High energy density and discharge efficiency polypropylene nanocomposites for potential high-power capacitor

Journal Pre-proof High energy density and discharge efficiency polypropylene nanocomposites for potential high-power capacitor Biao Liu, Minhao Yang, Wen-Ying Zhou, Hui-Wu Cai, Shao-Long Zhong, Ming-Sheng Zheng, Zhi-Min Dang PII:

S2405-8297(19)31084-0

DOI:

https://doi.org/10.1016/j.ensm.2019.12.006

Reference:

ENSM 1012

To appear in:

Energy Storage Materials

Received Date: 22 October 2019 Revised Date:

19 November 2019

Accepted Date: 5 December 2019

Please cite this article as: B. Liu, M. Yang, W.-Y. Zhou, H.-W. Cai, S.-L. Zhong, M.-S. Zheng, Z.-M. Dang, High energy density and discharge efficiency polypropylene nanocomposites for potential highpower capacitor, Energy Storage Materials, https://doi.org/10.1016/j.ensm.2019.12.006. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

High energy density and discharge efficiency polypropylene nanocomposites for potential high-power capacitor Biao Liu

a,b

, Minhao Yang c,*, Wen-Ying Zhou b, Hui-Wu Cai

b,*

, Shao-Long Zhong a, Ming-

Sheng Zheng a, Zhi-Min Dang a,* a

State Key Laboratory of Power System and Department of Electrical Engineering, Tsinghua

University, Beijing, 100084, China b

School of Chemistry and Chemical Engineering, Xi’an University of Science and Technology,

Xi’an, 710054, China c

Institute of Advanced Materials, North China Electric Power University, Beijing, 102206,

China

High

energy

polypropylene

density

and

discharge

nanocomposites

high-power capacitor

1

for

efficiency potential

Abstract The film capacitor, one typical type of electrostatic capacitors, exhibits its unique advantages in the high-power energy storage devices operating at a high electric field due to the high electrical breakdown strength (Eb) of the polymeric films. However, the development of film capacitor towards high energy storage density is severely hindered by the low dielectric constant (ε) and low charge-discharge efficiency (η) of the polymeric films. The film of polypropylene (PP), the most used polymeric film with a market share of 50%, owns a high η due to its low inherent hysteresis loss. Yet the low ε (2.2 at 103 Hz) impedes the increase of its energy storage density (1-2 J/cm3). Here we demonstrate that the discharged energy density (Ue) of PP film could be largely increased from 1.40 J/cm3 of pure PP film to 3.86 J/cm3 of PP nanocomposite film by incorporating a small loading of core-shell structured PMMA@BaTiO3 (PMMA@BT) nanoparticles (2.27 vol%) into PP matrix. The obtained Ue of 3.86 J/cm3, to the best of our knowledge, is the highest reported value for the PP based films using commercial PP resin. Meanwhile, the η merely undergoes a slight decrease from 99.5% to 94.1%. Similarly, the obtained η of 94.1%, to the best of our knowledge, is also the highest value for the polymeric films reported in the previous works. The significant increase of Ue (175.7%) and negligible decrease of η (5.4%) are mainly attributed to the increase of ε from 2.2 to 3.7 and Eb from 361 MV/m to 448 MV/m. Compared with the films of raw BT/PP nanocomposites, the Ue exhibits a significant increase of 365.1% from 0.83 J/cm3 to 3.86 J/cm3 and the η also displays an increase of 7.4% from 87.6% to 94.1%. The significant improvement of energy storage performance achieved in the film of PMMA@BT/PP nanocomposite suggests that the organic PMMA shell plays an important role in the amelioration of compatibility of BT nanoparticles and PP matrix and the alleviation of local electric field concentration. Moreover, the PMMA shell could also provide a robust scaffold to hinder the early breakdown failure of nanocomposites due to its high Eb (425 MV/m). Thus, high ε, low loss and high Eb were simultaneously achieved. Specifically, the high hot stretch ratio (1:4) of nanocomposite film indicates the strong feasibility of 2

industrialized film processing. Apart from the experimental analysis, theoretical analysis using the simulation of finite element is also carried out to figure out the influence of organic PMMA shell on the dielectric performance of PP nanocomposites films. These findings enable the development of film capacitor using non-polar PP based film towards high energy storage density.

Keywords： dielectric properties energy storage polypropylene nanocomposites high discharge efficiency core-shell structure

3

1. Introduction The ever-increasing demand for the development of high-power energy storage devices has been driven by the rise of high-power applications such as electric vehicles, advanced weapons and equipment, and ultra-high voltage direct current (HVDC) transmission systems [1-8]. The electrostatic capacitor, which stores and releases energy electrostatically, possesses the highest power density among various energy storage technologies [3]. Among various types of electrostatic capacitors, film capacitor exhibits its unique strengths in the high electric field environment due to the high electrical breakdown strength (Eb) of the dielectric polymer inside. However, the application of film capacitor in those high-power fields is severely hindered by its low energy storage density [6, 9, 10]. The energy storage density of a film capacitor is generally determined by the energy storage density of the dielectric polymer sandwiched between two electrodes. In general, the maximum energy storage density (Um) of a linear dielectric layer scales quadratically with its Eb and linearly with its dielectric constant (ε) according to the following the equation

=

, where

ε0 represents the vacuum dielectric constant [6]. Although Eb seems to be the most critical parameter in determining Um, the biaxially oriented polypropylene (BOPP) film with a high Eb of 600 MV/m, the state-of-the-art commercially available dielectric polymer, can only exhibit an energy storage density of 1-2 J/cm3 due to the low intrinsic ε (2.2) of PP [11, 12]. Recently, T. C. Mike Chung et al. reported that the energy storage density of PP based film could be significantly improved by using specially designed PP copolymer or cross-linkable PP copolymer [13-15]. For instance, a giant energy density of 7.42 J/cm3 was achieved in the PP copolymer with an appropriate amount of -OH side groups [13]. The formation of H-bondings between adjacent -OH groups not only offered a high crystallinity for PP copolymer but also produced a stable network for the reversible polarization-depolarization of segments [13]. Thus, high polarizability and good reversibility were simultaneously achieved in the PP copolymer. Additionally, the stable network used for the reversible electronic polarization in PP copolymer could also be constructed by cross-linking the 4

