High-Energy-Density Polymer Nanocomposites Composed of Newly

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High-energy-density polymer nanocomposites composed of newly-structured one-dimensional BaTiO3@Al2O3 nanofibers Zhongbin Pan, Lingmin Yao, JiWei Zhai, Dezhou Fu, Bo Shen, and Haitao Wang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b13663 • Publication Date (Web): 09 Jan 2017 Downloaded from http://pubs.acs.org on January 12, 2017

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ACS Applied Materials & Interfaces

High-energy-density polymer nanocomposites composed of newly-structured one-dimensional BaTiO3@ Al2O3 nanofibers

Zhongbin Pana†, Lingmin Yaob†, Jiwei Zhaia∗, Dezhou Fuc, Bo Shena, Haitao Wanga a

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

Shanghai 201804, China. b

Institute of Applied Physics and Materials Engineering, Faculty of Science and

Technology, University of Macau, Macao SAR 999078, China c

Department of Electrical Engineering, Tongji University, Shanghai 201804, China

ABSTRACT Flexible electrostatic capacitors are potentially applicable in modern electrical and electric power systems. In this study, flexible nanocomposites containing newly-structured one-dimensional (1D) BaTiO3@Al2O3 nanofibers (BT@AO NFs) and the ferroelectric polymer poly(vinylidene fluoride) (PVDF) matrix were prepared and systematically studied. The 1D BT@AO NFs, where BaTiO3 nanoparticles (BT NPs) were embedded and homogenously dispersed into the AO nanofibers, were successfully synthesized via an improved electrospinning technique. The additional AO layer, which has moderating dielectric constant, was introduced between BT NPs and PVDF matrices. To improve the compatibility and distributional homogeneity the nanofiller/matrix, the dopamine is coated onto nanofiller. The results show that energy density due to high dielectric polarization is about 10.58 J cm-3 at 420 MV m-1 and the fast charge-discharge time is 0.126 µs. A finite element simulation of electric field and electric current density distribution revealed the novel-structured 1D BT@AO NFs significantly improved the dielectric performance of the nanocomposites. The large extractable energy density and high dielectric breakdown strength suggest the potential applications of the BT@AO-DA NFs/PVDF nanocomposite films in electrostatic capacitors and embedded devices.

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Keyword: Nanocomposites; Capacitors; BaTiO3 nanoparticles; One-dimensional; Energy density; Dielectric properties

1. INTRODUCTION Flexible electrostatic capacitors, which are designed as power sources for modern electrical and electronic devices, have been applied to insulated-gate bipolar transistor snubbers, artificial muscles and particle accelerators.1-12 Classically, the energy density (Us= ∫ EdD) of a dielectric capacitor is decided simultaneously by the E (the applied electric field) and D (the electric displacement).11,

13

The D of dielectric

materials under an applied electric field is determined by the dielectric constant (εr).14 For this reason, most approaches to enhance Us of high-performance dielectric capacitors are categorized by improving the dielectric breakdown strength (Eb) and εr. The 1-3 polymer nanocomposites are a promising strategy to largely improve εr and Us compared with the 0-3 counterpart.15, 16 One-dimensional (1D) large aspect-ratio ferroelectric materials, including nanofibers, nanowires, nanorods, nanobelts and nanotubes, could more significantly enhance the εr at lower content (< 5 vol.%), which is ascribed to the large local electric fields.17-19 Additionally, the large aspect-ratio has the smaller surface region, which is advantageous to decreasing the surface energy and thus improving the fillers dispersion in the polymeric matrices.20, 21

Therefore, introducing the 1D large εr dielectric materials (eg. BaTiO3 nanofibers

