Decorating TiO2 Nanowires with BaTiO3 Nanoparticles: A New

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Decorating TiO2 Nanowires with BaTiO3 Nanoparticles: A New Approach Leading to Substantially Enhanced Energy Storage Capability of High-k Polymer Nanocomposites Da Kang, Guanyao Wang, Yanhui Huang, Pingkai Jiang, and Xingyi Huang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b16409 • Publication Date (Web): 04 Jan 2018 Downloaded from http://pubs.acs.org on January 4, 2018

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ACS Applied Materials & Interfaces

Decorating TiO2 Nanowires with BaTiO3 Nanoparticles: A New Approach Leading to Substantially Enhanced Energy Storage Capability of High-k Polymer Nanocomposites Da Kanga, Guanyao Wanga, Yanhui Huangb, Pingkai Jianga and Xingyi Huanga* aDepartment

of Polymer Science and Engineering, Shanghai Key Laboratory of Electrical Insulation and

Thermal Aging, Shanghai Jiao Tong University, Shanghai 200240, China bDepartment

of Materials Science and Engineering, Rensselaer Polytechnic Institute, Troy, New York 12180,

United States *Email: [email protected] Abstract: The urgent demand of high-energy-density and high-power-density devices has triggered significant interest in high dielectric constant (high-k) flexible nanocomposites comprising dielectric polymer and high-k inorganic nanofiller. However, the large electrical mismatch between polymer and nanofiller usually leads to earlier electric failure of the nanocomposites, resulting in undesirable decrease of electrical energy storage capability. A few studies show that the introduction of moderate-k shell onto high-k nanofiller surface can decrease the dielectric constant mismatch and thus the corresponding nanocomposites can withstand high electric field. Unfortunately, the low apparent dielectric enhancement of the nanocomposites and high electrical conductivity mismatch between matrix and nanofiller still result in low energy density and low efficiency. In this study, it is demonstrated that encapsulating moderate-k nanofiller with high-k but low electrical conductivity shell is effective to significantly enhance the energy storage capability of dielectric polymer nanocomposites. Specifically, using BaTiO3 nanoparticles encapsulated TiO2 (BaTiO3@TiO2) core-shell nanowires as filler, the corresponding poly(vinylidene fluoride-co-hexafluoropylene) nanocomposites exhibit superior energy storage capability in comparison with the nanocomposites filled by either BaTiO3 or TiO2 nanowires. The nanocomposite film with 5 wt % BaTiO3@TiO2 nanowires possesses an ultrahigh discharged energy density of 9.95 J cm-3 at 500 MV m-1, much higher than that of commercial biaxial-oriented polypropylene (BOPP) (3.56 J cm-3 at 600 MV m-1). This new strategy and corresponding results presented here provide new insights into the design of dielectric polymer nanocomposites with high electrical energy storage capability. Keywords: TiO2 nanowires, BaTiO3 nanoparticles, Electrical energy storage, Polymer nanocomposites, Dielectric constant 1 / 21