PP copolymer, which contributed to the increase of ε and Eb [14]. Therefore, a high energy storage density (>5 J/cm3 at 650 MV/m) was achieved [14]. However, the complex molecular design and polymerization of comonomer make it difficult to replace the commercially used PP homopolymers with the reported PP copolymer shortly. As a result, numerous efforts have been made to address the issue of the low ε of commercially used PP film [2]. Generally, the strategies employed can be mainly divided into two categories. A straightforward approach is replacing the non-polar PP with ferroelectric polymers, represented by poly(vinylidene fluoride) (PVDF) and its copolymers, as a novel dielectric layer for the high energy density film capacitors due to their high ε (~10.0) [16-18]. However, owing to their high intrinsic ferroelectric losses and remnant polarizations, the charge-discharge efficiency (η) achieved in previous works is usually lower than 80%. Besides, the process of ferroelectric polymer pellets into a thin film is more difficult than the process of PP and the applied melt-stretching or shearing force will inevitably induce the partial conversion of non-polar α-phase into polar β-phase [19]. Such a conversion will cause an increase of ferroelectric loss and remnant polarization, accordingly further decreasing the η. Therefore, the unique advantages of non-polar PP at low cost, ultralow dielectric loss, high Eb and easy processability make it challenging to replace PP with PVDF or its copolymers in film capacitor field [20]. Another promising strategy for the increase of ε is to prepare nanocomposites by incorporating inorganic nanofillers with a high dielectric constant into the PP matrix [20]. In such a way, the nanofillers contribute to the enhancement of the ε, while the polymer matrix protects against the electrical breakdown [21]. Nevertheless, the ultrahigh surface energy of nanofillers caused by the enormous specific surface area encourages the inhomogeneous distribution of nanofillers and impairs the uniformity of their nanocomposites, thus resulting in the severe deterioration of Eb [22]. Additionally, the large difference in the ε of nanofillers and polymer matrix leads to the electrical field intensification surrounding the nanoparticles, thus giving rise to the early electrical breakdown of nanocomposites [23-25]. Consequently, a variety of 5

surface modifications strategies for the amelioration of nanofillers dispersion in the polymer matrix have been intensively attempted by linking the small organic molecules or long polymeric chains on the surface of nanofillers [7]. However, the method using small molecular modifier has some limitations in improving the Eb of nanocomposites because the modifiers themselves usually do not possess a high inherent Eb. For this reason, coating the nanofillers with a polymer shell to prepare core-shell structured nanohybrids provides an intriguing approach to ameliorate the dispersion of nanofillers [7, 9, 26, 27]. In addition, in order to alleviate the huge difference of ε between nanofillers and polymer matrix, the ε of the selected polymer shell should be larger than that of the host polymer matrix [28]. Apart from the structure and component of nanofillers, the concentration of fillers should be carefully considered due to the tradeoff between the difficulty in film processing and the enhancement of dielectric properties [2]. In terms of film processing, the loading of fillers should be kept at a small value to ensure the formation of a nanocomposite film without any defect. Although ε increases with the increase of fillers loading, the improvement in Eb only occurs when the nanocomposite is loaded with a slight fraction of fillers (< 3 vol%) and the further increase in the loading of fillers will lead to the deterioration of Eb and formation of voids [20]. Thus, the optimum loading of fillers for achieving the maximum energy density should be below the onset of the decrease in Eb due to the quadratic relationship between Eb and Um. Therefore, a small loading of fillers in the polymer matrix is desirable and reasonable for the scalable production of nanocomposites films and the achievement of a high energy density in the film capacitor industry. In this contribution, we report that a giant energy density and η could be simultaneously achieved by incorporating the core-shell structured PMMA@BaTiO3 nanoparticles into the PP matrix. PP is chosen as the host polymer because of its commercial availability in film capacitor, high η, and easy processability of film. Commercially available BaTiO3 nanoparticles are selected as the nanofillers due to its advantages of high ε and capability of scalable production. PMMA is chosen as the polymer shell on the surface of BaTiO3 nanoparticles owing to its intermediate ε 6

between PP matrix and BaTiO3 nanoparticles. In addition, its high Eb, strong ability to promote the dispersion of BaTiO3 nanoparticles in the PP matrix and easy synthesis by atom transfer radical polymerization (ATRP) method make it as a prominent polymer shell. Apart from the experimental analysis, the finite element analysis is also employed in this work to theoretically demonstrate the influence of PMMA shell on the electric field distribution. The principle towards the direction of industrialization during the selection of each material and the energy storage performance achieved in this work play a significant role in the development of high energy storage PP based film capacitors.

7

2. Experimental 2.1. Materials BaTiO3 nanoparticles with an average diameter of 60 nm, methyl methacrylate (MMA), N, N, N’, N’’, N’’-pentamethyldiethylenetriamine (PMDETA), acetone and copper bromide (CuBr) were purchased from Shanghai Aladdin Bio-Chem Technology Co., Ltd. Polypropylene (HC312BF) was obtained from Borealis Company

(Finland).

An

aqueous

solution

of

H2O2

(30

wt%),

N,

N-dimethylformamide (DMF), toluene, acetic acid and dichloromethane (CH2Cl2) were supplied by Sinopharm Chemical Reagent Co., Ltd. γ-Aminopropyl triethoxysilane (APS) was provided by Beijing Shenda Chemical Company. 2-bromoisobutyric acid was purchased from Acros Organics. Triethylamine was obtained from Shanghai Yuanye Bio-Technology Co., Ltd. All the chemicals were used as received without further purification. 2.2. Methods Synthesis of core-shell structured PMMA@BaTiO3 nanoparticles. The typical procedures for the synthesis of core-shell structured PMMA@BaTiO3 nanoparticles were carried out as follows. Surface hydroxylation of as-received BaTiO3 nanoparticles: Briefly, 20 g of BaTiO3 (BT) nanoparticles were homogeneously dispersed in 100 mL of an aqueous solution of H2O2 (30 wt%) under the ultrasonic treatment for 30 min. Subsequently, the mixture was refluxed at 105 °C under magnetic stirring for 5 h. After that, the hydroxylated BT (HO-BT) nanoparticles were collected by centrifugation at 5000 rpm for 5 min and washed with deionized water for three times. Finally, HO-BT nanoparticles were freeze-dried for 9 h, followed by being dried in a vacuum oven at 80 °C for 5 h. Functionalization of HO-BT nanoparticles by APS: In brief, 15 g of HO-BT nanoparticles were added into 100 mL of toluene and the solution was sonicated for 30 min. Acetic acid was added dropwise into the solution described above to adjust the pH to 4, and then 4 g of APS was added slowly through a dropping funnel. The 8

