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High Energy Density Performance of Polymer Nanocomposites Induced by Designed Formation of BaTiO@Sheet-like TiO Hybrid Nanofillers 3

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Ran Su, Zhengdong Luo, Dawei Zhang, Yang Liu, Zhipeng Wang, Junning Li, Jihong Bian, Yanxi Li, Xinghao Hu, Jinghui Gao, and Yaodong Yang J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b01853 • Publication Date (Web): 25 May 2016 Downloaded from http://pubs.acs.org on May 31, 2016

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The Journal of Physical Chemistry C is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

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The Journal of Physical Chemistry

High Energy Density Performance of Polymer Nanocomposites Induced by Designed Formation of BaTiO3@sheet-likeTiO2 Hybrid Nanofillers Ran Su ,1 Zhengdong Luo,2 Dawei Zhang,1 Yang Liu,3 Zhipeng Wang,1 Junning Li,1 Jihong Bian,1 Yanxi Li,4 Xinghao Hu,5Jinghui Gao,5 Yaodong Yang*,1 a)

1

Frontier Institute of Science and Technology, State Key Laboratory for Mechanical

behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, China

2

Department of Physics, University of Warwick, Coventry CV47AL, UK

3

Laboratoire Structures, Propriétés et Modélisation des Solides, CentraleSupélec,

CNRS-UMR8580,

Université

Paris-Saclay,

Grande

Voie

des

Vignes,

Châtenay-Malabry Cedex 92295, France

4

Department of Materials Science and Engineering, Virginia Tech, Blacksburg,

Virginia 24061, USA

5

State Key Laboratory of Electrical Insulation and Power Equipment, Xi’an Jiaotong

University, Xi’an 710049, China

a) [email protected]; Tel: +86-029-83395132

1

ACS Paragon Plus Environment


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ABSTRACT: Nanocomposites incorporating inorganic nanoparticles (NPs) within ferroelectric polymeric matrices have great potential for high density energy capacitors. In this strategy, employing the nanostructures with specially designed morphology as fillers would notably improve the energy storage. However, this strategy is extremely challenging and has not been largely explored. Here, the BaTiO3@sheet-likeTiO2 core-shell NPs have been successfully synthesized, and can be well-dispersed into PVDF matrices. The nanocomposites with the volume fraction 2.5 % BT@TiO2 NPs have higher the electric displacement (6.0 µm/cm2) than that of PVDF films with the 2.5 vol % BaTiO3 (BT) NPs (5.1 µm/cm2) under the same electric field of 350 kV/mm, which is mainly ascribed to the hierarchical interfacical polarization induced by the large surface area of TiO2 sheet assembled on BT NPs in the nanocomposites. Simultaneously, the medium permittivity TiO2 with medium Ɛr as a buffer layer between the BT NPs and polymer matrix could minimize the inhomogeneous electric field in the nanocomposites, which results in the enhancement of the breakdown strength (490 kV/mm with 2.5 vol % BT@TiO2 NPs) compared to the pristine PVDF. As a result of a gratifying EB and D, the energy density of the nanocomposites with 2.5 vol % BT@TiO2 NPs reached 17.6 J/cc.

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INTRODUCTION

Capacitors with high energy density play a key role in advanced pulsed-power electronic devices and electric power systems.1 As a result, the development of materials with high energy density becomes an inexorable trend.2-3 The current surge of focus on ferroelectric polymer is fueled by the significant advantage as a pivotal energy density material compared to other capacitors materials such as conventional ceramics and ceramic thin films, including high breakdown strength, good mechanical flexibility, cheap price, light weight and desired level of process.4-5 However, the intrinsic electric permittivity in polymers is low, which limits the realization of higher performances.6-7 Thus the critical issue is to enhance the electric permittivity while maintaining other excellent features of the ferroelectric polymers, particularly the high breakdown strength. The design of proper nanocomposites, by incorporating inorganic nanoparticles into polymer matrices, represents one of the most promising avenues to solve the above conundrum.8-11 Indeed, plenty of studies unambiguously have demonstrated that the high dielectric permittivity of the polymer could be straightforward improved by adding inorganic fillers with the high dielectric permittivity such as TiO2, ZrO2, BaTiO3, calcium copper titanate (CCTO) and (Ba,Sr)TiO3.12-16 However, as aforementioned, the dielectric breakdown strength of those nanocomposite films decreases with increasing the doping amount of inorganic fillers. Recently, it is found that the weakness can be overcome by adjusting the special structure and morphology of the fillers. For instance, H. Tang et al. reported that the quenched PVDF and Ba0.2Sr0.8TiO3 NWs can 3