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Introduction High-energy-density devices are receiving more and more attention due to their wide applications in electronic and electrical industry1-2. Due to the high power density and fast charge-discharge speed, dielectric-based capacitors have been considered as potential candidates for the next-generation pulse power systems and DC link power applications for electric vehicles3. However, the low energy density of state-of-the-art dielectrics, such as commercially available biaxial-oriented polypropylene (BOPP), can’t meet the ever-increasing urgent demand nowadays4-6. Generally, for dielectric materials, the enhancement of electrical energy density can be achieved by increasing the dielectric constant (k) and/or improving the breakdown strength (Eb). Therefore, the polymer nanocomposites comprising high-k inorganic nanofiller and dielectric polymer matrices with high Eb hold great promise for the fabrication of modern energy storage devices. Nevertheless, the improvement of k is usually at the cost of descending Eb resulted from the large k mismatch between nanofillers and polymers7-9. Thus, it is vital to seek the appropriate strategies for resolving the paradox of k enhancement and Eb decrease in dielectric polymer nanocomposites10-14. Among the numerous polymers with excellent dielectric properties, poly(vinylidene fluoride) (PVDF) and its copolymers are often utilized as the polymer matrices of nanocomposites because of their high k and Eb6, 15-16. One-dimensional (1D) high-k nanofillers, such as BaTiO3 nanowires (NWs) and nanofibers (NFs), have proven to be superior to the nanoparticle counterparts in increasing the apparent dielectric constant and energy storage capability of polymer nanocomposites17-18. However, the large contrast in k between polymers (k < 10) and the nanofillers (k > 200) gives rise to undesirable electric field intensification and distortion in the nanocomposites19-21, leading to decreased capability of withstanding electric field. A common approach to address the issue is to introduce a shell layer with moderate k onto the surface of high-k, which acts as an outer buffer layer to reduce the concentration of electric field22. Thus, extensive efforts have been devoted to enhancing the electrical energy storage capability by introducing the appropriate buffer layers, such as SiO2, Al2O3 and TiO2, on the surface of high-k BaTiO3 NWs23-25. Nevertheless, as opposed to this strategy, the encapsulation of relative high-k outer shell onto the surface of nanofillers with moderate k was rarely investigated. Apart from the extensively used high-k BaTiO3 NWs, moderate-k NWs such as TiO2 have been also used to enhance the electrical energy storage capability of dielectric polymer nanocomposites because of their ease of large scale preparation26-27. Their moderate k (33, 40 and 115 for amorphous, anatase and rutile films, 2 / 21


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respectively28-29) can decrease the dielectric constant mismatch with the polymer matrix, making the nanocomposites theoretically withstand high electric field30. However, the semiconductor nature of TiO2 NWs can lead to an electrical conductivity higher than 10-4 S/m, which is depending on the preparation methods and heat treatment conditions31. In spite of the low dielectric constant mismatch, the high electrical conductivity contrast between polymer and nanofiller also results in significant interfacial polarization because of the space charge accumulation between the polymer/nanofiller interface32. A strong interfacial polarization also causes high dielectric loss and low breakdown strength in the nanocomposites, leading to low energy density and low energy efficiency19. Herein, core-shell structured BaTiO3@TiO2 NWs by decorating high-k but low electrical conductivity BaTiO3 NPs onto TiO2 NWs were prepared and used to increase the electrical energy storage capability of poly(vinylidene fluoride-co-hexafluoropylene) (PVDF-HFP) nanocomposites. The encapsulation of BaTiO3 outer shell not only suppressed the undesirable effects (e.g., high electrical conductivity) of TiO2 NWs on nanocomposites, but also brings more significant improvement in k and breakdown strength than the employment of bare TiO2 NWs or BaTiO3 NWs, resulting in significantly enhanced electrical energy storage capability of the PVDF-HFP nanocomposites. This strategy provides new insights into the design of high-k polymer nanocomposites with improved dielectric properties and enhanced electrical energy storage capability.

Experimental section Materials Titanium dioxide (TiO2, P25, ≥99.5%) nanoparticles were purchased from Sigma-Aldrich and used as received. Barium hydroxide octahydrate (Ba(OH)28H2O), diethylene glycol and tetrabutylammonium hydroxide solution (TBAH, 40 wt %) were purchased from Tansoole (Shanghai, China) and used as received. PVDF-HFP containing 15% HFP was kindly provided by Solvay Plastics (Shanghai, China). N,N-Dimethyl formamide (DMF) and other reagents/solvents were obtained from Tansoole (Shanghai, China). Preparation of BaTiO3@TiO2 NWs The TiO2 (Anatase) NWs were prepared using the hydrothermal method reported in our previous work26. Then, BaTiO3@TiO2 NWs were fabricated by converting TiO2 NWs surface via a hydrothermal reaction. Typically, 6 mmol TiO2 NWs (0.48 g) were firstly immersed in a 90 mL Teflon autoclave that filled with a solution of 3 mmol Ba(OH)28H2O (0.95 g) in 10 mL diethylene glycol, 10 mL ethanol, 3 mL 2-propanol, 1.2 g TBAH (40 wt %), and 3 / 21