mixture was refluxed subsequently at 80 °C for 24 h. After that, the amino functionalized BT nanoparticles (APS-BT) were recovered by centrifugation at 5000 rpm for 5 min and washed with deionized water for two times. At last, APS-BT nanoparticles were dried in an oven at 80 °C for 24 h. Grafting of PMMA onto the surface of APS-BT: 13 g of APS-BT nanoparticles were dispersed into 100 mL of CH2Cl2 solvent homogeneously under the ultrasonic treatment for 30 min, and then 1 mL of triethylamine as a stabilizer was dropwise added into above solution. After that, a homogeneous solution containing 50 mL of CH2Cl2 solvent and 0.56 g of 2-bromoisobutyric acid as initiator was added slowly into above dispersion. Afterward, the obtained mixture was stirred at 0 °C for 3 h, followed by being stirred at room temperature for 12 h. After the reaction, the obtained nanoparticles were collected through centrifugation and washed with CH2Cl2 solvent for 2 times. Finally, the obtained Br-APS-BT nanoparticles were dried in a vacuum oven at 65 °C for 24 h. 4 g of synthesized Br-APS-BT nanoparticles were added into 50 mL of DMF solvent and the solution was sonicated for 30 min. 0.5 g of CuBr was subsequently added into dispersion. After that, a certain amount of MMA monomer was added into the above mixture. In order to obtain the core-shell structured PMMA@BT nanoparticles with different PMMA shells thicknesses, the weight ratio of Br-APS-BT nanoparticles/MMA monomer was settled as 1/5, 1/4, 1/2 and 3/4. Afterward, the obtained mixture was stirred magnetically at 60 °C for 30 min under the protection of nitrogen. 0.6 g of PMDETA solution was introduced into the obtained mixture and the solution was stirred magnetically at 60 °C for 24 h under the protection of nitrogen. After the reaction, the treated nanoparticles were recovered by centrifugation at 5000 rpm for 5 min and washed with acetone for several times. Finally, the obtained core-shell structured PMMA@BT nanoparticles were freeze-dried for 10 h, followed by being dried in a vacuum oven at 65 °C for 12 h. Preparation of the films of core-shell structured PMMA@BaTiO3/PP and raw BaTiO3/PP nanocomposites. The films of PMMA@BaTiO3/PP nanocomposites with different loadings of PMMA@BaTiO3 nanoparticles were prepared by the combination of melt blending and hot pressing methods. Different loadings of 9

PMMA@BaTiO3 nanoparticles and PP matrix were mixed in a Haake torque rheometer with a speed of 60 rpm at 190 °C for 10 min. The obtained mixtures were cut into small pieces and these pieces were subsequently hot pressed into thin films with a thickness of 25 µm-30 µm at 210 °C under the pressure of 125 MPa for 10 min. The films of BaTiO3/PP nanocomposites were prepared by a similar method. The derivation of the volume fraction of core-shell structured PMMA@BT nanoparticles is displayed in the Supporting Information. 2.3. Characterizations Scanning electron microscopy (SEM) images of BaTiO3 and PMMA@BaTiO3 nanoparticles and PP nanocomposites were taken on a field-emission SEM (SU8010, HITACHI) operated at 5 kV. Transmission electron microscopy (TEM) images of PMMA@BaTiO3 nanoparticles were obtained from a TEM instrument (HT7700, HITACHI) operated at 100 kV. The chemical structure of PMMA@BaTiO3 nanoparticles was characterized by Fourier-transform infrared (FTIR) spectrum with a spectrometer (Nicolet 6700, Thermo Fisher Scientific). X-ray diffraction (XRD) patterns were collected on a powder diffractometer (SmartLab 9, Rigaku) with a range from 10° to 90°. Thermogravimetric analysis (TGA) was carried out using an instrument (SDT Q600, TA Instruments) with a heating rate of 10 °C min−1 from 30 °C to 800 °C under an argon atmosphere. Differential scanning calorimetry (DSC) analysis was performed on an instrument (TA-60, Shimadzu) with a heating rate of 10 °C min−1 from 30 °C to 200 °C under a nitrogen atmosphere. The tensile tests were conducted on a tensile testing equipment (TS-2000, GOTECH). The dielectric properties were measured using an impedance analyzer (Concept 40, Novocontrol GmbH) at room temperature from 102 Hz to 106 Hz. The measurement of the dielectric breakdown strengths was performed on a high voltage tester (DC-2010, Jiangsu Shenghua) with 20 mm ball to ball electrode. The polymer films were placed between the two-sphere electrodes immersed in the silicone oil at room temperature. A DC voltage ramp of 500 V/s was applied to the electrodes until the polymer films were punctured. Twelve specimens were tested for each sample. Copper electrodes with a diameter of 4 mm were sputtered on both sides of pure PP and PP 10

nanocomposites films for the measurement of Electric displacement-electric field (D-E) loops, which were measured by a ferroelectric analyzer with a high voltage source of 10 KV (Precision Multiferroic, Radiant) at 10 Hz. 2.4 Finite element simulation of electric potential, electric field and current density. In order to theoretically investigate the effect of PMMA shell on the electric field and current density distribution surrounding the BT nanoparticles, the films of raw BT/PP and PMMA@BT/PP nanocomposites were further analyzed through finite element simulation (COMSOL Multiphysics). The rectangle size of the simulated sample was 1.6 µm × 1.6 µm. The dielectric constants of BT, PMMA and PP were set as 3500, 3.7 and 2.2, respectively. The conductivities of BT, PMMA and PP were assigned as 10-8 S/m, 10-16 S/m, 6.35 × 10-15 S/m, respectively [29-31]. The diameter of BT nanoparticles was set as 60 nm and the thickness of PMMA shell was assigned as 8 nm. A voltage of 10-9 KV was applied to the simulated sample, which decreases gradually from the top to the bottom.