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obtain high breakdown strength (450 kV/mm) and good energy density (14.86 J/cc) due to the use of fillers with high aspect ratio rather than those of the equiaxial particles.17 Moreover, the core-shell structure nanoparticles as dielectric fillers, consisting of medium-Ɛr shells (such as Al2O3 or TiO2) and high-Ɛr cores (such as BaTiO3 or PbTiO3), were employed to obtain the enhanced breakdown strength.18-19 As a result, the dielectric breakdown strength of polymer nanocomposites could approach and even surmount the intrinsic dielectric breakdown strength of the corresponding polymer matrix. More speculatively, those composites approach not only takes full advantage of the impressive permittivies and breakdown strength of nanoparticles, but also could introduce other novel properties, such as the exchange coupling effect through a dipolar interface layer between the nanoparticles and the polymer matrix.11, 13 Hence the dielectric permittivity could be further improved by the interfacial polarization at the interfaces between the nanoparticles and the polymer matrix. The area of the interfaces is largely determined by the structure of nanoparticles,

proportional

to

the

interfacial

polarization

of

polymer

nanocomposites.20-22 In general, to achieve desired permittivity and breakdown strength for high energy densities of polymer nanocomposites, the rational design and precise control of the structure and component of nanofillers is crucial. In this study, we report a novel method for the controllable synthesis of BaTiO3@sheet-like TiO2 core-shell nanoparticles as fillers for nanocomposite capacitors with high energy density. It is found that the significantly enhanced polarization and high breakdown strength are achieved simultaneously, leading to 4


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excellent energy densities of more than 17.6 J/cc. These outstanding results can be attributed to the special hetero-nanostructure of fillers with high intrinsic permittivity in the polymer matrix, and the special hetero-nanostructure of fillers will induce well-dispersed, hierarchical interfaces and a buffer layer effect.

METHODS

Preparation of BaTiO3 NPs. A mixture of 22 g NaOH and 18g KOH was put into a 50 mL Teflon-lined stainless steel autoclave, followed by addition of 0.3g Ba(NO3)2 and 0.08g TiO2. The autoclave was then sealed and heated at 210 °C for 48 h. After cooling the reaction to room temperature, the products were washed with deionized water and dilute hydrochloric acid to remove excessive hydroxide and barium carbonate. Preparation of BaTiO3@sheet-like TiO2 Core-Shell NPs. The BaTiO3 nanocubes (0.04 g) were dispersed in 30 mL isopropanol, followed by the addition of 0.03mL DETA. After being stirred gently, 2 mL TIP was added to the solution. The mixture was then transferred into a Teflon-lined stainless steel autoclave with a capacity of 50 mL and kept at 200 °C for 24 h. Finally, the products were washed with ethanol and collected by centrifugation. The annealed BT@TiO2 NPs were obtained at 400 °C in static air. The densities of BaTiO3 NPs (5.8 g/cm3) and BaTiO3@TiO2 nanosheets core-shell NPs (5.3 g/cm3) are directly obtained by traditional dewatering method. Preparation of PVDF-based Nanocomposite Films. The nanocomposite films were prepared by dispersing the nanofillers into a 5% weight fraction of PVDF in 5



