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Cite This: ACS Appl. Energy Mater. XXXX, XXX, XXX−XXX

Multiple Interfacial Modifications in Poly(vinylidene fluoride)/ Barium Titanate Nanocomposites via Double-Shell Architecture for Significantly Enhanced Energy Storage Density Lingyu Zhang,† Yao Wang,*,†,‡ Meiyu Xu,† Wentian Wei,† and Yuan Deng*,†,‡ †

School of Materials Science and Engineering, Beihang University, Beijing 100191, China Beijing Advanced Innovation Center for Biomedical Engineering, Beihang University, Beijing 100083, China

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‡

ABSTRACT: Capacitors with high energy density are pressingly demanded in pulsed power systems and recent achievements in polymer-based nanocomposites with increasingly high energy storage capacity demonstrate their great potential in this field. Poly(vinylidene fluoride) (PVDF)-based composites with barium titanate (BT) nanoparticles as fillers are one of the most studied material systems. Here, we demonstrated in BT/PVDF composite that balance between high breakdown strength, moderately high dielectric constant, and low dielectric loss could be controlled via rational interfacial modification between fillers and polymer matrix. Insulating shell layer constructed from coherent and dense amorphous Al2O3 (AO) and organic polydopamine (PDA) has been encapsulated outside BT nanoparticles to gradually mitigate the large disparity in dielectric constants between BT and PVDF. Good dispersion of nanoparticles in the PVDF matrix is another important merit resulting from interfacial modification. The effects of different shell layer on the crystallinity, microstructures, transmittance of light, and dielectric performances were studied comprehensively. A suppressed dielectric loss of 0.016 with a high discharged energy density of 20.6 J cm−3 has been achieved at 659.1 MV m−1 in BT@AO@PDA/PVDF nanocomposite with 1 vol % loading. Three-dimensional finite element analysis was employed to analyze the effects of shell layers on local electric field distribution. This study shows that a significant increase in energy storage capacity of the nanocomposites incorporating a very small fraction of fillers could be realized via successful interface construction which enables them to be good candidates for pulse power system. KEYWORDS: Polymer-based nanocomposites, core−shell structure, dielectric properties, energy storage, finite element analysis

1. INTRODUCTION Polymer-based dielectric composites with high energy storage density have received intensive attention in the past few years. The superior power density stored via electrostatic capacitors over other energy storage techniques gives them unique applications in military and civilian systems, such as scalar weapons and electric vehicles.1−5 The energy storage density for nonlinear dielectrics could be calculated from the equation U=

∫ E dD = ∫0

Eb

E d(ε0K ·E)

The choice of polymers (over ceramics) in many applications is motivated by the need for “graceful failure” of the dielectric at high electric fields.7 However, the commercial energy storage capacitors employing biaxially oriented polypropylene (BOPP) as dielectric medium possess the maximum energy storage density around 5 J cm−3 at breakdown,8 which is far from the growing demand, due to the low dielectric constant of about 2.2. Therefore, how to increase the dielectric constant of the polymer without reducing its operating voltage is the core factor to develop dielectric with high capacitive energy density.

(1)

where E, D, Eb, ε0, and K represent the electric field, displacement, the breakdown strength, the vacuum permittivity (8.85 × 10−12 F m−1) and dielectric constant, respectively.6 © XXXX American Chemical Society

Received: May 28, 2019 Accepted: July 8, 2019 Published: July 8, 2019 A

DOI: 10.1021/acsaem.9b01052 ACS Appl. Energy Mater. XXXX, XXX, XXX−XXX

Article

ACS Applied Energy Materials Scheme 1. Schematic Illustrationsa

Preparation of (a) BT@AO@PDA nanoparticles and (b) BT/PVDF nanocomposite films with different core-shell structures.