(BT NFs), BaTiO3 nanowires (BT NWs), BaxSr1-xTiO3 nanofibers (BST NFs), BaxSr1-xTiO3 nanowires (BST NWs) and PbZrTiO3 nanowires (PZT NWs)), into the polymer matrix could essentially increase εr.22-26 As reported, the large-aspect-ratio and majority of in-plane-orientation of the BaTiO3@TiO2 nanofibers introduced into of the PVDF-based matrix improved εr of 35 with 7 vol.% filler loading and a maximum Us with 3 vol.% filler loading (31.2 J cm-3 at 797.7 MV m-1).27 The nanocomposite filled with PZT NWs (higher aspect ratio) shown smaller loss tangent (tanδ) and larger εr than the loading with PZT nanorods (lower aspect ratio), which increased Us.25 Large-surface-energy nanofillers highly incline to aggregate in the polymeric 2 ACS Paragon Plus Environment

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matrices, thus reducing the Eb and US of the nanocomposites. Therefore, one efficient strategy is to surface-functionalize and thereby stabilize the nanofillers in the polymer matrices, which could improve the compatibility and distributional homogeneity. This is beneficial for minimizing the distortion degree of the local field caused by the large contrast in εr between the high-εr fillers and low-εr polymeric matrices and decreasing the amount of microstructural defects (e.g. voids and flaws) at the filler-matrix interface. For instance, dopamine has strong interfacial bonding strength to the surface by chemical bonds, which is potential to be extremely useful.28, 29 Song et al.30 prepared P(VDF-TrFE)-based nanocomposites using dopamine-modified BT NFs as fillers and indicated that the use of fillers functionalized with suitable dopamine led to good distribution of the fillers, thus improving εr and US, due to its strong interfacial bonding to the fillers surfaces. Based on this strategy, different surface-modifying agents were employed to functionalize the high-εr fillers for polymer nanocomposite. For example, silane/titanate coupling agent, H2O2, phosphonic acid, epoxy and poly(vinyl pyrrolidone) (PVP), could effectively improve the distribution of fillers in the polymeric matrices and thereby moderately modify the εr and Eb of nanocomposites.3, 5, 31-34 Another approach to improve the dispersion of nanofillers is the “graft to” and “graft from” method, which could directly connect the nanofillers to the molecular chain of the polymer matrix by chemical bonding.3, 35, 36 The dense polymer surface layer eliminates the filler accumulation and limits the charge carrier migration on the nanofillers/matrices interface, thus largely increasing the amount of Eb and Us. Maxwell indicate that the introduction of nanofillers (high εr) improves the εr of composite films.37 The high εr of the nanofillers tends to induce high nonuniform electric field in the nanocomposites, which essentially decreases the Eb of their nanocomposites, perhaps because of the large εr contrast in fillers/matrix. One effective approach to alleviate the εr contrast is to incorporate a mediating-εr shell layer, such as zirconia (εr ≈ 25 of bulk), titania (εr ≈ 30~55 of 20 nm nanoparticles), hafnium oxide (εr ≈ 20~35 of bulk), and alumina (εr ≈ 10 of films) onto the surface of nanofillers.

38-41

As reported, the εr, Eb and US could be significantly improved by 3 ACS Paragon Plus Environment


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introducing an intermediate-εr TiO2 shell layer, which effectively promotes uniform dispersion and reduces the local electric-field, thus improving the Us to 12.2 J cm-3 at 340 MV m-1.

42

In addition, the insulating intermediate-εr AO layer can effectively

enhance the energy storage capacity of the resulting nanofillers. 4, 43, 44 In this article, nanocomposites comprising BT@AO NFs and the ferroelectric PVDF matrices were prepared and systematically studied. The novel structure of the 1D BT@AO NFs, where BT NPs were homogeneously dispersed into the AO nanofibers, were successfully synthesized via a facile improved electrospinning method. The additional AO layer, which has moderating εr, was introduced between BT NPs and PVDF matrices. To improve the compatibility and distributional homogeneity of the BT@AO NFs into the polymeric matrices, the dopamine is coated onto BT@AO NFs. The results show that the Us was significantly enhanced up to 10.58 J cm-3 at 420 MV m-1 compared with the PVDF (4.8 J cm-3 at 350 MV m-1).

