Significantly Enhanced Energy Density by Tailoring the Interface in

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Applications of Polymer, Composite, and Coating Materials

Significantly Enhanced Energy Density by Tailoring the Interface in a Hierarchical-Structured TiO2-BaTiO3-TiO2 Nanofillers in PVDF Based Thin Film Polymer Nanocomposite Prateek Prajapati, Ritamay Bhunia, Shahil Siddiqui, Ashish Garg, and Raju Kumar Gupta ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b01359 • Publication Date (Web): 20 Mar 2019 Downloaded from http://pubs.acs.org on March 22, 2019

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ACS Applied Materials & Interfaces

Significantly Enhanced Energy Density by Tailoring the Interface in a HierarchicalStructured TiO2-BaTiO3-TiO2 Nanofillers in PVDF Based Thin Film Polymer Nanocomposite Prateek,† Ritamay Bhunia,‡ Shahil Siddiqui,† Ashish Garg‡ and Raju Kumar Gupta*,†,‼ †

Department of Chemical Engineering, Indian Institute of Technology Kanpur, Kanpur

208016, Uttar Pradesh, India ‡

Department of Materials Science and Engineering, Indian Institute of Technology Kanpur,

Kanpur 208016, Uttar Pradesh, India ‼

Center for Environmental Science and Engineering, Indian Institute of Technology Kanpur,

Kanpur 208016, Uttar Pradesh, India *

Corresponding author. Tel: +91-5122596972; Fax: +91-5122590104.

E-mail address: [email protected]

ABSTRACT: Dielectric polymer nanocomposites with high breakdown field and high dielectric constant have drawn significant attention in modern electrical and electronic industries due to their potential applications in dielectric and energy storage systems. The interfaces of the nanomaterials play a significant role in improving the dielectric performance of polymer nanocomposites. In this work, polydopamine (dopa) functionalized TiO2-BaTiO3TiO2 (TiO2-BT-TiO2@dopa) core@double-shell nanoparticles have been developed as novel nanofillers for high energy density capacitor application. The hierarchically designed nanofillers help in tailoring the interfaces surrounding the polymer matrix as well as act as individual capacitors in which core and outer TiO2 shell functions as capacitor plate because of their high electrical conductivity while the middle BT layer functions as a dielectric medium due to high dielectric constant. Detailed electrical characterizations have revealed that TiO2BT-TiO2@dopa/PVDF possess maximum relative dielectric permittivity (εr), breakdown strength (Eb), as well as energy densities in comparison to PVDF, TiO2/PVDF, TiO2@dopa/PVDF, TiO2-BT@dopa/PVDF polymer nanocomposites. The εr and energy 1 ACS Paragon Plus Environment

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density of TiO2-BT-TiO2@dopa/PVDF was 12.6 at 1 kHz and 4.4 J cm-3 at 3128 kV cm-1, respectively, which was comparatively much higher than commercially available biaxially oriented polypropylene (BOPP) having εr of 2.2 and the energy density of 1.2 J cm-3 at much higher electric field of 6400 kV cm-1. It is expected that these results will further open new avenues for the design of novel architecture for high-performance polymer nanocompositesbased capacitors having core@multishell nanofillers with tailored interfaces. KEYWORDS: dielectrics, BaTiO3 nanoparticles, core-shell nanomaterials, polymer nanocomposites, capacitors

◼ INTRODUCTION In recent decades, novel dielectric materials for energy harvesting and storage systems are highly desirable due to their capability of ultrafast charging-discharging ability as well as ultrahigh power densities and their applications in numerous electrical and electronics industries such as pulsed power supply technology, inverters, hybrid vehicles, portable electronic devices, power grids, etc.1-10 Dielectric capacitors, commonly known as electrostatic capacitors, need to be developed with high energy density as well as efficiency, flexibility, low dielectric loss, long life time, lightweight, wide operating temperature range, reliability under high operating voltages, low-cost as well as environmentally friendly.11-14 In general, the energy density (Ud) of the capacitors depends on the applied electric field (E) and polarization (P) as equation 1: 0

U d =  EdP

(1)

Pmax

while for the linear dielectrics can be expressed as equation 2: Ud =

1  o r Eb2 2

2 ACS Paragon Plus Environment

(2)

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where εo and εr are the vacuum permittivity (8.854 × 10-12 F m-1) and relative permittivity, respectively and Eb is the breakdown strength. The present state-of-the-art of the commercially available electrostatic capacitor is biaxially oriented polypropylene (BOPP), which has high breakdown strength of 6400 kV cm-1, but suffers from low dielectric constant of 2.2 as well as low energy density of 1.2 J cm-3, thereby making it unsuitable for the next-generation dielectric materials to meet the growing demand for efficient and high energy density storage devices.15 Usually, ceramics such as barium titanate (BT), barium strontium titanate (BST), BZT−BCT, CCTO or Pb(ZrTiO3) (PZT) have substantially high dielectric constant while polymers such as polyvinylidene fluoride (PVDF), polystyrene (PS), polypropylene (PP) exhibit lower εr but high Eb.5,6,8,16-22 Also, there is always a trade-off between εr and Eb, related as E b  ε −r 0.5 for monophasic dielectric materials, which further restricts improvement in the energy density.3 Thus, polymer nanocomposite-based capacitors have been researched extremely due to possibility of simultaneous enhancement in εr and Eb provided by exploring high-εr of dielectrics and high Eb of polymers, respectively.3 The introduction of high-εr nanomaterials makes the electric field inhomogeneous at the interface due to interfacial polarization in the polymer matrix leading to early breakdown. Moreover, nanomaterials have high surface energy which lead to agglomeration. Further, these also create a large number of interfaces in the polymer matrix thereby introducing a large number of defects as well as porosity and electron conduction responsible for high dielectric loss tangent (tan δ) and hence lowering Eb of the polymer nanocomposites.23-26 Thus, tailoring the interface and surface modification of the nanomaterials play a crucial role in enhancing Eb as well as energy density of the polymer nanocomposites.26-31 Generally, nanomaterials have been modified as coreshell/core-multishell structures based on two major reasons in order to improve the dielectric performances. Firstly, a high-εr nanomaterial is rationally modified in such a manner that the εr of each layer follows a decreasing trend from core to shell. This alleviates the sharp εr contrast 3 ACS Paragon Plus Environment