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dimethylformamide (DMF) solution by stirring for 10 hours to make it stable and homogenous. The films were then dried at 70 °C under vacuum overnight followed by heating at 200 °C for 5 min, then immediately quenched in an ice-water bath. The final quenched films were dried at 70 °C and peeled from the substrates. Materials Characterization. The morphology and crystalline structure of the nanowires were characterized by Scanning Electron Microscopy (FE-SEM; SU-8010, HITACHI), Transmission Electron Microscopy (TEM, 2100 JEOL) at an acceleration voltage of 200 kV and powder X-ray diffractometry (XRD, Shimadzu 7000) with a Cu Kα (0.15418nm) source. The piezoelectric properties and I-V curves of nanoparticles were investigated by an atomic force microscope (AFM, Cypher, Asylum Research). Thermogravimetric analysis (TGA) was carried out under a flow of air with a temperature ramp of 5 °C/min. Fourier-transform infrared (FTIR) spectroscopy of BT@TiO2 NPs was performed with a NICOLET 6700 FTIR. Frequency-dependent dielectric permittivity constant and loss were measured using an Agilent 4980A LCR meter with a frequency range from 100 Hz to 1 MHz at 1 Vrms with a parallel equivalent circuit. The D-E loops of the nanocomposites were measured by a Sawyer-Tower circuit (Radiant Workstation), which allows direct computation of the energy density.

RESULTS AND DISCUSSION

Herein, the uniform BaTiO3 nanocubes were synthesized by a hydrothermal reaction (Figure 1a) and the ferroelectric properties of BaTiO3 nanocubes were measured by 6


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the piezoresponse force microscope (PFM).23 Local piezoelectric hysteresis loops of the BaTiO3 nanocubes present 180˚ phase switching (Figure S1a), indicating the polarization switching between two antiparallel polarization states along with external field. The piezoelectric amplitude also illustrates that BaTiO3 nanocubes have stable ferroelectric state. And then, the BaTiO3@TiO2 NPs were prepared.24-25 In this process, the surfactant of tridentate DETA (diethylenetriamine) played a critical role in the successful synthesis of the core-shell nanostructure. Transmission electron microscopy (TEM) observations in Figure 1b reveal that the BT nanocubes are well-wrapped by a large number of TiO2 nanosheets. Also the rippled and crimpled structure indicates that TiO2 sheets are only a few layers’ thick. It is found that the hybrid nanoparticles (NPs) have a uniform size and morphology as shown in the scanning electron microscopy (SEM) images (Figure S2). Moreover, Local piezoelectric hysteresis loops of the TiO2 nanosheets/BT nanocube also present 180˚ phase switching (Figure S1b). However, the polarization switching voltage of the TiO2 nanosheets/BT nanocube is much higher than that of the BT nanocube. This behavior is due to the fact that the switching behavior of the polarization is limited by the interfaces between TiO2 and BT. It is worth mentioning that the as-obtained antase TiO2 shells were stabilized by amine groups due to the existence of DETA in the reaction and poor crystallinity. The TiO2 shell with high crystallinity can be obtained by the calcination at 400 °C. Unfortunately, the TiO2 sheets shrink considerably as the consequence of the loss of amine groups, and the BaTiO3@TiO2 core-shell NPs are slightly agglomerated in the 7



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process of annealing (Figure S3). The thermogravimetric analysis (TGA) analysis revealed that BaTiO3@TiO2 core-shell NPs contained ~16 wt % of surfactant. By contrast, only few organic species remain (~3 wt %) for the nanoparticles annealed (Figure 2a). Fourier-transform infrared (FITR) spectrums of BaTiO3@TiO2 NPs before and after annealing were shown in Figure 2b. The loss of the amine groups on the surface of BaTiO3@TiO2 NPs after annealing is also evidenced by the disappearance of the peak at 1450 cm-1. The examination of individual TiO2 nanosheet with high-resolution (HR) TEM shows that the lattice fringes with interplanar spacing d001=1.90Å is consistent with anatase crystal structure (Figure 1c). The crystal structure of TiO2 shell is also verified via the selected-area diffraction pattern (SAED). Diffraction rings, from inside out, can be indexed to the (101), (105), and (204) facets of anatase TiO2, respectively. Figure 1d shows XRD patterns of the as-synthesized products of the BaTiO3 nanocubes, BaTiO3@TiO2 core-shell nanoparticles before and after calcination at 400 oC, respectively. The peaks of bottom curve could be easily indexed to perovskite (tetragonal phase) crystal structure of BaTiO3. As shown in the middle plot in Figure 1d, the characteristic peaks of TiO2 phase were vaguely observed, indicating that the solvothermal synthesized TiO2 shell is poorly crystalline. After the annealing, the emergence of characteristic peaks of the anatase-TiO2 indicates the mixture phases of BT and TiO2. The 10 µm nanocomposites with nanofillers were prepared by a solution-casting process. To further improve the ferroelectric properties of nanocomposites, the γ-PVDF phase was obtained by annealing and quenching in ice water.17 The surface 8