2. EXPERIMENTAL Materials. All chemicals of analytic grade except the Poly(vinylidene fluoride) (PVDF, Arkema, Kynars 301F) were supplied by Shanghai Aladdin Industrial Inc. Synthesis of BT@AO NFs. The BT@AO NFs were synthesized via a facile improved electrospinning technique. The precursor solution A was prepared by dissolving aluminum isopropoxide (3.8g) in ethanol. Then the BaTiO3 nanoparticles (BT NPs) (1g) were added into the solution A and stirred to form a homogeneous precursor solution B. The 50-nm-diameter BT NPs were acquired from GuoTeng Co. Ltd., Shandong. The viscosity of solution B was adjusted by the addition of PVP. The solution B was then moved to a syringe for electrospinning at 1.7 kV cm-1. The co-electrospinning nanofibers were air-calcinated at 700 °C for 3 h to form BT@AO NFs. To effectively improve the filler-matrix compatibility, we employed dopamine as a surface filler. The BT NPs and BT@AO NFs were dispersed in 0.02 mol L-1 dopamine hydrochloride solutions under stirring at 60 °C for 12 h, followed by drying at 80 °C for 10 h and named as dopamine-coated BT@AO NFs (BT@AO NFs-DA). The dense and robust surface layers of dopamine were observed by HRTEM and 4 ACS Paragon Plus Environment

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further evidenced by XPS, FT-IR and TGA. The composition and morphology of BT@AO NFs were characterized by HRTEM and FESEM. Fabrication of Nanocomposite Films. The nanocomposite films were prepared as follows. The BT@AO NFs-DA and PVDF were proportionally dissolved into DMF by stirring at 70 °C for 24 h, forming a stable suspension C. Then C was deposited through spin-coating onto an ITO substrate to form nanocomposite films. The thickness of the nanocomposite films was controlled by the spin-coating velocity and time (at 3000 rpm min-1 for 30 s).45 The nancomposite films were vacuum-dried at 45 °C for 10 h to volatilize the solvent and eliminate structural defects (eg. voids and probes). To obtain non-polar γ phase and a smooth and flat surface, all the nanocomposite films were subjected to heat at 210 °C for 10 min and then immediate quenching in ice-water. Then it was dried in the vacuum at 70 °C for 10 h to volatilize the residual water. For comparison, the pristine PVDF film was also deposited via spin-coating on an ITO substrate using the same treatment. The final nanocomposite films were ≈ 10 µm thick and their morphology were characterized by FESEM, TGA and FT-IR. Measurement of Dielectric Properties. The aluminum electrodes (60 nm in thickness and 2 mm in diameter) were sputtered on the nanocomposite films for electrical measurements. Their εr and tanδ were measured by an E4980A LCR meter in the frequency range 20 Hz to 2 MHz with 1000 mV at ambient temperature (various). The leakage current density was characterized by a Keithley 2400 source meter (Keithley Instruments, Inc.). The D-E hysteresis of the nanocomposite films were characterized by a Premier II ferroelectric test system at 10 Hz.

3. RESULTS AND DISCUSSION The BT@AO NFs, where BT nanoparticles were embedded and homogeneously dispersed into the AO nanofibers, were successfully synthesized via an improved electrospinning technique (Fig. 1a). The precursor solution for electrospinning contained ~50-nm-diameter BT NPs and aluminum isopropoxide. Then the electrospun fibers were heated at 700 °C to completely remove the organic solvent. As 5 ACS Paragon Plus Environment