11

3. Results and discussion

Fig. 1 (a) Schematic diagram illustrating the preparation of core-shell structured PMMA@BT nanoparticles by ATRP method. SEM images of (b) raw BT and (c) PMMA@BT nanoparticles. TEM images of PMMA@BT nanoparticles with different thicknesses of PMMA shells: (d) PMMA@BT-4, (e) PMMA@BT-6, (f) PMMA@BT-8 and (g) PMMA@BT-10 nanoparticles. (h) FT-IR spectra of raw BT, HO-BT, APS-BT and PMMA@BT nanoparticles. (i) TGA curves and (j) XRD patterns of raw BT and PMMA@BT with different thicknesses of PMMA shells. 12

Fig. 1a illustrates the synthesis process of core-shell structured PMMA@BT nanoparticles by the method of atom transfer radical polymerization (ATRP). ATRP method has been widely used for growing polymer on a variety of nanofillers surface owing to its advantages of uniform growth, high grafting density and high capping efficiency [7, 32]. The commercially available BT nanoparticles were used in this work because of its high ε and scalable production ability. The organic PMMA shell was chosen to reduce the surface energy of BT nanoparticles and improve the compatibility between BT nanoparticles and PP matrix. Firstly, a surface hydroxylation process of raw BT nanoparticles in H2O2 solution was employed to introduce functional hydroxyl groups (-OH) on the surface of BT nanoparticles. As shown in Fig. 1h, the appearance of a broad absorption peak at 3450 cm-1 confirms the existence of -OH groups on the surface of BT after the treatment of H2O2 solution [20, 32]. After that, HO-BT nanoparticles were reacted with APS to produce APS functionalized BT nanoparticles (APS-BT). The presence of absorption peaks at 2925 cm-1 and 2854 cm-1 suggests that APS is successfully introduced on the surface of BT nanoparticles [33]. Besides, the difference in the weight loss of raw BT and Br-APS-BT nanoparticles, calculated from the TGA curves in Fig. 1i, further indicates the successful introduction of APS on the surface BT nanoparticles [32]. The last step was grafting polymerization of MMA on the surface of BT nanoparticles. The characteristic absorption peaks at 1730 cm-1, 1190 cm-1 and 1150 cm-1 demonstrate that PMMA is successfully grafted on the surface of BT nanoparticles [32]. Additionally, the SEM images of raw BT and PMMA@BT nanoparticles show that the surface morphology of BT nanoparticles changes significantly after the polymerization, which further suggests the successful grafting of PMMA on the surface of BT nanoparticles (Fig. 1b and 1c) [20]. More importantly, the TEM images of PMMA@BT nanoparticles clearly indicate that a stable and dense polymer shell without any non-covered sites was directly coated on the surface of the BT nanoparticles, which could be precisely controlled from 4 nm to 10 nm by adjusting the concentration of MMA monomer (Fig. 1d, 1e, 1f and 1g). In addition, both the gradual increase in the weight loss of PMMA@BT nanoparticles (inset of Fig. 1i) and 13

gradual decrease in the intensity of XRD characteristic peaks of PMMA@BT nanoparticles (Fig. 1j) suggest the increase of PMMA shell thickness in the PMMA@BT nanoparticles [34, 35]. As shown in Fig. S1b, obvious aggregations of raw BT nanoparticles in the PP matrix could be easily observed. However, PMMA@BT nanoparticles disperse well in the PP matrix for the nanocomposite with the same loading of fillers (5 wt%) as shown in Fig. S1c. Moreover, even for the nanocomposite with a high loading of PMMA@BT fillers (10 wt%), the dispersion of PMMA@BT nanoparticles in the PP matrix is still homogeneous and no obvious aggregations could be observed. The results indicate that the organic PMMA shell plays a crucial role in optimizing the dispersion of BT nanoparticles in the PP matrix. The coating of PMMA shell on the surface of BT nanoparticles not only decreases the surface energy of BT nanoparticles but also improves the compatibility between BT nanoparticles and PP matrix.7 From DSC cooling curves (Fig. S2b), it is interesting to note that the presence of raw BT nanoparticles in the PP matrix provides active sites for heterogeneous nucleation, which increases the crystallization temperature of PP matrix. After coating the BT nanoparticle with a PMMA shell, the physical entanglement of PMMA shell chains with PP matrix chains facilitates the movement and alignment of PP chains into the active sites of heterogeneous nucleation (BT nanoparticles) during the crystallization process, accordingly moving the crystallization temperature to higher temperature. The entanglement of PMMA chains with PP chains is beneficial to reduce the defect density of the interface between BT nanoparticles and PP matrix, accordingly improving the Eb of PP nanocomposites.

14

Fig. 2 (a) The dielectric properties, (b) Weibull distributions for the electrical breakdown strengths and (c) D-E loops at 250 MV/m of PP nanocomposites with 10 wt% loading of raw BT nanoparticles and PMMA@BT nanoparticles with different thicknesses of PMMA shells. (d) The dependences of charged and discharged energy densities at 250 MV/m of PP nanocomposites with 10 wt% loading of fillers on the thickness of PMMA shell. The inset in (d) shows the dependence of charge-discharge efficiency at 250 MV/m of PP nanocomposites with 10 wt% loading of fillers on the thickness of PMMA shell. In order to figure out the influence of the thickness of PMMA shell on the dielectric and energy storage performances and obtain the optimum shell thickness, the films of PMMA@BT/PP nanocomposites with different thicknesses of PMMA shells at the same loading of fillers (10 wt%) were prepared by the combination of melt blending and hot pressing methods. As shown in Fig. 2a and Fig. S3a, both the ε and loss of pure PP and PP nanocomposites exhibit a stable frequency-dependent behavior in the range of 102-106 Hz. With incorporating 10 wt% loading of raw BT fillers into the PP matrix, the ε at 103 Hz of PP increases from 2.2 to 3.0. Compared 15

