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High Energy and Power Density Capacitors from SolutionProcessed Ternary Ferroelectric Polymer Nanocomposites Qi Li, Kuo Han, Matthew Robert Gadinski, Guangzu Zhang, and Qing Wang* Electrical energy storage plays an essential role in advanced electronics and electrical power systems.[1] Among the currently available electrical energy storage devices, dielectric capacitors possess the highest power density because of their fast charge–discharge capability.[2,3] Yet they are limited by energy densities that are at least an order of magnitude lower than those of electrochemical devices such as batteries[4] and electrochemical capacitors (ECs).[5,6] For instance, the energy densities of most commercially available ECs to date are in the range of 18 to 29 J cm−3 (5 to 8 W h L−1),[7] while that of the best commercial capacitor film represented by biaxially oriented polypropylenes (BOPP) is only about 1.2 J cm−3 (0.3 W h L−1).[8] Since capacitors can contribute more than 25% of the volume and weight to power electronics and pulsed power systems, dramatic improvement of the energy density of capacitors would be essential to realize their full potential as an enabling technology. For example, high-energydensity polymer capacitors would help to reduce the volume, weight, and cost of the electric power system in hybrid electric vehicles. The energy density of capacitors is governed by the dielectric material that separates the opposite static charges between two electrodes and is given by: U = ∫EdD, where U is the total stored energy density, E is the applied electric field and D is the electric displacement (see Supporting Information Figure S1). For linear dielectrics, U = ½DE = ½Kε0E2, where K is the dielectric constant and ε0 is the vacuum permittivity. Therefore, U is strongly dependent on both K and E, with E limited by the breakdown strength (Eb). Compared with ceramic and electrolytic capacitors, polymer capacitors are easily processed, can be operated under high voltages, and fail gracefully with an open circuit, making them the preferred high-energy-density capacitors.[2,8,9] However, the state-of-the-art dielectric polymers such as biaxially oriented polypropylenes (BOPP) have low values of K (ca. 2.2), which substantially limits U, to, for example, less than 4 J cm−3, at the material level. To raise K, a variety of ferroelectric oxides such as BaTiO3 and Pb(Zr,Ti) O3 with K ranging from hundreds to thousands have been introduced into the polymer to form nanocomposites.[10–17] Unfortunately, a general characteristic of the dielectric nanocomposites is that the improvements in K are obtained at the Dr. Q. Li, K. Han, M. R. Gadinski, Dr. G. Zhang, Prof. Q. Wang Department of Materials Science and Engineering The Pennsylvania State University University Park Pennsylvania 16802, USA E-mail: [email protected]
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expense of sharp reduction in Eb,[10] which is primarily attributed to the presence of highly heterogeneous field distribution in the composites as a result of a large contrast in K between ceramic filler and polymer matrix phases.[3,18] As U scales quadratically with E, a marked decrease in Eb negates any potential substantial increase in U under high fields. The composite approach has thus been hindered by the challenge of resolving the seeming paradox of enhancing K while maintaining high Eb of polymer matrix. Although elevated Eb has been observed in polyethylenes with SiO2 nanoparticles[19,20] and oriented montmorillonite nanoplates,[22], 22 the composites show reduced K relative to the polymer matrix. In this work, we present and demonstrate a new design of the filler system in dielectric polymer nanocomposites for concurrent enhancements in both Eb and K that in turn give rise to greatly improved energy densities that are comparable to ECs, and outstanding power densities that are more than 8 times that of BOPP. Different from the current dielectric nanocomposites, which are usually binary systems containing a single filler, two fillers with different geometries and functions have been readily incorporated into a ferroelectric polymer matrix to obtain ternary polymer nanocomposites. The benefits of the two fillers—viz., two-dimensional (2-D) hexagonal boron nitride nanosheets (BNNSs) used to improve Eb and barium titanate (BT) nanoparticles used to promote K—lead to unprecedented energy storage capability of polymer nanocomposites, including large discharged energy density, high charge– discharge efficiency, and great power density. We first carried out a systematic investigation on the ferroelectric polymer/BNNS nanocomposites as no study regarding the dielectric properties of BNNS-filled polymer composites has been reported before this work. Ferroelectric poly(vinylidene fluoride-co-chlorotrifluoroethylene), P(VDF-CTFE), was chosen as the matrix because of its highest U among the known polymers, arising from the highly polar C–F bonds and spontaneous orgnization of dipoles in the crystalline domains.[23] Moreover, bulky CTFE stabilizes the trans–gauche conformations to realize a reversible change in conformations between the non-polar and polar phases, which accounts for the small remnant polarization (D at zero field) and large change in D under the applied electric field (e.g., ca. 6.0 µC cm−2 at 350 MV m−1).[9] Hexagonal boron nitride (h-BN), a layered structure similar to graphite, is a wide bandgap (ca. 6 eV) insulator with a dielectric strength of 800 MV m−1.[24,25] The breakdown voltage of h-BN can be further improved by decreasing the thickness of h-BN sheet down to a few layers.[26] In this study, ultra-thin BNNSs are prepared through the chemical exfoliation of bulk h-BN powders,[27] and characterized by transmission electron microscopy (TEM; Figure 1a, b and Supporting Information Figure S2) and atomic force microscopy (AFM; Supporting Information Figure S3).