between the nanofillers and polymer matrix interface and hence mitigates the electric field concentration and current density. For example, researchers have modified high-εr BT surface with alumina (Al2O3) shell,32-36 titanium dioxide (TiO2) shell37-39 as well as sheets,40 TiO2@Al2O341 double-shell, polymers42 like poly(methyl methacrylate) (PMMA), PS,43 polydopamine (dopa), polylactic acid (PLA),44 etc. to improve the energy density of the polymer nanocomposites based capacitors. The second strategy involves selection of low-εr nanomaterials such as TiO2 as core,1,45-47 so that the electric field is relatively more homogenous as compared to high- εr nanomaterials in the polymer matrix. However, the main concern about using TiO2 as nanofillers is their large electrical conductivity (σ) (10-4 S m-1) as compared to BT (10-10 S m-1) which inherently leads to high-conductivity mismatch between nanofillers and polymer, thereby causing significant interfacial polarization and consequently high dielectric loss and low Eb.1,48 The literature suggests that the exact phenomenon occurring at the interface of TiO2 (high-σ, low-εr) and BT (high-εr, low-σ) in the polymer matrix is still not clear. Thus, the rationale engineering of the nanomaterial interface is required to modulate the interfacial polarization between the nanofillers and polymer matrix as well as forming electrical barriers across the film in order to enhance the dielectric performance of the capacitors. In this work, polymer nanocomposite-based capacitors comprising of TiO2-BTTiO2@dopa as core-multishell nanofillers and PVDF as polymer matrix is developed and their dielectric properties are systematically studied. Here, the nanofiller TiO2-BT-TiO2 is hierarchically designed to work as individual capacitor in which TiO2 NP and shell work as capacitor plates as it has comparatively high electrical conductivity as compared to BT while the middle BT layer serves as the dielectric medium because of high-εr. Moreover, in the multishell BT-TiO2-dopa layer; BT, TiO2 and dopa are treated as polarization layer, buffer layer, and dispersion layer, respectively. The results show that the present novel interfacial


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architecture simultaneously improved the dielectric constant as well as breakdown strength and offers a promising approach in substantially improving the dielectric performance of the capacitors.

◼ RESULTS AND DISCUSSION The surface modification of TiO2 NPs was carried out using the steps mentioned in experimental section (Scheme 1).

Scheme 1. Schematics for the surface modification of TiO2 NPs. The thermal stability and difference in the composition of pristine TiO2 and surface modified TiO2 NPs was studied by TGA (Figure S1). Without dopa surface modification, TiO2, TiO2-BT, and TiO2-BT-TiO2 NPs have negligible weight losses of 2.5, 4.4, and 3.4%, respectively in the temperature range of 200-600 oC indicating that very little hydroxyl groups are bonded on the surface.49 With the dopa functionalization, the corresponding weight losses of TiO2@dopa, TiO2-BT@dopa, TiO2-BT-TiO2@dopa NPs are 9.8, 9.2, and 13.1%, respectively. Also, in the temperature range of 25 – 600 oC, the TGA curves show that the residual weights of 4.7, 6.3 and 7.1% for TiO2, TiO2-BT, TiO2-BT-TiO2 NPs, respectively while for TiO2@dopa, TiO2-BT@dopa, TiO2-BT-TiO2@dopa NPs are 13.3, 13.0 and 18.1%, respectively. Thus, the dopa content in TiO2@dopa, TiO2-BT@dopa, and TiO2-BTTiO2@dopa are 8.6, 6.7, and 11.0%, respectively responsible for thermal degradation and higher weight losses as compared to without dopa surface modified NPs.50



The surface morphologies of the NPs were studied using FESEM as shown in Figure 1a-c. The images show that the all the NPs are of uniform shapes and sizes. Figure 1d-f shows HRTEM images of TiO2, TiO2-BT, and TiO2-BT-TiO2 NPs. The HRTEM image of pristine TiO2 NPs shows the lattice fringes with interplanar spacing of 3.5 Å which represents (101) plane of anatase phase (ICDD file number: 21-1272) (Figure 1d). The coating of BT shell on the TiO2 NPs shows (111) plane of interplanar distance 2.3 Å of tetragonal BT (ICDD file number: 05-0626) (Figure 1e). Further, the TiO2 planes of (101) and (110) show anatase and rutile (ICDD file number: 21-1276) phases, respectively confirming TiO2 coating on the TiO2BT NPs (Figure 1f). The average thickness of each layer was determined from TEM images as shown in Figure S2. The average diameter of the purchased TiO2 NPs was ~21 nm and the thickness of BT and TiO2 in TiO2-BT and TiO2-BT-TiO2 was ~2 nm and ~5 nm, respectively (Figure S2 a-f). Also, the average thickness of dopa on TiO2 NPs was 2.4 nm (Figure S2 g, h).

