Stretchable quaternary phasic PVDF-HFP nanocomposite films containing graphene-titania-SrTiO3 for mechanical energy harvesting

  • Deepalekshmi Ponnamma
  • Alper Erturk
  • Hemalatha Parangusan
  • Kalim Deshmukh
  • M. Basheer Ahamed
  • Mariam Al Ali Al-Maadeed
Original Article


Integrating efficient energy harvesting materials in to soft, flexible, and eco-friendly substrates could yield significant breakthroughs in wearable and flexible electronics. Substantial advances are emerged in fabricating devices which can conform to irregular surfaces in addition to integrating piezoelectric polymer nanocomposites in to mechanical generators and bendable electronics. Here, we present a tri-phasic filler combination of one-dimensional titanium dioxide (TiO2) nanotubes, two-dimensional reduced graphene oxide, and three-dimensional strontium titanate (SrTiO3), introduced in to a semi-crystalline polymer, poly(vinylidene fluoride-co-hexafluoropropylene). Simple mixing method was adopted for the composite fabrication after ensuring a high interaction between the various fillers. The prepared films were tested for their piezoelectric responses and mechanical stretchability. The results showed that the piezoelectric constant has increased due to the change in the filler concentration and reached a value of 7.52 pC/N at 1:2 filler combination. The output voltage obtained for the same filler composition was about 10.5 times that of the voltage generated by the neat polymer. Thus, we propose integration of these materials in fabricating energy conversion devices that can be useful in flexible and wearable electronics.

Graphical abstract


Flexoelectric Energy harvesting Foldable films Nanocomposites 

1 Introduction

Polyvinyledene fluoride (PVDF) is a semi crystalline polymer exhibiting excellent piezoelectric and pyroelectric properties [1, 2, 3]. Its functional applications in electronic, structural, and biomedical fields are remarkable due to the thermoelectric properties, mechanical strength, chemical resistance, and biocompatibility. It has widespread applications in supercapacitors, ultrafiltration membranes, rechargeable batteries, and optoelectronics [4, 5, 6, 7]. The –CH2 and –CF2 groups on the skeletal PVDF are arranged in specific ways and this is responsible for the PVDF unique properties [8, 9]. Depending on the processing conditions of this polymer, the polymer crystallizes to five different conformations-α, β, γ, δ, and ε [10, 11]. The major α form has stability and possesses good mechanical and thermal properties. The significance of β and γ forms lies in its ability to transfer and even store ionic charge. Manufacturing method has great influence on the degree of crystallinity and amount of each phase [12, 13, 14, 15].

PVDF and its co-polymers, poly(vinylidene fluoride-co-hexa-fluoropropylene) (PVDF-HFP) and poly(vinylidene fluoride-cotrifluoroethylene) (PVDF-TrFE) possess high electric displacement polarization [16, 17] and their nanocomposites are reported for high dielectric constant. Glycidyl methacrylate functionalized PVDF-HFP nanocomposites containing amino silane modified BaTiO3 were prepared by Xie et al. through a grafting to approach [18]. With 50 wt.% of BaTiO3, 3.9 times enhancement in dielectric constant was observed, at the same time achieving a 2.67 fold increase in thermal conductivity. Prabakaran et al. varied the concentration of silane treated TiO2 in PVDF-HFP by 4–20 wt.% and observed a two fold increase in dielectric constant [19]. With the emergence of carbonaceous fillers such as carbon nanotubes and graphene, such composites of PVDF-HFP are also well reported for their increased electrical conductivity [15, 16, 20, 21, 22].

Energy harvesting materials and devices based on flexible polymer nanocomposites have significant attention in recent times, due to their ability to convert mechanical vibrations to electrical energy [23]. The harvested electric power density of 464 and 561% respectively were reported for the neat PVDF-HFP films using ac and dc circuits [24]. The significance of carbon black filler particles in regulating the α-β phase transformation by acting as nucleating agents was well demonstrated. Graphene is a two-dimensional carbon nanostructure and as such it is a little difficult to get it dispersed well within a polymer and thus to achieve required properties [25, 26]. Thus, graphene oxide (GO), reduced graphene oxide (rGO), or functionalized forms were used to increase the electrical and dielectric properties of the nanocomposites [27, 28, 29]. PVDF composites containing GO and rGO were prepared by Rahman et al. and achieved remarkable dielectric and ferroelectric properties with 0.1 wt.% loading [30]. The energy harvesting device made from the PVDF/rGO film showed 36 nW power against 704 kΩ load resistance.

In this study, we take the advantage of ceramic filler particles of high piezoelectric coefficient, graphene and TiO2 of good dielectric performance and the flexible ferroelectric PVDF-HFP. The hybrid architecture of graphene-TiO2 was achieved by hydrothermally growing the TiO2 nanotubes in presence of GO [31]. The nanosheet-nanotube hybrid structure imposed good dielectric properties, at the same time, the third ceramic SrTiO3 filler particles were added to induce the piezoelectric effect. The poled composite films showed enhanced energy harvesting responses in terms of generated voltage and piezoelectric coefficient values. The electrostatic charges in the hybrid graphene-TiO2 architecture, crystalline regions of SrTiO3, and the good interfacial connection formed in between the quaternary nanocomposite (PVDF-HFP/graphene/TiO2/SrTiO3) altogether lead to good piezoelectric response.