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COMMUNICATION Figure 1. a) Low-resolution TEM image of BNNSs. b) High-resolution TEM image of BNNSs with exposed edge showing a four-layer feature. Inset, electron-diffraction pattern displaying a hexagonal symmetry. c–e) SEM images of P(VDF-CTFE)/BNNS nanocomposite films with 12 wt% BNNSs. f) Weibull plots for P(VDF-CTFE) and P(VDF-CTFE)/BNNS nanocomposites indicating the failure distribution. g) Weibull breakdown strength of P(VDFCTFE)/BNNS nanocomposites as a function of filler content. h) Temperature-dependent dielectric loss of P(VDF-CTFE)/BNNS nanocomposites. Inset, room-temperature dielectric loss of P(VDF-CTFE)/BNNS nanocomposites at 1 kHz as a function of BNNS content. i) Electrical resistivity of P(VDFCTFE)/BNNS nanocomposites with various filler concentrations.
The so-obtained BNNSs are able to form a stable dispersion in N,N-dimethyl formamide (DMF) because of the polar surface of BNNSs induced by B–N bonds, which also enables a homogeneous distribution of BNNSs in polar polymers[27] such
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as P(VDF-CTFE). The P(VDF-CTFE)/BNNS nanocomposites are fabricated through the solution-casting method and present as uniform and flexible films. A set of scanning electron microscopy (SEM) images of the nanocomposite films shown
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in Figure 1c–e validates the homogeneous distribution of BNNSs in the polymer matrix. The characteristic Eb of the P(VDF-CTFE)/BNNS nanocomposites with different concentrations of BNNSs is analyzed within the framework of the Weibull statistics (Figure 1f, g and Supporting Information Table S1). Clearly, Eb is greatly enhanced upon the incorporation of BNNSs, from 387 MV m−1 for the solution-processed pristine P(VDF-CTFE), which is consistent with the literature,[28] to 649 MV m−1 for the nanocomposite with 12 wt% BNNSs. The dramatic improvement in Eb has been rationalized by dielectric spectroscopy (Figure 1h and Supporting Information Figure S6), electrical resistivity (Figure 1i), and mechanical property measurements (Supporting Information Figure S7). The significantly decreased dielectric loss, especially at high temperatures in the dielectric spectra, and nearly an order-of-magnitude increase in electrical resistivity underline the fact that 2-D BNNSs function as an efficient insulating barrier against the leakage current and the space–charge conduction.[24] Meanwhile, the superior Young’s modulus and higher tensile stresses displayed in the composites signify that BNNSs create a robust scaffold hampering the onset of electromechanical failure, one of the main breakdown mechanisms of ferroelectric polymers.[29] In accordance with the experimental results, calculations based on the electromechanical model[30] also reveal the Eb is maximized at 12 wt% of BNNSs in the nanocomposites (Supporting Information Figure S7). Nonetheless, the K value of the P(VDF-CTFE)/
BNNS nanocomposites decreases monotonically with the increase of BNNS content (Supporting Information Figure S8) because of the lower K (i.e., 3–4) of BNNSs[24] relative to the ferroelectric polymer matrix (ca. 10 at 1 kHz). To achieve a simultaneous enhancement in both Eb and K, BT nanoparticles with an average diameter of 100 nm were included in the composites along with BNNSs to create ternary polymer nanocomposites. Interestingly, the bare BT nanoparticles are found to be well-dispersed throughout the polymer matrix, as evidenced in the SEM images (Figure 2e, f) and the energy-dispersive X-ray spectroscopy (EDS) mapping of B and Ba elements (Figure 2h, i) of the ternary nanocomposites. Comparatively, considerable nanoparticle aggregation has been seen in the binary nanocomposites of P(VDF-CTFE)/BT (Figure 2b, c). It is postulated that 2-D nanostructures of BNNSs with high specific surface area can help with the dispersion of the second filler (e.g., BT) by physically dividing the matrix into numerous sub-blocks and placing impervious barriers to prevent BT from aggregating. The role of BNNSs in prevention of BT aggregation has also been verified in the compositional dependence of mechanical properties of the nanocomposites (Figure 2j and Supporting Information Figure S11). The tensile strength of the ternary nanocomposites is seen to increase consistently with increasing BT content at a fixed concentration of BNNSs (12 wt%), which suggests a reinforcement by BT nanoparticles. In stark contrast, the tensile strength of the P(VDF-CTFE)/BT binary nanocomposites declines when the BT concentration
Figure 2. a–f) Schematic and cross-sectional SEM images of P(VDF-CTFE)/BT composite film with 15 wt% BT, and of P(VDF-CTFE)/BNNS/BT ternary nanocomposite film with 12 wt% BNNSs and 15 wt% of BT. g) A photograph of the ternary nanocomposite film. h) Ba, and, i) B element mapping on the cross-section of the ternary nanocomposite established using EDS. j) Tensile strength of P(VDF-CTFE)/BT composites and P(VDF-CTFE)/BNNS/ BT ternary nanocomposites with 12 wt% BNNSs as a function of weight-percentage of BT.