Figure 1. FESEM and HRTEM images of (a,d) TiO2 NPs, (b,e) TiO2-BT NPs, and (c,f) TiO2BT-TiO2 NPs. 6 ACS Paragon Plus Environment

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The crystallinity of TiO2, TiO2-BT, TiO2-BT-TiO2 NPs was further confirmed from the XRD data as shown in Figure 2a. The peaks of TiO2 NPs are matched with anatase phase (ICDD file number: 21-1272) along with presence of small amount of rutile phase (ICDD file number: 21-1276). After the BT NPs were grown hydrothermally, the XRD data of TiO2-BT NPs confirms the presence of both TiO2 as well as tetragonal phase of BT (ICDD file number: 05-0626). A clear insight near 2θ = 45o suggests that the peak can be deconvoluted into two separate peaks at 44.8o and 45o corresponding to (002) and (200) planes, respectively as shown in inset of Figure 2a. However, a small amount of barium carbonate (BaCO3) is also formed during the BT NPs synthesis as deduced from ICDD file number 05-0378 which may be due to the reaction between Ba2+ in solution and CO2 in atmosphere (Figure S3).51

Figure 2. XRD spectra of (a) TiO2 NPs, TiO2-BT NPs and TiO2-BT-TiO2 NPs, and (b) PVDF film and TiO2-BT-TiO2@dopa/PVDF polymer nanocomposites. It has been previously reported that BaCO3 usually forms as an impurity during BT synthesis.52-55 During the reaction, the atmospheric CO2 is dissolved as CO32- in the solution and reacts with Ba2+ to form BaCO3 as per reactions (3-5):56

Ba (OH )2  Ba 2+ + 2OH −

(3)

CO2 + 2OH −  CO32− + H 2O

(4)



Ba 2+ + CO32−  BaCO3

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(5)

However, López et al. deduced that the most effective way to extract the BaCO3 contamination from as-synthesized BT is acid wash followed by water wash (reaction 6).57 BaCO3( s ) + 2 HCl( aq) → BaCl2( aq) + H 2O(l ) + CO2( g )

(6)

XRD confirms that BaCO3 is removed by acid wash (see experimental methods) as shown in Figure S3. Further, the XRD spectra of TiO2-BT-TiO2 NPs show similar peaks of TiO2 and BT suggesting that no additional product is formed during the TiO2 shell synthesis on the TiO2-BT NPs (Figure 2a). In addition, the BaCO3 content is approximately calculated by measuring the weights of dried samples before and after acid and water wash. It is found that the BaCO3 content in TiO2-BT (without acid wash) and TiO2-BT-TiO2 (without acid wash) NPs are ~23 and ~8%, respectively. The comparatively smaller content in TiO2-BT-TiO2 (without acid wash) NPs suggests that some amount of BaCO3 might be removed during TiO2 shell synthesis. Thus, XRD, TEM as well as HRTEM images confirm successful synthesis of BT and TiO2 shell on the TiO2 NPs. Also, XRD spectra of PVDF film and TiO2-BT-TiO2@dopa/PVDF confirm that PVDF is present in α-phase (2θ = 17.8°, 18.3° and 20.1° referred to the planes (100), (020), (110), respectively) and the phase remains same after addition of nanofillers in the polymer nanocomposites (Figure 2b). Additional peaks in TiO2-BT-TiO2@dopa/PVDF are indexed to indium tin oxide (ITO) electrode coated on glass substrate, as confirmed from ICDD file number 44-1087. The surface modification of different nanomaterials was examined using FTIR spectroscopy (Figure S4). The Figure S4a shows the FTIR spectra of TiO2, TiO2@dopa, TiO2-BT, TiO2-BT@dopa, TiO2-BT-TiO2, and TiO2-BT-TiO2@dopa NPs. In all the spectra, the peak at 3440 cm-1 is associated with the vibration of surface hydroxyl or the adsorbed water while the OH bending vibration peak at 1635 cm-1 corresponds to chemisorbed water. Other peaks at 1400 cm-1 as well as in the range of 394 - 984 cm-1 correspond to Ti–O stretching peaks.58-60 A careful investigation of TiO2-BT (without acid wash) and TiO2-BT-TiO2 (without 8 ACS Paragon Plus Environment

