Dramatically improved piezoelectric properties of poly(vinylidene fluoride) composites by incorporating aligned TiO2@MWCNTs

Composites Science and Technology 123 (2016) 259e267

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Composites Science and Technology journal homepage: http://www.elsevier.com/locate/compscitech

Dramatically improved piezoelectric properties of poly(vinylidene fluoride) composites by incorporating aligned TiO2@MWCNTs Lu Yang a, b, Hongli Ji a, c, Kongjun Zhu a, c, Jing Wang a, c, Jinhao Qiu a, c, * a

State Key Laboratory of Mechanics and Control of Mechanical Structures, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, China College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, China c College of Aerospace Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, China b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 12 June 2015 Received in revised form 16 November 2015 Accepted 30 November 2015 Available online 8 December 2015

Composites comprising TiO2@MWCNTs (shell@core) nanoparticles and poly(vinylidene fluoride) (PVDF) were prepared through a two-step method, including solution cast and mechanical rolling, inducing a highly oriented structure with both PVDF lamellar and TiO2@MWCNTs. A systematic study was performed to investigate the crystalline structure and piezoelectric properties of composites. The obtained results reveal that greatly enhanced breakdown strength and piezoelectric properties can be achieved in composites by the incorporation of highly aligned TiO2@MWCNTs. A maximum piezoelectric coefficient d33 of ~41 pC/N can be achieved in composites at a loading level of 0.3wt.%, which is nearly double that of the pure PVDF. © 2016 Elsevier Ltd. All rights reserved.

1. Introduction Polyvinylidene fluoride (PVDF) and its copolymer have been widely used in the fields of sensing and actuating devices due to their excellent piezoelectric and ferroelectric properties among the polymeric materials [1,2]. PVDF is a linear fluorinated hydrocarbon with a repeat unit (CH2eCF2), which is semi-crystaline and exhibits two dominant crystalline forms, namely, a-phase and b-phase. The a-phase with a TGTG’ (T-trans, G-gaucheþ, G’-gauche-) dihedral conformation arrangement in a centrosymmetric unit cell, resulting in a nonpolar crystal structure, while the b-phase exhibits all TTTT conformation, giving rise to a polar, non centrosymmetric crystal with a large dipole when electrically poled [3,4]. According to previous investigations [5,6], the remarkable piezoelectricity in PVDF is attributed to high contents of b-phase, thus most of the work has been directed toward obtaining high contents of b-phase to obtain a good piezoelectric and ferroelectric material. As reported [7,8], the crystal phases obtained in PVDF are dependent on the processing history, including casting solvent, thermal history, mechanical history and electric field history. The bphase can be obtained directly by melt crystallization under high pressure or by crystallization from solution at temperatures below

* Corresponding author. College of Aerospace Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, China. E-mail address: [email protected] (J. Qiu). http://dx.doi.org/10.1016/j.compscitech.2015.11.032 0266-3538/© 2016 Elsevier Ltd. All rights reserved.

70  C. Moreover, the b-phase can be obtained by phase transformation from a-phase under the application of strong electric field or mechanical stretching. Recently, incorporating particles (e.g., clay, carbon nanotubes, ferrite nanoparticles) into PVDF has been proven to be one of the facile ways to induce the b-phase formation [9e12]. Due to the large aspect ratio, remarkable mechanical, electrical, and thermal properties of multi-walled carbon nanotubes (MWCNTs), the compounding of MWCNTs into PVDF is especially attractive [13]. By introducing a small amount of MWCNTs into PVDF matrix, one can realize enhancements in both mechanical and piezoelectric properties, which enables the use of resulting composites as actuators and sensors [14]. However, one issue of the final MWCNTs/PVDF composites is the sharp decrease in breakdown strength because of the presence of conductive MWCNTs [15]. On the basis of fact that high poling electric field is critical to the piezoelectric performance of PVDF, the enhancement in piezoelectric properties of composites is largely limited as they cannot be poled under sufficiently high electric fields due to the diminished breakdown strength. Therefore, the fabrication of composites containing MWCNTs with high breakdown strength is still a challenge. One convenient way to improve the breakdown strength of composites containing MWCNTs is to align them in directions perpendicular to the applied electric field [16]. Up to now, several techniques including electrospinning [17,18], high speed melt-

