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Ultrafast Discharge and Enhanced Energy Density of Polymer Nanocomposites loaded with 0.5(Ba0.7Ca0.3)TiO3– 0.5Ba(Zr0.2Ti0.8)O3 One-dimensional Nanofibers Zhongbin Pan, Lingmin Yao, JiWei Zhai, Haitao Wang, and Bo Shen ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b01381 • Publication Date (Web): 04 Apr 2017 Downloaded from http://pubs.acs.org on April 7, 2017

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ACS Applied Materials & Interfaces

Ultrafast Discharge and Enhanced Energy Density of Polymer Nanocomposites loaded with 0.5(Ba0.7Ca0.3)TiO3–0.5Ba(Zr0.2Ti0.8)O3 One-dimensional Nanofibers Zhongbin Pana, Lingmin Yaob, Jiwei Zhaia*, Haitao Wanga, Bo Shena a

School of Materials Science & Engineering, Tongji University, 4800 Caoan Road,

Shanghai 201804, China. b

School of Physics and Electronic Engineering, Guangzhou University, Guangzhou,

510006, China.

ABSTRACT One-dimensional (1D) materials as fillers introduced into the polymer matrixes have shown great potential in achieving high energy storage capacity because of their large dipole moments. In this article, 1D lead-free 0.5(Ba0.7Ca0.3)TiO3–0.5Ba(Zr0.2Ti0.8)O3 nanofibers (BCZT NFs) were prepared via electrospinning and their formation mechanism was systematically studied. The polypropylene acyl tetraethylene pentamine (PATP) grafted into the surface of BCZT NFs was embedded in the polymer matrixes, which effectively improved the distribution and compatibility of the fillers via chemical bonding and confined the movement of the charge carriers in the interface filler/matrixes. The energy density at relatively low electric field 380 MV m-1 was increased to 8.23 J cm-3 by small loading fillers, far more than that of the biaxially-oriented polypropylene (BOPP) (≈ 1.2 J cm-3 at 640 MV m-1). Moreover, the nanocomposite loaded with 2.1 vol.% BCZT@PATP NFs exhibits a superior discharge speed ≈ 0.189 µs, which indicates the potential application in practice. The finite element simulation of electric potential and electric current density distribution revealed the PATP grafted into the BCZT NFs surface could significantly improve the dielectric performances. This work could provide a new design strategy for high-performance dielectric polymer nanocomposites capacitors.

KEYWORDS: Nanocomposites, Capacitors, Dielectric properties, Lead-free, Energy density

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■ INTRODUCTION High-performance dielectric capacitors have attracted ever-increasing attention because of fast charging-discharging speed and thus can be applied into the next-generation advanced electric devices and systems [1-5]. The discharged energy density of a capacitor (Uc) is expressed as Uc = ∫EdD, where D is the electric displacement and E is the dielectric breakdown strength [6]. D is described as D = εrE, which means large εr (dielectric constant) tends to increase D under the applied electric field. However, there are still critical challenges that maximize both dielectric breakdown strength (Eb) and dielectric constant (εr) for high-performance dielectric materials. Ferroelectric polymers, which possesses high Eb, low dielectric loss (tanδ), and maneuverability, are excellent candidates for dielectric capacitors. But pure polymers generally are limited by the low εr (about < 10) and low energy density [7]. For instance, the best commercially biaxial-oriented polypropylenes (BOPP) possess rather low Uc because of their low εr, which seriously limits their applications [8,9]. To address these problems, researchers have intensively explored the polymer nanocomposite approach, which combines the high ε of inorganic materials with the high Eb and low dielectric loss of polymers, leading to a high energy density. Compared with pristine polymer, BaTiO3-based ceramics have outstanding dielectric and piezoelectric properties [10-16]. In particular, the lead-free 0.5(Ba0.7Ca0.3)TiO3–0.5Ba(Zr0.2Ti0.8)O3 (BCZT) is widely studied

due to its large dielectric constant (εr ≈ 3200) and high piezoelectric coefficient