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shown in Fig. 1b, the BT@AO NFs are tens of micrometers long featured with large aspect-ratios and diameters from 150 to 300 nm. Remarkably, the XRD peaks of BT@AO NFs are consistent with the BT NPs (Fig. 1c). The line in XRD patterns is characteristic of the cubic perovskite structure without evident secondary phase (within the accuracy of XRD). This result obviously confirms that the AO may exist in the form of amorphous/microcrystalline phase.46, 47 The composition of Ba, Al, Ti and O was confirmed by elemental mapping, as corroborated in Fig. 1d. Clearly, the Ba, Al, Ti and O elements were homogeneously dispersed in the nanofibers during the annealing. The composition and architectural characteristics of the BT@AO NFs complex were further confirmed by HRTEM. As shown in Fig. 2a, the BT NPs are connected to chain and cluster inside the AO nanofibers. The element distribution identified by energy dispersive spectrometry (EDS) is presented in Fig. 2b. Point A is mainly composed by elemental Ba, Ti, Al and O while point B only has elemental Al and O. These results further indicate the successful synthesis of BT@AO NFs. As shown in Fig. 2c, the d spacing estimated from the lattice fringe is 3.99A ± 0.05, which corresponds well to the [100] crystal direction of the BT NPs. However, no lattice fringe at region C is observed. One possible reason is that AO could be formed as a distinct amorphous/microcrystalline phase under the annealing, which is in accordance with XRD. The filler-matrix dispersion and compatibility play a vital role in improving the Us of dielectric capacitors. The dopamine which acts as a surface modifier can effectively improve the BT@AO NFs and PVDF matrices distribution and achieve the compatibility between.29 As shown in Fig. 2d, The dense polymer surface layer with diameter ≈ 7 nm were coated on the surface of BT@AO NFs, which is ascribed to the adsorption of dopamine molecules and evidenced by XPS (Fig. 3). Compared with BT@AO NFs, the additional peak of N 1s in BT@AO-DA NFs appears at approximately 401 eV due to free amino-group, indicating the successful incorporation of dopamine onto the BT@AO NFs surfaces. The TGA and FT-IR further confirm the successful coating of dopamine onto surface of the BT@AO NFs 6 ACS Paragon Plus Environment

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(Fig. S1 and S2). The BT@AO-DA NFs/PVDF nanocomposite films (≈ 10 µm) were obtained via spin-coating and treated by annealing and quenching, which were beneficial for the formation of the nonpolar γ phase

24

(Fig. S3 and S4). The mass fractions of

nano-inclusions were attained from the TGA curves and identified nanofillers contents in the final nanocomposite films (Fig. S5). We found the experimental data were consistent with the theoretical data (see supporting information 5). The top-view and cross-section morphologies of the BT@AO-DA NFs/PVDF nanocomposites were investigated by SEM. Owing to the previous surface functionalization, the BT@AO-DA NFs were relatively good distributed in the polymer matrix, though fewer micro-structure defects (such as micro-pores, micro-cracks and micro-voids) (Fig. 2e and 2f) were available. The 3.6% BT@AO-DA NF film can be easily bended (inset of Fig. 2e), indicating good flexibility. Moreover, the majority of the embedded nanofibers have been distributed on the polymeric matrices with little agglomeration and the BT@AO-DA NFs inclined to the in-plane-oriented direction of the PVDF matrix (Fig. 2f), which are favorable for improvement of Eb and US. The εr of dielectric materials reflects the ability to be poled under high voltage. Figure 4a presents the εr and tanδ of the BT@AO-DA NFs/PVDF nanocomposites from 102 Hz to 106 Hz. It could be seen that the εr of the nanocomposite films rises with the further addition of fillers. This trend is attributed to the higher εr and aspect-ratio of nanofillers than pure PVDF. Remarkably, the 8 vol.% BT@AO-DA NFs nanocomposite films have higher εr of 16.34 at 103 Hz, with an increase of ~ 198% from that of pure PVDF (≈ 8.26). As corroborated in Fig.4a, the tanδ of all the nanocomposites is lower than pure PVDF at high frequency. The fillers are similar to “freezing”, which limits the macromolecular movement and thus results in low dielectric loss tangent.48 It is noted that all the nanocomposites have a current density below 3.83 × 10-6 A cm-2 (Fig.4b), which is preferable to improvement of Eb and Us. One

possible

reason

is

that

the

nanocomposites

have

relatively

low

Maxwell−Wagner−Sillars (MWS) interfacial polarization after the incorporation of moderate-εr AO between the nanofillers and polymeric matrices, within which the 7 ACS Paragon Plus Environment