a

2.2. Preparation of Double-Shell Structured BT@AO@PDA Nanoparticles. BT nanoparticles with an average diameter of 100 nm were first coated by an Al2O3 (AO) layer through a heterogeneous nucleation method based on our previous procedure.26 Then polydopamine (PDA) was coated outside the BT@AO particles by an autopolymerization approach. Two grams of BT@AO nanoparticles were dispersed in 150 mL of buffer solution (pH = ∼8.5, adjusted via the mass ratio between TRIS-Base and TRIS-HCl in deionized water) following by severe ultrasonic dispersion to form a uniform suspension liquid. Next, 0.3 g of dopamine hydrochloride was added into the liquid and stirred vigorously for 48 h at room temperature. After that, the liquid was filtered and the residue was carefully washed by deionized water then dried at 60 °C for 12 h. Finally, the particles were ground and collected. 2.3. Preparation of BT@AO@PDA/PVDF Nanocomposites. BT, BT@AO, and BT@AO@PDA nanoparticles were used as fillers to fabricate three series of nanocomposites, respectively. The nanoparticles were first dispersed in 20 mL DMF solvent for 1 h of severe ultrasonication; next, 1.5 g of PVDF powders were introduced. The mixture was then stirred vigorously for 24 h at room temperature. Next, the suspension was cast onto films on a glass substrate by Elcometer 4340 at 80 °C after being degassed in a vacuum chamber at 40 °C for 0.5 h. Then the films were dried at 80 °C for 12 h and thermal treated at 200 °C for 5 min, followed by quenching in cold atmosphere. Thickness of all of the as-prepared films is controlled around 10 μm. The whole fabrication process is shown in Scheme 1. 2.4. Characterization. The X-ray diffraction (XRD) patterns of all samples were collected at room temperature by a D/max-2200/PC Diffractometer using Cu Kα radiation (λ = 0.154056 nm) at a scanning rate of 6° min−1 in the angle range of 10−90°. Morphologies of the cross sections of films were observed by a field emission scanning electron microscope (FE-SEM, Quanta 250 FEG, FEI, Czech). Microstructures of the nanoparticles were observed by transmission electron microscopy (TEM, JEM-2100 JEOL, Japan). Differential scanning calorimetry analysis (DSC) and thermogravimetric analysis (TG) measurements of nanoparticles were performed on a thermal analysis instruments analyzer (STA-449F3) from 30 to 800 °C at a heating rate of 10 °C min−1. Fourier-transform infrared spectrometer (FT-IR, iN10MX) was used to check the functional groups on the surface of nanoparticles and the crystalline phases of PVDF. The ultraviolet−visible spectra of composite films were collected at room temperature by a UV−vis−NIR spectrophotometer (UV-3600, Shimadzu) in the wavelength range between 300 and 800 nm. The dielectric properties of nanocomposites were measured on HP4294A Precision Impedance Analyzer (Agilent) at room temperature in the frequency range of 1 kHz to 10 MHz. The values of breakdown strength were measured by a high voltage polarization device (MPD-20KV, Partulab) at a voltage increasing rate of 100 V s−1 and each sample was measured more than 15 times. Before the displacement-electric field (D-E) loops measurement, circular Cu electrodes with a diameter of 2 mm were sputtered on both sides of

Ferroelectric ceramics have dielectric constant as high as thousands and have been incorporated as fillers into the polymer matrix to form composites for a long time.9−16 For examples, BaTiO3 (BT) with various sizes, morphologies (i.e., nanoparticles, nanorods, nanofibers), loadings, and distributions have been widely studied.9−13 Simple incorporation of BT nanoparticles often requires a high concentration to show an obvious increase in dielectric constant, which results in dramatically reduced flexibility of the polymer, early polarization saturation and electric field concentration, leading to dramtically reduced breakdown strength of these composites.17 Recently, incorporation of one-dimensional BT nanomaterials with large aspect ratio11,12 or ultrafine nanoparticles13 has shown a promising solution to the dilemma of coexistance of high Eb and high ε. In the meantime, constructing core−shell architectures with surfactants,18,19 grafted polymer,20,21 or inorganic layers22,23 have been proven as an effective strategy to obtain high energy density also. Although the insulating shell layer could reduce energy loss due to the suppression of Maxwell−Wagner−Sillars interfacial polarization,24,25 the local field concentration due to the large disparity between the dielectric constants of the BT filler and polymer matrix still remains. One of the key issues is how to design composites with proper interface structure where the permittivities gradually decrease from the filler to the polymer.22 Here, we demonstrated the widely studied nanocomposite system, that is, poly(vinylidene fluoride)/BaTiO3, via elaborate core−shell structure design, that is, coating BT nanoparticles first with a thin layer of highly insulating Al2O3 (εr ∼ 10) to greatly suppress interfacial polarization and polydopamine as the second organic buffer layer to enhance dispersity of nanoparticles and strengthen the bonding between fillers and PVDF, would finally result in an ultrahigh discharged energy density of 20.6 J cm−3 obtained at 659.1 MV m−1 with only 1% volume fraction of filler.