(d33 ≈ 620 pC/N) [17,18]. However, the BCZT NFs have been rarely reported in application of high-energy-density nanocomposite materials. The dispersion and interaction of fillers into the polymeric matrixes are vital factors in enhancing the Uc of dielectric capacitors. Experimental practice and computational studies show that one-dimensional (1D) materials (eg. nanowires, nanofibers, nanorods, and nanotubes) have a larger aspect ratio than zero-dimensional (0D) materials (eg. nanoparticles) [19-22]. The large aspect ratio has two advantages. First, the 1-D fillers incorporation into the polymer matrixes are preferred to improve Uc at low concentration, which is ascribed to the enhancement of local fields. Second, 1D materials possess a smaller specific surface, which is beneficial for decreasing the surface energy and


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alleviating the agglomeration of the nanofillers in the polymeric matrixes. Besides the methods mentioned above, surface functionalization with polymeric coating, which affords favorable control over the interfacial performance of nanofillers, could significantly improve the distribution of the filler and interfacial bonding of the fillers/matrices. For example, the “grafting to” or “grafting from” method could directly connect the nanofillers to the host matrixes through chemically bonded. Therefore, the polymeric coating on filler surfaces to form core-shell structure can effectively reduce filler aggregation and limit the charge carriers migration upon the filler/polymer interfacial, thus largely improving the Eb and Uc [7, 23-26]. Huang and Jiang et.al [27] reported that the bio-inspired fluoro-polydopamine modification of BaTiO3 nanowires as fillers was introduced into P(VDF-HFP) matrixes, which has guaranteed both the increase of dielectric constant and the maintenance of Eb, resulting in significantly improved Us. The improvement reason is that the surface layer improves nanofillers from dispersion in the polymeric matrixes and confines the movement of the charge carriers in the interface filler/matrixes. In the present study, a novel design of nanocomposite films was proposed by employing poly(vinylidene fluoride) (PVDF) blends as the polymeric matrixes to form nanocomposites with high-aspect-ratio BCZT@(polypropylene acyl tetraethylene pentamine) nanofibers (BCZT@PATP NFs). The large-aspect-ratio 1D BCZT NFs were prepared via electrospinning and their formation mechanism was systematically studied. The PATP was grafted from the surface of BCZT NFs to improve the dispersion and compatibility of the nanofillers in the polymeric matrix. In addition, the PATP acting as shell layer confines the movement of the charge carriers in the interface filler/matrixes and thereby decreases the dielectric loss and current density. As a result, the energy density at the relatively low electric field of 380 MV m-1 was improved to 8.23 Jcm-3 by small loading fillers. Therefore, this new strategy permits insights into polymer electrostatic capacitors with high energy storage capability. ■ EXPERIMENTAL SECTION Materials. All chemicals of analytic grade were supplied by Shanghai Aladdin Industrial Inc., except PVDF (Arkema, Kynars 301F). They were used without any



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further purification. Preparation of BCZT NFs. The BCZT NFs were prepared via electrospinning. Precursor sols were prepared as follows. First, the Ba(COOH)2 (1.95 g) and Ca(COOH)2 (0.212 g) were dissolved in CH3COOH (10 ml) and stirred at 50 °C for 1 h to form a stable precursor solution A (Sol A). Second, the C16H36O4Ti (2.748 g) and C12H28O4Zr were dissolved in C5H8O2 (2.3 g) and stirred to form a stable precursor solution B (Sol B). Third, the Sol B was added into Sol A and stirred to form a stable precursor solution C (Sol C). The viscosity was adjusted by the addition of PVP. Then the Sol C was carried out in a syringe, followed by electrospun at 1.7 kVcm-1. As-prepared nanofibers were air-calcinated at 900 °C for 3 h, so as to completely remove the organics. The phase formation and morphology of the BCZT NFs were measured via XRD, Raman, FT-IR, DSC, TG, and FESEM. Preparation of Core-Shell Structured PATP@ BCZT NFs. First, the BCZT NFs (2 g) were dispersed in a H2O2 aqueous solution (30 %, 300 ml), ultrasonicated for 10 minutes and then stirred at 102 °C for 2 hours. Then, the modified nanofibers were repeatedly rinsed with purified water and CH3CH2OH, and dried under vacuum at 70 °C for 20 hours and named as h-BCZT NFs. Second, the h-BCZT NFs (2 g) were added into 150 ml of N,N-dimethylformamide (DMF) and sonicated for 30 minutes. Then the mixture doped in 1-ethyl-3-(3-dimethylaminopropyl) carbodiimide hydrochloride (EDC.HCl) (0.5 g), polyacrylic acid (PAA) (1 g) and 4-dimethylaminopyridine (DMAP) (0.05 g), ultrasonicated for 20 min and then stirred at 80 °C for 24 h. The mixture were washed with DMF for several times, vacuum-dried at 70 °C for 20 h and named as PAA@BCZT NFs. Third, the PAA@BCZT NFs (2 g) were added into 150 ml of DMF and ultrasonicated for 10 minutes, and followed by the mixture was added with EDC.HCl (0.25 g), tetraethylene pentamine (TEPA) (0.5 g) and DMAP (0.05 g), ultrasonicated for 10 minutes, followed by stirred at 70 °C for 24 hours. The mixture were repeatedly rinsed with DMF, and dried under vacuum at 70 °C for 20 hours and named as BCZT@PATP NFs. The surface layer were observed by HRTEM and further confirmed by FT-IR and TG. Fabrication of BCZT@PATP NFs/PVDF nanocomposites. The nanocomposite film was prepared as follows. First, the BCZT@PATP NFs and PVDF (2 g) were