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is over 10 wt%, which suggests that BT nanoparticles agglomerate at high loading ratios in the absence of BNNSs, since a decreasing trend of tensile strength with filler content implies agglomeration of the filler nanoparticles in a nanocomposite.[31] This result is significant because bare inorganic fillers can be intimately dispersed in the polymer matrix without the need for any surfactants, as is the case for conventional dielectric nanocomposites. Although surface modification of the fillers has proved to be effective in improving dispersibility of fillers within the polymer matrix,[10] the ligand with long hydrocarbon chains typically has the lowest K values when compared to ferroelectric matrix and fillers, and thus is the most vulnerable
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local region under high fields and deteriorates the overall dielectric strength.[18] A library of the ternary P(VDF-CTFE) nanocomposites with systematically varied concentrations of BNNS and BT fillers has been prepared. The dependence of crystalline structure, electrical resistivity, dielectric spectra, and breakdown strength of the nanocomposites on composite composition and filler content has been investigated (Supporting Information Figure S12–17). As expected, high contents of BT along with low proportions of BNNSs give rise to a larger K value, while low percentages of BT with a proper loading of BNNSs favor better voltage resilience. As shown in Figures 3a and b, Eb can reach
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up to 649 MV m−1, whereas K varies from 8.0 to 15.3 measured at 1 kHz with different contents of BT and BNNS fillers in the ternary nanocomposites. The results suggest that these key dielectric characteristics could be readily customized to suit different applications by simply tailoring the composite compositions. The energy-storage capability of the dielectric materials was evaluated from the D–E loops by a modified SawyerTower circuit. Typical D–E loops presented in Figure 3d and Supporting Information Figure S20 indicate clearly a concurrent and marked increase in both D and breakdown field from unfilled P(VDF-CTFE) to the ternary nanocomposites, which are in concert with the dielectric spectroscopic and Eb results. More notably, slimmer D–E hysteresis loops with lower remnant polarization (Pr) are evident in the nanocomposites, denoting their lower energy loss (Ul) and higher discharged energy density (Ue =U – Ul) (see Supporting Information Figure S18 for schematic representation). Figure 3c presents Ue measured at respective Eb as a function of filler contents in the ternary nanocomposites. As Ue values are defined by both K and Eb, an optimal combination of K and Eb across the 3-D mapping spectra shown in Figures 3a and b gives rise to a highest Ue of 21.2 J cm−3 of the ternary polymer nanocomposites at a K of 12 and an Eb of 552 MV m−1. This is obtained at a composition of 12 wt% BNNSs and 15 wt% BT, where a 20% increase in K along with a 43% improvement in Eb relative to the pristine matrix are found. This Ue is much greater than that of the solution-processed pristine P(VDF-CTFE), i.e., 7.1 J cm−3 at an Eb of 380 MV m−1 (see Zhang et al.[28], Figure 3e); to the best of the authors' knowledge, this is the highest value reported for dielectric polymer nanocomposites. Although the ferroelectric copolymers such as P(VDFCTFE) and P(VDF-HFP) manufactured via melt-stretching and extrusion-blowing processes can offer similar levels of energy density, ranging from 17–25 J cm−3,[9,32,33] these stretchingbased techniques promote a partial conversion of the non-polar α-phase in PVDF to the polar β-phase and, accordingly, result in high ferroelectric loss and low charge–discharge efficiency (η =Ue/U) that generates waste heat and reduces the lifetime of capacitors. For instance, the extrusion-blown and melt-stretched P(VDF-CTFE)s and P(VDF-HFP)s bear the charge–discharge efficiencies of ca. 66% and 57% at 300 MV m−1, respectively.[33] On the other hand, a much higher charge-–ischarge efficiency of ca. 82% at comparable electric field is seen in the ternary nanocomposite. Even at high electric fields, e.g., at 500 MV m−1, the charge–discharge efficiency of the ternary nanocomposite (78%) is around 10% superior to that of the extrusion-blown and melt-stretched P(VDF-CTFE)s (70%) or P(VDF-HFP)s (65%).