with raw BT/PP nanocomposite, all the core-shell structure PMMA@BT nanoparticles filled samples exhibit a significant improvement in ε at the same loading of fillers (10 wt%). The ε increases gradually from 3.0 to 3.8 when the thickness of PMMA shell increases to 10 nm. The organic PMMA shell not only improves the compatibility between BT nanoparticles and PP matrix, thus leading to a homogeneous dispersion of BT (Fig. S1), but also alleviates the huge difference of ε between BT nanoparticles and PP matrix due to the intermediate ε (3.1) of PMMA at 103 Hz (Fig. S3a). Thereby, the ε of PP nanocomposites undergoes a significant increase after the introduction of PMMA shell. For the part of dielectric loss, the remarkable suppression of the dielectric loss of raw BT/PP nanocomposite could be easily observed after grafting PMMA on the surface of raw BT nanoparticles. More importantly, the dielectric losses of the core-shell structured PMMA@BT nanoparticles filled nanocomposites are below a value of 0.005 and such a low value is helpful to increase the η and reduce the production of joule heating. Apart from the dielectric properties, another critical parameter that defines the Um of dielectric film is Eb as Um scales quadratically with Eb. The Eb could be calculated = 1−

using ( −(

a

two-parameter

Weibull

distribution

function,

⁄ ) ), where P is the cumulative probability of electric failure,

Eb is the measured breakdown strength for each specimen, α is the characteristic breakdown strength (characteristic Eb) that corresponds to a ~63.2% probability of failure, and β is the slope parameter that evaluates the scatter of data [36]. Herein, the characteristic Eb is calculated from a linear fitting using Weibull failure statistics across 10 specimens per sample. As displayed in Fig. 2b and Fig. S3b, the Eb of PP decreases rapidly from 361 MV/m to 256 MV/m after incorporating 10 wt% loading of raw BT nanoparticles into the PP matrix. Yet the Eb of PP nanocomposite at the same loading of fillers increases significantly to 385 MV/m when a thin PMMA shell (4 nm) were grafted on the surface of raw BT nanoparticles. More interestingly, the Eb of PP nanocomposite at the same loading of fillers increases from 385 MV/m to 448 MV/m as the thickness of PMMA shell increases from 4 nm to 8 nm. With further increasing the thickness of PMMA shell to 10 nm, the Eb decreases to 434 MV/m. The 16

enhancement of Eb in PP nanocomposites using core-shell structure PMMA@BT nanoparticles can be attributed to the following factors. Firstly, the flexible polymer chains of PMMA shell can adjust its conformation easily and modify the interface between the inorganic BT cores and organic PP matrix, thus contributing to the reduction of defect density [7]. In addition, the high intrinsic Eb of PMMA shell (425 MV/m) favors the formation of a robust layer that can prevent the development of electric trees as shown in Fig. S3b. One more thing is that the PMMA shell with an intermediate ε acts as a buffer layer to minimize the local electric field concentration (Fig. 5), consequently leading to the enhancement of Eb [23]. The energy storage performance of the films of PP nanocomposites with 10 wt% loading of fillers at 250 MV/m was analyzed by the electric displacement-electric field (D-E) loops (Fig. 2c). As shown in Fig. S4, the energy density (U) and discharged energy density (Ue) of dielectric film could be calculated from the D-E loops using the integral and subtraction equations, respectively.6 A higher maximum polarization (Pm) and a lower energy loss (Ul) at the same electric field usually contribute to a higher Ue. The Pm of PP at 250 MV/m increases from 0.58 µC/cm2 (Fig. S5a) to 0.73 µC/cm2 after incorporating 10 wt% loading of raw BT nanoparticles into the PP matrix. For the core-shell structured PMMA@BT nanoparticles filled nanocomposites at the same loading of fillers, the Pm increases gradually from 0.88 µC/cm2 to 1.05 µC/cm2 when the thickness of PMMA shell increases from 4 nm to 10 nm. The charged energy density of PP at 250 MV/m, calculated from the charge curve (Fig. S5a), increases from 0.72 J/cm3 to 0.95 J/cm3 when the PP was filled with 10 wt% loading of raw BT nanoparticles. The dependence of charged energy density of PP nanocomposites at the same loading of fillers (10 wt%), calculated from the charge curves (Fig. 2c), on the thickness of PMMA shell displays a similar tendency that the charged energy density increases gradually from 0.95 J/cm3 to 1.62 J/cm3 as the thickness of PMMA shell increases from 0 nm to 10 nm. However, the discharged energy density, calculated from the discharge curves (Fig. 2c), increases from 0.72 J/cm3 for pure PP to 1.51 J/cm3 for the PP nanocomposite with PMMA@BT-8 nanoparticles, followed by being decreased to 1.45 J/cm3 for the PMMA@BT-10/PP 17

nanocomposite. According to the values of charged and discharged energy density, the

η could be obtained using the equation of

=

⁄

as shown in the inset of Fig. 2d.

The η decreases from 100 % for pure PP to 87.6% for the PP nanocomposite with 10 wt% loading of raw BT nanoparticles. When the raw BT nanoparticles were coated with a thin PMMA shell (4 nm), the η increases sharply to 92.0% and reaches a maximum value of 95.1% as the thickness of PMMA shell increases to 8 nm. However, with further increasing the thickness to 10 nm, the η decreases rapidly to 89.9%. Even though the highest charged energy density (1.62 J/cm3) is achieved for the PP nanocomposite with PMMA@BT-10 nanoparticles, such a low η (below 90%) is not suitable for the development of high-performance dielectric film.2 Therefore, the optimum thickness of PMMA shell is 8 nm and the PMMA@BT-8/PP nanocomposites were systematically investigated in the following parts to obtain an appropriate loading of fillers.

