Enhancement of dielectric and energy density properties in the PVDF-based copolymer/terpolymer blends.
Poly(vinylidene difluoride) (PVDF)-based polymers have attracted much attention due to their multifunctional electroactive polymers with superior mechanical properties. Their potential to be used as high permittivity dielectrics [1-6], organic ferroelectric memory devices [7, 8], electrostrictive actuators [9-12], and high energy density capacitors [13-20] has been widely studied over the last few decades.
Among various crystalline phases, the ferroelectric [beta]-phase is most likely to form in P(VDF-TrFE) due to the steric hindrance of the bulky TrFE group [21, 22]. The [beta]-phase is a highly ordered state, in which the polymer chains adopt the all-trans (TTTT) conformation. In this conformation, all the fluorine atoms are on one side of the chain and form dipoles perpendicular to the chain direction. P(VDF-TrFE) copolymer has Curie temperature ([T.sub.c]) around 111[degrees]C. The melting (Tm) and crystallization temperature are 152 and 142[degrees]C, respectively. PVDF-TrFE films can be generally annealed between the Curie and melting temperature because the chain mobility is higher in the paraelectric phase as compared to the ferroelectric phase. Moreover, chain mobility increases as a function of temperature. A higher chain mobility favors the lowest energy conformation (all-trans) thereby increasing its ferroelectric character. When the temperature increases, the ferroelectric [beta]-phase will undergo a phase transition and change into a paraelectric phase, consisting of a random sequence of trans-gauche (TG) bonds (e.g., TGTG') and [T.sub.3]G[T.sub.3]G' isomers , Correspondingly, the dipoles in the crystallites also change from an ordered state to a disordered state. When the temperature is just above the ferroelectric-paraelectric (FE-PE) transition, the dipoles can still be aligned by an external electric field. When an external electric field is applied, the random dipoles in the relaxor ferroelectric will reorient in the same direction as the electric field, causing a structure change in the polymer. Compared to normal ferroelectrics, relaxors exhibit a broad, diffuse dielectric peak at 25[degrees]C.
One of the greatest advantages of polymeric materials is their tunability. Apart from the chemical processes used to synthesize a material (e.g., copolymerization or irradiation), certain physical methods (e.g., changing processing conditions) are much easier to adjust in order to tune the properties of the polymers. For example, stretching the poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene) [P(VDF-TrFE-CFE)] terpolymer can lead to different electrocaloric responses . Casting terpolymer films at different temperatures also causes variations in the structure and dielectric response , Additionally, using different annealing temperatures can result in different electromechanical properties . Finally, adding nanoparticles or other polymers into the matrix polymer, to make a composite, is an easy way to tailor the properties or endow the composites with new properties [15-17]. The blend system provides a model system to study how random defects in the copolymer influence the polarization response in the copolymer. The blend also provides a direct way to understand the relationship between the polarization responses in the blends. When the defect content reaches 7-9%, the normal ferroelectric polymers will be converted to ferroelectric relaxors ; randomly distributed nanopolar regions will be embedded in a nonpolar matrix. These results demonstrate the promise of using a composite approach for tailoring the ferroelectric properties of PVDF-based ferroelectric polymers.
In a recent report, Chen et al.  investigated the electrocaloric effect of the P(VDF-TrFE-CFE) terpolymer (62.5/29/8.5 mol%) with a blended P(VDF-TrFE) (55/45 mol%) copolymer. They observed an enhancement in the ferroelectric properties with the P(VDF-TrFE) copolymer. As P(VDF-TrFE-CFE) terpolymer possesses a high dielectric constant at room temperature ~45 at 1 kHz  and the P(VDF-TrFE) copolymer shows a lower dielectric constant at room temperature (~12 at 1 kHz), polymer blends, which exploit the merits of both the base polymer and the additive polymer, offer a great opportunity to improve and tailor the properties of the base polymer [27, 28], Therefore, in this paper, we investigate the blends of P(VDF-TrFE) copolymer and P(VDF-TrFE-CFE) terpolymer. These blends exhibit the improvements in the dielectric constant and dielectric energy density, which are crucial to applications requiring capacitors with high energy density and good reliability under relatively low electric field.