[33] This can be understood from the lower Ul of the ternary nanocomposite that arises from improved electrical resistivity and reduced hysteresis loss. While the conduction loss is decreased by the presence of BNNSs as a result of the increased electric resistivity, the hysteresis loss is related to the crystalline structures of the polymer matrix.[34] In the ternary nanocomposites, the α-phase with an estimated relative fraction of 94.7% remains as the majority crystalline phase of the polymer matrix, as corroborated by the wide-angle X-ray diffraction (WAXD) (Supporting Information Figure S17 and Table S7) and the Fourier-transform infrared (FTIR) spectra (Supporting Information Figure S19). Furthermore, the crystallite sizes are reduced with
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the addition of BNNS and BT fillers (Supporting Information Table S6). The existence of small crystallites coupled with a stable α-phase facilitates dipole orientation and phase transitions to the non-polar state during the discharge process, which reduces Pr and the associated hysteresis loss of the nanocomposites. For example, in comparison to the melt-stretched P(VDF-CTFE) with Pr of 1.1 µC cm−2 and D of 8.4 µC cm−2 at 500 MV m−1, the ternary nanocomposite exhibits a lower Pr of 0.8 µC cm−2 and a higher D of 9.6 µC cm−2 (Supporting Information Figure S21). Therefore, a higher charge–discharge efficiency is found in the ternary nanocomposite. Besides, compared with the multi-step extrusion-blowing and zone-heated uniaxially stretching process, the preparation of the ternary polymer nanocomposite films using a simple solution blending and casting technique is straightforward, versatile, and costeffective. These results indicate that a similar level of Ue values as achieved in the traditionally processed (i.e., stretching-based methods) dielectric films are now attained in solution-processable dielectric materials with much higher charge–discharge efficiencies. In addition to energy density, power density is another crucial criterion in the evaluation of the performance of energy storage devices. To demonstrate the ultra-fast discharge rate of the ternary nanocomposites and thus their high power density, a typical fast-discharge experiment[9,35] (see the Experimental Section in the Supporting Information for details) was carried out on both the nanocomposite and BOPP at identical RC time constant (Figure 4). As expected, a microsecond discharge rate has been obtained in the nanocomposite with 12 wt% BNNSs and 15 wt% BT. It is clear that the ternary polymer nanocomposites with 12 wt% BNNSs and 15 wt% BT display a superior power density of 850 MW L−1 that is more than eight times that of BOPP. This result, along with the ultra-high energy density of 5.9 W h L−1 (21.2 J cm−3), which is comparable to existing ECs,[5,7] suggests that the ternary ferroelectric polymer nanocomposite capacitors are capable of storing as large amount of
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Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
Acknowledgements This work was supported by the US Office of Naval Research under grant number N00014–11–1–0342. Received: May 10, 2014 Revised: June 12, 2014 Published online: July 16, 2014
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electric energy as their electrochemical counterpart devices that can be delivered at a rate even greater than the commercial high power density dielectric capacitor films. The pristine P(VDFCTFE) and P(VDF-CTFE)/BT binary nanocomposite filled with 20 wt% BT exhibit relatively lower power densities of 550 and 630 MW L−1, respectively, compared to the ternary nanocomposite, which are attributable to their inferior energy densities. The excellent performance coupled with the simplicity and scalability of preparation makes the ternary polymer nanocomposites promising candidates for compact and flexible highenergy high-power capacitive energy storage devices. This successful design of synergistic dual-filler nanocomposites could also be applicable to other materials that require integration of antagonistic properties to enhance collective performance.