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and cross-sectional images of nanocomposites were simultaneously acquired by SEM (Figure 3), which clearly shows that the BaTiO3@TiO2 NPS are homogeneously dispersed in the polymer matrix and no voids exist in the polymer matrix. The effect of nanofillers on microstructure of the PVDF matrix was analyzed by differential scanning calorimetry (DSC). The DSC profiles of 2.5% vol nanocomposites with BT, BT@TiO2 NPs, BT@TiO2 annealled NPs and pure PVDF were presented during the heating scan (Figure 4). As can be seen, the endothermic peaks of the nanocomposites with BT@TiO2 NPs have slight shift and broad trend compared to that of BT NPs, BT@TiO2 annealed NPs and pure PVDF, which indicates that the nanofillers promote the crystallization process of PVDF matrix and improve the thermal stability of PVDF matrix.11 In general, the results benefit from the special morphology of BT@TiO2 NPs in that the nanosheets with high specific surface area can facilitate dispersion of nanofiller into polymer matrix.26 Let us focus on the dielectric response of nanocomposites containing BT NPs, BT@TiO2 NPs with different volume fractions. As expected, enhanced dielectric permittivity is observed as the volume fraction of fillers increases. This is mainly because the nanoparticles have a higher dielectric constant than the PVDF matrix (Figure 5).27-28 Figure S4 summarizes the dielectric permittivity of nanocomposites containing BT NPs, BT@TiO2 NPs with volume fraction of 2.5% fillers the pure PVDF thin films. It is shown that the dielectric permittivity of the nanocomposites filled with BT@TiO2 NPs is higher than that of nanocomposites filled with BT NPs. This is due to the result of the production of hierarchical interfacial polarization in 9



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nanocomposites by employing the topological structure of BT@TiO2 NPs. To be specific, these two kinds of interfaces may be induced among BT core, TiO2 shell and PVDF matrix according to the physical properties and placement contrast in nanocomposites, including the ‘soft interface’ (between TiO2 nanosheet and polymer matrix) and the ‘hard interface’ (between BT nanocubes and TiO2 nanosheet). Those two types of interfacial mechanisms simultaneously evoked the permittivity response in nanocomposites filled with BT@TiO2. On one hand, the electrical double-layer may be induced by diffusion of the mobile space charges in ‘soft interface’ zone between TiO2 nanosheets and PVDF.29-30 In nanocomposites of BT@TiO2 NPs, this effect can be reinforced due to the large surface area of sheets like TiO2, which is the dominating factor for the proposed mechanism. Moreover, the interfacial regions around adjacent TiO2 sheets on individual BT@TiO2 NPs may probably by forming percolation paths within the diffuse layers. The infiltration of mobile charges would greatly enhance the dielectric permittivity.31 On the other hand, internal strain, vacancies and dead layer generated by the additional interfaces (‘hard interface’) between TiO2 shell and BT core could induce the mobile space charges, which can further enhance the dielectric permittivity.32-33 Meanwhile the hard interface effect is indirectly indicated by the switching hysteresis behavior of I-V cures of BT@TiO2 nanocubes (Figure S1c). The electrical hysteresis behavior of I-V cures of BT@TiO2 nanocubes are attributed to a high oxygen vacancy concentration in BT@TiO2 nanocubes, which could lead to a nonlinearity response under different voltage.34-35 10