2. EXPERIMENTAL SECTION 2.1. Materials. PVDF was supplied by Shanghai 3F company, China. N,N-dimethylformamide (DMF) was obtained from Tianjin Kangkede Technology Co., Ltd. Barium titanate nanoparticles (∼100 nm in diameter), TRIS-Base (C4H11NO3) and TRIS-HCl were supplied by Shanghai Aladdin Biochemical Technology Co., Ltd. Analytical pure aluminum sulfate octadecahydrate, formic acid, and ammonium formate were supplied from Sinopharm Chemical Reagent Co., Ltd. Dopamine hydrochloride was supplied by Shanghai Macklin Biochemical Co., Ltd. All of the reagents were directly used without further purification. B


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ACS Applied Energy Materials the composite films using magnetron sputtering. The samples were immersed in silicone oil during D-E measurement. The test was carried out at 10 Hz in unipolar mode by a Premier II Ferroelectric Test System (Radiant Technologies, Inc.) at room temperature.

to around 10 nm, moreover, no distinguishable interface between AO and PDA was observed. Thermogravimetric analysis (TG) curves of BT, BT@AO, and BT@AO@PDA nanoparticles were shown in Figure 2a. The enlarged image of curves ranging from 50 to 250 °C shown in the inset was given to present the small loss in mass due small solvent molecular volatilization. The loss of mass remained 0.71% until 800 °C due to the thermostability of Al2O3. Although for the BT@AO@PDA nanoparticles, much rapid weight loss was observed, and with the increase of temperature after 300 °C 12.1% mass was lost due to pyrolysis of PDA. Figure 2b shows FT-IR spectra of nanoparticles. Peak at 651 cm−1 is attributed to the TiO bonding and peaks at 3494 and 1620 cm−1 are the stretching vibration and bending vibration of −OH as well as aluminum hydroxyl groups (Al OH) formed during the reaction process of coating, respectively.27,28 A new peak at 1106 cm−1 was observed after bare BT particles were coated with Al2O3, which corresponds to the AlOAl bending vibration. The introduction of the PDA shell led to new absorption peaks at 1670, 1500, and 1288 cm−1, corresponding to the stretching vibration of NH, CC, and p−OH (phenolic hydroxyl group), respectively,29 confirming the core−shell structure of BT@AO@PDA. 3.2. Microstructure and Crystallization Properties of Nanocomposite Films. The effect of interface modification on the microstructures of nanocomposites could be clearly demonstrated from cross-sectional SEM images. As seen in Figure 3, all of the films showed uniform thickness of ∼10 μm. The pure PVDF film presented in Figure 3a is dense and of high quality; in contrast, for the BT/PVDF composite without any interface modification, serious agglomeration occurred as shown in the inset of Figure 3b, and obvious fractures happened around the interfaces between BT nanoparticle clusters and PVDF. When the BT surface was modified by insulating and amorphous Al2O3, the agglomeration of the nanoparticles was reduced and less fractures were found near the filler/polymer interface. Further, when BT@AO@PDA core−shell was incorporated, homogeneous dispersion of fillers and good interfaces with PVDF has been achieved as shown in the inset of Figure 3d, indicating that successful shell architecture design would bring near-perfect microstructure of the composite. Crystallization properties of the nanocomposites with BT nanoparticles in different core−shell structures were studied XRD, FT-IR, and DSC thermal analysis, as shown in Figure 4. As seen from Figure 4a XRD patterns, diffraction peaks at 18.4° and 20.8° correspond to the (020) crystal plane of αphase and (220) crystal plane of β-phase PVDF, respectively.30,31 FT-IR absorption peaks (see Figure 4b) at 763, 795,

3. RESULTS AND DISCUSSION 3.1. Core−Shell Structure of Nanoparticles. X-ray diffraction patterns of pristine BT, BT@AO, and BT@AO@ PDA nanoparticles were shown in Figure 1a. The sharp