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High Energy and Power Density Capacitors from SolutionProcessed Ternary Ferroelectric Polymer Nanocomposites Qi Li, Kuo Han, Matthew Robert Gadinski, Guangzu Zhang, and Qing Wang* Electrical energy storage plays an essential role in advanced electronics and electrical power systems.[1] Among the currently available electrical energy storage devices, dielectric capacitors possess the highest power density because of their fast charge–discharge capability.[2,3] Yet they are limited by energy densities that are at least an order of magnitude lower than those of electrochemical devices such as batteries[4] and electrochemical capacitors (ECs).[5,6] For instance, the energy densities of most commercially available ECs to date are in the range of 18 to 29 J cm−3 (5 to 8 W h L−1),[7] while that of the best commercial capacitor film represented by biaxially oriented polypropylenes (BOPP) is only about 1.2 J cm−3 (0.3 W h L−1).[8] Since capacitors can contribute more than 25% of the volume and weight to power electronics and pulsed power systems, dramatic improvement of the energy density of capacitors would be essential to realize their full potential as an enabling technology. For example, high-energydensity polymer capacitors would help to reduce the volume, weight, and cost of the electric power system in hybrid electric vehicles. The energy density of capacitors is governed by the dielectric material that separates the opposite static charges between two electrodes and is given by: U = ∫EdD, where U is the total stored energy density, E is the applied electric field and D is the electric displacement (see Supporting Information Figure S1). For linear dielectrics, U = ½DE = ½Kε0E2, where K is the dielectric constant and ε0 is the vacuum permittivity. Therefore, U is strongly dependent on both K and E, with E limited by the breakdown strength (Eb). Compared with ceramic and electrolytic capacitors, polymer capacitors are easily processed, can be operated under high voltages, and fail gracefully with an open circuit, making them the preferred high-energy-density capacitors.[2,8,9] However, the state-of-the-art dielectric polymers such as biaxially oriented polypropylenes (BOPP) have low values of K (ca. 2.2), which substantially limits U, to, for example, less than 4 J cm−3, at the material level. To raise K, a variety of ferroelectric oxides such as BaTiO3 and Pb(Zr,Ti) O3 with K ranging from hundreds to thousands have been introduced into the polymer to form nanocomposites.[10–17] Unfortunately, a general characteristic of the dielectric nanocomposites is that the improvements in K are obtained at the Dr. Q. Li, K. Han, M. R. Gadinski, Dr. G. Zhang, Prof. Q. Wang Department of Materials Science and Engineering The Pennsylvania State University University Park Pennsylvania 16802, USA E-mail: [email protected]
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expense of sharp reduction in Eb,[10] which is primarily attributed to the presence of highly heterogeneous field distribution in the composites as a result of a large contrast in K between ceramic filler and polymer matrix phases.[3,18] As U scales quadratically with E, a marked decrease in Eb negates any potential substantial increase in U under high fields. The composite approach has thus been hindered by the challenge of resolving the seeming paradox of enhancing K while maintaining high Eb of polymer matrix. Although elevated Eb has been observed in polyethylenes with SiO2 nanoparticles[19,20] and oriented montmorillonite nanoplates,[22], 22 the composites show reduced K relative to the polymer matrix. In this work, we present and demonstrate a new design of the filler system in dielectric polymer nanocomposites for concurrent enhancements in both Eb and K that in turn give rise to greatly improved energy densities that are comparable to ECs, and outstanding power densities that are more than 8 times that of BOPP. Different from the current dielectric nanocomposites, which are usually binary systems containing a single filler, two fillers with different geometries and functions have been readily incorporated into a ferroelectric polymer matrix to obtain ternary polymer nanocomposites. The benefits of the two fillers—viz., two-dimensional (2-D) hexagonal boron nitride nanosheets (BNNSs) used to improve Eb and barium titanate (BT) nanoparticles used to promote K—lead to unprecedented energy storage capability of polymer nanocomposites, including large discharged energy density, high charge– discharge efficiency, and great power density. We first carried out a systematic investigation on the ferroelectric polymer/BNNS nanocomposites as no study regarding the dielectric properties of BNNS-filled polymer composites has been reported before this work. Ferroelectric poly(vinylidene fluoride-co-chlorotrifluoroethylene), P(VDF-CTFE), was chosen as the matrix because of its highest U among the known polymers, arising from the highly polar C–F bonds and spontaneous orgnization of dipoles in the crystalline domains.[23] Moreover, bulky CTFE stabilizes the trans–gauche conformations to realize a reversible change in conformations between the non-polar and polar phases, which accounts for the small remnant polarization (D at zero field) and large change in D under the applied electric field (e.g., ca. 6.0 µC cm−2 at 350 MV m−1).[9] Hexagonal boron nitride (h-BN), a layered structure similar to graphite, is a wide bandgap (ca. 6 eV) insulator with a dielectric strength of 800 MV m−1.[24,25] The breakdown voltage of h-BN can be further improved by decreasing the thickness of h-BN sheet down to a few layers.[26] In this study, ultra-thin BNNSs are prepared through the chemical exfoliation of bulk h-BN powders,[27] and characterized by transmission electron microscopy (TEM; Figure 1a, b and Supporting Information Figure S2) and atomic force microscopy (AFM; Supporting Information Figure S3).