The PVDF-TrFE copolymer with a VDF/TrFE molar ratio of 70/30 was obtained from Solvey Korea and P(VDF-TrFE-CFE) (62.5/29/8.5 mol%) was supplied by Piezotech (France). The 4-methyle-2-pentanone solvent was supplied by (Sigma-Aldrich) of purity [greater than or equal to] 99%. P(VDF-TrFE-CFE) terpolymer and P(VDF-TrFE) copolymer powders were separately dissolved in the solvent at 40[degrees]C. The two solutions were mixed at proper ratios for various blend compositions and then stirred at 50[degrees]C for 1 h. Pt/ Ti[O.sub.2]/Si[O.sub.2]/Si substrates were thoroughly cleaned via a series of ultrasonic baths (isopropanol, ethanol, and then distilled water) for 5 min each and then dried with [N.sub.2] gas. Before fabrication, the humidity, and atmospheric pressure were controlled. Solutions of P(VDF-TrFE)/P(VDF-TrFE-CFE) blends were pipetted on to the Pt/Ti[O.sub.2]/Si[O.sub.2]/Si substrate surface and then spin-coated at 2,500 rpm for 30 s. The spin-coated films were dried at 110[degrees]C for 30 min to remove solvent. Then, films were thermally annealed at 130[degrees]C in vacuum for 6 h to enhance film quality. This procedure produces films with a thickness of ~0.8 [micro]m. Finally, P(VDF-TrFE)/P(VDF-TrFE-CFE) blend capacitors with a metal/blend/metal structure were fabricated to investigate the ferroelectric properties. Pt top electrodes with an area of 4.0 X [10.sup.-3] [cm.sup.2] were deposited using DC sputtering (Cressington 108, Cressington, USA). P-E hysteresis loops were measured by a ferroelectric tester (PRECISON LC, Radiant Technologies, USA). The dielectric and impedance data for the P(VDF-TrFE)/ P(VDF-TrFE-CFE) blends were obtained in the frequency range of 0.01 Hz to 10 MHz over a temperature range of 20-120[degrees]C using an impedance analyzer (Hioki 3522-50 LCR HiTESTER). FTIR was performed by using a Fourier transformer (Agilent, USA).
RESULTS AND DISCUSSION
Typical hysteresis loops for P(VDF-TrFE) and P(VDF-TrFE-CFE) films are shown in Fig. la and b. A step-by-step enhancement in the applied electric field resulted in a systematic increase in the remnant polarization ([P.sub.r]). Well-saturated hysteresis loops were obtained at 125 MV/m. The typical [P.sub.r] and coercive field ([E.sub.c]) values were ~6.3 [micro]C/[cm.sup.2] and 65 MV/m, respectively. These results indicate that the polymer can sustain an electric field greater than 125 MV/m (which was the voltage limit of our polarization measurement set-up). The observed [P.sub.r] and maximum polarization (Pm) values agree with those reported by Chen et al. , For the P(VDF-TrFE-CFE) terpolymer, a very slim P-E loop with very low [P.sub.r] and [E.sub.c] values was observed; this is typical for a ferroelectric relaxor.
The ferroelectric responses of blends with [greater than or equal to] 20 wt% terpolymer were investigated. Figure 2a shows the bipolar polarization loops of the blends measured under an electric field of 125 MV/ m at room temperature. The copolymer film exhibited a near square, saturated P-E loop with a large [P.sub.r] of ~6.3 [micro]C/[cm.sup.2]. A very slim P-E loop with very low [P.sub.r] and [E.sub.c] values, which is typical for a ferroelectric relaxor, was observed for the terpolymer. The data show that both the [P.sub.r] and [E.sub.c] can be tuned by adjusting the blend compositions of the copolymer/terpolymer blends. Since the maximum polarization ([P.sub.m]) of the copolymer and terpolymer differ from one another, there is also a difference between the [P.sub.m] in the polarization hysteresis (P-E) loops among different blends as seen in Fig. 2a. On the other hand, significant variations were found in the [E.sub.c] and [P.sub.r], as can be seen in Fig. 2b. In the intermediate composition, both the [E.sub.c] and [P.sub.r] can be tuned between these two extremes, which are potentially attractive for device applications.
It is well known introducing defects into the crystal lattice of a traditional, normal ferroelectric ceramic can transform it to a relaxor ferroelectric with improved dielectric constant and electromechanical response [29, 30], Therefore, we conclude that the copolymer changed to a typical relaxor ferroelectric after the addition of the terpolymer, which resulted in the degradation of [E.sub.c] and [P.sub.r] under an electric field of 125 MV/m.