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spinning [19], injection molding [20], stretching [21] and roling [22] have been adopted to realize the alignment of MWCNTs in polymer matrix. Another route to enhance the breakdown strength of involves in utilizing insulative/semi-conductive shells outside the MWCNTs [23,24]. By the employment of insulative/semi-conductive shell such as TiO2 [25], the tunneling current of composites at high fields can be suppressed to some extent and thus enhance the breakdown strength of composite. Previously, we reported the breakdown strength of solution cast PVDF can be enhanced by incorporating TiO2 coated MWCNTs (TiO2@MWCNTs), i.e. the room temperature breakdown strength increase from ~170 V/mm for pure PVDF with a thickness of ~30 mme~210 V/mm for composites with a loading of 0.3wt.% [25]. However, even with enhanced breakdown strength, the obtained solution cast composites cannot be poled under electric field above 100 V/m with sufficient polarization time (>1 h), largely limiting the achievable piezoelectric performance. In fact, due to the limited poling electric field, the obtained maximum piezoelectric coefficient d33 in the solution cast composites is only ~14 pC/N. In this work, we will report a combined approach to dramatically raise the breakdown strength and piezoelectric properties of PVDF composites by using TiO2@MWCNTs and a mechanical rolling process. Composites comprising TiO2@MWCNTs (shell@core) nanoparticles and PVDF were prepared by a solution cast method, followed by a rolling process, with an aim to align TiO2@MWCNTs. A comprehensive investigation was performed to study the crystalline structure, breakdown strength and piezoelectric properties of composites comprising PVDF and aligned TiO2@MWCNTs. The results demonstrate that the significantly improved piezoelectric properties can be induced by the existence of aligned TiO2@MWCNTs, indicating the promise of using TiO2@MWCNTs in sensing and actuating devices. 2. Experimental 2.1. Materials Pristine MWCNTs (diameter: 40e60 nm, length: 5e15 mm) were purchased from Shenzhen Nanotech Port Co., China. The chemicals including tetrabutyl-orthotitanate (TBOT), absolute ethanol and nitric acid (HNO3) were purchased from China National Chemicals Corporation Ltd. PVDF powders were supplied by 3F Co. Ltd., Shanghai, China, under the trade name PR904. N, N-dimethylformamide (DMF; J T Baker, Inc., USA; 9222-01 PHOTREX_ Reagent) was used as solvent in the preparation of composite. 2.2. Preparation of TiO2@MWCNTs and TiO2@MWCNTs/PVDF composites The detailed synthesis of TiO2@MWCNTs was described in Ref. [25]. For composite preparation, a desired amount of TiO2@MWCNTs was added into DMF and sonicated for 1.5 h. Meanwhile, PVDF powder was added to DMF in a Teflon mechanical stirrer with ultrasound for 1.5 h for complete dissolution of the polymer. Then TiO2@MWCNTs suspension was slowly added into PVDF solution. After sonicating and vacuum mixing, the mixture was cast on a clean glass substrate, the solvent evaporated and the PVDF crystallized at 80  C. When the films were formed, they were placed in a vacuum oven for further heat treatment at 80  C for 16 he24 h. The thickness of solution cast composites was 60 ± 4 mm. The composites fabricated by solution casting were cut into 5 cm  5 cm pieces and rolled with a two-roller rolling machine. Reduction in thickness was accomplished in two passes through the rollers. The rolling temperatures was 50  C and the