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proportionally dispersed into dimethylformamide (10 ml) for 2 hours of ultrasonication and then stirred for at 70 °C for 5 hours, forming a stable suspension. The suspension was deposited via spin-coating onto an ITO substrate to form nanocomposite films, which were heated at 60 °C for 5 h to volatilize the solvent. The final nanocomposite films were ~10 µm thick and their morphologies were characterized by FESEM. Figure 1 shows the modification of BCZT NFs and the fabrication of the nanocomposites. Characterization. The nanocomposite films were sputtered with 60-nm-thick and 2-mm-diameter aluminum electrodes for electrical measurements. Their dielectric performance were characterized by an E4980A LCR meter from frequency 102 Hz to 106 Hz at 1000 mV and ambient (various) temperature. The D-E loops were characterized by a Premier II ferroelectric test system at 10 Hz. For expression convenience, the acronyms in this present work were indicated by abbreviations as listed in Table S1. ■ RESULTS AND DISCUSSION To study the effect of heating temperatures and confirm the phase formation of the BCZT NFs, we sent all the samples calcined at different temperatures to characterization by XRD (Figure 2(a)). Clearly, the broad amorphous phase appears at 300 °C, because the electrospun fibers are decomposed into amorphous inorganics [28]. Under further heating at 500 °C, the diffraction peaks of BaCO3 (CaCO3), TiO2 (ZrO2), BCZT and C phases appear, meaning that the BCZT crystalline phase has evolved and the decomposition is still incomplete [17]. At the heating temperature ≥ 700 °C, the line in XRD patterns can be indexed to the pseudo-cubic perovskite structure with no other phases present, which reveals the single phase (BCZT) formation. The BCZT NFs calcined at 900°C are tens of micrometers long featured with large aspect-ratio, with diameters from 200 to 350 nm (FESEM image in Figure 2(b)). The composition of Ba, Ca, Zr, Ti and O was confirmed by element mapping, as displayed in Figure 2(c). Noted that the Ba, Ca, Zr, Ti and O elements were uniform distributed in the nanofibers during the annealing. The FT-IR spectra of as-prepared samples with different heat-treatment temperatures are shown in Figure S1. Clearly, they are mixed of organometallics and the remaining



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solvent, whose peaks both gradually vanish with the increase of calcination temperature. To further understand the formation mechanism of BCZT NFs, we carried out TGA and DSC curves. Based on the DSC curve, the total weight loss occurs at three stages in the formation process. As shown in Figure 3, the TG curves have the first weight loss with 10.28% at period Ι, which may be caused by the evaporation of residual water and solvent. Also an exothermic effect appears on the DSC curve. An intensive exothermal peak is found at 95 °C, which can be ascribed to a loss of non-structural water molecules. The samples changed from pale yellow (as spun) to grey. The sharp weight loss is ~ 48.92% at periodⅡ. One possible reason is that the organic functional groups in the raw materials were destroyed, which was accompanied by an exothermal peak on the DSC curve. The samples turned black, indicating the decomposition into unburned organics. At the period Ⅲ, the heated samples turn black. Also we observed an endothermal effect with 19.12% weight loss. These results can be associated with the decomposition of the intermediates BaCO3 (CaCO3), CaO (ZrO) and C. Then the BCZT crystalline starts to evolve and the decomposition is still incomplete. With temperatures rising to period Ⅳ, the sample turned white. Organometallics and the remaining solvent were completely decomposed, suggesting that the pure BCZT NFS have formed. To confirm the phase formation of BCZT NFs, we tested the Raman spectra for different heating temperatures as-prepared samples (Figure 4). Clearly, the region ≤ 500 °

C is dominated by the amorphous phase, which well agrees with the XRD and FT-IR.