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acid wash) spectra reveals the presence of new peaks of BaCO3 which are absent in acid washed samples as shown in Figure S4b. The peaks at 857 and 693 cm-1 are attributed to the in-plane and out-of-plane bending of CO32-. The sharp and significant peak at 1451 cm-1 represent asymmetric stretching mode of C-O mode while one at 1059 cm-1 corresponds to symmetric C-O stretching vibrations.61 The similar observation of BaCO3 removal in TiO2-BT and TiO2BT-TiO2 NPs on acid treatment was confirmed by XRD (Figure S3). It is also evident from Figure S4a that there is a broad -OH peak present in the range of 3700-3000 cm-1, which is helpful in forming covalent bond with dopa. Lin et al. investigated that -OH group of dopamine hydrochloride (DP) is expected to bind with the nanomaterial surfaces.62 Further, the peaks in this region of 3700-3000 cm-1 becomes broader which suggest that the amount of surface hydroxyl or the adsorbed water on the nanomaterials surface further increased with dopa coating. Thus, the presence of -NH2 groups of dopa as well as -OH groups of adsorbed water forms the interface regions which is capable of strong coupling with C-F bond of PVDF (−C−F···H−N− or −C−F···H−O−) and has significant influence on the dielectric performance.50,63 Another important observation from the FTIR spectra (Figure S4a) can be seen that BaCO3 peaks are removed during dopa coating as in TiO2-BT@dopa as well as TiO2-BTTiO2@dopa NPs. Here, the DP used as precursor simultaneously helps in removing BaCO3 from the TiO2-BT (without acid wash) as well as TiO2-BT-TiO2 (without acid wash) NPs as well as forms dopa coating. The possible reason for BaCO3 removal during dopa coating is the presence of HCl in DP precursor which when dissolved in solution formed BaCl2 and CO2 as previously shown in reaction 6. After dopa functionalization, the colour of the nanomaterials changed from white to deep dark (the digital images of TiO2-BT-TiO2 and TiO2-BTTiO2@dopa NPs is shown in the Figure S5). The dopa functionalization is characterized by stretching vibrations of C–C/C=C aromatic ring at 1623 and 1490 cm-1, C–OH at 1344, 1305,



and 1262 cm-1, C–H at 2925 cm-1, stretching (3700-3200 cm-1) and bending (1618 cm-1) vibrations of N–H. The peak at around 3400 cm-1 is assigned to stretching vibration of –OH group of catechol, and the bending vibrations of C–OH at 1397, 1155, and 1127 cm-1. Also, the strong peak at 983 cm-1 corresponds to C–C–H bending of dopa.15,58,64-67 The peaks at 1066 and 1122 cm-1 are attributed to the –OH groups while the one at 1635 cm-1 is assigned to the bending mode of absorbed H–O–H molecules.68,69 Thus, FTIR confirms the successful attachment of dopa on TiO2, TiO2-BT, and TiO2–BT–TiO2 NPs surface. The BT shell is specifically designed to serve as dielectric medium as previously reported by Yang et al.70 It was studied that the hydrothermal reaction temperature (usually > 150 oC) plays a key role in local conversion of TiO2 NPs surface and growth on to BT NPs. The conversion of TiO2 as well as growth of BT is significantly accelerated at higher temperatures (170 to 210 oC) resulting in generation of large number of BT NPs on the TiO2 surface which led to formation of uniform BT shell. Thus, the BT shell thickness in TiO2-BT core-shell heterostructures can be tuned with the variation in reaction temperatures. Similar was also observed by Jang and Yang.71 Hence, we have carried out the reaction at 180 oC for 2 h to form a smooth surface of BT shell on the TiO2 NPs surface. The second layer of uniform TiO2 shell was synthesized by sol-gel approach using titanium (IV) butoxide (TBOT) and ammonia (NH3) as precursors.72-75 Li et al. have reported a simple, facile and kineticallycontrolled approach to construct a uniform TiO2 shell of desired thickness.72 In brief, the heterogeneous nucleation and growth of TiO2 shell is governed by a two-step reaction. Firstly, the amorphous TiO2 nucleation and growth is proceeded by the plentiful formation and diffusion of polymerized TiO2 networks or titanium oligomers through hydrolysis and condensation of TBOT (reaction 7) which is controlled by the amount of NH3 present in the reaction system.


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(7) With the increase in shell thickness, the shell growth rate slows down due to the hinderance in diffusion of large titanium oligomers to the nucleation sites and is controlled or limited by mass transfer or diffusion. However, excess NH3 content leads to further hydrolysis and condensation of TBOT, which results in formation of undesired isolated TiO2 NPs by homogeneous nucleation. Thus, TiO2 shell thickness can be controlled by NH3 content and the proper tuning of the reaction kinetics, which only allows heterogeneous nucleation and growth of uniform TiO2 shell on the TiO2-BT NPs surface and avoids homogeneous nucleation of titanium oligomers to form undesired TiO2 NPs in the solution. Also, the amorphous phase of TiO2 can easily be transformed to crystalline phase without disrupting the uniform TiO2 shell nanostructures. Thus, shell thickness of BT and TiO2 layers can be easily tuned via appropriate reaction conditions for desired applications. The dielectric properties of TiO2@dopa/PVDF polymer nanocomposites were investigated at different loadings and 10 vol% loading was selected based on the optimum results for further studies (Figure S6). The influence of TiO2 NPs based nanofillers (loading: 10 vol%) on the dielectric properties of PVDF based capacitors was carefully investigated. The Figure 3 shows that the addition of TiO2 NPs based fillers enhances the dielectric constant as compared to pure PVDF. The dielectric constant is increased from 9.9 for PVDF to 11.1 and 11.6 for TiO2/PVDF and TiO2@dopa/PVDF, respectively at 1 kHz. The high dielectric constant of TiO2 (≈ 40) helps in improving the dielectric constant of the polymer nanocomposites. Further, the incorporation of BT layer on the TiO2 NPs (TiO2-BT NPs) increased the dielectric constant to 11.9 at 1 kHz due to the high-k BT (> 200) and interfacial polarization at the NPs interfaces in TiO2-BT@dopa/PVDF. The maximum dielectric constant 11 ACS Paragon Plus Environment


of 12.6 at 1 kHz for TiO2-BT-TiO2@dopa/PVDF is associated with the formation of TiO2-BTTiO2 as an individual capacitor as well as improved interfacial polarization and interface modulation.