18

Fig. 3 (a) Dielectric constants, (b) dielectric losses, (c) electrical breakdown strengths and (d) D-E loops of the films of pure PP, raw BT/PP and nanocomposites. (e) Comparison of the discharged energy densities of the films of pure PP, raw BT/PP and PMMA@BT/PP nanocomposites at different electric fields. (f) The maximum energy densities and charge-discharge efficiencies of the films of pure PP and PMMA@BT/PP nanocomposites. Considering the factors of discharged energy density and charge-discharge efficiency, the optimum thickness of PMMA shell for PMMA@BT nanoparticles is 8 nm. Hence, the dielectric and energy performances of the films of PMMA@BT-8/PP 19

nanocomposites with different loadings of fillers are needed to be investigated systematically to obtain an optimum loading of fillers. For convenience, the PMMA@BT-8/PP nanocomposites used in the following parts will be denoted as the PMMA@BT/PP nanocomposites. As shown in Fig. 3a, the ε of the films exhibits a stable behavior over a range of 102-106 Hz and the ε increases monotonously with increasing the loading of PMMA@BT nanoparticles. For the dielectric loss part, as displayed in Fig. 3b, it also presents a monotonous increasing tendency. It should be noted that the dielectric loss at 103 Hz increases abruptly from 0.002 to 0.009 as the loading of nanoparticles increases from 10 wt% to 15 wt%. Although the nanocomposite with 15 wt% loading of fillers exhibits a higher ε (3.86 at 103 Hz) than that (3.75 at 103 Hz) of the nanocomposite with 10 wt% loading of fillers, such a high loss (above 0.005) would lead to a lower Eb and a lower η. The Eb increases significantly from 361 MV/m of pure PP to 448 MV/m of the nanocomposite with 10 wt% loading of nanoparticles, followed by being decreased to 326 MV/m of the nanocomposite with 15 wt% loading of nanoparticles (Fig. 3c). The enhancement of Eb with increasing the loading of fillers from 0 wt% to 10 wt% could be partially attributed to the increase of maximum tensile stress of PP nanocomposites (Fig. 4a and b), which indicates that the PMMA shell acts as a robust scaffold hampering the onset of electromechanical failure [37, 38]. According to the D-E loops of the films at different electric fields (Fig. S5), the D-E loops of the films at Eb were selected and plotted in Fig. 3d. Besides, the discharged energy densities at different electric fields were calculated from the D-E loops of the films in Fig. S5 and plotted in Fig. 3e. The discharged energy density can reach up to a maximum value of 3.86 J/cm3 from the films of nanocomposite with 10 wt% loading of PMMA@BT nanoparticles at 367 MV/m. As shown in Fig. 3f, the maximum discharged energy density at Eb increases from 1.40 J/cm3 for pure PP to 3.86 J/cm3 for the nanocomposite with 10 wt% loading of PMMA@BT nanoparticles. With further increasing the loading of PMMA@BT nanoparticles to 15 wt%, the maximum discharged energy density decreases to 2.66 J/cm3 due to its high loss and low Eb. Besides, the η at Eb decreases gradually from 99.5% for pure PP to 89.3% for the nanocomposite with 15 wt% loading of fillers. 20

The η of the nanocomposite with 10 wt% loading of fillers decreases monotonously from 98.6% at 50 MV/m to 94.1% at Eb of 367 MV/m (Fig. S6a). Compared with the raw BT filled nanocomposites, the core-shell structured PMMA@BT filled nanocomposites exhibit a higher Eb and a slimmer loop with a higher Pm, accordingly leading to a higher discharged energy density and a higher η (Fig. S6b-d). Although a slight decrease of η from 99.5% for pure PP to 94.1% for core-shell structured nanoparticles filled nanocomposite is observed, the achieved discharged energy density exhibits a significant increase from 1.40 J/cm3 to 3.86 J/cm3 (Fig. S6d). Therefore, the organic PMMA shell on the surface of BT nanoparticles is of great importance to improve the dielectric and energy performances of polymer nanocomposites.

21

Fig. 4 (a) The stress-strain curves and (b) the maximum tensile stresses and strains of the films of pure PP, raw BT/PP and PMMA@BT/PP nanocomposites. The pictures of the film of PP nanocomposite with 10 wt% loading of PMMA@BT nanoparticles (c) before and (d) after hot-stretching at 165 °C. The pictures of the film of PP nanocomposite with 10 wt% loading of PMMA@BT nanoparticles (e) before and (f) after bending. In order to figure out the influences of organic PMMA shell and loadings of PMMA@BT nanoparticles on the mechanical properties of corresponding nanocomposites, the tensile test was carried out and the typical stress-strain curves were displayed in Fig. 4a. Subsequently, the corresponding maximum tensile stress and maximum tensile strain were presented in Fig. 4b by analyzing those stress-strain curves. Compared with the PP nanocomposite with 10 wt% loading of raw BT nanoparticles, both the maximum stress and strain of PMMA@BT/PP nanocomposite at the same loading of fillers were improved, which suggests that the organic PMMA shell is helpful to optimize the interface between BT nanoparticles and PP matrix. The increased stress of the nanocomposites using core-shell structured nanoparticles plays 22

a positive role in improving its Eb (Fig. 2b) [37]. More importantly, the significant enhancement of strain from 60.52% for raw BT filled nanocomposite to 94.28% for PMMA@BT filled nanocomposite is extremely important in the films winding process at room temperature [2]. For the dependence of mechanical properties on the loadings of fillers, the PMMA@BT nanocomposites exhibit a monotonous increase of maximum stress and a monotonous decrease of strain with increasing the loading of PMMA@BT nanoparticles. In addition to the tensile test performed at room temperature, a simple hot stretching experiment was also carried out in an oven at 165 °C. As shown in Fig. 4d, the film of PMMA@BT nanocomposite with 10 wt% loading of fillers exhibits a high extension ratio of 1:4 without producing any defect and such a high extension ratio is of great importance in the process of the biaxial stretching of PP nanocomposites. The film of PMMA@BT nanocomposite also displays a good light transmission property and excellent flexibility (Fig. 4e and f).