Figure 1. (a) XRD patterns of BT, BT@AO, and BT@AO@PDA nanoparticles with standard PDF #74-1967 for comparison. (b) SEM image of bare BT nanoparticles with the inset showing the microstructure of a single BT particle observed from TEM. (c,d) TEM images of BT@AO and BT@AO@PDA particles, respectively.

diffraction peaks at 22.1°, 31.4°, 39.3°, 45.2°, 50.9°, 56.1°, and 65.8° correspond to the (100), (110), (111), (200), (210), (211), and (220) crystal planes of BT in space group Pm3̅m (PDF #74−1967). No new diffraction peaks other than those of BT appeared, suggesting the shell layers are amorphous. As seen in Figure 1b, bare BT nanoparticles are spherical with diameter uniformly distributed around 100 nm. The morphologies of the BT@AO and BT@AO@PAD core− shell structure are clearly seen in Figure 1c,d. As compared with the inset in Figure 1b showing the morphology of a single BT nanoparticle, Al2O3 (see Figure 1c) formed a coherent and compact shell layer outside BT with a thickness of around 5 nm. No indications of crystallized Al2O3 were observed, further confirming the amorphous phase of the AO layer. The thickness of AO@PDA shell layer shown in Figure 1d increases

Figure 2. (a) TG measurements and (b) FT-IR spectra of BT, BT@Al2O3, and BT@Al2O3@PDA nanoparticles. C


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ACS Applied Energy Materials

Table 1. Summary of Degree of Crystallinity and Melting Temperature of Composites with Different Core-Shell Structured BT Fillers at 4% Loading samples

χc (%)

Tm (°C)

pure PVDF 4%-BT/PVDF 4%-BT@AO/PVDF 4%-BT@AO@PDA/PVDF

44.05 42.92 40.81 26.09

164.0 165.0 164.7 165.1

Pictures of films with BT nanoparticles in different core− shell structures presented in Figure 5a showed obvious color

Figure 3. SEM images of cross sections of (a) neat PVDF film and composite films with (b) BT, (c) BT@AO, and (d) BT@AO@PDA nanoparticles at 4% loading with insets showing the enlarged image indicated by square area of each sample, respectively.

Figure 5. (a) Photographs of each sample. (b) UV−vis spectra of pure PVDF, BT/PVDF, BT@AO/PVDF, and BT@AO@PDA/PVDF with the loading of 4 vol %.

976, 1211, 1382 cm−1 correspond to α-phase, 840, 1070, 1279 cm−1 correspond to β-phase, and 1180 cm−1 corresponds to γphase PVDF.32 It is seen that incorporating BT nanoparticles greatly reduced the content of α-phase, whereas β-phase, and γ-phase were less effected. The degree of crystallinity in PVDF matrix was calculated from DSC curves according to the equation χc (%) =

ΔHf (1 − ϕ)ΔHm100

changing from transparent to dark brown. To further understand this color variation, the ultraviolet−visible spectra (UV−vis) of composite films were collected in Figure 5b. As seen, incorporation of BT fillers would increase light absorption, and the obvious absorption of BT/PVDF and BT@AO/PVDF composites near 3.25 eV corresponds to band gap of BT.34 BT@AO@PDA/PVDF showed much stronger light absorption in the full photon energy range due to the π−π* transition in the conjugated system of polydopamine,35 leading to the dark brown appearance. 3.3. Dielectric and Electrical Characteristics of Nanocomposite Films. Frequency-dependent dielectric constant and dielectric loss of BT/PVDF, BT@AO/PVDF, and BT@ AO@PDA/PVDF nanocomposites with various filler loadings are displayed in Figure 6. Dielectric constants of all the composites increase gradually with the increase of filler content and showed the same frequency dependence behavior as PVDF. Variation of dielectric constants and losses at 1 kHz with the content of the composites with different core−shell structured fillers are compared in Figure 6d. As seen, insulating Al2O3 shell layer slightly lowers the dielectric constant due to the intrinsic low dielectric constant (∼10) and greatly reduces