The mechanics and aggregation characteristics of polymeric chains can differ when different crystalline phases are formed, resulting in different surface morphologies. This can be studied using atomic force microscopy (AFM). Figure 3 shows surface and 3D 1 nm X 1 nm AFM images of a typical films made from P(VDF-TrFE), P(VDF-TrFE-CFE), and their blends. The rod-like shape of the grains was attributed to [beta] phase crystallites. The morphology of the grains and the roughness of the surface were affected by terpolymer defects. The average length of the grains was approximately 180 nm with root mean square (RMS) roughnesses of 1.4, 6.7, 5.5, and 5.0 nm for P(VDF-TrFE), P(VDF-TrFE-CFE) 20%, P(VDF-TrFE-CFE) 60%, and P(VDF-TrFE-CFE) 100%, respectively.
Molecular vibration analysis is a key to understanding the dynamics of a material. Fourier transform infrared spectroscopy (FTIR) can be used to detect the vibrational mechanics of a material system by monitoring the absorption of infrared energy. FTIR was employed to detect local conformational structure changes of the PVDF-based copolymers. The normalized FTIR absorption spectra for the copolymer, terpolymer, and with blends are shown in Fig. 4. Here, we only focus on the three intense bands associated with the [beta] phase of P(VDF-TrFE) (850,1285, and 1400 [cm.sup.-1]). Reynolds et al.  described the 850 and 1285 [cm.sup.-1] band as belonging to the CF2 symmetric stretching with the dipole moments parallel to the polar b axis. The 1400 [cm.sup.-1] band is assigned to the C[H.sub.2] vibration with the dipole moment along the c axis. The copolymer shows a chain of conformations that is consistent with what has been reported in the literature [11, 32], With the addition of the terpolymer, the crystal bands at 1285 and 850 [cm.sup.-1], which correspond to the all trans ([T.sub.m > 3]) conformation, gradually grow downward. It can be seen that at a terpolymer content of 20%, the TGTG' conformation is increased. This means that the terpolymer embedded in the blends has adopted a mixed structure. As the terpolymer content becomes greater than 20%, a quick decrease of the crystal bands is observed. This indicates that the influence of the terpolymer has grown. The random defects in the copolymer have influenced the crystallite in the copolymer and confirmed that the structure has transformed from a normal ferroelectric into a ferroelectric relaxor.
Figure 5a and b shows the frequency dependent dielectric constant of the pure copolymer, terpolymer, and their blends. The dielectric constant of the pure P(VDF-TrFE) copolymer was observed to be ~11 at 1 kHz. By blending with only 20% P(VDF-TrFE-CFE), the dielectric constant was increased to 24 and achieved ~55 when the terpolymer doping level reached 60% as shown in Table 1 of this manuscript. At low teropolymer content, the random defects in the teipolymer influence the ferroelectric response in the copolymer very little through interfaces between the two polymers, and consequently, the whole blends exhibit relaxor ferroelectric response. These combined effects of crystallinity and interfacial couplings lead to an enhanced dielectric constant compared with that in the neat terpolymer. Therefore, the enhancement observed in blends with the terpolymer is believed to be related to interfacial couplings between the terpolymer and the copolymer . While, at high copolymer content, the two phases become more distinguished, and each component tends to keep their own structures, leading to intermediate electric properties between pure terpolymer and copolymer. The copolymer is converted to relaxor through interfacial couplings and thus enhances the overall dielectric response of the blends. When the terpolymer weighs lower 60%, the dielectric constant of the blends becomes lower than that of pure terpolymer. However, it is noted that for the blend with 60 wt% terpolymer, its dielectric constant was found to increase than that of pure terpolymer indicating that the interface effect still play its role in the blends. The dielectric constant was found to increase with terpolymer but making a shoulder at 40% sample which is almost same to that of sample 30%. This may be due to not the optimize level for the increase of dielectric constant and the interface effects plays little low role to increase the dielectric constant. However, the dielectric constant was found to increase for the sample 50%. Defects in the copolymer can hinder the formation of large ferroelectric domains in the P(VDF-TrFE) copolymer, converting the copolymer into the relaxor ferroelectric, leading to a higher dielectric constant. Conversely, at high copolymer content, the copolymer is a normal ferroelectric with a lower dielectric constant. Consequently, the total dielectric response increases. The dielectric loss was observed for the pure copolymer around 8% and was increased to 10% for the sample with terpolymer 60%.