23

Fig. 5 The distribution of the (a, b) electric potential, (c, d) electric field and (e, f) current density of the films of raw BT/PP and PMMA@BT/PP nanocomposites with 10 wt% loading of nanoparticles. Apart from the experimental analysis, the theoretical analysis using the simulation of finite element is also conducted to investigate the influence of organic PMMA shell on the dielectric performance as shown in Fig. 5. The distribution of the electric potential was simulated as displayed in Fig. 5a and the electric potential decreases gradually from the top to the bottom. It is well known that the huge ε contrast between fillers and polymer matrix could result in the electric field concentration [39]. Compared with the raw BT filled nanocomposite, the electric field 24

is more uniformly distributed in the core-shell structured PMMA@BT filled nanocomposite under the same applied electric field (Fig. 5a-b), accordingly leading to a higher ε at the same loading of fillers. More importantly, the PMMA@BT filled nanocomposite displays a relatively low electric field strength surrounding the fillers compared with the raw BT filled one (Fig. 5c-d), which suggests that the PMMA@BT filled nanocomposite could withstand a higher electric field than the raw BT filled one. Additionally, the current density is another important factor that affects the loss and Eb of nanocomposites. The contrast of the local electric current density mainly results from the difference of the local electric resistivity between fillers and polymer matrix [25, 40] The PMMA@BT filled nanocomposite displays a lower local electric current density compared with raw BT filled one (Fig. 5e-f), which indicates that the organic PMMA shell could effectively confine the mobility of free electrons and excessive current percolation [25]. Thus, a higher ε, lower dielectric loss and higher Eb are simultaneously obtained and such a change would lead to a higher discharged energy density and η. The simulation results further demonstrate that the organic PMMA shell plays a critical role in improving the dielectric performance of PP nanocomposites.

25

Fig. 6 The comparison diagram of the increase rate of discharged energy density, charge-discharge efficiency and loading of fillers of this work and the related dielectric nanocomposites reported in previous literatures. As illustrated in Fig. S7, compared with the film of pure PP, the maximum discharged energy density and charge-discharge efficiency of the film of PP nanocomposite with raw BT nanoparticles exhibit a significant decrease of 12.0% and 40.7%, respectively. However, for the film of core-shell structured PMMA@BT nanoparticles filled nanocomposite, the maximum discharged energy density displays a remarkable increase of 175.7% and the charge-discharge efficiency shows a negligible decrease of 5.4%. The results suggest that the organic PMMA shell is of great importance in improving the energy storage performance of PP nanocomposites film. Besides, we also compare the resultant energy storage performance achieved in this work with the energy storage performances of related dielectric nanocomposites reported in previous literatures in Fig. 6 [3, 10, 18, 20-24, 37, 39, 41-44]. The ideal sample is expected to appear in the direction of the arrow indicated in Fig. 6. Such a position, in which the ideal sample is located, indicates that high energy storage density, high charge-discharge efficiency and low loading of fillers are indispensable for high-performance dielectric nanocomposites films. In comparison with those 26

reported systems, the film of PMMA@BT/PP nanocomposite in this work possesses a totally competent energy storage performance.

4. Conclusion In summary, the core-shell structured PMMA@BT nanoparticles with adjustable shell thickness were successfully prepared with the aid of ATRP method and the films of PP nanocomposites filled with PMMA@BT nanoparticles were fabricated by the combination of melt blending and hot pressing methods. The influence of the thickness of PMMA shell on the dielectric and energy performances of PP nanocomposites films was systematically investigated and the optimum thickness of PMMA shell (8 nm) was selected for the subsequent fabrication of PMMA@BT nanoparticles and corresponding nanocomposites films. A maximum discharged energy density of 3.86 J/cm3 was obtained by optimizing the loading of PMMA@BT nanoparticles in the PP matrix. Compared with the film of raw BT/PP nanocomposite at the same loading of fillers, the maximum discharged energy density of the film of PMMA@BT/PP nanocomposite exhibited a significant increase of 365.1% from 0.83 J/cm3 to 3.86 J/cm3 and the charge-discharge efficiency also displayed an increase of 7.4% from 87.6% to 94.1%. In comparison with the film of pure PP, although a slight decrease of charge-discharge efficiency (5.4%) from 99.5% to 94.1% was observed, a significant increase of discharged energy density (175.7%) from 1.40 J/cm3 to 3.86 J/cm3 was achieved. The significant increase of discharged energy density is mainly attributed to the enhancement of dielectric constant and electrical breakdown strength and the high charge-discharge efficiency is attributed to the suppression of dielectric loss. The optimization of the interface between BT nanoparticles and PP matrix and the mitigation of local electric field concentration after grafting organic PMMA shell on the surface of BT contribute to the simultaneous enhancement of dielectric constant and breakdown strength at a small loading of fillers. The results of the simulation of finite element provide theoretical evidence to prove that the organic PMMA shell is beneficial to alleviate the electric field concentration. The principle of 27

materials selection and strategy employed here could have a profound impact on the development of non-polar PP nanocomposites films towards high energy density and high charge-discharge efficiency.

Acknowledgment The authors gratefully acknowledge the support from State Grid Corporation Science Technology Project (5202011600UK), the National Natural Science Foundation of China (Nos. 51425201 and 51622701), and the National Basic Research Program of China (973 Program) (Grant No. 2015CB654603). This project is also supported by State Key Laboratory of New Ceramic and Fine Processing Tsinghua University (No. KF201907).

References [1] Q. Li, L. Chen, M. R. Gadinski, S. Zhang, G. Zhang, H. U. Li, E. Iagodkine, A. Haque, L.-Q. Chen, T. N. Jackson, Q. Wang, Nature 523 (2015) 576-579. [2] D. Q. Tan, Adv. Funct. Mater. (2019), https://doi.org/10.1002/adfm.201808567. [3] Q. Li, G. Zhang, F. Liu, K Han, M. R. Gadinski, C. Xiong, Q Wang, Energy Environ. Sci. 8 (2015) 922-931. [4] S. Luo, Y. Shen, S. Yu, Y. Wan, W.-H. Liao, R. Sun, C.-P. Wong, Energy Environ. Sci. 10 (2017) 137-144. [5] Z.-M. Dang, J.-K. Yuan, J.-W. Zha, T. Zhou, S.-T. Li, G.-H. Hu, Prog. Mater. Sci. 57 (2012) 660-723. [6] Z.-M Dang, J.-K. Yuan, S.-H. Yao, R.-J Liao, Adv. Mater. 25 (2013) 6334-6365. [7] X. Huang, P. Jiang, Adv. Mater. 27 (2015) 546-554. [8] Z. Pan, L. Yao, J Zhai, X. Yao, H. Chen, Adv. Mater. 30 (2018) 1705662. [9] X. Huang, B. Sun, Y. Zhu, S. li, P. Jiang, Prog. Mater. Sci. 100 (2019) 187-225. [10] J. Jiang, Z. Shen, X. Cai, J. Qian, Z. Dan, Y. Lin, B. Liu, C.-W. Nan, L. Chen, Y. Shen, Adv. Energy Mater. 9 (2019 ) 1970047. 28