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According to the results by the aforementioned TGA analysis and FITR spectrums, the surface area shrinks considerably after annealing. This result in a smaller dielectric permittivity of annealed nanocomposites with annealed BT@TiO2 NPs compared to that in the pristine unannealed counterparts. Interestingly, it is found that of dielectric permittivity of the nanocomposites with annealed BT@TiO2 NPs is higher than that in the nanocomposites with unannealed BT NPs. However, the dielectric permittivity of BT@TiO2 NPs/PVDF has a significant decline than that of BT NPs/PVDF and PVDF under the high frequencies (Figure S4c), which is attributed to the slow response of space charges with the high frequencies. This indicates the existence of the ‘hard interface’ in hybrid annealed BT@TiO2 NPs. Moreover, it is found that the promotion of dielectric permittivity in BT@TiO2/PVDF is not at the cost of the dielectric loss (Figure 5). The low dielectric loss in BT@TiO2/PVDF may benefit from the homogeneous good dispersion of BT@TiO2 in polymer matrix, compared to that in BT/PVDF and BT@TiO2 annealed/PVDF. The typical electric displacement (D) versus electric field loops were measured under a unipolar field at 10 HZ with various peak electric fields as shown in Figure 6 (and Figure S5). Notably, the polarization of the nanocomposites was greatly enhanced with the increasing concentration of the nanoparticles as a result of the high dielectric permittivity of the inorganic fillers (Figure S5).27-28 It is found that the electric displacement of nanocomposites containing BT@TiO2 NPs are much higher than that of the nanocomposites of BT NPs with the same concentration and fields (Figure S5a and S5b), which is also due to a stronger interfacial polarization induced 11



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by the topological and composite nanostructure. For instance, the electric displacement of nanocomposites with 2.5 vol % loading of BT@TiO2 reaches 6.0 µm/cm2, while the displacement of BT/PVDF film is 5.1 µm/cm2 under the same electric field of 350 kV/mm. Moreover, while the nanocomposites of annealed BT@TiO2 NPs have a higher (comparable) electric displacement than the nanocomposites containing BT NPs under the same electric field, the maximum value of electric displacement is constrained due to relatively larger dielectric losses. Let us concentrate on the electric breakdown strength (EB), which is also a key parameter in determining the energy density. It is known that the breakdown strength of nanocomposites usually decreases with the increasing concentration of fillers (see Figure S5).27-28 The typical results are summarized in Figure 6b. It is shown that the dielectric breakdown strength of BT NPs/PVDF nanocomposites decreases from ca. 440 to ca. 370 kV/mm as the volume fraction of BT NPs increases from 1% to 5%. On the contrary, the BT@TiO2 NPs/PVDF nanocomposites could maintain relatively high electrical breakdown strength even with inorganic fillers. The electric breakdown strength of nanocomposites with volume fraction 2.5% BT@TiO2 NPs (ca. 490 kV/mm) is only slightly smaller than the pure polymer (ca. 500 kV/mm), while electric breakdown strength of the BT NPs/PVDF nanocomposites drops down drastically to ca. 350 kV/mm at (with same volume fraction). The stabilization of EB is analyzed using a two-parameter Weibull cumulative probability function: P(E)=1-exp[1-(E/EB)β], where P(E) is the cumulative probability of failure occurring at the electric field; EB is a scale parameter referring to the breakdown strength for a 12


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63.2%; E is the experimental breakdown strength; β is the shape parameter associated with linear regressive fit of distribution. The β of the PVDF/ BT@TiO2 NPs is similar with that of the PVDF/ BT NPs and slightly below that of the pure PVDF. It is known that the dispersion of inorganic fillers in the polymer matrix and its giant permittivity in contrast with the matrix are problematic, both of which could lead to a reduction of the breakdown strength.14,