Figure 3. Dielectric constant (solid) and loss tangent (tan δ) (hollow) variation with frequency for pristine PVDF and different 10 vol% TiO2 based nanofillers/PVDF capacitors. Also, the dielectric constant for all the devices follows a decreasing trend with frequency, which is expected as dipoles fail to rotate at higher frequencies which further lowers down the polarization and hence dielectric constant. The corresponding loss tangent of PVDF, TiO2/PVDF, TiO2@dopa/PVDF, TiO2-BT@dopa/PVDF and TiO2-BT-TiO2@dopa/PVDF at 1 kHz were 0.03, 0.05, 0.04, 0.06 and 0.05, respectively. A closer look into values of loss tangent reveals that at lower frequencies, pure PVDF has the lowest loss tangent which is due to the lower conduction losses as well as segmental motion of the amorphous phase, while the 12 ACS Paragon Plus Environment

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addition of TiO2 NPs increases the loss tangent showing that additional free charges carriers are incorporated in the PVDF matrix. However, insulating dopa layer in TiO2@dopa/PVDF helps in reducing loss tangent due to the strong interface adhesion as compared to that in TiO2/PVDF which suppresses charge carrier movement. The BT layer in TiO2BT@dopa/PVDF slightly increased loss tangent than that of TiO2@dopa/PVDF which might be due to increased interfacial polarization and defect formation responsible for charge migration. The loss tangent of TiO2-BT-TiO2@dopa/PVDF was decreased as compared to TiO2-BT@dopa/PVDF which shows that the additional TiO2 layer helps in free charge mitigation and improved interfacial polarization and thus, improved the dielectric performance of the capacitors. Also, at higher frequency range, only PVDF has maximum loss tangent which further suppressed with the nanofillers addition due to restricted segmental motion of the polymer chains as well as fewer molecular dipoles.1,15,24 The breakdown failure of the devices was studied by using two-parameter Weibull distribution function.11 Briefly, the electric failure can be expressed in terms of cumulative probability P as given by equation 8,

  E   P = 1 − exp −      Eb    

(8)

where, E is the experimental breakdown strength, Eb is a scale parameter known as characteristic breakdown strength and represents the cumulative failure probability of 63.2%, β is the shape parameter which indicates the scatter of the breakdown strength data, a higher β value indicates less scattered experimental data. Figure 4a and b represents the Weibull analysis and Eb of pristine PVDF and different TiO2 nanofillers filled in PVDF based capacitors.



Figure 4. (a) Weibull distribution, (b) characteristic breakdown strength (Eb) of different NPs filled polymer nanocomposites (nanofiller loading: 10 vol%). (c) Different breakdown paths (path A and B) as well as movement of charges (path C) in the polymer nanocomposites, and (d) schematic distribution of electric potential under the applied electric field. For pristine PVDF, Eb and β were 2837 kV cm-1 and 20.3, respectively which suffers a severe decrease with TiO2 addition in TiO2/PVDF to 1966 kV cm-1 and 18.1, respectively. The comparatively lower β in TiO2/PVDF suggests the formation of microdefects in the film which was absent in pristine PVDF. However, dopa coating improved the TiO2@dopa NPs interface as well as adhesion with the PVDF polymer matrix resulting in enhanced Eb (2164 kV cm-1) and β (35.2). Eb was further increased to 2631 kV cm-1 for TiO2-BT@dopa/PVDF which was due to the reduction in the electric percolation pathways associated with the mitigation of 14 ACS Paragon Plus Environment

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charges in the nanocomposites.1 The sudden decrease in the β (8.3) value shows that more defects might be introduced in the polymer matrix. In TiO2-BT-TiO2@dopa/PVDF polymer nanocomposites, Eb showed an increment to 3128 kV cm-1, the maximum among all the devices, with β value of 26.4. The high β shows that the film quality was improved with core@multishell NPs and interface modulation helps in enhancing the Eb.15 Further, leakage current density vs electric field (J vs E) measurements were carried out to clearly understand the breakdown behaviour of the samples and the results are shown in Figure S7. The leakage current was found to be maximum for TiO2/PVDF and further decreased with dopa coating. Thus, dopa coating helps in minimizing the leakage current density and improving the breakdown strength.76 FTIR data confirmed earlier that peaks of -OH as well as -NH groups of dopa enhanced the interaction as well as compatibility of the functionalized nanofillers with the PVDF matrix and suppressed the inhomogeneity at the interface and led to improved dielectric performance. It can also be observed from Figure S8 that the dispersion of the NPs after dopa modification is improved while TiO2 NPs without dopa modification forms agglomerates (Figures S8 c, d), thereby deteriorates the dielectric performance of TiO2/PVDF as compared to other devices. The current density of different polymer nanocomposites follows the order: TiO2/PVDF > TiO2@dopa/PVDF > PVDF > TiO2-BT@dopa/PVDF > TiO2-BTTiO2@dopa/PVDF which is opposite to that of the breakdown strength. Also, log (J) at higher fields is limited by space–charge-limited conduction mechanism which is minimum for TiO2BT-TiO2@dopa/PVDF,77 suggesting that the free charges get trapped in the enclosed interfacial zone as well as vicinity of the NPs interface and enhances the breakdown strength.1 The enhancement in the Eb of different surface modified TiO2 nanofillers can be further understood in another way. In a polymer nanocomposite system, the breakdown strength is dependent on the charge transportation through nanofillers as well as polymer matrix. On the application of electric field, two phenomena occur in the polymer nanocomposites. Firstly, a