rolling speed of the two rollers for each pass was set at 30 and 27 r/ min, respectively. The final thickness of rolled composites was 20 ± 2 mm. For comparison, solution cast composites with a thickness of 20 ± 3 mm were also fabricated. Thereafter, solution cast composites and rolled composites will designated as S-X and R-X, respectively, where X represents (0, 0.1, 0.3, 0.5, 0.7, 1) the weight fraction of MWCNTs with respect to the PVDF matrix. Poling was performed in a silicon oil bath with high-voltage DC supply at 70  C for 1.2 h. For the poling process, gold electrodes (area of 1 cm  1 cm) were vacuum-evaporated onto both surfaces of the films. 2.3. Characterization The distribution of TiO2@MWCNTs in composites was visualized by transmission electron microscopy (TEM; JEOL JEM 2100) operated at an accelerated voltage of 200 kV. Fourier transform infrared spectroscopy (FTIR; Nicolet 6700, Thermoscientific, USA) was employed to record the wave numbers from 400 cm1e1500 cm1. X-ray diffraction (XRD; Bruker D8 Advance System, Germany) was performed at room temperature with Cu-target Ka radiation (l ¼ 0.154 nm). Small-angle X-ray scattering (SAXS) measurements were carried out on a pinhole NanoSTAR SAXS camera (Bruker AXS, Germany), using Cu Ka radiation (l ¼ 0.154 nm) and a HiStar 2D detector. The X-ray beam was perpendicular to the surface of samples and the exposure time was 1 h for each measurement. The Lorentz corrections were performed after subtraction of the air scattering. Crystallinity and melting point of the samples were determined through differential scanning calorimetry (DSC; DSC7020, Japan) at 40  Ce200  C at a heating rate of 10  C/min. The dielectric properties of the samples were measured at room temperature using an Impedance Analyzer (HP4294, Agilent, USA). The electric breakdown strength was tested by a dielectric withstand voltage test (Beijing Electromechanical Research Institute Supesvoltage Technique) at a ramping rate of 200 V s1 and a limiting current of 5 mA. Polarization-electric field (P-E) hysteresis loops of samples were obtained with a ferroelectric testing system (RT66A, Radiant Technologies, USA) connected with a high voltage interface (Trek 609B, Trek, USA). The piezoelectric strain coefficients (d33) were measured at ~20  C using a piezo d33 m (ZJ-3A, Institute of Acoustics, Chinese Academy of Sciences, China). 3. Results and discussion 3.1. Morphology characterization of composites As already demonstrated in our previous works [22], rolling can induce a highly oriented structure of samples due to the provided shear force. Shown as examples in Fig. 1a are the 2D-SAXS patterns for rolled pure PVDF(inset) and composites with 0.7wt.% loadings (R-0.7), the corresponding pattern for solution cast pure PVDF and composite with 0.7wt.% loadings (S-0.7) are also included for comparison. For the solution cast samples, a well-defined isotropic ring can be observed, implying the presence of well-developed and randomly distributed lamellae. However, in the case of rolled samples, two blob-like reflections on the meridian can be observed, indicating an oriented lamellar structure perpendicular to the rolling direction. Remarkably, the similar SAXS patterns in rolled pure PVDF and composite is clear manifestation of the similar orientation degree of lamellar structure in these rolled samples. In an effort to examine the dispersion and orientation of TiO2@MWCNTs in composites, TEM was adopted. Presented in Fig. 1c and d are the TEM images of samples containing 0.7 wt.% loadings. To better identify the nanofillers in PVDF matrix, the TEM

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Fig. 1. SAXS patterns of solution cast (a) and rolled (b) pure PVDF (inset) and composites with 0.7 wt.% loadings; TEM images of solution cast (c) and rolled (d) composites containing 0.7wt.% loadings, wherein the arrow points to the rolling direction. The inset in (c) represents the TEM image of solution cast pure PVDF.

image of solution cast pure PVDF is also provided (inset in Fig. 1c). Clearly, as reported in our previous work [25], for the solution cast composites (S-0.7), randomly dispersed and distributed TiO2@MWCNTs nanoparticles can be identified. However, in the case of rolled composites (R-0.7), a preferred parallel orientation of TiO2@ MWCNTs to the rolling direction was clearly observed. Additionally, a comparison between Fig. 2 a and b reveals that rolling of solution cast composites results in disentanglement of some small TiO2@ MWCNTs bundles, inducing more individual TiO2@MWCNTs and increased separation between TiO2@MWCNTs. 3.2. Crystalline structure of composites XRD was performed to establish the PVDF crystalline phases of the rolled samples and the obtained patterns are displayed in Fig. 2a. As shown, the peaks at 20.7 and 36.7 are characteristic of the b-phase and the peak at 18 characterizes the b-phase [26,27]. Apparently, all samples were all characterized primarily by strong peaks corresponding to the b-phase, indicative of the predominant role of b-phase. Meanwhile, an intensified peak at 20.7 can also be detected in composites as compared with that of pure PVDF, which signifies the enhanced content of b-phase in composites. FTIR was adopted here to further identify the crystalline structure of PVDF in composites, as illustrated in Fig. 2b. A strong transition band corresponding to b-phase (840 cm1) and a weak characteristic band of a-phase (771 cm1) can be detected in all the rolled samples, suggesting that the b-phase is dominant in the rolled samples, which is in accordance with the XRD analysis. Furthermore, despite the fact that the b-phase transition bands of

all samples remain almost unchanged, the corresponding a-phase transition band of composites is weakened compared with that of the pure PVDF, which is clear manifestation of the suppressed aphase content (or increased b-phase content) in composites. In aiming to quantify the relative amounts of b-phase in samples, Equation. (1) was adopted [28,29].