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COMMUNICATION Figure 1. a) Low-resolution TEM image of BNNSs. b) High-resolution TEM image of BNNSs with exposed edge showing a four-layer feature. Inset, electron-diffraction pattern displaying a hexagonal symmetry. c–e) SEM images of P(VDF-CTFE)/BNNS nanocomposite films with 12 wt% BNNSs. f) Weibull plots for P(VDF-CTFE) and P(VDF-CTFE)/BNNS nanocomposites indicating the failure distribution. g) Weibull breakdown strength of P(VDFCTFE)/BNNS nanocomposites as a function of filler content. h) Temperature-dependent dielectric loss of P(VDF-CTFE)/BNNS nanocomposites. Inset, room-temperature dielectric loss of P(VDF-CTFE)/BNNS nanocomposites at 1 kHz as a function of BNNS content. i) Electrical resistivity of P(VDFCTFE)/BNNS nanocomposites with various filler concentrations.
The so-obtained BNNSs are able to form a stable dispersion in N,N-dimethyl formamide (DMF) because of the polar surface of BNNSs induced by B–N bonds, which also enables a homogeneous distribution of BNNSs in polar polymers[27] such
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as P(VDF-CTFE). The P(VDF-CTFE)/BNNS nanocomposites are fabricated through the solution-casting method and present as uniform and flexible films. A set of scanning electron microscopy (SEM) images of the nanocomposite films shown
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in Figure 1c–e validates the homogeneous distribution of BNNSs in the polymer matrix. The characteristic Eb of the P(VDF-CTFE)/BNNS nanocomposites with different concentrations of BNNSs is analyzed within the framework of the Weibull statistics (Figure 1f, g and Supporting Information Table S1). Clearly, Eb is greatly enhanced upon the incorporation of BNNSs, from 387 MV m−1 for the solution-processed pristine P(VDF-CTFE), which is consistent with the literature,[28] to 649 MV m−1 for the nanocomposite with 12 wt% BNNSs. The dramatic improvement in Eb has been rationalized by dielectric spectroscopy (Figure 1h and Supporting Information Figure S6), electrical resistivity (Figure 1i), and mechanical property measurements (Supporting Information Figure S7). The significantly decreased dielectric loss, especially at high temperatures in the dielectric spectra, and nearly an order-of-magnitude increase in electrical resistivity underline the fact that 2-D BNNSs function as an efficient insulating barrier against the leakage current and the space–charge conduction.[24] Meanwhile, the superior Young’s modulus and higher tensile stresses displayed in the composites signify that BNNSs create a robust scaffold hampering the onset of electromechanical failure, one of the main breakdown mechanisms of ferroelectric polymers.[29] In accordance with the experimental results, calculations based on the electromechanical model[30] also reveal the Eb is maximized at 12 wt% of BNNSs in the nanocomposites (Supporting Information Figure S7). Nonetheless, the K value of the P(VDF-CTFE)/
BNNS nanocomposites decreases monotonically with the increase of BNNS content (Supporting Information Figure S8) because of the lower K (i.e., 3–4) of BNNSs[24] relative to the ferroelectric polymer matrix (ca. 10 at 1 kHz). To achieve a simultaneous enhancement in both Eb and K, BT nanoparticles with an average diameter of 100 nm were included in the composites along with BNNSs to create ternary polymer nanocomposites. Interestingly, the bare BT nanoparticles are found to be well-dispersed throughout the polymer matrix, as evidenced in the SEM images (Figure 2e, f) and the energy-dispersive X-ray spectroscopy (EDS) mapping of B and Ba elements (Figure 2h, i) of the ternary nanocomposites. Comparatively, considerable nanoparticle aggregation has been seen in the binary nanocomposites of P(VDF-CTFE)/BT (Figure 2b, c). It is postulated that 2-D nanostructures of BNNSs with high specific surface area can help with the dispersion of the second filler (e.g., BT) by physically dividing the matrix into numerous sub-blocks and placing impervious barriers to prevent BT from aggregating. The role of BNNSs in prevention of BT aggregation has also been verified in the compositional dependence of mechanical properties of the nanocomposites (Figure 2j and Supporting Information Figure S11). The tensile strength of the ternary nanocomposites is seen to increase consistently with increasing BT content at a fixed concentration of BNNSs (12 wt%), which suggests a reinforcement by BT nanoparticles. In stark contrast, the tensile strength of the P(VDF-CTFE)/BT binary nanocomposites declines when the BT concentration
Figure 2. a–f) Schematic and cross-sectional SEM images of P(VDF-CTFE)/BT composite film with 15 wt% BT, and of P(VDF-CTFE)/BNNS/BT ternary nanocomposite film with 12 wt% BNNSs and 15 wt% of BT. g) A photograph of the ternary nanocomposite film. h) Ba, and, i) B element mapping on the cross-section of the ternary nanocomposite established using EDS. j) Tensile strength of P(VDF-CTFE)/BT composites and P(VDF-CTFE)/BNNS/ BT ternary nanocomposites with 12 wt% BNNSs as a function of weight-percentage of BT.