charged interfacial region is generated on each layer of the multishell nanostructures as well as in the surrounding polymer matrix, both collectively responsible for internal micro-electric field at the interfaces which increases with the electric field and hence forms micro-capacitors. The direction of applied electric field opposite to the internal micro-electric field significantly nullifies the electric field concentration thereby making the electric field relatively more homogeneous in the nanocomposites. Secondly, electron injection from the electrodes as well as de-trapping from the composites result in generation of more space charges and carriers in the device (path A in Figure 4c). Some of the space charges get trapped in the nanocomposites. However, remaining large number of space charges prefer to move through the interface region of the nanofillers due to relatively high charge densities near the interface (path B in Figure 4c), thus, responsible for comparatively lower resistance as compared to that of polymer matrix. With more space charges accumulating near the interface, the charge densities and hence interface thickness (or interfacial volume) further increases which restricts the movement of space charges at the nanofiller interface (path C in Figure 4c). Thus, large number of accumulated interfacial charges due to the formation of internal micro-capacitors increases the surface potential (ψo) of the nanofillers to ψ’o as shown in Figure 4d. As a consequence, the mechanical stress in interfacial region and hence interfacial thickness (t) also increases to Δt which is beneficial for restricting more charges in that region (path C in Figure 4c). Thus, formation of internal-micro-capacitors and restriction of space charge movement through the nanofiller interfaces are responsible for improved breakdown strength.24 A detailed explanation for improved dielectric performance of TiO2-BT-TiO2@dopa/PVDF polymer nanocomposites over other devices is given in the next sub-section. The polarization test was performed to calculate the discharged energy density (Ud) of the devices (equation 1). The energy density values of different devices were calculated from P-E loop as shown in Figure 5a. The Figure 5b represents P-E loops of different devices at their


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Eb while at different electric fields is shown in Figure S9. The calculated energy density values at different electric fields are summarized in Figure 5c.

Figure 5. (a) Schematic representation of P-E loop, (b) P-E loops of devices at Eb. (c) Energy density and (d) efficiency of different NPs filled polymer nanocomposites (nanofiller loading: 10 vol%). Energy density of the polymer nanocomposite is increased with the addition of nanofillers as compared to pristine PVDF. The addition of high-εr nanofillers in the PVDF polymer matrix results in increase in the energy densities similar to the dielectric constant (Figure 3): PVDF < TiO2/PVDF < TiO2@dopa/PVDF < TiO2-BT@dopa/PVDF < TiO2-BTTiO2@dopa/PVDF. The maximum energy density of 4.4 J cm-3 at Eb was obtained for TiO2BT-TiO2@dopa/PVDF which was 83.3% higher than that of pristine PVDF. The dielectric 17 ACS Paragon Plus Environment


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performance is also evaluated by an important criterion termed as energy efficiency, η given by equation 9:

=

Ud U d + Ul

(9)

where Ul is the loss energy density (Figure 5a). The variation of efficiency with electric field of different devices is summarized in Figure 5d. The efficiency of TiO2-BT-TiO2@dopa/PVDF is maximum among all the polymer nanocomposite-based capacitors except for the pristine PVDF. Noticeably, the efficiency follows a decreasing trend with electric field. The lower efficiency of TiO2/PVDF, TiO2@dopa/PVDF, TiO2-BT@dopa/PVDF as compared to TiO2BT-TiO2@dopa/PVDF is attributed to higher leakage current density or conduction loss and possible defects formation responsible for space charge migration in the films.33 The advantage of TiO2-BT-TiO2 structure is the presence of more number of interfaces than other nanofillers which constraints the charge conduction pathways and reduces the leakage current and thus, helping in enhancing the efficiency. Further, to study the effect of BaCO3 in as-synthesized TiO2-BT (without acid wash) NPs on the dielectric performance, the dielectric results are compared with acid washed TiO2BT NPs as shown in Figure S10. The comparative dielectric study reveals that BaCO3 is undesirable for improved dielectric performance and should be removed from NPs before device fabrication. The improved dielectric performance of TiO2-BT-TiO2@dopa/PVDF is explained as follows.


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Role of the novel TiO2-BT-TiO2@dopa tri-layered architecture nanofiller in improved dielectric performance The novel tri-layered architecture of the nanofillers plays a crucial role in enhancing the dielectric performance of the capacitors as explained in Figure 6. In order to obtain high energy density of the polymer nanocomposites, it is necessary to improve the dielectric constant (or polarization) as well as breakdown strength (equations 1 and 2).

Figure 6. Schematic illustration of the TiO2-BT-TiO2@dopa/PVDF device (left) and enlarged view of the nanomaterial interface (right) under the application of electric field.