With the heating temperature increasing to 700 °C and above, bands appear at 185, 305, 520 and 726 cm-1, which may be caused by A1/TO, B1/E(TO+LO), A1/E(TO) and A1/E(LO) modes, respectively [17, 28, 29]. The band at 520 cm-1 indicates the formation of cubic and tetragonal phases. The bands at 305 and 726 cm-1 confirm the ferroelectric pseudo-cubic phases of the BCT-BZT NFs [29]. Based on the above analysis, we proposed a possible reaction mechanism of formation BCZT NFs as follows. The schematic diagram is shown in Figure 5. PeriodⅠ,Ⅱ, PVP → C 2Zr(C5H7O2)4 + (46 + x)O2 → 2ZrOx + 40CO2↑+ 28H2O


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2C16H36O4Ti + (46 + x)O2 → 2TiOx + 32CO2↑+ 36H2O Ba(OOCCH3)2 [Ca(OOCCH3)2] → Ba(OOC)2 [Ca(OOC)2] + 2CH3↑ Period Ⅲ, C + O2 → CO2↑ 2ZrOx (TiOx) + (2 - x)O2 → 2ZrO2 (TiO2) Ba(OOC)2 [Ca(OOC)2] → BaCO3 [CaCO3] + CO2↑ BaCO3 (CaCO3) + ZrO2 (TiO2) → 0.5Ba(Zr0.2Ti0.8)O3-0.5(Ba0.7Ca0.3)TiO3 Period Ⅳ, BaCO3 (CaCO3) + ZrO2 (TiO2) → 0.5Ba(Zr0.2Ti0.8)O3-0.5(Ba0.7Ca0.3)TiO3 0.5BZT-0.5BCT crystal grows

The fillers uniformly dispersion polymeric matrixes formation perfect interfacial combination filler-matrixes could significantly enhance the Eb and Uc. The morphology of BCZT@PATP NFs was characterized by TEM, as shown in Figure 7(a) and (b). A dense polymer surface layer appears around the BCZT NFs with ~ 22 nm thickness, which further confirms the successful spontaneous graft of PATP upon the surface of BCZT NFs. The crystal structures of the BCZT@PATP NFs and BCZT NFs did not still change from XRD patterns (Figure S2). The TGA curves show that the BCZT@PATP NFs have a larger weight loss at 800 °C (Figure 6(a)), in which ~13.5 wt. % PATP was grafted. The BCZT NFs, h-BCZT NFs, PAA@BCZT NFs and BCZT@PATP NFs were measured by FT-IR, as presented in Figure 6(b). Compared with BCZT NFs, addition peak appears at ~ 3450 cm-1, indicating that –OH was introduced upon the surface of BCZT NFs [16]. The PAA@BCZT NFs show the infrared absorbance signals at 2925 and 2854 cm-1 (−CH2 stretching vibrations), 2960 cm-1 (of C-H stretching vibration), 1565 and 1415 cm-1 (–COO- stretching vibrations) [30, 31]. These differences of absorbance signals indicate that PAA has been introduced successfully onto the surfaces of h-BCZT NFs. In comparison with BCZT@PAA NFs, the new absorbance signals around 1664 cm-1 corresponds to the N-H stretching vibration [32]. The above results suggest that the



