It can be seen that the unmodied BT nanoparticles are poorly compatible with the P(VDF-CTFE) matrix as most llers are exposed outside of the polymer matrix. In addition, the obvious agglomeration will induce the defects and disordered structures inside the lms. In contrast, PDFMA@BT nanoparticles have a uniform dispersion in the P(VDF-CTFE) matrix without any agglomeration, which indicates that the excellent compatibility between PDFMA@BT and P(VDF-CTFE) matrix is constructed. The improvement of the compatibility can be attributed to the fact that PDFMA shells can reduce the surface energy of BT nanoparticles, increasing the distance between the nanoparticles. On the other hand, the strong interchain force between the uoropolymer PDFMA and P(VDF-CTFE) matrix can undeniably enhance the interfacial adhesion between the llers and polymer matrix. 27 Fig . 4 shows the XRD patterns of the PDFMA@BT/P(VDF-CTFE) nanocomposite lms. The diffraction pattern of the pure     05-0626) and conrm the existence of inorganic BT nanoparticles in the polymer matrix. As the volume fraction of PDFMA@BT increases, the diffuse reection of P(VDF-CTFE) in the nanocomposite lm weakens and even disappears, and the intensity of the BT peak increases, which can be attributed to the compositional changes. 28 Breakdown strength is an important factor affecting the performance of dielectric materials. Fig. 5 shows the characteristic breakdown strength of the nanocomposite lms, which is analyzed via two-parameter Weibull distribution function: where P(E) is the cumulative probability of electrical faults. b is the shape parameter tted by the distributed linear regression. E is the experimental breakdown strength, and E b is the characteristic breakdown strength obtained when P(E) reaches 63.2%. This distribution function can convert into its logarithmic form: when the value of log[Àln(1 À P)] reaches zero, the E is equal to E b . Besides, for every specic value of E, P is calculated as follows: where i indicates that this E value ranks the ith in the ascending order of breakdown strength data and n is the number of total data points. 34 The E b of the nanocomposite lms with different llers shows the same tendency as the volume fraction of the ller increases ( Fig. 5a and b). As the volume fraction of BT nanoparticles increases, the llers are more accessible to be agglomerated, causing the inevitable cracks and voids inside the nanocomposite lms. The defects will sharply decrease the E b . The mismatch of the relative permittivity or conductivity between the inorganic ller and the polymer matrix usually leads to an uneven distribution of the electric eld inside the composite, resulting in a signicant reduction in the breakdown strength of the dielectric composite. 30 It worth noting that PDFMA@BT/P(VDF-CTFE) nanocomposite lms possess higher b value under the same volume fraction of llers. It also proves the uniform dispersion of the llers and the fewer defects inside the lms from another aspect.
In order to study the effect of the llers loading on the E b , the E b at different volume fractions are calculated as shown in Fig. 5c. The E b of the pure P(VDF-CTFE) matrix is 284.1 mV m À1 , however, the E b is drastically reduced to 136 mV m À1 for BT/ P(VDF-CTFE) nanocomposite lms when the llers loading is 20 vol%. As for PDFMA@BT/P(VDF-CTFE) nanocomposite lm, the PDFMA polymer shells greatly improve the compatibility of BT nanoparticles with P(VDF-CTFE) matrix and reduce the generation of leakage current channels, which are benecial to increasing the E b . As shown in Fig. 5c, the curve of normalized energy density factor changing 35 with the volume fraction of conductors also indicates that the PDFMA shell exhibits the ability to decrease the permittivity contrast between the BT ller and polymer matrix in the nanocomposite, and lessen the local electric eld concentration, which improve the E b of the PDFMA@BT/P(VDF-CTFE) nanocomposite lm. As a result, the E b of PDFMA@BT/P(VDF-CTFE) nanocomposite lm with 20 vol% llers loading achieves 188 mV m À1 . Fig. 6 shows the dielectric