14 mL deionized water. The mixture was stirred for 2 h and the autoclave was put into an oven of 180 oC for another 2 h. The obtained products were extensively washed to remove any unreacted excess reagent with deionized water and ethanol and then dried in air. The finally obtained products were denoted as BaTiO3@TiO2 NWs. Preparation of Nanocomposite Films The preparation of nanocomposite films is described as follows: Firstly, the given amount of BaTiO3@TiO2 NWs was homogeneously dispersed in DMF and ultrasonicated for 2 h. After stirring for another 24 h, P(VDF-HFP) was added into the mixture. The corresponding composite mixture was stirred vigorously for 48 h and then casted into films on a smooth and clean glass. In order to accelerate the DMF evaporation and thoroughly remove the trace solvent, the nanocomposite films were dried at 40 oC in air and vacuum for several hours and 12 h, respectively. Before peeling from the glass substrates, the corresponding films were heated at 200 oC for 5 min and then quenched in ice water immediately in order to increase the nonpolar γ-phase in the polymer matrix. The obtained films were dried at 40 oC in vacuum for another 12 h. Nanocomposites with different weight fractions (5%, 10%, 15%, and 20%) of BaTiO3@TiO2 NWs were fabricated, respectively. Then the nanocomposite films with 20 wt % raw BaTiO3 and TiO2 NWs NWs were also prepared for comparison, respectively. The thickness of film samples mentioned in this article was all about 15 µm. Characterization The morphology of the synthesized NWs and the corresponding nanocomposite films were characterized by a Nova NanoSEM scanning electron microscope (SEM, 450, FEI, USA) and a JEM-2010 transmission electron microscope (TEM, JEOL, Japan). X-ray diffraction (XRD) patterns were collected using Rigaku D/max-2200/PC X-ray powder diffractometer (Cu Kα irradiation, 40 KV, 20 mA). We conducted X-ray photoelectron spectrum (XPS) measurements to analyze the elements composition of the NW surface using an Axis UltraDLD spectrometer with a monochromated Al Kα source, which is a product of Shimadzu-Kratos Analytical (UK) Furthermore, the dielectric spectra of the samples were obtained from a Novocontrol Alpha-A high resolution dielectric analyzer at a wide range of temperatures (-50-150 oC) and frequency (10-1-106 Hz). A layer of gold was evaporated on both surfaces of the films to serve as electrodes (the diameter is 12 mm and the thickness is 50-100 nm). The DC electric breakdown tests were conducted at room temperature with a ramping rate of 500 V s-1. The electric displacement-electric field (D-E) loops and leakage current tests were both performed by Precision Multiferroic Materials Analyzer (Radiant Inc., USA). A layer of copper was evaporated on both sides of 4 / 21


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films samples to serve as electrodes (the diameter is 3 mm and the thickness is 50-100 nm) during the measurements of electric breakdown, D-E loops and leakage current tests.

Results and discussion Preparation and Characterization of the Nanowires The preparation of TiO2/BaTiO3 core/shell NWs was accomplished by BaTiO3 surface conversion. After converting partial TiO2 into BaTiO3 at 180

for 2 h, numerous BaTiO3 NPs with diameter of ~100 nm yielded a

thick layer on the surface of TiO2 NWs. Furthermore, Figure 1 demonstrated the successful encapsulation of TiO2 NWs by BaTiO3 NPs. As shown in Figure 1, the pristine TiO2 NWs show smooth surface. However, after the high-temperature hydrothermal growth of BaTiO3, the smooth surface was covered by aggregated nanoparticles, implying the successful injection of Ba2+.

Figure 1. TEM images of (a-b) pristine TiO2 and (c-d) BaTiO3@TiO2 NWs. Schematic illustration of the (e) pristine TiO2 and (f) BaTiO3@TiO2 NWs.