× 100 (2)

where ΔHf and ΔHm are the melting enthalpy of composite and 100% crystalline PVDF (104.6 J g−1), whereas ϕ represents the mass fraction of fillers, respectively.33 Table 1 summarizes the changes in degree of crystallinity (χc) and melting temperature (Tm) of the films with various core−shell structured BT fillers. A slight decrease in crystallinity was caused by BT and BT@AO nanoparticles, whereas the doubleshell structured particles greatly reduced the crystallinity of the composite to 26.09%. This phenomenon is related to the greatly improved dispersity of BT@AO@PDA nanoparticles, because good dispersion brings high interface volume fraction which is influenced by the PDA layer. The result suggests that PDA inhibits PVDF crystallization. The melting temperature raised slightly for the composites. 100

Figure 4. (a) XRD patterns, (b) FT-IR spectra, and (c) DSC thermal analysis of pure PVDF, BT/PVDF, BT@AO/PVDF, and BT@AO@PDA/ PVDF composites with 4% loading. D


ÅÄÅ ÑÉ ÅÅ ij E yz β ÑÑÑ Å P(E) = 1 − expÅÅÅ−jjj zzz ÑÑÑÑ ÅÅ jk E b z{ ÑÑ ÅÅÇ ÑÑÖ


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(6)

where P is the cumulative probability of electrical failure, and E is the breakdown strength measured of each repetition. Eb and β are the shape parameters where Eb is the characteristic breakdown strength (breakdown strength at the cumulative failure probability of 63.2%) and β represents the slope of least-squares fitted line of the distribution. The cumulative probability is defined as i − 0.44 P= (7) n + 0.25 where i represents the sequence number of measured breakdown strength values ranked in an ascending order and n is the total number of repeats. The Weibull distribution and fitting straight-line of BT/PVDF, BT@AO/PVDF, and BT@ AO@PDA/PVDF films with different filler loadings are shown in Figure 7a−c, respectively. As seen, pure PVDF film has the

Figure 6. Dielectric constant and dielectric loss of (a) BT/PVDF, (b) BT@AO/PVDF, and (c) BT@AO@PDA/PVDF nanocomposites as a function of frequency from 1 kHz to 10 MHz at room temperature. (d) Variation of dielectric constants and losses with filler loadings at 1 kHz.

the dielectric loss due to largely suppressed MWS interfacial polarization. PDA layer further lowers the dielectric constant and loss a little, due to the much mitigated aggregation of nanoparticles. To quantitatively estimate the interfacial polarization among different interfaces, that is, BT-PVDF, BT-AO, BT@AO-PVDF, and BT@AO@PDA−PVDF, calculations were carried out according to our previous deduction36 ε0(γBTεPVDF − γPVDFεBT)

Q BT − PVDF =

γBTdPVDF + γPVDFdBT εPVDF − εBT

=

kf

−1/3

E ̅ (dPVDF + dBT)

γPVDF γBT

−1

kf −1/3 ε0E ̅ (3)

εPVDF − εAO Q BT@AO − PVDF =

(kf −1/3 − 1)dBT +

(

Figure 7. Weibull distribution of breakdown strength of (a) BT/ PVDF, (b) BT@AO/PVDF, and (c) BT@AO@PDA/PVDF nanocomposites. (d) Comparison of characteristic breakdown strengths of composites with various core−shell structures at different filler contents.

γPVDF γAO γPVDF γAO

)

dBTkf −1/3 ε0E ̅

− 1 dAO

(4)

highest breakdown strength (496.0 MV m−1) and Eb gradually decreases with the filler content increasing. However, it is evident that as summarized in Figure 7d, both core−shell structures retard the breakdown with increasing filler content, and moreover the PDA layer contributes more than the AO layer to retain the Eb. Although the insulating AO layer significantly suppresses the interfacial polarization, the resistance to high electric field is still not strong. Therefore, double-shell structure with PDA solves the electric field concentration via enhancing dispersity, which retains the high breakdown strength of PVDF for composites. 3.4. Energy Storage Performance of Nanocomposite Films. Displacement-electric field (D-E) loops of BT@AO@ PDA/PVDF composites with different loadings at maximum breakdown strength were shown in Figure 8a. Composite film with 1% loading showed the lowest polarization but could withstand the highest Eb, which is the competing result between large decrease in crystallinity of PVDF (i.e., the decrease of polar phase) and the polarization brought by

Q BT@AO@PDA − PVDF εPVDF − εPDA = (kf

−1/3

− 1)dBT +

(

γPVDF γPDA

γPVDF γPDA

)