The specially designed TiO2-BT-TiO2 NPs helps in improving the polarization as follows. The TiO2-BT-TiO2 NPs is supposed to work as an individual capacitor with TiO2 (both core and outer shell) and BT serve as capacitor plates and dielectric medium, respectively. The reason for selecting TiO2 as capacitor plates is that TiO2 has comparatively higher electrical conductivity of 10-4 S m-1 as compared to BT (10-10 S m-1) while BT has higher dielectric constant (> 200) than TiO2 (≈ 40).1,39 As starting TiO2 nanomaterial has very small size (21 nm), the number of TiO2-BT-TiO2-dopa NPs and hence the number of capacitors per unit volume also increases which leads to the enhancement in the dielectric constant as well as energy density.12 Also, modulating the interfacial polarization at the interfaces inside the NPs as well as space charge accumulation at the NPs interface play a key role in enhancing the 19 ACS Paragon Plus Environment


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dielectric properties. As TiO2 is an n-type semiconductor and Fermi level difference between TiO2 and BT is > 0.5 eV, this results in the accumulation of positively and negatively space charges on the TiO2 and BT. The surface charge densities (σ) in different layers of TiO2-BTTiO2 NPs follows the trend as σ(TiO2)core > σ(BT)middle

layer

> σ(TiO2)outer

layer

in order to

maintain the charge neutrality during electrostatic induction. The positive charges in TiO2 layer in turn develop negative charges on the dopa layer due to polar interaction. The interfacial charges present in the polymer matrix form the Gouy–Chapman–Stern layer at the NPs interface which enhances the interfacial polarization at the interface. Usually, the thickness of Stern layer and diffused layer is ~ 1 and ~ 2-9 nm.3,26,29,39,78,79 This NPs architecture also enhances the breakdown strength in the manner as follows. The BT, TiO2, and dopa layers are also considered to play the role of polarization layer, buffer layer, and dispersion layer, respectively. The BT polarization layer provides high dielectric constant to the NPs while TiO2 buffer layer has moderate dielectric constant in between the BT and dopa layer which helps in suppressing the local electric field at the interface. The dopa layer helps in homogeneous dispersion of NPs in the PVDF polymer matrix. NH2 groups of dopa and C-F groups of PVDF form H-bond which enhances the compatibility of the NPs with the polymer matrix.62 The leakage current density is minimum for the TiO2-BT-TiO2/PVDF polymer nanocomposites (Figure S7), which shows that the conduction mechanism is due to space–charge-limited conduction,77 in which interface entraps the space charges and restricts the formation of conduction channels. Also, TiO2-BT-TiO2 NPs has two different interfacial zones (TiO2-BT and BT-TiO2) which form an enclosed interfacial architecture. This restricts the migration of charges and improves the dielectric performance as compared to that in TiO 2 NPs based device in which there is no enclosed interfacial zone to restrict the charge migration while TiO2-BT has only one interface zone. The similar behaviour was observed by Kang et al.1 They showed that the breakdown of TiO2 nanowires was comparatively lower than that of


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TiO2-BT nanowires in poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP) based polymer nanocomposites which was due to the restricted migration of charges in the enclosed interfacial zone of TiO2-BT as compared to the bare TiO2 nanowires. Thus, the interfacial engineered TiO2 NPs improved the overall dielectric performance of the dielectric capacitors. The present work shows improved dielectric properties due to the interface tuning as compared to earlier reported works (Figure S11). We expect that this work will augment the exploration of new pathways via design of nanofillers with novel architectures for highperformance polymer nanocomposite-based capacitors.

◼ CONCLUSIONS In conclusion, a novel hierarchical architecture of core@multishell TiO2-BT-TiO2@dopa NPs as nanofillers were designed to study the effect of interfaces modulation on the dielectric properties of the polymer nanocomposite-based capacitors. The results showed significantly improved dielectric performances as compared to other TiO2 nanofillers and PVDF based capacitors. The maximum energy density of 4.4 J cm-3 was obtained for TiO2-BTTiO2@dopa/PVDF which was 83.3% higher than that of pure PVDF. Also, the device showed maximum dielectric constant, breakdown strength as compared to other devices. The results show that each TiO2-BT-TiO2 NPs works as an individual capacitor with improved interface surrounding the polymer matrix resulting in improved dielectric performance as compared to PVDF, TiO2/PVDF, TiO2@dopa/PVDF, TiO2-BT@dopa/PVDF. Apart from the increase in dielectric constant or polarization, the breakdown strength also improved due to formation of enclosed interfacial zones (TiO2-BT and BT-TiO2) as well as reduced leakage current which help in charge entrapment in both NPs as well as polymer matrix. This work reveals for the first time that the newly designed core@multishell nanofillers could function as an individual capacitor and their interfacial modulation improves polarization as well as breakdown strength.



Further, this work provides a promising and novel approach in maximizing the dielectric properties of the polymer nanocomposite-based capacitors.