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PATP has been successfully grafted onto the surface of the BCZT@PAA NFs. The grafting of PATP to the BCZT NFs surface could be used as a bridge-linking role between BCZT NFs and PVDF matrix. The PVDF nanocomposite films were obtained through spin-coating and treated by annealing and quenching, which is beneficial for formation of the nonploar γ-phase (Figure S3). The morphologic top-view and freeze-fractured cross-sectional of the nanocomposite films were studied by SEM. Compared with the BCZT NFs/PVDF nanocomposites, the BCZT@PATP NFs possessed a homogeneous distribution in the PVDF matrix (Figure 7 (c)-(d)), suggesting their excellent compatibility with the host matrixes (Figure 7(e)-(f)). Majority of the nanofibers are inclined to the in-plane direction in the polymeric matrixes [2,33], which is advantageous to improve the Eb and Uc. The εr and tanδ of nanocomposites loading with different contents of BCZT@PATP NFs and BCTZ NFs fillers are shown in Figure 8(a) (details in Figs. S4 and S5). Clearly, the dielectric constant increases gradually with increased of fillers, due to the higher dielectric constant of BCZT NFs than that of pristine polymer [33-35]. At the same loading, the BCZT@PATP NFs/PVDF nanocomposite films exhibit inferior elevation of dielectric constant compared with the BCTZ NFs/PVDF nanocomposite films, which is ascribed to the good dispersion and interaction of the fillers into the polymer matrix. Specially, the dielectric constant grows up to 25.5 at 103 Hz when the volume fraction of BCZT@PATP NFs is 9.2 vol.%, which is ≈ 320% that of pristine PVDF (8.26). In addition, the BCZT@PATP NFs/PVDF nanocomposites have lower dielectric loss than the BCTZ NFs/PVDF nanocomposites. Because the PATP shell layer directly links to the surface of the fillers, which further helps to improve the adhesion of the fillers and thereby reduces the structure defects (e.g. pores, cracks). In addition, the AC conductivity is very minor with increasing amount of BCZT@PATP NFs up to 106 Hz (Figure S6). This reason is that the PATP shell layer increases the insulation of the BCZT NF fillers, which effectively restricts the charge carrier migration in the filler-matrixes space. To

further

illustrate

these

properties

of

the

proposed

nanocomposites,

temperature-dependent dielectric spectroscopy was employed. The intuitional variation tendencies of the εr and tanδ with 6.8 vol.% BCZT@PATP NFs/PVDF, BCZT NFs/PVDF nanocomposites and pure PVDF are presented in Figure 8 (b)-(d). With an increase in


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frequency from 1 to 1000 kHz, the εr and tanδ peaks significantly shift to the high temperature, which is a typical thermally-excited relaxation [36]. Compared to the bare PVDF nanocomposites, the BCZT@PATP NFs/PVDF and BCZT NFs/PVDF nanocomposites have a larger dielectric constant along with the temperature rising, due to the higher dielectric constant, dipole polarization, and interfacial polarization of BCZT NFs than that of pristine polymer. The BCZT@PATP NFs/PVDF nanaocomposites have larger dielectric constant than the BCZT NFs/PVDF nanocomposites, since the BCZT@PATP NFs are well-dispersed and adhere into the PVDF matrixes. BCZT@PATP NFs/PVDF nanocomposites have a lower dielectric loss than BCZT NFs/PVDF nanocomposites. This is because the PATP shell layer increases the insulation of the BCZT NF fillers, which effectively restricts the charge carrier migration in the filler-matrixes space. Specially, the dielectric loss peak of the 6.8 vol.% BCZT@PATP NFs/PVDF nanocomposites significantly shifts to the low temperature compared with the bare PVDF at the same frequency. This shift is a hint of the Maxwell-Wagner-Sillars (MWS) polarization and an increase of trap density in the nanocomposites [37]. In practical application, Eb is also tremendous important in dielectric materials and analyzed using Weibull distribution function [38]. From the Figure 9(a), The slightly increased Eb with 2.5 vol.% BCZT@PATP NFs/PVDF nanocomposites, a further increase of the loading would give rise to the decrease of Eb, because of the incorporation of more structure defects in the host matrixes. Noticeably, the BCZT@PATP NFs/PVDF nanocomposites maintain a higher Eb than the BCZT NFs/PVDF nanocomposites at the same loading (Figure 9(b)), which can be explained as follows. First, the dispersion and compatibility of BCZT@PATP NFs are greatly improved through strong chemical bonds. Second, the PATP shell layer increases the insulation of the BCZT NFs fillers, which effectively restricts the charge carrier migration in fillers-matrixes space. To investigate the dielectric response under a high electric field, we measured the energy storage capability of the BCZT@PATP NFs/PVDF nanocomposites and electric displacement dependence of the electric field (D-E) (Figure S7). The energy density of the composites derived from the D-E loops and equation Uc = ∫EdD is summarized in Figure 10(a). The Uc rises with increased content of BCZT@PATP NFs at the same electric field, which is mainly attributed to the relatively more content of BCZT@PATP