Typical EDX elemental mapping images (Figure 2a) and elemental analysis (Figure 2b) were introduced to validate the composition of BaTiO3@TiO2 NWs. Apparently, the appearance of element Ba revealed the coating 5 / 21



of BaTiO3 upon the surface of TiO2, further identifying the results of electron microscopy. Besides, the generation of BaTiO3 was also confirmed by the comparison of XPS spectra shown in Figure 2c. With the BaTiO3 conversion, the Ba peak intensity shows an apparent enhancement, while those of Ti and O peaks were just the opposite. XRD patterns shown in Figure S1 (Supporting Information) also verified the change of the composition and crystallography of TiO2 NWs after BaTiO3 conversion. As shown in Figure S1, the obvious peaks of BaTiO3 (JCPDS 81-2203) evidenced the consumption of TiO2 (JCPDS 21-1272) into BaTiO3. These characterizations provided consistent and convincing evidences of the successful transition of BaTiO3 shell layers upon the surface of TiO2 NWs.

Figure 2. (a) EDX elemental mapping images of BaTiO3@TiO2 NWs. Ba mapping in cyan, Ti mapping magenta, O mapping in yellow. (b) EDX elemental analysis of BaTiO3@TiO2 NWs. (c) XPS spectra of pristine BaTiO3, pristine TiO2 and BaTiO3@TiO2 NWs.

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Dielectric Properties of Polymer Nanocomposites It has been demonstrated in previous work that compared with high-k BaTiO3 NWs, the incorporation of moderate-k TiO2 NWs can give rise to high breakdown strength of polymer nanocomposites because of the relatively low dielectric constant mismatch between polymer matrices and NWs26. However, the improvement of breakdown strength only takes place at a low concentration. In addition, the enhancement of apparent dielectric constant is lower in comparison with the BaTiO3 NWs based nanocomposites, resulting in lower discharged energy density below the breakdown electric field26. Thus, it is of vital importance to give full play to the role of TiO2 NWs while effectively avoiding its disadvantage. Herein, high-k but low electrical conductivity BaTiO3 NPs (about 10-10 S/m according to Ref.19 ) was introduced as the encapsulation shell to restrain undesirable high electrical conductivity and to increase k values of NWs. As shown in Figure 3a, the k of the nanocomposites comprising BaTiO3@TiO2 NWs showed an increasing trend with raising the content of nanofillers, since the proposed nanowires possessed much higher k than the polymer matrix. In order to evaluate the influence of BaTiO3 encapsulation on the performance of those nanocomposites, the comparison between core-shell structured nanowires and the pristine ones was presented. As shown in Figure 3b, the nanocomposite of TiO2 NWs possessed lower k than that with BaTiO3 NWs due to the lower k of anatase TiO2 (k ≈ 40) to in comparison with BaTiO3 (k > 200). In general, the k of the nanocomposites of BaTiO3@TiO2 NWs might lie in the range of these two kinds of aforementioned nanocomposites with bare NWs according to the effective medium theory such as Maxwell-Garnett formula33. Nevertheless, the core-shell structured nanowires exhibited more effective enhancement of k in the nanocomposites at the same weight fractions loading (Figure 3b). As shown in Figure S2, the nanocomposites comprising BaTiO3@TiO2 NWs show higher k when compared with the nanocomposites with raw BaTiO3 NWs or TiO2 NWs at the same loading. This abnormity could be attributed to the specific core-shell structure of BaTiO3@TiO2 NWs, since additional polarization might be induced in the interfacial region between BaTiO3 NPs and TiO2 NWs. Namely, the lattice distortion in the interfacial area, resulted from the different structured lattice of BaTiO3 and TiO2, would give rise to additional micro defects to facilitate the formation of carriers and channels for the charge shifting over long distance. The interfacial polarization stemmed from the charge migration through those defects would finally increase the k of those nanocomposites. In fact, interfacial polarization within the nanofiller induced dielectric enhancement has been documented in different composites systems, including PVDF nanocomposites with TiO2 encapsulated BaTiO3 NWs or nanoparticles34-36. 7 / 21