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The inhomogeneous mixture between

nanoparticles and polymer could yield excessive defects.14 We argue that the striking feature of the greatly strengthened breakdown field found here mainly arises from the topological nanostructure of fillers. Firstly, the strong interaction between amino groups absorbed on the fillers and PVDF will be induced. As a result, the BT@TiO2 NPs could be more homogenously dispersed in the polymer matrix.36 In addition, the 2D structure of TiO2 shells can also be effective in improving the dispersion of fillers within the polymer matrix.26, 37 After being annealed, the BT@TiO2 NPs loose the majority of amino groups and can even be reunited. Therefore, the dielectric breakdown strength sharply decreases with the increased concentration due to the introduction of more defects (see Figure S5c). Second and most importantly, the TiO2 nanosheet could be regarded as a buffer layer owing to a middle permittivity between PVDF and BT.18-19, 30 When BT NPs with a high permittivity are directly in contact with PVDF matrix with a low permittivity, the large contrast of permittivity between the two phases could create an even larger distortion and inhomogeneity of electric fields. The hetero-nanostructure of BT@TiO2 NPs with grading permittivities from the center to border could prevent this effect and thus generally give rise to the 13



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significant reduced current leakage.38 Our results in Figure 6 demonstrate that the enhancement in electrical displacement and the retention in breakdown strength could be achieved by using the novel structure of BT@TiO2 nanosheet fillers. The introduction

of

well-designed

nanoparticles

into

polymer

matrices

could

simultaneously achieve higher values of D and EB, which indicates the promise of BT@TiO2 nanosheets fillers in high energy density storage applications. As shown in Figure 7a, a remarkable energy density of 17.6 J/cc is obtained in the nanocomposite filled with 2.5 vol % of BT@TiO2 under an electric of 490 kV/mm, which is higher than that of pure polymer matrix (11.5 J/cc) and 1366% greater than the reported commercial capacitor (biaxial oriented polypropylene) with the energy density of 1.2 J/cc at 640 kV/mm.17 Based on the above observation and investigations, we present the schematic of the origin of the ultrahigh energy density for BT@TiO2 NPs (as illustrated in Figure 7b and Figure S6). The large specific surface area of TiO2 nanosheets coated on BT core in the nanocomposite could induce much stronger interfacial polarization by diffusion of space charges. Furthermore the motions of mobile space charges between neighboring layers of TiO2 significantly increase the charge polarization. In addition, the enhanced dielectric polarization could also be attributed to the introduction of additional interfaces between TiO2 shell and BT core. The formation of amino surface groups greatly enhances the dispersibility of the fillers in organic media. Even more essentiality, by establishing a buffer ceramic layer of medium permittivity between the BT NPs and polymer matrix, the grading permittivity from the center of filler to the border could minimize the 14


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inhomogenous electric field in the polymer matrix. To further understand the mechanism of dielectric breakdown, we employ the finite element method to simulate the electric field distribution in the nanocomposites. With an applied vertical electric field of 500 kV on the nanocomposites with a thickness of 10 µm, it is clearly found that the distribution of electric field turns inhomogeneous when doping the BaTiO3 NPs into PVDF film (Figure 7c and Figure S7a, b), and the high electric field appears primarily in the interface between the BaTiO3 NPs and PVDF matrix. When coating TiO2 nanosheets on BT NPs, the distribution of electric field of the PVDF nanocomposites becomes more homogeneous than that of the BaTiO3/PVDF (Figure 7d and Figure S7c, d). Compared with the smooth TiO2 shell, TiO2 nanosheets could lead to point discharge. Furthermore, the intensified yellow area (high electric field distribution) between the neighboring TiO2 nanosheets in the PVDF matrix is induced, which is more easily connected to form percolated interfacial regions. In general, the precise control of morphology and component of fillers that enhances D and maintains EB should be responsible for the high energy density.