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NF nanocomposites with higher dielectric constant and electric displacement. As shown in Figure 10(a), with the loading of 2.1 vol.% BCZT@PATP NFs at a slightly lower electric field of 380 MVm-1, the energy-storage density is maximized to about 8.23 J cm-3, which is an enhancement of ≈ 171% over the pure PVDF (4.81 J cm-3 at 350 MV m-1) and ≈ 686% over the benchmark BOPP (≈ 1.2 J cm-3) [8, 9]. The reasons are the higher electric displacement, good compatibility and well attachment onto the polymer matrix. Our work exhibits excellent breakdown strength and energy density, which seems to rival or exceed some reported advanced PVDF-based nanocomposites at 380 MVm-1 (Table 1). This could be interpreted from three aspects. (i) The BCZT NFs have a larger aspect ratio and high dipole displacement. (ii) The PATP shell layer directly links to the surface of the fillers, which further helps to improve the adhesion of the fillers and thereby reduces the formation of structure defects (e.g. pores, cracks). (iii) The PATP shell layer increases the insulation of the BCZT NF fillers, which effectively restricts the charge carrier migration in the filler-matrixes space, thus reducing the dielectric loss and interface polarization. For dielectric applications, another important criterion that evaluates the property of dielectric capacitors is the discharge energy efficiency (η = Ue/(Ue + Uloss)) [2]. Figure 10(a) shows the η of the nanocomposites containing different contents of BCZT@PATP NFs at various electric fields. Clearly, at the same electric field, the η decreases gradually with increased of fillers, which may be attributed to the increment in ferroelectric loss and conduction loss. Specially, the 2.1 vol.% BCZT@PATP NFs nanocomposites have a relatively high η ≈ 58% at 380 MVm-1, which is mainly attributed to the low interfacial polarization and electrical conduction loss [48]. Compared with the BCZT NFs/PVDF nanocomposites, the BCZT@PATP NFs/PVDF nanocomposites exhibit higher energy density and efficiency at the same condition (Figure 11(b)), which could be ascribed to the large electric displacement & high breakdown strength (Figure S(8)) and the lower dielectric loss, respectively. Given the practical applications in the pulsed power sources, the power energy density (P) and discharge energy density (W) of the nanocomposites loaded with 2.1 vol.% BCZT@PATP NFs were evaluated by RLC circuit (Figure 11). [1, 49-51] This capacitor has a fast discharge speed of ≈ 0.189 µs (Figure 11(b)). When the BCZT@PATP NFs/PVDF nanocomposites capacitor was charged at 200 MV m-1, the discharge energy density was about 2.38 J cm-3, which agrees well with the result


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computed from the D-E curves. Accordingly, the nanocomposite shows a superior power density of 2.21 MW cm-3. The distributions of electric potential and flux density were simulated to further understand the effect of the PATP shell layer on the Uc of BCZT NFs/PVDF nanocomposites. Majority of nanofillers were located along the in-plane-oriented directions into the polymeric matrixes, as evidenced by SEM. Clearly, the electric potential decreases gradually from top to bottom (Figure 12(a)-(b)). The applied voltage between the top and bottom in the simulation system decreased gradually from 1 to 0 kV. The contrast of current density is mainly due to the difference of local resistivity between the fillers and the matrixes [52-55]. The local current density strength of the BCZT@PATP NFs decreased gradually and uniformly dispersed in the matrixes (Figure 12(c)), as exhibited by the enlargement from blue-purple area to pine-green area through the transitional green area (Figure 12(e)). This is because the PATP shell layer significantly improves the distribution and compatibility of the fillers via chemical bonding in the matrixes. Another reason is that the PATP shell layer acts as an interphase and insulation layer, which increases the insulation of the BCZT NFs fillers and confines the accumulation and mobility of the filler-matrix space charge. More interestingly, the local current density of some adjacent nanoparticles is chained to form a channel along the electric field direction (Figure 12(d), (g)), as exhibited by the enlarged pink area. Therefore, this channel of current density is potentially the origin of breakdown. The local current density is maximized at the neighboring fillers, as proved by the enlarged red area (Figure 12(d), (g)). This area is most likely the region of breakdown due to the largest heating [56, 57]. These results further indicate that the PATP shell layer could effectively improve the breakdown strength and energy density. ■ CONCLUSIONS BCZT NFs were prepared via electrospinning and their formation mechanism was systematically studied. The promising strategy was proposed to prepare core-shell structured PATP grafted to BCZT NFs/PVDF nanocomposites by spin-coating method. The nanocomposites loading with BCZT@PATP NFs could more effectively improve the dielectric performance, Eb and Uc compared with the BCZT NFs. The PATP could