− 1 dPDA − dAO

dBTkf −1/3 ε0E ̅

(5)

where ε, γ, and d represent the dielectric constant, conductivity, and thickness of each component, respectively. ε0 is the vacuum permittivity, k is the distribution factor, f is the volume fraction of filler, and E is the applied electric field. Therefore, take 4% composite films being applied to an external electric field of 100 MV m−1 as examples, the interfacial polarization of each composite is 0.013 C m−2 for QBT‑PVDF, −0.199 C m−2 for QBT@AO‑PVDF, and −0.03 C m−2 for QBT@AO@PDA−PVDF where the minus sign indicates that the polarization direction is opposite to the external electric field. Breakdown strength of the composite is a determining parameter for achieving high energy density as shown in eq 1, and it is analyzed by two-parameter Weibull distribution function E


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Figure 9. (a) Comparison in discharged energy densities of BT/ PVDF with different filler shell structure at 1% volume fraction changing with electric field. (b) Comparison in discharged energy densities among BT/polymer-based composites with various morphologies and loadings13,37−46 and this work.

compared with recent reported values in various BT/polymerbased composites with different morphologies and contents as presented in Figure 9b. It is clear that much higher values have been achieved in this study at quite low BT loading without sacrificing the flexibility of polymer. 3.5. Three-Dimensional Finite Element Simulations of Electrical Properties of Composites with Double-Shell Coated Nanoparticles. To further analyze the dielectric and electrical behaviors of composites with different core−shell structure under external electric field, finite element analyzation (FEA) was carried out by COMSOL Multiphysics 5.3a. Three dimensional models were established for nanocomposites with 4 vol % bare BT, BT@AO, and BT@AO@ PDA nanoparticles, as shown in Figure 10a−c. The distributions of fillers in three models were generated from the stochastic matrix by MATLAB. Dielectric constants of BaTiO3, Al2O3, PDA, and PVDF were assigned as 1000, 9.54, 4, and 8.5,47,48 and the corresponding conductivities are set as 10−6, 10−16, 10−14, and 10−13 S m−1, respectively. The intensity of applied external electric field was set to be 500 MV m−1 and the direction was applied from the top surface to the bottom in all the simulations. The side length of the cube, diameter of the BaTiO3 nanoparticles, thickness of Al2O3, and PDA layers were unified to 1 μm, 100 nm, 5 nm, and 5 nm, respectively, based on the experimental observation. The local electric field distributions in the composites are displayed in Figure 10d−f. It is clearly seen that the distributions of the electric field are significantly changed by the fillers’ shell structure. For the composite with bare BT particles as shown in Figure 10d, field concentration is around the BT nanoparticles along the field direction, which provides the high possibility that these regions would connect as the breakdown path. Because of the depolarization field, the electric field inside the BT nanoparticles is low. When the insulating Al2O3 shell layer is employed as shown in Figure 10e, only the Al2O3 layer suffers a large local field whereas the field concentration near the Al2O3/PVDF interfaces along the applied field direction is greatly alleviated, which is due to the largely reduced interfacial polarization.36−46 It is also interesting to note that electric fields perpendicular to the fillers begin to increase. When the PDA layer was added, as shown in Figure 10f, the PDA layer did not change the local field distribution around the fillers/ PVDF interface but a higher local field tends to distribute perpendicular to the external field direction further which delays the electric breakdown along the external field direction. Cross profiles of simulated distributions of leakage current density in the BT/PVDF, BT@AO/PVDF, and BT@AO@ PDA/PVDF composites are shown in Figure 11a−c. For the

Figure 8. (a) D-E loops of BT@AO@PDA/PVDF nanocomposites with different filler contents. (b) Discharged energy density and (c) charge−discharge efficiency of BT@AO@PDA/PVDF as a function of external electric field. (d) Variation of discharged energy density and efficiency at Eb of composite films with different filler contents.

ferroelectric BT nanoparticles. Discharged energy densities of composites films with increasing electric field until the breakdown strength were calculated according to eq 1 and displayed in Figure 8b. Discharged energy densities showed little variation among samples in the low electric field range (under 400 MV m−1) and 1% composite film reached the highest value of 20.6 J cm−3 at its maximum Eb of 659.1 MV m −1 . The charge−discharge efficiency is defined as η=