Ab  FðbÞ ¼
Ab þ 1:26Aa

(1)

where F (b) represents the relative mass fraction of the b-phase and Aa and Ab are the areas of the peaks at 771 and 840 cm1 corresponding to the a- and b-phases, respectively. Presented in Fig. 2 c is the evolution of the fraction of b-phase in rolled samples. For comparison, the fraction of b-phase in solution cast samples is also presented. As reported in previous work [25], the existence of TiO2@MWCNTs can promote the formation of bphase due to the hydroxyl groups on TiO2 surface. Hence, an uptrend of the fraction of b-phase is seen in the samples containing TiO2@MWCNTs. Moreover, the rolled samples present enhanced bphase content over that of these solution cast samples, which may arise from the a-phase /b-phase transition induced by the applied tension during the rolling process [3,22]. Of particular interest is that a higher increment in F(b) induced by rolling is found in the composites when compared with the pure PVDF. This implies that the presence of TiO2@MWCNTs may be advantageous to the aphase/b-phase transition, which is also observable in the MWCNTs/PVDF system [22]. With an applied external stress, the phase transition only takes place in regions where local stress is

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Fig. 2. (a) XRD patterns of the rolled composites; (b)FTIR spectra of rolled composites; (c) fraction of b-phase of solution cast and rolled composites.

higher than the critical stress. The well-dispersed TiO2@MWCNTs nanoparticles are likely to assist stress transfer within the polymer matrix and lead to stress concentrations around them, resulting in an efficient transition in the interfacial areas. With a more efficient transition as compared to pure PVDF, the fraction of b-phase in composites is maximized at a loading of 0.3wt.% (~99%). The changes in melting temperature peak (Tm) as well as the crystallinity of the samples were detected by DSC scans, as shown in Fig. 3. The crystallinity is calculated according to the following equation (2) [30]:

Xc ¼

DHc  100% DH

(2)

where DHc is the melting heat of samples, DH is the melting heat of 100% crystalline PVDF, which is 93.07 and 103.40 J/g for pure a- and b-PVDF, respectively. DH ¼ 93.07  F (a)þ103.40  F (b) J/g is used for the calculation (F (a) ¼ 1  F (b)). It is well accepted that the Tm is related with the crystalline phase and lamellae thickness. Differernt crystalline phases possess varied Tm and a decreased lamellae thickness is expected to result in a lower Tm [31]. From the DSC thermograms (Fig. 3a), it is evident that all the rolled samples exhibit one single melting peak linked to the b-phase, confirming the predominate role of b-phase. Additionally, a progressive shift of Tm toward high temperature with increasing nanofillers concentration is observable. On the basis of fact that that the lamellae thickness is reduced by the

inclusion of nanofillers, as verified in the following SAXS analysis, the increased Tm in composites may be attributable to the strong interactions between nanofillers and polymer chains [32]. The calculated total crystallinity (Xc) of samples is detailed in Fig 3b. Generally, the existence of nanofillers can significantly affect crystallization behavior of crystalline polymers, enhancing the nucleation while hindering the crystallization process [33]. Thus, it is hard to predict the trend for crystallinity of composites due to the competition between these two effects. From our data, the later one prevails as the Xc of samples decreased with increasing loading content. Compared with the solution cast samples, the rolled composites, particularly these with loadings below 0.7wt.% exhibit enhanced Xc, suggesting that part of the amorphous area may turn into the crystalline area during the rolling process [34]. The crystallinity of b-phase (Xcb), which is critical to the resultant ferroelectric and piezoelectric response of PVDF, is also adopted here to make quantitative comparison. It can be seen that the Xcb of samples, determined from the product of F (b) and total crystallinity (Xcb ¼ Xc* Fb), shows a downward trend with increasing loadings except in the case of 0.3 wt.% loadings. With the slightly reduced crystallinity while higher b-phase content relative to these of pure PVDF, the rolled composites containing 0.3wt.% loadings display the maximum Xcb. To further investigate the bulk morphology of rolled composites, one-dimensional electron density correlation function g (z)

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Fig. 3. (a) DSC melting traces of rolled composites; (b) the calculated total crystallinity (Xc) and the crystallinity of b-phase (Xcb) of solution cast and rolled composites.