One can also see that the nanocomposites with BaTiO3@TiO2 NWs show suppressed low frequency and high frequency dielectric losses in comparison with those nanocomposites of BaTiO3 or TiO2 NWs. For PVDF based polymers, the low frequency and high frequency dielectric losses originate from the electrical conduction process and segmental motions of amorphous phase32, 37, respectively. The lower dielectric loss at low frequencies indicates the decreased electrical conductivity in the BaTiO3@TiO2 nanocomposites, which is consistent with the leakage current results shown below. The decreased high frequency dielectric loss should originate from the stronger restriction effect of BaTiO3@TiO2 NWs on the segmental motions of amorphous phase of PVDF-HFP38. It would be shown later that the decrease of low frequency loss is of particular importance to achieve higher electrical energy storage efficiency.

Figure 3. Frequency dependence of dielectric constant and dielectric loss of P(VDF-HFP)-based nanocomposites with (a) different weight fractions of BaTiO3@TiO2 NWs and (b) 20 wt % of BaTiO3@TiO2, TiO2, 8 / 21


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and BaTiO3 NWs.

The temperature-dependent dielectric spectroscopy was obtained to further investigate the electrical properties of those proposed nanocomposites. As shown in Figure 4, the dielectric loss of the neat polymer at high temperature and low frequency was significantly reduced by introducing these NWs, especially BaTiO3@TiO2 and BaTiO3 NWs. Furthermore, it should be noted that, when compared with the TiO2 NWs, the BaTiO3@TiO2 NWs remarkably reduced the dielectric loss of P(VDF-HFP) nanocomposites due to the encapsulation of TiO2 core NWs by BaTiO3 shell. As discussed earlier, these results indicate that compared with the TiO2 NWs, BaTiO3 is more effective to suppress the electrical conductivity of the nanocomposites. For demonstrating the influence of nanofillers on the high frequency loss of the polymer matrix, Figure S2 presented the frequency and temperature dependent imaginary electric modulus (M″) of the nanocomposites with 20 wt % NWs, as well as pure P(VDF-HFP). The electric modulus (M″) analyzing is useful to understand the relaxation processes because that the large variations of dielectric constant and dielectric loss of materials at low frequencies can be minimized39. As we can see in Figure S3 (ESI†), two relaxation peaks show up both in the pure P(VDF-HFP) and the nanocomposites. The relaxation peak at low frequency and 50 oC is attributed to the interfacial polarization. As the temperature increases, the mobility of the carrier enhances, the relaxation time decreases, and the relaxation peak shifts to high temperature and high frequency. Another relaxation peak at low temperature and high frequency is attributed to the segment movement in the amorphous region of the polymer, named αa relaxation peak. As the temperature increases, the intensity of the relaxation peak is constantly increasing and the relaxation peak shifting to high frequency40-41. As shown in Figure S3, the αa relaxation at low temperatures and high frequencies and the interfacial polarization at high temperatures and low frequencies were significantly suppressed due to the incorporation of NWs. However, the M″ of the nanocomposites seemed to be insensitive to different types of NWs, as the intuitional variation tendencies of these nanocomposites were quite similar.

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Figure 4. Temperature-dependent dielectric spectra of (a) pure P(VDF-HFP) and nanocomposites with 20 wt % of (b) BaTiO3@TiO2, (c) TiO2, and (d) BaTiO3 NWs.