◼ EXPERIMENTAL METHODS Materials. Titanium (IV) oxide (TiO2, 21 nm, ≥ 99.5%), barium hydroxide octahydrate (Ba(OH)2.8H2O, ≥ 98%), tetrabutylammonium hydroxide solution (TBAH, 40% in water), titanium (IV) butoxide (TBOT, reagent grade, 97%), dopamine hydrochloride (DP), tris(hydroxymethyl)aminomethane (ACS regent, ≥ 99.8%), and polyvinylidene fluoride (PVDF, MW = ~534 000 by GPC) were supplied from Aldrich. Ethanol (absolute for analysis), 2-propanol (≥ 99.0%), ammonium peroxodisulfate (AP), and N, N-dimethylformamide (DMF) were purchased from Merck. Diethylene glycol (98.5%) was obtained from Loba Chemie Pvt. Ltd. Ammonia (NH3, 25%) was supplied from Fisher Scientific. ITO coated glass substrate (15 Ω) was purchased from Lumtec, Taiwan. Synthesis of Nanofillers. TiO2 NPs were decorated with BT using the procedure reported earlier.1,70 Firstly, 0.48 g (6 mmol) TiO2 NPs were well-dispersed in 5 mL ethanol through sonicating for 1 h. Another solution was prepared by mixing 0.95 g (3 mmol) Ba(OH)2.8H2O in 10 mL of diethylene glycol, 5 mL of ethanol, 3 mL of 2-propanol, 1.2 g of TBAH, and 14 mL of deionized (DI) water. Both the solutions were mixed and stirred for 2 h and transferred in 50 mL autoclave. The reaction was carried out at 180 oC for 2 h. After the completion of reaction, cooled product was centrifuged, washed with DI water and ethanol repeatedly at 10000 rpm for 10 min to remove any impurities as well as unreacted excess reagents. The product was dried at 60 oC in hot air oven overnight. To further remove the BaCO3 formed inherently during the reaction as impurities, the as-synthesized TiO2-BT NPs were further washed with 0.1 M HCl (three times) and DI water (three times) and were dried under vacuum.57 In order to encapsulate TiO2-BT NPs with TiO2 shell,73 50 mg of TiO2-BT NPs were


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dispersed in 80 mL ethanol through sonicating for 1 h. Further, 0.625 mL of NH3 was added and stirred for 15 min. Subsequently, a solution of 0.73 mL TBOT in 20 mL ethanol was added dropwise to the solution and the reaction mixture was vigorously stirred at 45 oC for 10 h. The final product was washed with ethanol several times through centrifuging at 10000 rpm for 10 min. The product (TiO2-BT-TiO2) was dried in oven at 60 oC overnight. To further improve the crystallinity of the TiO2 layer, the product was annealed at 400 oC for 5 h in air.80 Lastly, the functionalization of the nanomaterials with dopa was carried out as per literature.81,82 Firstly, 90 mg (0.475 mmol) DP was dissolved in 45 mL DI water followed by addition of 60.57 mg (10 mM) of tris(hydroxymethyl)aminomethane. The pH of the resulting solution was 8.5. A homogenous solution of 146 mg of each of the nanomaterials (TiO2, TiO2-BT and TiO2BT-TiO2) in 5 mL DI water was prepared by sonicating for 30 min. Both the solutions were mixed and 54.2 mg (0.2375 mmol) of AP was added and stirred for 10 min (AP:DP molar ratio = 1:2). The reaction was carried out under mild stirring for 6 h in dark condition. The reaction product was centrifuged washed repeatedly with DI water at 10000 rpm for 10 min and dried in oven at 60 oC overnight (Scheme 1). Fabrication of Devices. The devices were fabricated as follows. Firstly, nanofillers of desired loadings were homogeneously dispersed in DMF through sonicating for 1 h. Then, PVDF of concentration 0.1 g mL-1 was added and the final mixture was kept for vigorous stirring at 60 oC for 24 h. The solution was spin-coated on ITO coated glass substrate at 1000 rpm for 1 min. The film was dried at 60 oC for 4 h on hot-plate followed by vacuum-drying for 24 h. The film was annealed at 180 oC for 10 min. The resulting film thickness was ~1-3 µm. The aluminium (Al) electrodes were deposited on the top side of the film using a mask of 1 mm eyelets which served as the top electrode while the ITO was used as bottom electrode. Characterization. The thermogravimetric analysis (TGA) was performed using SDT Q600 in the following conditions: temperature range = 25 oC – 800 oC, heating rate = 10 oC min-1,



carrier gas = air, sample weight = 8–10 mg. The chemical structures of the nanomaterials were examined using PerkinElmer Fourier transform infrared (FTIR) spectrometer in the range of 400 - 4000 cm-1 at ambient conditions using KBr pellet method. The crystal structures of the nanofillers were determined using an X-ray diffraction (XRD) pattern obtained from Two Circle Diffractometer, Rigaku, MiniFlex600 (angle range = 10o to 90o, step size = 2o min-1, radiation = Cu Kα, λ = 1.5406 Å). The surface morphologies were imaged by field emission scanning electron microscope (FESEM, Mira3, Tescan). The morphologies of the nanomaterials were observed using transmission electron microscope (TEM, Tecnai 20G2) and high-resolution transmission electron microscope (HRTEM, FEI Titan G2 60 − 300). The thickness of the films was measured using DektakXT, Bruker surface profilometer. The dielectric constant and loss tangent were measured using E4980AL LCR meter at room temperature (frequency range = 102 – 106 Hz, voltage = 1 V). The P-E hysteresis loops for energy density, efficiency measurements as well as breakdown strengths and leakage currents were characterized using Precision 4 kV HVI test system (Radiant Technologies, Inc.) at 10 Hz.

◼ SUPPORTING INFORMATION Supporting Information is available from ACS Publications website free of charge.

◼ ACKNOWLEDGEMENTS Financial support from Department of Atomic Energy (DAE), BRNS, India for grant 34/14/14/2014−BRNS and Department of Science and Technology (DST) Grant No. DST/TMD/CERI/C140(G) under Clean Energy Research Initiative is acknowledged. RKG acknowledges financial assistance from DST, India, through the INSPIRE Faculty Award (Project No. IFA-13 ENG-57).


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◼ CONFLICT OF INTEREST The authors declare no conflict of interest.

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Significantly Enhanced Energy Density by Tailoring the Interface in

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