Figure 6 shows plots of the leakage current density versus the applied electric field (J-E) for pure and P(VDF-TrFE)/ P(VDF-TrFE-CFE) blend films under a DC applied electric field measured at room temperature. The leakage current densities of the copolymer were observed to be large and decreased gradually with increasing terpolymer content.
The measured leakage current density of the copolymer thin film was 8.3 x [10.sup.-8] A/[cm.sup.2] at an applied electric field of 10 MV/m. At the same applied electric field, the measured leakage current density of the 60% terpolymer thin film was 3.7 x [10.sup.-8] A/[cm.sup.2]. The leakage current density of the 60% terpolymer thin film is an order of magnitude lower than that of the copolymer thin film. At electric field 10 MV/m the leakage current density for all compositions is almost same, because the electric field is very low. In the case of PVDF-based polymers need relatively large electric field. At electric field above 20 MV/m the change in the leakage current density is seem to be clear and the leakage current density is decreased with doping of terpolymer as in shown in Fig. 6b.
In order to understand the origin of the leakage current, we investigated the leakage current mechanism of the thin films. A linearity of the curves was observed over the entire region of the applied electric field for the 60% terpolymer blend and the pure terpolymer thin films, which indicates an Ohmic conduction mechanism (S ~ 1) . The log (J) is nearly proportional to log(E); however, the change in the slope from 1 to 2 for the copolymer and the 20% terpolymer-doped thin films indicates a transition of the conduction mechanism from an Ohmic to a space-charge-limited conduction mechanism . In the high electric field region, the copolymer and [less than or equal to] 30% terpolymer blend thin films show a change in the slope combination of the space-charge-limited conduction mechanism with another unknown conduction mechanism .
The unipolar polarization hysteresis (P-E) loops of the copolymer, terpolymer, and their blends are presented in Fig. 7a. For blends with 20-60 wt% terpolymer, the polarization response is greatly increased compared to the pure copolymer because the polarization level of the copolymer is lower than that of the terpolymer as shown in the Fig. 7a. The blend with 60 wt% copolymer was 6.6 [micro]C/[cm.sup.2] and also higher than the polarization response of the terpolymer 5.8 [micro]C/[cm.sup.2]. The energy density of copolymer/terpolymer blends calculated from the unipolar P-E loops is presented in Fig. 7b. This data shows an increase in the energy density for 20-60 wt% terpolymer blends. The possible structural changes in the terpolymer or copolymer components in the blends were investigated under 125 MV/m to discover the mechanisms for the enhancement of the polarization response and energy density. A high energy density of ~4.2 J/[cm.sup.3] was obtained in these blends at 60 wt%, in contrast to ~3.6 J/[cm.sup.3] which was seen in the terpolymer, at an applied electric field of 125 MV/m. Blending the pure copolymer with 60 wt% terpolymer increased the energy density to 4.2 J/[cm.sup.3] at 125 MV/m; this increase is larger than what was observed by Chu et al.  for their terpolymer P(VDF-TrFE-CFE) 63/37/7.5 composition which had an energy density of 3.2 J/[cm.sup.3] at 125 MV/m.
We can explain the observed polarization and energy density enhancement at relatively low fields is the polarization contribution from the interfaces between these two polymer components, which has been observed in many multicomponent systems (blends and nanocomposites) [1, 35-38], For example, in PVDF/ nylon blends and 2-2 composites, observation of a higher polarization response than both pure PVDF and nylon in certain composition range was reported. This phenomenon is obviously not consistent with the consideration based on the traditional mixing rule of the composite materials. An additional polarization contribution from the interfaces was proposed in these multicomponent systems even though the origin of the interfacial polarization is still unclear (probably due to the different conductivity between two components and the resulting charge accumulation at interface) [35, 37], The results demonstrate the promise of blend approaches for tailoring and enhancing the dielectric properties of ferroelectric polymers.
The high dielectric constant and energy storage density of the P(VDF-TrFE)/P(VDF-TrFE-CFE) blend system can be used for practical application in flexible field-effect transistor (FET), nanogenerators and ferroelectric random access memory (FeRAM) like devices."