Energy Storage Capability of Polymer Nanocomposites The energy density (Ue) of dielectric materials can be calculated by the equation: U e = ∫ EdD

(1)

where E and D represent the applied electric field and corresponding electric displacement, respectively. Thus, the energy storage capability of the proposed nanocomposites could be acquired by typical D-E loops of an optimized Sawyer-Tower circuit. Compared with the slim loops of neat polymer (Figure S4), the nanocomposites with BaTiO3@TiO2 NWs shown in Figure S5 depicted fatter circuit as the nanofiller loading increased gradually. That is, the incorporation of nanowires gave rise to higher remnant polarization, which usually resulted in lower energy density and efficiency. Notably, benefited from the special core-shell structure of BaTiO3@TiO2 NWs, the corresponding nanocomposites possess slimmer D-E loops (Figure S5d) in comparison with those consisting of raw TiO2 (Figure S5e) and (d) BaTiO3 NWs (Figure S5f) at the same loading. The slimmer D-E loops means a higher charge-discharge efficiency, which originates from lower dielectric loss and/or lower electrical conduction 10 / 21


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(the leakage current density will be discussed in the following). The summarized discharged energy densities and efficiencies shown in Figure 5 demonstrated a more visual approach. Along with the increasing electric field, both the total stored and discharged energy density exhibited a rising trend (Figure 5a and Figure S6). On the contrary, charge-discharge efficiency decreased gradually except a little fluctuation at higher electric filed. As shown in Figure 5a, the nanocomposite film with 5 wt % BaTiO3@TiO2 NWs discharged an energy density of 9.95 J cm-3 at 500 MV m-1, much higher than that of pure polymer (7.60 J cm-3 at 440 MV m-1). Nevertheless, further increasing the loading of NWs would bring negative effect to the energy storage capability, since those nanocomposites withstand decreased electric fields with increased nanofiller loading. For instance, the nanocomposite films with 10 wt % BaTiO3@TiO2 NWs possessed a lower discharged energy density of 8.34 J cm-3 at a reduced electric field of 440 MV m-1. Moreover, by increasing the loading to 15 wt % and 20 wt %, the value dropped to 7.05 J cm-3 at 400 MV m-1 and 5.13 J cm-3 at 340 MV m-1, respectively. The prominent effect of BaTiO3 NPs encapsulation upon TiO2 NWs was further evidenced by the comparison of energy storage capability and efficiency between these nanocomposites with core-shell structured nanowires and bare ones. The highest energy density obtained from those nanocomposites with 20 wt % pristine nanowires was ~2.01 J cm-3 at 300 MV m-1 because of their high remnant polarization. Conversely, the incorporation of 20 wt % BaTiO3@TiO2 NWs led to a much higher energy density of 5.13 J cm-3 at 340 MV m-1. Furthermore, the charge-discharge efficiency of the nanocomposite comprising BaTiO3@TiO2 NWs could be still kept upon 45%, while the nanocomposites with other two kinds of raw nanowires exhibited much lower efficiencies of ~10%. Namely, the sharp decline of energy densities and efficiencies was significantly suppressed by the employment of BaTiO3@TiO2 NWs. Compared to the other two kinds of bare nanowires, most of the charges in the nanocomposites with BaTiO3@TiO2 NWs could only transfer in the enclosed interfacial zone between the core TiO2 NWs and the shell BaTiO3 NPs. Thus, the electric percolation pathway was remarkably reduced as the migration of charges in the nanocomposites was restricted inside the core-shell structured nanowires34-35, 42. The leakage current densities discussed below further validated the above hypothesis.

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Figure 5. (a) Discharged energy densities and (b) charge-discharge efficiencies of P(VDF-HFP)-based nanocomposites with different weight fractions of BaTiO3@TiO2 NWs. For comparison, the properties of composites comprising 20 wt % of raw TiO2 and BaTiO3 NWs were also included.