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electric-field concentration might be alleviated.49 Another reason is that the AO, which functions as an insulated layer, efficiently restricts the charge carrier migration in the filler-matrices space. In addition, the εr () was decreased, and the tanδ as well as current density (Fig. 4c and 4d) increased in the nanocomposite filled with BT@AO NFs under the same condition, as compared with the nanocomposites filled with dopamine modified BT@AO NFs (more details in Fig. S6). These trends may be ascribed to the poor dispersion of the BT@AO NFs without dopamine modification and poor interfacial adhesion resulting in microstructure defect (such as micropores and microcracks) in the nanocomposites. The D-E hysteresis of both the 3.6 vol.% BT-DA NPs and BT@AO-DA NFs nanocomposites at 300 MV m-1 are shown in Fig. 5a. The BT@AO-DA NFs nanocomposites exhibit a narrower D-E curve, lower Dmax and smaller Dr compared with the BT-DA NPs nanocomposites at the same condition, which can be explained as follows. First, the MWS interfacial polarization of the nanocomposites could be reduced by the incorporation of AO into the surface of BT NPs, resulting from the gradient verities of εr. Second, the AO buffer layer increases the insulation of the BT NPs, which effectively restricts the charge carrier migration in the filler/matrices space, thus reducing the space charge polarization of the nanocomposites. Specially, the BT@AO-DA NFs nanocomposites have a larger electric displacement (≈ 7.7 µC cm-2 at 420 MV m-1) and still maintain a relatively lower remnant polarization than that of BT-DA NPs nanocomposites (Fig. 5b), which are beneficial for improving the Eb and Us. We investigated how the incorporation of BT-DA NPs or BT@AO-DA NFs would affect the MWS interfacial polarization relaxation of the PVDF matrix. Figure 6a shows the frequency-dependence of M″ in the 3.6 vol.% BT@AO NF nanocomposites. Clearly, the nanocomposite films show the same MWS(PVDF) interfacial polarization, which originates from the restricted charges at the boundaries between the lamellar crystals and interlamellar amorphous regions.50 Obviously, the peak of fmax moves to higher frequency with temperature rise, which is due to the MWS(PVDF)

interfacial

polarization.50,

51

The

nanocomposites


loaded

with

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BT@AO-DA NFs have a relatively lower relaxation intensity of MWSPVDF polarization compared with the BT-DA NFs nanocomposites, suggesting the introduction of the AO buffer layer confined the aggregation of space charges on the crystal/amorphous boundaries. The 1000/T versus ln(fmax) is presented in Fig. 6b. The activation energy (Ea) could be computed based on Arrhenius equation,52,

53

as

displayed in Fig. 6b. The lower Ea of BT@AO NFs/PVDF nanocomposite films versus the BT NPs/PVDF nanocomposite films (79.06 vs. 85.51 kJ mol-1) means a lower energy is required for space charge flow barriers and aggregates in the crystal/amorphous boundary in the polymer matrix. Equivalently, the addition of the Al2O3 buffer layer drives the space charges to move away rather than gathering on the interface. Therefore, the mobility of charges can be confined by introducing the insulating AO buffer layer, which reduces the space charge polarization. The experimental data suggest that the AO introduction onto the surface of BT NPs as interphase and an insulator restricts the migration of free electrons and excessive current percolation effects, thus reducing the leakage current and dielectric loss. The Eb of the nanocomposites is described using a Weibull distribution function. Based on the Weibull distribution,54 the BDS data of the nanocomposite films with the varying filler contents are given in Fig. 7a. The highest BDS of the 3.6 vol.% BT@AO-DA NFs/PVDF nanocomposites is 420.6 MV m-1, which is larger than pure PVDF (≈ 350.8 MV m-1). One possible explanation is that the Weibull modulus β (β is the slope of the fitted line) is higher than that of pure PVDF (12.06 vs. 9.89). Because β quantifies the scattering of experimental data, a higher β means the experimental data are less scattered.55 The β of the nanocomposite films is higher with a smaller load of BT@AO-DA NFs, meaning the fillers were uniformly distributed into the polymeric matrices and exhibited relatively fewer structure imperfections. The decrease in Eb with the increase of filler content is mainly contributed to the more structure defects (e.g. micro-voids and cracks) of the nanocomposite films. The high Eb nanocomposite films may be polarized to the higher field, which suggests that nanocomposite films with higher electric displacement (D) can be obtained. As shown in Fig. 7b, the 3.6 vol.% BT@AO NFs-DA/PVDF nanocomposites have a larger 9 ACS Paragon Plus Environment