obtained from SAXS profiles was employed here, given by Refs. [35,36]:

Z gðzÞ ¼

0

∞

IðqÞq2 cosðqzÞdq Z ∞ IðqÞq2 dq

(3)

0

where I(q) is the scattering intensity obtained from the SAXS measurements, q¼(4p/l)*sin(q/2) (q is the scattering angle, and z is the direction normal to the lamellar interface (rolling direction). Accordingly, the correlation function curves of samples investigated are illustrated in Fig. 4a, with the inset demonstrating how the structural parameters including the long spacing (L), average thickness of the crystalline lamellae (lc) and interphase thickness (E) were obtained. From the structural parameters listed in Fig. 4b, it is distinctly seen that rolling of the solution cast composites leads to decreased L and lc, which may be interpreted by a stress-induced fragmentation process during deformation. Most notably, compared with the pure PVDF, the L together with the E of composites shows negiliable change while the lc was considerablely

reduced, suggesting that smaller crystallites are being nucleated through the incorporation of nanofillers. The inner crystallinity of the lamellae as defined by the ratio of lc to the L is also shown in Fig. 4b [37,38]. Obviously, the inner crystallinity, which represents the volume fraction of crystallites within the lamellae stacks, gradually decreases with increasing loadings in the solution cast samples. Rolling of the solution cast samples leads to decreased L and lc, but inducing enhanced inner crystallinity. The above results demonstated that the inclusion of TiO2@MWCNTs has the effect of reducing the crystallinity of composites while rolling can enhance the crystallinity of composites, which are consistent with the DSC analysis.

3.3. Dielectric properties and breakdown strength It is generally accepted that doping conductive component such as MWCNTs into polymer matrix allows for significant increment in dielectric constant due to the formation of micro-capacitancestructure and interfacial polarization effect (also named MaxwellWagner-Sillars (MWS) effect) [39,40]. As expected, upon the addition of TiO2@MWCNTs, enhanced dielectric constant can be

Fig. 4. (a) One-dimensional electron density correlation function g (z) of solution cast and rolled composites containing 1wt.% loadings; (b) average long spacing (L), interphase thickness (E), thickness of the crystalline lamellae (lc) and inner crystallinity of solution cast and rolled composites.

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achieved in composites relative to the pure PVDF (see Fig. 5a and b). More interestingly, the rolling process is capable of yielding dielectric constant increment for the pure PVDF whereas it has the opposite effect on that of composites, as summarized in Fig. 5b, which describes the dielectric properties of the two types of composites measured at 1000 Hz. With regard to pure PVDF, the increased dielectric constant induced by rolling is likely to be ascribed to the improved dipole density caused by the enhanced crystallinity, b-phase content and chain alignment, as confirmed in the above analysis. On the other hand, in the case of composites, the enlarged isolation distance between nanofillers induced by rolling may be responsible for the decline in dielectric constant [41,42]. According to the microcapacitance-structure model, the introduced nanofillers within the matrix are isolated by polymer layer, forming a number of microcapacitors, and thus contribute to enhanced dielectric constant. Due to the enlarged isolation distance between TiO2@MWCNTs and TiO2@MWCNTs, as evidenced in TEM, the average capacitance of the microcapacitors in the rolled composites diminishes, resulting in reduced dielectric constant [43]. Simultaneously, the depression of dielectric constant in rolled composites is also accompanied by a drastically suppressed dielectric loss when compared with the solution cast samples, as shown in the inset of Fig. 5b. In general, for the solution cast composites, there may exist some cavities left by the evaporated solvent, which is presumably to result in high dc conduction and consequently high dielectric loss. Rolling can be an efficient way to

induce a more dense structure for composites, and thus eliminate the dielectric loss originating from dc conduction in the lowfrequency range (G 103 Hz). Additionally, the dielectric loss in the low-frequency range also involves contribution from dipole relaxation in the crystalline region [25]. The improved chain order of crystalline region in these rolled samples may favor packing of the macromolecules and reduce free volume, restraining the chain mobility and thus leads to reduced dielectric loss under weak electric field. It should be note that for the two types of composites, the dielectric loss varies marginally with increasing TiO2@MWCNTs loadings and is comparable to that of pure PVDF, suggesting the effectiveness of TiO2@MWCNTs coreshell structure in depressing conduction loss at weak field [23,25,44]. The two-parameter Weibull analysis was adopted here to evaluate the breakdown strength of composites, in which the cumulative probability of breakdown P for a sample was expressed as

i h P ¼ 1  exp  ðE=Eb Þb

(4)

where E is experimental breakdown strength, Eb is the characteristic breakdown field at the cumulative failure probability of 63.2% and b is a shape parameter that evaluates the scattering of data [45,46]. The Weibull cumulative distribution function can then be expressed as:

lnð  lnð1  PÞÞ ¼ b ln E  b ln Eb

(5)