2 Experimental details

2.1 Materials

PVDF-HFP pellets of Mw ≈ 400,000 were purchased from Sigma-Aldrich. N,N-dimethylformamide (DMF) and acetone (BDH Chemicals, Qatar) were used as solvents for the polymer. All other chemicals required for the synthesis of hybrid nanomaterial (graphene-TiO2) as per the followed protocol and SrTiO3 of < 100 nm particle size were also obtained from Sigma Aldrich.

2.2 Synthesis of graphene-TiO2 filler

The graphene-TiO2 hybrid structure was synthesized according to the previously reported protocol [31]. By this procedure, a dispersion of GO in distilled water was made (0.1 g in 10 ml) to which TiO2 dispersion (1.2 g in 10 ml) was added. The whole suspension after stirring was transferred to an autoclave and kept at 130 °C for 10 h. The cooled resultant material was later filtered and the precipitate obtained was washed in water and 0.1 M HCl until it achieves neutrality. The final powder, TiO2 nanotubes grown rGO (G-T) was obtained after drying the precipitate at 80 °C for 4 h.

2.3 Synthesis of PVDF-HFP composites

G-T was mixed with PVDF-HFP in specific ratio by solvent casting. A ternary filler composition was also prepared by mixing the G-T together with SrTiO3 and thereafter reinforcing the PVDF-HFP at various ratios. The DMF/Acetone solvent mixture in 1:1 ratio has been used as the common solvent and the nano powders were dispersed in the solvent mixture prior to nanocomposite fabrication by bath sonication (30 min). The final composites were synthesized by magnetic stirring followed by casting, drying and hot pressing at 170 °C. The sample compositions are provided in Table 1.
Table 1

Contact angle measurements for the PVDF-HFP composites



Contact angle





109.01 ± 1.05




PVDF-HFP with 1 wt.% G-T

110.06 ± 1.68




PVDF-HFP with 1 wt.% G-T and 1 wt.% SrTiO3

95.66 ± 2.49




PVDF-HFP with 1 wt.% G-T and 2 wt.% SrTiO3

99.84 ± 3.08




PVDF-HFP with 1 wt.% G-T and 3 wt.% SrTiO3

97.48 ± 1.95



2.4 Characterization methods

Sample morphology was analyzed by SEM and TEM. The SEM studies were conducted by SEM, XL-30E Philips Co., Holland and the TEM by transmission electron microscope FEI TECNAI G2. The surface contact angles were measured on a drop shape analysis system (OCA 35- Dataphysics) using deionized water to study the hydrophobic effect of the composite surface. FTIR spectra of the samples were recorded with PerkinElmer Spectrum 400 spectrophotometer in the range, 400–4000 cm−1 with a resolution of 2 cm−1. X-ray diffraction studies were performed by XRD diffractometer (Mini Flex 2, Rigaku). Nickel-filtered CuKα radiation (λ = 0.1564 nm) operated at 30 V and 15 mA served as the source. The patterns were recorded in the 2θ range of 10–30 ° at a scanning speed of 1.8 °/min. Broadband dielectric/impedance spectroscopy-Novocontrol was used for measuring the dielectric properties of the samples. The dielectric constant, dielectric loss, tan δ and conductivity values of the samples were checked during the frequency range from 10−2 to 106 Hz. Piezoelectric studies were done using an assembled set up consisting of a vibrating shaker, amplifier, frequency generator and accelerometer [32, 33]. Specific weights were placed on the sample in such a way that variable forces were imposed on it in the transverse direction. The resistance values were also adjusted to get maximum output. All samples were corona poled for 7 s by subjecting to an applied voltage of 6 kV prior to piezoelectric measurement.

3 Results and discussion

3.1 Morphology

Schematic representation of the composite film fabrication is shown in Fig. 1. Both the filler morphologies are clearly observed as TEM images. The mixed nanosheet-nanotube morphology (TiO2 nanotubes grown among the rGO nanosheets) of G-T and spherical morphology of SrTiO3 are expected to create a uniformity in filler dispersion within the polymer as the scheme represents. In addition, the poling process aligns the filler units and the polymer chains, which contributes towards the enhanced piezoelectric property of the composite.
Fig. 1

Schematic representation of the preparation of PVDF-HFP composites

Distribution of fillers throughout the PVDF-HFP matrix is clear from the SEM images of Fig. 2. The graphene-TiO2 hybrid architecture provides a continuous web like pattern within the nanocomposites, and the spherical SrTiO3 nanoparticles are embedded within the web. This leads to a very good homogeneity and uniformity in filler dispersion within the polymer, and thus the nanocomposite is capable of showing good properties even at low filler concentrations (as will be seen later). Hydrothermal synthesis of G-T nanohybrid ensures the formation of nanotubes on the surface of rGO sheets as well as the folded sheets around the tubes [31]. Good interfacial interaction was observed for such hybrid architecture due to the presence of additional C-Ti bond [34]. The uniform dispersion of spherical SrTiO3 particles hide the layers of G-T as the latter is present in small concentration (1 wt.%). However, the combination of one-, two-, and three-dimensional nanostructures contributes homogeneous filler dispersion in the nanocomposite and causes strong network structures in the PVDF-HFP matrix. But the usual agglomeration tendency of the nanoparticles is observed here also when the filler combination ratio became 1:3 (Fig. 2d). The uniform surface conductivity coming from the uniform filler distribution helps the nanocomposite to drain the surface charge and thus to achieve better device performance [35].
Fig. 2