[11] D. Wang, Z.-M. Dang, in Dielectric Polymer Materials for High-Density Energy Storage. William Andrew 12 (2018) 429-446. [12] T. D. Huan, S. Boggs, G. Teyssedre, C. Laurent, M. Cakmak, S. Kumar, R. Ramprasad, Prog. Mater. Sci. 83 (2016) 236-269. [13] X. Yuan, Y. Matsuyama, T. C. Mike Chung, Macromolecules 43 (2010) 4011-4015. [14] X. Yuan, T. C. Mike Chung, Appl. Phys. Lett. 98 (2011) 062901. [15] G. Zhang, H. Li, M. Antensteiner, T. C. Mike Chung, Macromolecules 48 (2015) 2925-2934. [16] G. Wang, X. Huang, P. Jiang, ACS Appl. Mater. Interfaces 9 (2017) 7547-7555. [17] Z. Pan, M. Wang, J. Chen, B. Shen, J. Liu, J. Zhai, Nanoscale 10 (2018) 16621-16629. [18] X. Zhang, J. Jiang, Z. Shen, Z. Dan, M. Li, Y. Lin, C.-W. Nan, L. Chen, Y. Shen, Adv. Mater. 30 (2018) 1707269. [19] P. Martins, A. C. Lopes, S. Lanceros-Mendezab, Prog. Poly. Sci. 39 (2014) 683-706. [20] M.-S. Zheng, Y.-T. Zheng, J.-W. Zha, Y. Yang, P. Han, Y.-Q. Wen, Z.-M. Dang, Nano Energy 48 (2018) 144-151. [21] K. Bi, M. Bi, Y. Hao, W. Luo, Z. Cai, X. Wang, Y. Huang, Nano Energy 51 (2018) 513-523. [22] Y. Hao, X. Wang, K. Bi, J. Zhang, Y. Huang, L. Wu, P. Zhao, K. X, M. Lei, L. Li, Nano Energy 31 (2017) 49-56. [23] X. Lin, P. Hu, Z. Jia, S. Gao, J. Mater. Chem. A 4 (2016) 2314-2320. [24] X. Zhang, Y. Shen, B. Xu, Q. Zhang, L. Gu, J. Jiang, J. Ma, Y. Lin, C.-W. Nan, Adv. Mater. 28 (2016) 2055-2061. [25] Z. Pan, J. Zhai, B. Shen, J. Mater. Chem. A 5 (2017) 15217-15226. [26] K. Yang, X. Huang, Y. Huang, L. Xie, P. Jiang, Chem. Mater. 25 (2013) 2327-2338. [27] J. Chen, Y. Wang, H. Li, H. Han, X. Liao,R. Sun, X. Huang, M. Xie, Chem. Mater. 30 (2018) 1102-1112. 29

[28] Y. Huang, X. Huang, L. S. Schadler, J. He, P. Jiang, ACS Appl. Mater. Interfaces 8 (2016) 25496-25507. [29] S. Shankar, O. P. Thakur, M. Jayasimhadri, Mater. Chem. Phys., 234 (2019) 110-121. [30] K. Brandt, C. Neusel, S. Behr, G. A. Schneider, J. Mater. Chem. C, 1 (2013) 3129-3137. [31] K. A. Imran, J. Lou, K. N. Shivakumar, J. Appl. Polym. Sci., 135 (2018) 45833. [32] L. Xie, X. Huang, C. Wua, P. Jiang, J. Mater. Chem. 21 (2011) 5897-5906. [33] M. Zhu, X. Huang, K. Yang, X. Zhai, J. Zhang, J He P. Jiang, ACS Appl. Mater. Interfaces 6 (2014) 19644-19654. [34] M. Yang, C. Hu, H. Zhao, P. Haghi-Ashtiania, D. He, Y. Yang, J. Yuan, J. Bai, Carbon 132 (2018) 152-156. [35] M. Yang, H. Zhao, C. Hu, P. Haghi-Ashtiani, D. He, Z.-M. Dang, J. Bai, Phys. Chem. Chem. Phys. 20 (2018) 2777-2786. [36] K. Han, Q. Li, C. Chanthad, M. R. Gadinski, G. Zhang, Q. Wang, Adv. Funct. Mater. 25 (2015) 3505-3513. [37] Q. Li, K. Han, M. R. Gadinski, G. Zhang, Q. Wang, Adv. Mater. 26 (2014) 6244-6249. [38] M. Yang, H. Zhao, D. He, C. Hu, H. Chen, Jin. Bai, Materials, 10 (2017) 741. [39] X. Zhang, Y. Shen, Q. Zhang, L.Gu, Y. Hu, J. Du, Y. Lin, C.-W. Nan, Adv. Mater. 27 (2015) 819-824. [40] J. Chen, X. Wang, X. Yu, L. Yao, Z. Duan, Y. Fan, Y. Jiang, Y. Zhou, Z. Pan, J. Mater. Chem. C 6 (2018) 271-279. [41] D. Kang, G. Wang, Y. Huang, P. Jiang, X. Huang, ACS Appl. Mater. Interfaces 10 (2018) 4077-4085. [42] Q. Chi, T. Ma, Y. Zhang, Y. Cui, C. Zhang, J. Lin, X. Wang, Q. Lei, J. Mater. Chem. A 5 (2017) 16757-16766. [43] G. Zhang, D. Brannum, D. Dong, L. Tang, E. Allahyarov, S. Tang, K. Kodweis, J.-K. Lee, L. Zhu, Chem. Mater. 28 (2016) 4646-4660. [44] H. Tang, H. A. Sodano, Nano Lett. 13 (2013) 1373-1379. 30

Conflict of interest The authors declared that they have no conflicts of interest to this work.