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maximum D than bare PVDF (7.14 vs. 4.3 µC cm-1). The rise in D with the increasing volume fraction of fillers is mainly attributed to the incorporation of high εr fillers, yet the earlier breakdown of nanocomposite films is due to the presence of more defects in these nanocomposites. This result agrees well with β (β = 8.49 at 8 vol.% BT@AO NFs-DA/PVDF nanocomposites). It is worthwhile to note that the BT@AO-DA NFs/PVDF nanocompoaires show higher Eb compared with the BT@AO NFs/PVDF composites films at the same condition (Fig. 8a) (more details in Fig. S7 and Fig. S8). This may be due to the strongly interfacial adhesion by the surface layer between nanofillers and matrices, thus alleviating the mobility of the polymer chains and suppressing the transfer of charge carries, which gave rise to higher Eb. The US of these composite films was determined from the D-E hysteresis, as presented in Fig. 7c (more details in Fig. S8). Clearly, the Us rises with the volume fractions of nanofillers rising from 0 vol.% to 8 vol.% at the same condition. This trend is mainly ascribed to the large εr of the fillers that induces the large electric displacement. It is noteworthy that the 3.6 vol.% BT@AO-DA NFs/PVDF nanocomposites have a higher US of ≈ 10.58 J cm-3 at 420 MV m-1, which is 220% higher than that of pristine PVDF (4.8 J cm-3 at 350 MV m-1) and 882% over the BOPP (≈ 1.2 J cm-3 at 640 MV m-1).55 For comparison, the Us of BT@AO-DA NFs/PVDF nanocomposites are much larger compared with the BT@AO NFs/PVDF nanocomposites at the same condition (Fig. 8b) (more details in Fig. S7 and Fig. S8). This could be attributed to relatively good dispersion and strongly adhesion of BT@AO-DA NFs in the polymer matrix, which would in a certain way modify the Eb, dielectric loss and current density to improve the Us. Nan et al.56 reported the Us of PVDF-based nanocomposites was 20 J cm-3 at 650 MV m-1 with the presence of 3 vol.% BaTiO3@TiO2 NFs, but the Us is about 8 J cm-3at the relatively applied electric field of 420 MV m-1. For further comparison, the Us and Eb of some other PVDF-based nanocomposites are summarized in Table S2. This study possesses the relatively larger Eb and higher Us compared with some previous report, which could be interpreted from three-fold. (I) the BT@AO NFs have a large aspect ratio and high dipole displacement. (II) strong interfacial adhesion by the dopamine surface layer 10 ACS Paragon Plus Environment

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nanofillers-matrices, thus alleviating the mobility of the polymer chains and suppressing the transfer of charge carries. (III) the AO shell layer increases the insulation of the BT NFs and restricts the accumulation and migration of the fillers/matrix space charge, thus reducing the dielectric loss and interface polarization. Particularly, for dielectric applications, another important criterion that evaluates the performance of dielectric capacitors is the energy efficiency (η =