properties of the different types of nanocomposite lms with different llers loading at room temperature. The dielectric constant of the nanocomposite lms increases with the increase of llers loading. For example, when the frequency is 100 Hz, the dielectric constant of PDFMA@BT/ P(VDF-CTFE) nanocomposite lm with 20 vol% llers loading is 27.3, while the pure P(VDF-CTFE) lm is 11. The reason is that BT nanoparticles have higher dielectric constant than that of the pure matrix. In addition, the increased content of the BT nanoparticles in the nanocomposite lm can induce interfacial polarization and enhance the Maxwell-Wagner-Silars effect. 36,37 As the frequency increases, the dielectric constant of the nanocomposite lm gradually decreases and maintains relatively high stability at low frequency (10 2 to 10 5 Hz). In contrast, the dielectric constant of the nanocomposite lm shows a signicant decrease in the high-frequency range (10 5 to 10 7 Hz). It is because the speed of establishing interface polarization at high frequency cannot keep up with the change of the applied electric eld, which leads to the sharp decrease of dielectric constant. Unfortunately, the PDFMA@BT/P(VDF-CTFE) nanocomposite lm has a lower dielectric constant than that of BT/(PVDF-CTFFE) nanocomposite lm under the same doping amount, especially in the low-frequency range. In addition, the PDFMA@BT/P(VDF-CTFE) nanocomposite lm exhibits a weaker frequency dependency under the same condition compared with the BT/P(VDF-CTFE) nanocomposite lm. Due to the weak polarization ability and low dependence on the frequency of the C-F bond in PDMFA, the modied BT nanoparticles can reduce the sensitivity to the electric elds due to the presence of the insulating PDMFA shells. Besides, the PDMFA shells can also enhance the interfacial adhesion between the llers and matrix, which inhibits the interfacial polarization.
The polymer shells also show a positive effect on the dielectric loss of the nanocomposite. As shown in Fig. 6c and d, the dielectric loss of all nanocomposites are kept at a low level (<0.25). Compared with the pure P(VDF-CTFE) lm, the dielectric loss of all nanocomposite lms gradually decreases at the range of 10 4 to 10 7 Hz when the doping ller content increases. In the low-frequency range (10 2 to 10 4 Hz), although the dielectric loss of PDFMA@BT/ P(VDF-CTFE) nanocomposite lm has a slight increase than P(VDF-CTFE) lm, there is a signicant suppression compared with BT/P(VDF-CTFE) nanocomposites doping untreated BT. The inferior dielectric loss of BT/P(VDF-CTFE) nanocomposites is derived from the easy agglomeration of untreated BT nanoparticles in the polymer matrix and the poor compatibility between the two components, causing the stronger space-charge polarization and high dielectric loss. 17,31 For the PDFMA@BT/P(VDF-CTFE) nanocomposite lms, the modied BT nanoparticles with PDFMA shells can effectively improve the dispersibility of the BT nanoparticles in the P(VDF-CTFE) matrix and enhance the interfacial adhesion, which limits the macromolecular chain motion of the polymer matrix and the accumulation of charges. 38

Conclusions
In summary, a facile "thiol-ene" click reaction between thiol groups functionalized PDFMA (PDFMA-SH) and olenic bonds contained BT nanoparticles is used to reduce the surface energy of BT NPs and the corresponding PDFMA@BT/P(VDF-CTFE) nanocomposite lms were prepared. The results show that the core-shell structured PDFMA@BT nanoparticles are more uniformly dispersed in the P(VDF-CTFE) matrix and exhibit good compatibility compared to the unmodied BT nanoparticles. When the content of nanoparticles is 20 vol%, the breakdown strength of PDFMA@BT/P(VDF-CTFE) nanocomposite lm is 1.3 times that of BT/P(VDF-CTFE) nanocomposite lm. In addition, PDFMA@BT/P(VDF-CTFE) nanocomposite lms exhibit a significant suppression of dielectric loss. It is proved that the uoropolymer shells have a positive effect on the improvement of the dielectric properties and breakdown strength.

Conflicts of interest
Author declares no conict of interest for this research paper.