Breakdown Strength and Leakage Current Density of Nanocomposites Generally, for the linear dielectric materials, the electric displacement (D) varies linearly versus the electric field (E). Therefore, the energy storage capability of linear dielectrics can be calculated by a simplified form of equation (1): 1 U e = ε r ε 0 Eb 2 2

(3)

where Eb is the breakdown strength, and εr and ε0 (8.8542 × 10-12 F m-1) are the relative and vacuum dielectric constants, respectively. Therefore, a high level of Eb plays a key role in achieving a high energy storage 12 / 21


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capability. As can be seen in Figure 6a, at lower loading of BaTiO3@TiO2 NWs (e.g., 5 wt %), the nanocomposite film possessed comparable Eb (~480 MV m-1) to the pure polymer. However, higher loadings would inevitably result in the decrease of Eb. For instance, the nanocomposite with 20 wt % BaTiO3@TiO2 NWs could withstand the electric field of about 327.5 MV m-1. It's worth noting that those nanocomposites with raw nanowires could only tolerate the electric field of ~300 MV m-1. Their inferior Eb exactly reflected the outstanding performance of BaTiO3@TiO2 NWs due to restriction of outer BaTiO3 shell to the inner TiO2 core with semiconductive characteristic26. Leakage current density was provided to further investigate the superior performance of BaTiO3@TiO2 NWs to the pristine ones. Apparently, the leakage current densities at 100 MV m-1 shown in Figure 6b demonstrated that more structural defects and field concentration on polymer matrix resulted from higher loading of nanowires would give rise to higher leakage current densities. Besides, the nanocomposite film consisting of core-shell BaTiO3@TiO2 NWs exhibited a much lower density in comparison with those with same loading of bare TiO2 or BaTiO3 NWs, further verifying the suppression of BaTiO3 shell upon the TiO2 core. Namely, the charges in the nanocomposites could rarely go through the core-shell structured nanofillers into the polymer matrix.

Figure 6. (a) Breakdown strength and (b) leakage current densities of P(VDF-HFP)-based nanocomposites with different weight fractions of BaTiO3@TiO2 NWs. For comparison, the properties of composites comprising 20 wt % of raw TiO2 and BaTiO3 NWs were also included.

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Insights from the Computer Simulation The dielectric property and energy stoarge capability of dielectric composites are closely associated with the applied electric field intensity and local electric field distribution within the composites19, 43. Here finite element simulation was used to analyze the local electric field distribution in the polymer matrix after considering the dielectric constant contrast between the polymer matrix and the NWs, the NW aspect ratio, a random distribution and the NW loading fractions44. Figure 7a shows the 3D distribution of the normalized local electric field in the nanocomposites with 20 wt% NWs. For a qualititive analysis, the probability distribution function of the normalized local electric field in the nanocomposites were also given in Figure 7d. For the nanocomposites with TiO2 NWs, the electric field distribution peak strongly nears the normlized electric field of 1.03, then drops quickly, resulting in a weak tail at high normlized electric field. In the case of nanocomposites with BaTiO3 and BaTiO3@TiO2 NWs, the main distribution peaks decrease and the tails are more prominent in comparison with the nanocomposites with TiO2 NWs. In addition, the nanocomposites with BaTiO3 NWs shows a stronger main distribution peak and a slightly lareger tail. Such results suggest that, compared with the nanocomposites with TiO2 NWs, a larger portion of polymer matrix sustains stronger electric fields in the the nanocomposites with BaTiO3 or BaTiO3@TiO2 NWs45. Thus the higher dielectric enhancent in the corresponding BaTiO3 or BaTiO3@TiO2 NWs nanocomposites can be easily understood. On the other hand, in the case of the nanocomposites with BaTiO3 NWs, the slightly stronger electric field enhancement is not consistent with the slightly weak dielectric enhancent. As discussed ealier, however, such a result further demontrated the additional effect of interfacial polarization on dielectric enhancement of the BaTiO3@TiO2 nanocomposites, which was induced in the interfacial region between BaTiO3 NPs and TiO2 NWs but was not considered in the electric field simulation. The probability distribution function of the calulated local electric field in the nanocomposites with BaTiO3 or BaTiO3@TiO2 NWs can explain the difference of breakdown strength. Namely, a higher electric field enhancement results in a lower breakdown strength of the nanocomposites with BaTiO3 NWs. However, Figure 7 can’t not fully illustrate the breakodwn strength results of all the nancomposites since those with TiO2 NWs have the lowest breakodwn strength but exhibit a weak electric field enhancement. Such results show that, apart from the dielectric cosntant contrast, the electrical conductivity of the filler might also play a vital role in influencing the breakdown strength of the nanocmposites. The lower breakdown strength of the nanocomposites with TiO2 NWs should be mainly attributed to the highest leakage currents originating from the 14 / 21