௎೐ ௎೐ ା௎೗೚ೞೞ

). Figure

7d presents the η of the composite films containing different contents of BT@AO-DA NFs at varying electric fields, as computed from the D-E hysteresis. Noticeably, the η at the same electric field decreases with an increase in filler contents, which may be ascribed to the rise of ferroelectric loss and conduction loss. Another reason is that the presence of more structure defects (e.g. micro-voids and cracks) in the nanocomposite films would improve energy loss. Particularly, the 3.6 vol.% BT@AO-DA NFs/PVDF nanocomposites maintain a relatively high η of 63.85% at 420 MV m-1, which is mainly ascribed to the incorporation of the AO buffer layer in between the BT NPs and PVDF matrix that reduced the leakage current and dielectric loss. Another important reason is the quenching process that generated the γ-PVDF phase.24 Taking into the actual application in pulsed power systems, the power density (P) and discharge energy density (W) of the 3.6 vol. % BT@AO NFs/PVDF nanocomposites capacitors was further evaluated by an RLC circuit1, 57, 58 as shown in Fig. S9 and Fig. 9. This capacitor compared with the commercial BOPP capacitor has a faster discharge speed (0.126 µs vs. 2.6 µs) (Fig. 9) and about 650% larger W under the charge at 200 MV m-1 (2.21 J cm-3 vs. 0.34 J cm-3) (Fig. 9(b)).59 The maximum P is about 1.47 MW cm-3 at 0.126 µs, which is seven times higher than the BOPP capacitor at 0.62 µs (0.197 MW cm-3) (Fig. 9(a)).59 The distributions of electric field and electric flux density, which largely affect the US and η of nanocomposites, were simulated in the configuration that the majority of nanofillers were located along the in-plane-oriented direction of the polymer matrices, as evidenced by SEM. The local electric-field strength is traceable to charge agglomeration interface, due to the large contrast of εr matrices/fillers.13, 60 As a result, 11 ACS Paragon Plus Environment


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the local electric-field strength of BT-DA NPs nanocomposites is much higher than BT@AO-DA NFs nanocomposites in between the adjacent fillers as indicated by the intensified pink region (Fig. 10 a1 and Fig. 10a3). As shown in Fig. 10a1, the local electric-field of some adjacent nanoparticles is chained to form a channel along the electric-field direction, 29 as exhibited by the intensified yellow region. Therefore, this channel of local electric-field is potentially origin of breakdown. This may be an electrical breakdown that originates from electronic instability.61, 62 It is interesting to observe the maximum local electric-field strength, as exhibited by the intensified red region (Fig. 10a2). This region has the largest electric-field strength, which will electronically destabilize to cause breakdown.62 Thus, this region is the most likely region of breakdown, which can be further evidenced by electric field vector (Fig. S10a1 and Fig. S10b1). The difference of the electric current density is mainly due to the contrast of local electric resistivity between the fillers and polymeric matrices. Moreover, the BT@AO-DA NFs/PVDF nanocomposites have a much lower local electric current density than the BT-DA NPs nanocomposites in between the adjacent fillers (Fig. 10b1 and Fig. 10b3). One possible reason is that the AO, which functions as an insulated layer, efficiently restricts the charge carrier migration in the space between the fillers and matrices.63, 64 In addition, the local electric current density of some adjacent nanoparticles is chained to form a channel along the electric field direction, as exhibited by the intensified yellow region. Therefore, this channel of electric current density is potentially the starting place of breakdown, which may be a thermal breakdown that originates from thermal instability.61, 62 It is interesting to observe the maximum local electric current density strength, as exhibited by the intensified red region (Fig. 10b2). This region has the largest heating, which will melt insulation results in breakdown.62 Thus, this region is the most likely region of breakdown, which is further evidenced by the electric field vector (Fig. S10a2 and Fig. S10b2). However, the BT@AO-DA NFs/PVDF nanocomposites have larger US and η than the BT-DA NPs/PVDF nanocomposites.