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Figure 3. Composition dependence of: a) K, b) Eb, and, c) Ue of P(VDF-CTFE)/BNNS/BT ternary nanocomposites. The Ue value for each composite sample is obtained at respective Eb plotted in panel b. d) D–E loops measured under unipolar electric fields of 10 Hz for pristine P(VDF-CTFE) and P(VDF-CTFE)/BNNS/BT ternary nanocomposite with 12 wt% BNNSs and 15 wt% BT. e) Ue, and, f) charge–discharge efficiency of solution-processed pristine P(VDF-CTFE) and P(VDF-CTFE)/BNNS/BT ternary nanocomposite with 12 wt% of BNNSs and 15 wt% of BT as a function of electric field.
is over 10 wt%, which suggests that BT nanoparticles agglomerate at high loading ratios in the absence of BNNSs, since a decreasing trend of tensile strength with filler content implies agglomeration of the filler nanoparticles in a nanocomposite.[31] This result is significant because bare inorganic fillers can be intimately dispersed in the polymer matrix without the need for any surfactants, as is the case for conventional dielectric nanocomposites. Although surface modification of the fillers has proved to be effective in improving dispersibility of fillers within the polymer matrix,[10] the ligand with long hydrocarbon chains typically has the lowest K values when compared to ferroelectric matrix and fillers, and thus is the most vulnerable
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local region under high fields and deteriorates the overall dielectric strength.[18] A library of the ternary P(VDF-CTFE) nanocomposites with systematically varied concentrations of BNNS and BT fillers has been prepared. The dependence of crystalline structure, electrical resistivity, dielectric spectra, and breakdown strength of the nanocomposites on composite composition and filler content has been investigated (Supporting Information Figure S12–17). As expected, high contents of BT along with low proportions of BNNSs give rise to a larger K value, while low percentages of BT with a proper loading of BNNSs favor better voltage resilience. As shown in Figures 3a and b, Eb can reach
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up to 649 MV m−1, whereas K varies from 8.0 to 15.3 measured at 1 kHz with different contents of BT and BNNS fillers in the ternary nanocomposites. The results suggest that these key dielectric characteristics could be readily customized to suit different applications by simply tailoring the composite compositions. The energy-storage capability of the dielectric materials was evaluated from the D–E loops by a modified SawyerTower circuit. Typical D–E loops presented in Figure 3d and Supporting Information Figure S20 indicate clearly a concurrent and marked increase in both D and breakdown field from unfilled P(VDF-CTFE) to the ternary nanocomposites, which are in concert with the dielectric spectroscopic and Eb results. More notably, slimmer D–E hysteresis loops with lower remnant polarization (Pr) are evident in the nanocomposites, denoting their lower energy loss (Ul) and higher discharged energy density (Ue =U – Ul) (see Supporting Information Figure S18 for schematic representation). Figure 3c presents Ue measured at respective Eb as a function of filler contents in the ternary nanocomposites. As Ue values are defined by both K and Eb, an optimal combination of K and Eb across the 3-D mapping spectra shown in Figures 3a and b gives rise to a highest Ue of 21.2 J cm−3 of the ternary polymer nanocomposites at a K of 12 and an Eb of 552 MV m−1. This is obtained at a composition of 12 wt% BNNSs and 15 wt% BT, where a 20% increase in K along with a 43% improvement in Eb relative to the pristine matrix are found. This Ue is much greater than that of the solution-processed pristine P(VDF-CTFE), i.e., 7.1 J cm−3 at an Eb of 380 MV m−1 (see Zhang et al.[28], Figure 3e); to the best of the authors' knowledge, this is the highest value reported for dielectric polymer nanocomposites. Although the ferroelectric copolymers such as P(VDFCTFE) and P(VDF-HFP) manufactured via melt-stretching and extrusion-blowing processes can offer similar levels of energy density, ranging from 17–25 J cm−3,[9,32,33] these stretchingbased techniques promote a partial conversion of the non-polar α-phase in PVDF to the polar β-phase and, accordingly, result in high ferroelectric loss and low charge–discharge efficiency (η =Ue/U) that generates waste heat and reduces the lifetime of capacitors. For instance, the extrusion-blown and melt-stretched P(VDF-CTFE)s and P(VDF-HFP)s bear the charge–discharge efficiencies of ca. 66% and 57% at 300 MV m−1, respectively.