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effectively improve nanofillers from dispersion in the polymeric matrixes and confine the movement of the charge carriers in the interface filler/matrixes. As a result, the energy density under the electric field of 380 MV m-1 was improved to 8.23 J cm-3 by a small load of fillers, which was an enhancement of ≈ 171% over the pure PVDF (4.81 J cm-3 at 350 MV m-1) and ≈ 686% over the benchmark BOPP (≈ 1.2 J cm-3). Moreover, the nanocomposite loaded with 2.1 vol.% BCZT@PATP NFs exhibits a superior discharge speed ≈ 0.189 µs, which indicates the potential application in practice. The finite element simulation of electric potential and current density distributions revealed that the introduction of BCZT@PATP NFs into the polymeric matrix significantly improved the breakdown strength and energy density. This article will open up a new design strategy to significantly enhance the energy density of polymeric nanocomposites at relatively low applied electric field, which makes safer working conditions. ■ ASSOCIATED CONTENT *Supporting Information FT-IR spectra of BCZT NFs with different heat-treated temperatures and nanocomposites before and after heat-treated. XRD patterns of BCZT NFs and BCZT@PATP NFs. Frequency versus the εr, tanδ, AC conductivity, and D-E hysteresis of pristine PVDF and their nanocomposites. D-E loops of 2.1 vol.% BCZT@PATP NFs/PVDF and BCZT NFs/PVDF nanocomposites. ■ AUTHOR INFORMATION Corresponding Authors* E-mail: [email protected] ORID Zhongbin Pan: 0000-00002-7522-5840 Notes The authors declare no competing financial interest. ■ ACKNOWLEDGEMENTS


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This work was supported by the Ministry of Science and Technology of China through 973-project under Grant (2015CB654601). ■ REFERENCES [1] Chu, B. J.; Zhou, X.; Ren, K. L.; Neese, B.; Lin, M. R.; Wang, Q.; Bauer, F.; Zhang, Q. M. A Dielectric Polymer with High Electric Energy Density and Fast Discharge Speed. Science 2006, 313, 334-346. [2] Dang, Z. M.; Yuan, J. K.; Zha, J. W.; Zhou, T.; Li, S. T.; Hu, G. H. Fundamentals, Processes and Applications of High-Permittivity Polymer Matrix Composites. Prog. Mater. Sci. 2012, 57, 660-723. [3] Li, Q.; Chen, L.; Gadinski, M. R.; Zhang, S. H.; Zhang, G. Z.; Li, H. Y.; Haque, A.; Chen, L. Q.; Jackson, T.; Wang, Q. Flexible High-Temperature Dielectric Materials from Polymer Nanocomposites. Nature 2015, 523, 576-579. [4] Zhang, X.; Shen, Y.; Xu, B.; Zhang, Q. H.; Gu, L.; Jiang, J. Y.; Ma, J.; Lin, Y. H.; Nan, C. W. Giant Energy Density and Improved Discharge Efficiency of Solution-Processed Polymer Nanocomposites for Dielectric Energy Storage. Adv. Mater. 2016, 28, 2055-2061. [5] Li, Q.; Zhang, G. Z.; Liu, F. H.; Han, K.; Gadinski, M. R.; Xiong, C. X.; Wang, Q. Solution-Processed Ferroelectric Terpolymer Nanocomposites with High Breakdown Strength and Energy Density Utilizing Boron Nitride Nanosheets. Energy Environ. Sci. 2015, 8, 922-931. [6] Tian, Y.; Jin, L.; Zhang, H.; Xu, Z.; Wei, X.; Politova, E. D.; Stefanovich, S. Yu.; Tarakina, N.; Abrahams, I.; Yan, H. High Energy Density in Silver Niobate Ceramics. J. Mater. Chem. A 2016, 4 , 17279-17287. [7] Huang, X. Y.; Jiang, P. K. Core-shell Structured High-k Polymer Nanocomposites for Energy Storage and Dielectric Applications. Adv. Mater. 2015, 27, 546-554. [8] Zhang, L.; Cheng, Z. Y. Development of Polymer-Based 0-3 Composites with High Dielectric Constant. J. Adv. Dielectr. 2011, 1, 389–406.. [9] Barshaw, E.; White, J.; Chait, M.; Cornette, J.; Bustamante, J.; Folli, F.; Biltchick, D.; Borelli, G.; Picci, G.; Rabuffi, M. High Energy Density (HED) Biaxially-Oriented Poly-Propylene (BOPP) Capacitors for Pulse Power Applications. IEEE T. Magn. 2007,