4. CONCLUSIONS 12 ACS Paragon Plus Environment

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Flexible nanocomposites containing the newly-structured 1D BT@AO NFs and the ferroelectric PVDF matrix were prepared and systematically studied. The 1D BT@AO NFs, where BT NPs were embedded and homogeneously dispersed into the AO nanofibers, were successfully synthesized via a facile improved electrospinning method. The additional AO layer, which has moderating εr, was introduced between BT NPs and PVDF matrices. The results show that energy density due to high dielectric polarization is about 10.58 J cm-3 at 420 MV m-1 and the fast charge-discharge time is 0.126 µs. A finite element simulation of electric field and electric current density distribution revealed the significant implications of the newly-structured 1D BT@AO NFs on the dielectric performance of the nanocomposites. This study may develop a new approach to design and enhance the energy density of nanocomposites, which are potentially applicable for electrostatic capacitors and embedded devices.

■ ASSOCIATED CONTENT *Supporting Information FT-IR and TGA curves of the BT@AO NFs and BT@AO NFs-DA and their nanocomposites; Frequency versus the εr and tanδ, electric field versus the current density of the BT@AO NFs/PVDF nanocomposites; D-E hysteresis of the composite films filled with BT@AO NFs and BT@AO-DA NFs; The distribution of electric field vector and electric flux density vectors simulated for the 5.1 vol.% BT-DA NPs/PVDF and BT@AO-DA NFs/PVDF nanocomposites.

■ AUTHOR INFORMATION Corresponding Authors* E-mail: [email protected] Authors Contributions† Zhongbin Pan and Lingmin Yao contributed equally to this work. Notes The authors declare no competing financial interest. 13 ACS Paragon Plus Environment


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■ ACKNOWLEDGEMENTS This work was supported by the Ministry of Science and Technology of China through 973-project under Grant (2015CB654601).

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Fig. 1(a) Schematic illustration of improved electrospinning, (b) SEM images of BT@AO NFs, (c) XRD patterns of BT NPs and BT@AO NFs, (d) elemental mapping of BT@AO NFs.

Fig. 2 (a) TEM images, (b) EDS analysis and (c) HRTEM of BT@AO NFs, (d) TEM images of dopamine-coated BT@AO NFs; (e) top-view, (f) cross-section SEM and photo images (inset of (e)) of the 5.1 vol.% BT@AO NFs/PVDF nanocomposites.



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Fig. 3 XPS spectra of the BT@AO NFs and BT@AO NFs-DA.

Fig. 4 (a) Frequency-dependence of the dielectric constant (εr) and dielectric loss (tanδ), (b) electric field dependence of the current density of the BT@AO NFs-DA/PVDF nanocomposites; (c) εr, tanδ and (d) current density of BT@AO NFs/PVDF and BT@AO NFs-DA/PVDF nanocomposites loading with different volume fractions of the fillers. 22 ACS Paragon Plus Environment

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Fig. 5 (a) D-E curves at 300 MV m-1, (b) maximum and remnant polarization for nanocomposites with 3.6 vol.% BT-DA nfs and BT@AO-DA nfs.

Fig. 6 (a) Frequency-dependence of imaginary electric modulus (M″), and (b) Arrhenius plots ln(fmax) vs. reciprocal of temperature (1/T) for the 3.6 vol.% BT-DA NP nanocomposites and BT@AO-DA NF nanocomposites.



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Fig. 7 (a) Weibull plots, (b) D-E loops, (c) energy density and (d) efficiency of the nanocomposites with different contents of the BT@AO-DA NFs.


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Fig. 8 (a) breakdown strength and (b) energy density of BT@AO NFs/PVDF and BT@AO NFs-DA/PVDF nanocomposites loaded with different volume fractions of the fillers.

Fig. 9 (a) the power density (b) discharge energy density of the 3.6 vol.% BT@AO NFs/PVDF nanocomposites capacitors as a function of time. The load resistor R0 is 200 Ω, and the electrical field is 200 MV m-1.



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Fig. 10 Distributions of electric field and electric flux density simulated for the 5.1 vol.% BT-DA NPs/PVDF nanocomposites (a1 and b1, respectively) and the 5.1 vol.% BT@AO-DA NFs/PVDF nanocomposites (a3 and b3, respectively); (a2) and (b2) local magnifications of (a1) and (b1), respectively.


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