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high electrical conductivity (>10-4 S/m) of anatase TiO2,26, 31 as shown in Figure 6b. Keep this insight in mind, the lowest breakdown strength of the nanocomposites with BaTiO3 NWs can be understood by the high dielectric constant contrast between NWs and matrix and the high electrical conductivity of TiO2 NWs. On one hand, the BaTiO3 NWs have a higher dielectric constant contrast with the matrix when compared with the TiO2 NWs, resulting in larger electric field enhancement in the nanocomposites of BaTiO3 NWs. On the other hand, their nanocomposites also have the comparable high leakage current with the nanocomposites of TiO2 NWs (Figure 6), resulting in signicant heat generation under high electric field. In brief, this simulation shows that both the dielectric constant and electical conductivity of the shell layers of the core-shell filler play important roles in detemining the dielectric properties of the composites, finally resulting in significant influence on the electrical enegy storage capbility of the composites.

Figure 7. 3D distribution of normalized local electric field in the P(VDF-HFP)-based nanocomposites with 20 wt% TiO2 (a), BaTiO3 (b) and BaTiO3@TiO2 (c) NWs. (d) The corresponding probability distribution function of the normalized local electric field in the nanocomposites. The normalized local electric field is defined as the ratio of the computed local electric field to the applied electric field. 15 / 21



Conclusions In summary, BaTiO3 nanoparticles decorated TiO2 NWs were successfully synthesized by hydrothermal conversion of TiO2 NWs. The outer shell consisting of aggregated BaTiO3 NPs simultaneously has high dielectric constant and low electrical conductivity, which not only suppressed the undesirable effects caused by the high electrical conductivity of TiO2 NWs but also enhanced k and breakdown strength of the nanocomposites, resulting in significantly enhanced electrical energy storage capability of the PVDF-HFP nanocomposites. For instance, the nanocomposite film with 5 wt % BaTiO3@TiO2 NWs discharged an ultrahigh energy density of 9.95 J cm-3 at 500 MV m-1, much higher than that of BOPP (3.56 J cm-3 at 600 MV m-1) and pure polymer (7.60 J cm-3 at 440 MV m-1). Moreover, the nanocomposite film with BaTiO3@TiO2 NWs also possessed higher k, Eb and energy density than those consisting of raw nanowires. The outstanding performance of these proposed nanocomposites could be attributed to the utilization of the designed core-shell structured nanowires. The results and method presented here demonstrates that both the dielectric constant and electical conductivity of the filler outer layer are important in detemining the dielectric properties of the nanocomposites, which should be simultaeously considered when designing new filler for dielectric polymer nanocomposites with high electrical enegy storage capbility.

ASSOCIATED CONTENT

Supporting Information. XRD spectra patterns of BaTiO3@TiO2 NWs, dielectric constant of different nanocomposites at 1000 Hz, frequency dependent of imaginary electric modulus at various temperature for different nanocomposites, TGA spectra, additional dielectric spectra, and D-E loops of P(VDF-HFP) and the nanocomposites, total stored energy densities of different nanocomposites. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION INFORMATION Corresponding Author

*Email: [email protected] (X.Y.H.) Notes 16 / 21


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The authors declare no competing financial interest. Acknowledgements The financial support from National Natural Science Foundation of China (nos. 51522703, 51477096) and Special Fund of the National Priority Basic Research of China (Grant 2014CB239503) was acknowledged.

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Decorating TiO2 Nanowires with BaTiO3 Nanoparticles: A New

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