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Figure 1 Schematic illustration for the fabrication of BCZT@PATP NFs and the nanocomposites processes.

Figure 2 (a) XRD patterns of as-spun BCZT NFs calcined at different temperatures; (b) FESEM image and (c) element mapping images of the BCZT NFs calcined at 900 °C.



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Figure 3 DSC and TGA curves for the as-spun BCZT NFs measured at 10 °C min-1 in air atmosphere.

Figure 4 Raman spectra of as-spun BCZT NFs at different heating temperatures.


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Figure 5 Schematic diagram about the formation of BCZT NFs.

Figure 6 (a) TGA curves and (b) FT-IR spectra of the BCZT NFs, h-BCZT NFs, BCZT@PAA NFs and BCZT@PATP NFs.



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Figure 7 Typical TEM images (a) and enlarged view (b) of BCZT@PATP NFs; SEM images of 4.5 vol.% BCZT NFs/PVDF (c: surface; e: cross-section) and 4.5 vol.% BCZT@PATP NFs/PVDF nanocomposites (d: surface; f: cross-section).


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Figure 8 (a) Frequency dependence of the dielectric constant and dielectric loss of BCZT@PATP NFs/PVDF and BCZT NFs/PVDF nanocomposites with different volume fractions; temperature-dependent dielectric properties for the nanocomposite films with (b) bare PVDF, (c) 6.8 vol.% BCZT@PATP NFs and (d) 6.8 vol.% BCZT NFs.



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Figure 9 (a) Weibull plots of the BCZT@PATP NFs/PVDF nanocomposites and (b) breakdown strength of the BCZT@PATP NFs/PVDF and BCZT NFs/PVDF nanocomposites with different volume fractions.


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Figure 10 Energy density and efficiency of (a) BCZT@PATP NFs/PVDF nanocomposites with different volume fractions and (b) 2.1 vol.% BCZT@PATP NFs/PVDF and BCZT NFs/PVDF nanocomposites.



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Figure 11 Charge-discharge properties of 2.1 vol.% BCZT@PATP NFs/PVDF nanocomposites at 200 MV m-1: (a) energy density and (b) power energy density.


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Figure 12 Distributions of electric potential and electric flux density simulated for the 2.1 vol.% BCZT@PATP NFs/PVDF nanocomposites (a and c, respectively) and BCZT NFs/PVDF nanocomposites (b and d, respectively); (e)-(f) and (g): local magnification of (c) and (d), respectively.



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Table 1 Comparison of breakdown strength (Eb) and (Uc) of PVDF-based nanocomposites with our work and literature. Materials

1

Modifier

Eb

Uc

(MV m-1)

(J cm-3)

Ref.

BOPP

None

650

1.2

[9]

2.5 vol.% BZT NFs

PVP1

380

6.3

[39]

2 vol.% BT NPs

SiO2

340

6.28

[26]

2

4 vol.% BCZT NPs

PDA

170

1.3

[40]

30 vol.% BT NPs

Phthalic acid

300

6.7

[41]

5 vol.% BST NCs

H2O2

250

3.9

[42]

2.5 vol.% ST NFs

PVP

380

6.8

[43]

1 vol.% BT NPs

PDA

340

6.37

[44]

50 vol.% PZT NWs

None

80

1.158

[45]

10 vol.% BT NPs

DN-1013

260

4.31

[46]

2.5 vol.% BT NPs

Al2O3

360

6.19

[47]

5 vol.% BT NPs

Al2O3

312

4.7

[48]

2.1 vol.% BCZT NFs

PATP

380

8.23

Our work

Polyvinylpyrrolidone; 2Polydomapine; 3Titanate coupling agent.


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