[33] On the other hand, a much higher charge-–ischarge efficiency of ca. 82% at comparable electric field is seen in the ternary nanocomposite. Even at high electric fields, e.g., at 500 MV m−1, the charge–discharge efficiency of the ternary nanocomposite (78%) is around 10% superior to that of the extrusion-blown and melt-stretched P(VDF-CTFE)s (70%) or P(VDF-HFP)s (65%).[33] This can be understood from the lower Ul of the ternary nanocomposite that arises from improved electrical resistivity and reduced hysteresis loss. While the conduction loss is decreased by the presence of BNNSs as a result of the increased electric resistivity, the hysteresis loss is related to the crystalline structures of the polymer matrix.[34] In the ternary nanocomposites, the α-phase with an estimated relative fraction of 94.7% remains as the majority crystalline phase of the polymer matrix, as corroborated by the wide-angle X-ray diffraction (WAXD) (Supporting Information Figure S17 and Table S7) and the Fourier-transform infrared (FTIR) spectra (Supporting Information Figure S19). Furthermore, the crystallite sizes are reduced with
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the addition of BNNS and BT fillers (Supporting Information Table S6). The existence of small crystallites coupled with a stable α-phase facilitates dipole orientation and phase transitions to the non-polar state during the discharge process, which reduces Pr and the associated hysteresis loss of the nanocomposites. For example, in comparison to the melt-stretched P(VDF-CTFE) with Pr of 1.1 µC cm−2 and D of 8.4 µC cm−2 at 500 MV m−1, the ternary nanocomposite exhibits a lower Pr of 0.8 µC cm−2 and a higher D of 9.6 µC cm−2 (Supporting Information Figure S21). Therefore, a higher charge–discharge efficiency is found in the ternary nanocomposite. Besides, compared with the multi-step extrusion-blowing and zone-heated uniaxially stretching process, the preparation of the ternary polymer nanocomposite films using a simple solution blending and casting technique is straightforward, versatile, and costeffective. These results indicate that a similar level of Ue values as achieved in the traditionally processed (i.e., stretching-based methods) dielectric films are now attained in solution-processable dielectric materials with much higher charge–discharge efficiencies. In addition to energy density, power density is another crucial criterion in the evaluation of the performance of energy storage devices. To demonstrate the ultra-fast discharge rate of the ternary nanocomposites and thus their high power density, a typical fast-discharge experiment[9,35] (see the Experimental Section in the Supporting Information for details) was carried out on both the nanocomposite and BOPP at identical RC time constant (Figure 4). As expected, a microsecond discharge rate has been obtained in the nanocomposite with 12 wt% BNNSs and 15 wt% BT. It is clear that the ternary polymer nanocomposites with 12 wt% BNNSs and 15 wt% BT display a superior power density of 850 MW L−1 that is more than eight times that of BOPP. This result, along with the ultra-high energy density of 5.9 W h L−1 (21.2 J cm−3), which is comparable to existing ECs,[5,7] suggests that the ternary ferroelectric polymer nanocomposite capacitors are capable of storing as large amount of
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Supporting Information Supporting Information is available from the Wiley Online Library or from the author.
Acknowledgements This work was supported by the US Office of Naval Research under grant number N00014–11–1–0342. Received: May 10, 2014 Revised: June 12, 2014 Published online: July 16, 2014
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electric energy as their electrochemical counterpart devices that can be delivered at a rate even greater than the commercial high power density dielectric capacitor films. The pristine P(VDFCTFE) and P(VDF-CTFE)/BT binary nanocomposite filled with 20 wt% BT exhibit relatively lower power densities of 550 and 630 MW L−1, respectively, compared to the ternary nanocomposite, which are attributable to their inferior energy densities. The excellent performance coupled with the simplicity and scalability of preparation makes the ternary polymer nanocomposites promising candidates for compact and flexible highenergy high-power capacitive energy storage devices. This successful design of synergistic dual-filler nanocomposites could also be applicable to other materials that require integration of antagonistic properties to enhance collective performance.
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