Fig. 5. (a) Dielectric constants (εr) and dielectric loss(tand) of rolled composites with frequency range between 102 and 107 Hz at room temperature; (b) dielectric constant and dielectric loss of solution cast and rolled composites measured at 1000Hz; (c)Weibull analysis of rolled composites and (d) breakdown strength of solution cast and rolled composites.

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In our experiments, fifteen specimens were used to perform the breakdown measurements of one sample. The Weibull plots of rolled composites are presented in Fig. 5c with the characteristic breakdown strength of each sample summarized in Fig. 5 d. The characteristic breakdown strength of solution cast composites is also depicted in Fig. 5d for comparison. One can see that the introduction of TiO2@MWCNTs can improve the breakdown strength of solution cast samples to some extent, as has proven in previous work [25]. Such enhancement in breakdown strength is presumably attributed to three factors: (1) due to the existence of TiO2 layer outside MWCNTs, the direct contact between neighboring conductive particles can be effectively cut off, thus suppressing the tunneling current at high fields; (2) the semiconductive TiO2 can act as a buffer layer between conductive MWCNTs and insulating PVDF matrix, which is anticipated to benefit the breakdown stability by mitigating field contrast between MWCNTs and PVDF [47,48]; (3) the strong interfaces induced by the good compatibility and dispersion of TiO2@MWCNTs in PVDF, as confirmed in our previous report [25], is beneficial to localize electrons, ions and polymer chains, providing more stable potential energy states to reduce the breakdown probability. Notably, the solution cast samples with a thickness of ~20 mm exhibit higher breakdown strength as compared to those reported before [25], which may be caused by the diminished thickness. Compared with the solution cast samples, the rolled samples present raised breakdown strength, presumably as a consequence of a more dense structure induced by rolling and the increased dispersion and orientation state of TiO2@MWCNTs. Due to the increased separation between TiO2@MWCNTs, less percolative pathways will be formed to transfer charge. Furthermore, the well dispersed TiO2@MWCNTs oriented in rolling direction (perpendicular to the applied electric field) can provide tortuous paths for treeing and more scattering centers for the hot electrons, thus restraining the electric treeing throughout the material and leading to enhanced breakdown strength [49]. 3.4. Piezoelectric properties The ferroelectric properties of the solution cast and composites were studied by room-temperature P-E hysteresis loop measurements, presented in Fig. 6a. It is interesting to note that compared with the solution cast samples, these rolled samples present lower values of coercive field (Ec) and higher values of remnant polarization (Pr). Kepler and Anderson [50] have suggested that dipoles in the crystallitc reorient by rotation of polymer chain segments about the chain axis through a step of 60 with the applied electric field. The mechanism of dipolar reorientation can be explained in terms of kink propagation with the 60 model [51], which involves the formation of kinks in the chain at the surface of the crystallite. The enhanced crystallinity while reduced crystallite size in those samples may lead to more surface area, increasing the probability of formation of kink bands and consequently decreasing the coercive field. Furthermore, the improved lamellar orientation as well as the chain orientation of rolled samples may be partly contribute to the drop in coercive field and enhanced Pr [52]. As for the rolled composites, shown in Fig. 6 b, the enhanced Pr value over that of the rolled pure PVDF may be explained by the interfacial effect. The value of Pr reaches a peak at a loading of 0.3 wt.%, further increasing content results in a decreased value, which is assumed to be caused by the observed decreased crystallinity of b-phase and saturation of interfacial effect due to the coalescing and overlapping of interfaces [53,54]. Nevertheless, further work should be performed to elucidate this phenomenon. The measured piezoelectric constants d33 for the poled samples