CONCLUSION

In summary, we have successfully designed a novel dielectric nanocomposites filled with special structure of fillers, which is formed by self-assembled TiO2 nanosheets onto BT NPs. It is demonstrated that the significant enhancement of dielectric polarization could be induced by the modulation of hierarchical interfaces besides high intrinsic permittivity, benefiting from the existence of larger area interfaces 15



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between BT@TiO2 NPs and polymer matrix. Meanwhile, the TiO2 nanosheets as a buffer layer coated on BT NPs is for the purpose of facilitating dispersibility and decreasing the permittivity contrast between filler and matrix, and thus it increases the breakdown strength. As a result of the gratifying EB and D, the nanocomposites with 2.5 vol % BT@TiO2 NPs exhibited a higher energy density (17.6 J/cc).

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Associated Content Supporting Information Available: Piezoresponse phase, amplitude and I-V curves for BaTiO3 nanocubes and BT@TiO2 core-shell NPs. TEM images of TiO2 nanosheets and BT@TiO2 NPs annealed. SEM images of the BT@TiO2 core-shell NPs. Frequency-dependent dielectric constant and dielectric loss of PVDF and nanocomposites. The P-E loops of composites films. Energy density of PVDF nanocomposites of BT@TiO2, BT and BT@TiO2 annealed NPs with different volume fraction. Data on the finite element simulation.

Acknowledgements

The authors gratefully acknowledge Chuansheng Ma, Guang Yang from International Center for Dielectric Research at Xi’an Jiaotong University for his help to take TEM images. We would like to thank Linglong Li from Xi’an Jiaotong University for the useful discussion and help during electrical properties characterization too. This work was supported by the Ministry of Science and Technology of China through a 973-Project under Grant No. 2012CB619401, Natural Science Foundation of China (Grant Nos. 11204233, 51431007, 51321003), the fundamental research funds for the central universities and MOE innovation team (Grant No.IRT13034) both from Ministry of Education. Y.L. acknowledges the China Scholarship Council (CSC) for funding Y.L.'s stay in France.

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Figure 1. (a) TEM image of the BT nanocubes (scale bar=500 nm). Inset: HRTEM image of individual BT NPs, scale bar=5 nm. (b, c) TEM and HRTEM images of BT@TiO2 NPs (scale bar=100 nm) and TiO2 nanosheet (scale bar=5 nm). Inset: the SAED pattern of TiO2 nanosheets, scale bar=2 1/nm. (d) XRD patterns of BT, BT@TiO2 NPs and BT@TiO2 annealed NPs.

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Figure 2. (a) TGA curve of the BT@TiO2 and BT@TiO2 annealed NPs. (b) FTIR spectra of the BT@TiO2 NPs and BT@TiO2 NPs annealed.

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Figure 3. Surface and cross-sectional SEM images of composites films with BT@TiO2 core-shell NPs volume ratio of 10% (scale bar=5 µm). Inset: Photo images of PVDF/BT@TiO2 core-shell NPs nanocomposites from with BT@TiO2 core-shell NPs volume ratio of 2.5%.

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Figure 4. DSC curves of the polymer and 2.5 vol% nanocomposites with BT, BT@TiO2 and BT@TiO2 annealed NPs.

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Figure 5. Frequency-dependent dielectric constant (a, b, c) and dielectric loss (d, e, f) for pure PVDF film and the nanocomposites with the different nano-inclusion volume fraction.

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Figure 6. (a) Unipolar electric displacement-electric field loops of composites films of BT and BT@TiO2 NPs with different volume fraction of 2.5%. (b) Variation of maximum electric displacement and characteristic breakdown strength with different volume fraction of nanofillers for PVDF nanocomposites embedded with BT and BT@TiO2 NPs. (c) Weibull distribution and observed dielectric breakdown strength of nanocomposites with BT NPs (2.5 vol%), BT@TiO2 NPs (2.5 vol%), and pure PVDF.

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Figure 7. (a) Energy density of PVDF nanocomposites embedded with BT@TiO2 NPs, BT NPs and pure PVDF films as a function of electric field calculated from D-E loops. (b) Schematic illustration of the origin of the ultrahigh energy density for PVDF/BT@TiO2 NPs film. (c, d) Electric field distribution in the nanocomposites with BaTiO3 NPs, BT@TiO2 nanosheets NPs.

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