Udischarged Udischarged + Uloss

, where Udischarged is obtained by integration

of the area between the D-E loop curve and the corresponding displacement ordinate, and Uloss is the integration area of the D-E loop. Figure 8c shows the efficiency of each film changing with electric field. The 1% composite film shows higher efficiency than pure PVDF before its breakdown and an efficiency higher than 0.63 is retained when the field continues to increase. Especially, the efficiency of the 1% composite film even increased when electric field increased from 500 MV m−1 to Eb. This phenomenon is because the polarization gradually reaches saturation with increasing electric field, which means the remnant polarization (Pr) and the saturated polarization (Ps) do not change whereas the polarization increases slightly with increasing electric field with the small linear part. In this case, as the electric field increased to more than 500 MV m−1, the polarization reached saturation, and Udischarged continued increasing whereas the Uloss barely increased; thus, the efficiency increased slightly with electric field. The summary of discharged energy density and efficiency of these composite films at Eb changing with filler content is shown Figure 8d, and 1% BT@AO@PDA/PVDF shows the largest discharged energy density of 20.6 J cm−3, which is 52.6% higher than 13.5 J cm−3 for the pure PVDF film; the efficiency at Eb is 0.688, not lower than that of PVDF. Further, to study the effect of the core−shell structure on the energy storage capacity, discharged energy density composites with 1% volume fraction of various core−shell structured BT nanoparticles are shown in Figure 9a. As seen, the decisive factor for high energy density is the high Eb, which is achieved in BT@AO@PDA/PVDF composite. The discharged energy density values obtained in this work are F


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Figure 10. (a−c) Three-dimensional finite element models and (d−f) cross-sectional (highlighted in red in (a−c)) profiles of simulated distribution of local electric field of composites with 4% loading of (a,d) BT, (b,e) BT@AO, and (c,f) BT@AO@PDA nanoparticles.

Figure 11. (a−c) Cross-sectional profiles of simulated distribution of leakage current density in the composites with different core−shell structures. (d−f) Three-dimensional images of the major conductive path of composites at 4% loading of (a,d) BT, (b,e) BT@AO, and (c,f) BT@AO@PDA nanoparticles.

dielectric constant (9.5) could be achieved via rational interface construction between BT and PVDF. Double-shell structure constructed via dense and coherent inorganic Al2O3 and organic polydopamine is effective in suppressing the large interfacial polarization between BT and PVDF, and more importantly, leading to good dispersion of BT nanoparticles inside PVDF matrix, which greatly changes the local electric field distribution from along the external electric field direction to perpendicular direction. Finally, a significantly enhanced highly discharged energy density of 20.6 J cm−3 has been achieved in BT@AO@PDA/PVDF nanocomposite at quite low filler loading, that is, 1 vol %, which is superior to the state of the art value of the BT/PVDF system. The study provides a practical solution for the dielectric polymer-based nanocomposite achieving a high energy storage capacity to be applied in the pulse power system.

composite with bare particles, large areas with high current density were found inside the filler and tend to link each other, especially between adjacent fillers. The major leakage regions in the matrix are mainly distributed on the top and bottom sides of the particles, forming a conductive channel along the direction of external electric field, as demonstrated in the 3D major conductive path in Figure 11d, which results in a high dielectric loss and low breakdown strength. The introduction of shell layer not only reduces the current density inside the particles but also prevents formation of conduction paths between the neighbor fillers. Interestingly, the direction of the leakage current route changed from vertical to horizontal and formed a prolonged path perpendicular to the applied field in the composite with the core−shell structure [Figure 11e,f], which prevents direct breakdown of the composites.

■

■

CONCLUSION BT/PVDF composites have been the most extensively studied dielectric polymer-based nanocomposite system. Herein, we demonstrated in this composite that high breakdown strength (659.1 MV m−1), low dielectric loss (0.016), and moderate

AUTHOR INFORMATION

Corresponding Authors

*Email: [email protected] (Y.W.). *E-mail: [email protected] (Y.D.). G


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Yao Wang: 0000-0002-3849-9607 Yuan Deng: 0000-0002-1454-2965 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The work was supported by National Natural Science Foundation of China (Grant 51872009), and the Fundamental Research Funds for the Central Universities.

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