265

are shown in Fig. 6c where data are plotted as a function of poling voltage ranging from 60 V/mm to 120 V/mm. After poling at a 60 V/ mm field at 70  C for 1.2 h, the solution cast pure PVDF films acquires a piezoelectric coefficient of approximately 5 pC/N; that of the solution cast samples containing TiO2@MWCNTs is higher (maximum of 9 pC/N). Increasing poling electric fields can boost piezoelectricity. However, the poling electric field of the solution cast composites was limited below 90 V/mm, thus the maximum piezoelectric coefficient obtained is 14 pC/N. It is noteworthy that higher d33 values in rolled composites were measured against the solution cast composites when poling under the same condition. Moreover, given the improved breakdown strength, the rolled composites can be polarized at a higher field, giving rise to super piezoelectric properties. With a poling electric field of 120 V/mm, the piezoelectric coefficient d33 for the rolled composites reaches a maximum of ~41 pC/N at a loading level of 0.3wt.%. The b-phase crystallinity achieved through mechanical deformation is crucial for the piezoelectricity after poling [55]. With improved b-phase crystallinity, it is reasonable for the rolled samples to achieve super piezoelectricity against the solution cast samples. However, the increment in piezoelectricity is more significant than that in b-phase crystallinity when comparing these two types of samples, i.e., compared with the pure PVDF, the rolled pure PVDF yields a 10% increment in b-phase crystallinity while the increment in piezoelectricity can reach 40% (at a poling electric field of 60 V/mm). Thus, it is deduced that the super piezoelectricity of rolled samples may also originate from the highly oriented crystalline structure. As rolling induce the preferential alignment of the lamellar as well as polymer chains, the dipole moments within the oriented crystallites in the rolled specimens are likely to be distributed on the plane perpendicular to the rolling direction, whereas those in the solution cast samples are mostly randomly oriented. Considering the electric field was applied perpendicular to the film surface, the dipoles in the crystallites tended to rotate toward the direction of the electric field, thereby narrowing the distribution of crystallite orientation. The orientational distribution of the dipoles in the rolled samples with a better preliminary orientation of crystallites was biased much more strongly toward the poling direction than that of the solution films after poling [56]. Thus, the piezoelectric activity of the rolled samples improved. Meanwhile, despite the fact that the Xcb of composites displays a downtrend with loadings above 0.3 wt.%, these rolled composites still perform super piezoelectric properties against pure PVDF under the same poling condition. This suggests that the considerable enhancement in piezoelectric properties may also involve contributions from other factors. Through the addition of TiO2@MWCNTs, smaller crystallites and quantity interfaces were induced, which can significantly facilitate interfacial charge storage, allowing for an increased efficiency in dipole polarization and consequently improved piezoelectric response. Besides that, the existence of well-dispersed TiO2@MWCNTs within matrix may enhance the local electric field and create localized stress points, increasing domain mobility and facilitate polarization to some extent [57]. 4. Conclusion Composites consisting TiO2@MWCNTs and PVDF were prepared through the solution cast route, followed by a rolling process. Rolling of the solution cast composites give rise to enhanced bphase crystallinity, reduced crystallite size and a highly oriented structure, which are advantageous to yield improved ferroelectric and piezoelectric response. Additionally, the rolled composites possess significantly improved breakdown strength relative to the solution cast composites, presumably as a consequence of a more

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Fig. 6. (a) P-E loops of solution cast and rolled composites containing 0.5wt.% loadings and (b) P-E loops of rolled composites; (c) piezoelectric constants d33 of the solution cast (inset) and rolled composites as a function of poling electric field.

dense structure, increased dispersion and orientation state of TiO2@MWCNTs. Most importantly, smaller crystallites and quantity interfaces are introduced into the PVDF based composites upon the addition of TiO2@MWCNTs, resulting in considerable improvement in ferroelectric as well as piezoelectric performance. Acknowledgments This work was supported by the Major State Basic Research Development Program of China (973 Program, Grant No. 2015CB057501), National Nature Science Foundation of China (NSFC No. 51172108), Fundamental Research Funds for the Central Universities (NE2015101 & NE2015001), A Project Funded by the Priority Academic Program Development of Jiangsu Higher Education Institutions. References [1] G. Casar, X. Li, Q.M. Zhang, V. Bobnar, Influencing dielectric properties of relaxor polymer system by blending vinylidene fluoride-trifluoroethylenebased terpolymer with a ferroelectric copolymer, J. Appl. Phys. 115 (2014) 104101. [2] Xue JM. Alamusi, L.K. Wu, N. Hu, J.H. Qiu, C. Chang, et al., Evaluation of piezoelectric property of reduced graphene oxide(rGO)- poly(vinylidene fluoride) nanocomposites, Nanoscale 4 (2012) 7250.

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