We investigated the microstructure, dielectric, and ferroelectric properties of P(VDF-TrFE)/P(VDF-TrFE-CFE) blends. In AFM images, the average length of the grains was approximately 180 nm with RMS roughnesses of 1.4, 6.7, 5.5, and 5.0 nm for P(VDF-TrFE), P(VDF-TrFE-CFE) 20%, P(VDF-TrFE-CFE) 60%, and P(VDF-TrFE-CFE) 100% were, respectively. FTIR confirmed that the addition of the terpolymer caused the crystal bands at 850, 1285, and 1400 [cm.sup.-1] (corresponding to the all trans ([T.sub.m > 3]) conformation) to gradually grow downwards. This indicates a transition from a normal ferroelectric to a relaxor. Furthermore, a well-saturated hysteresis loop for P(VDF-TrFE) is obtained at 125 MV/m with typical measured Pr and Ec values of ~6.3 [micro]C/ [cm.sup.2] and 65 MV/m, respectively. The dielectric constant with only 20% terpolymer was increased to 24 and reached ~55 when the terpolymer doping level reached 60%. The measured leakage current density of the copolymer thin film was 1.007 X 10 7 A/ [cm.sup.2] at an applied electric field of 10 MV/m. At the same applied electric field, the measured leakage current density of the 60% terpolymer blend was decreased to 4.25 x [10.sup.-8] A/[cm.sup.2]. Consequently, a higher energy density of ~4.2 J/[cm.sup.3] was obtained in these blends, in contrast to ~ 3.6 J/[cm.sup.3] observed in the terpolymer, at an applied electric field of 125 MV/m.
The dielectric and energy storage density behavior of this system can make it a highly appropriate for practical application in flexible FET, nanogenerators, and FeRAM like devices.
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Amir Ullah, (1,2) Ata ur Rahman, (3) Chang Won Ahn, (2) Muneeb-ur Rahman, (1) Aman Ullah, (4) Zia-ur Rehman, (5) Muhammad Javid Iqbal, (6) III Won Kim (2)
(1) Department of Physics, Islamia College, Peshawar, Peshawar 25120, KP, Pakistan
(2) Department of Physics and EH SRC, University of Ulsan, Ulsan 680-749, Republic of Korea
(3) Institute of Chemical Sciences, University of Peshawar, Peshawar 25120, KP, Pakistan
(4) Department of Physics, University of Science and Technology, Bannu, KP, Pakistan
(5) Department of Chemistry, Quaid-i-Azam University, 45320 Islamabad, Pakistan
(6) Department of Physics, University of Peshawar, Peshawar 25120 KP, Pakistan
Correspondence to: I. W. Kim; e-mail: email@example.com Contract grant sponsor: Ministry of Education (Basic Science Research Program through the National Research Foundation [NRF], Korea); contract grant numbers: 2012R1A1A2005922, 2014R1A1A4A01004404, and 2009-0093818.
Published online in Wiley Online Library (wileyonlinelibrary.com).
TABLE 1. Electrical data of PVDF-TrFE/PVDF-TrFE-CFE blend system. [P.sub.r] [E.sub.c] Material ([micro]C/[cm.sup.2]) (MV/m) PVDF-TrFE 6.3 77.0 PVDF-TrFE-CFE20% 4.6 66.4 PVDF-TrFE-CFE30% 4.4 52.6 PVDF-TrFE-CFE50% 4.1 49.8 PVDF-TrFE-CFE60% 2.4 39.8 P(VDF-TrFE-CFE) 100% 1.1 5.9 Energy density [epsilon] Material (J/[cm.sup.3]) (10 kHz) Loss PVDF-TrFE i.i 12 0.08 PVDF-TrFE-CFE20% 2.0 21 0.09 PVDF-TrFE-CFE30% 2.7 29 0.11 PVDF-TrFE-CFE50% 3.3 40 0.12 PVDF-TrFE-CFE60% 4.2 53 0.10 P(VDF-TrFE-CFE) 100% 3.6 47 0.13 Leakage current Material (A/[cm.sup.2]) at 40 MV/nt PVDF-TrFE 5.6E-6 PVDF-TrFE-CFE20% 1.0E-6 PVDF-TrFE-CFE30% 3.8E-7 PVDF-TrFE-CFE50% 2.2E-7 PVDF-TrFE-CFE60% 1.4E-7 P(VDF-TrFE-CFE) 100% 8.6E-8
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|Author:||Ullah, Amir; Rahman, Ata ur; Ahn, Chang Won; Rahman, Muneeb-ur; Ullah, Aman; Rehman, Zia-ur; Iqbal,|
|Publication:||Polymer Engineering and Science|
|Date:||Jun 1, 2015|
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