FESEM images of the fracture surface morphology of PVDF-HFP nanocomposites a PHP/GT1, b PHP/GS11, c PHP/GS12, d PHP/GS13

The contact angle values for the sample films towards water medium are also investigated and included in the Table 1. The values show that the binary filler combination of graphene-TiO2 enhanced the hydrophobicity whereas the ceramic filler SrTiO3 reduces the hydrophobicity [36, 37]. This evidences the pronounced influence of ceramic filler particles within the ternary nanocomposites.

3.2 Crystallinity

Crystalline β and γ phases of PVDF-HFP nanocomposites were identified by XRD and FTIR studies. Figure 3 represents the FTIR spectra and XRD diffraction patterns for all nanocomposites studied. The various bands shown in the FTIR spectra give clear evidence for the presence of α, β and γ crystal planes of the samples. The bands near 610, 760, 790, 870, 970, 1160, 1380, and 1420 cm−1 are due to the α-crystalline phases of the samples, whereas the bands near 840, 1270, and 1420 cm−1 correspond to the β-crystalline phases [38, 39]. In addition the bands near 810, 835, and 1220 cm−1 are due to the γ-phase of the nanocomposites. The bands at 840 cm−1 become sharper as the filler composition enhances, which is attributed to the nucleation and stabilization of the electroactive γ and β phases [38, 39]. A quantitative estimation of the crystallinity is done to understand the influence of hybrid nanofillers on the crystalline behavior of the PVDF-HFP. The relative amount of β and γ-phases of the samples are calculated [38, 40] by following the two similar equations given below.
$$ F\left(\beta \right)=\frac{{\mathrm{X}}_{\beta }}{{\mathrm{X}}_{\alpha }+{\mathrm{X}}_{\beta }}=\frac{A_{\beta }}{\left(\raisebox{1ex}{${K}_{\beta }$}\!\left/ \!\raisebox{-1ex}{${K}_{\alpha }$}\right.\right){A}_{\alpha }+{A}_{\beta }}\kern0.75em $$
$$ F\left(\gamma \right)=\frac{{\mathrm{X}}_{\gamma }}{{\mathrm{X}}_{\alpha }+{\mathrm{X}}_{\gamma }}=\frac{A_{\gamma }}{\left(\raisebox{1ex}{${K}_{\gamma }$}\!\left/ \!\raisebox{-1ex}{${K}_{\alpha }$}\right.\right){A}_{\alpha }+{A}_{\gamma }}\kern0.5em $$
where Kα (0.365 μm−1), Kβ (0.132 μm−1) and Kγ (0.150 μm−1) are the absorption coefficients at the particular wavenumber. Xα, Xβ, and Xγ are mass fractions of α, β, and γ crystalline phases. Aα, Aβ, and Aγ are the area of absorption bands at 764, 840 and 832 cm−1.
Fig. 3

FTIR spectra and XRD diffraction patterns for the PVDF-HFP nanocomposites

In the XRD spectra (Fig. 3b), all samples showed similar peaks with variable intensities corresponding to the reflection planes of PVDF-HFP. The main peaks at 17.7, 18.3, 19.9, and 26.5° correspond to the (100), (020), (110), and (021) crystal planes of the non-polar α phase [41, 42]. The intensity of the peak around 20° is high since this peak is overlapped with the diffraction from the (002) plane coming from the polar γ phase. The β phase also contributes to have a peak at this position. The reduced peak intensity of the α-phase at 17.7 and the enhanced peak intensity near 20 suggest the conversion of non-polar α-phase to polar γ-phase which favors the piezoelectric behavior of the sample [38]. This is also an indication of the good filler-polymer interactions existing within the samples.

Crystallization behavior of PVDF-HFP by the addition of ternary hybrid nanofillers is also studied by DSC. For this, the second heating and cooling characteristics of the samples were monitored and the curves obtained are depicted in Fig. 4. Compared with the neat matrix, all composites showed an increase in melting and crystallization temperature as marked in the figure. The melting temperature was enhanced from 108.5 to 110.5 °C and then to 113.9 °C as the samples were changed from pure polymer to PVDF-HFP/GS11 and then to PVDF-HFP/GS13 respectively. Similarly crystallization temperature became 145.8 °C for PVDF-HFP/GS12 from 142.5 °C of pure polymer. This enhancement in the temperature values is attributed to the nucleating action of the fillers within the polymer [41, 43]. However, the filler dispersion significantly can affect the uniform crystallite formation, so at higher loading (PVDF-HFP/GS13), a decrease in crystallization temperature is observed. Other than the filler dispersion, the filler dimension and surface area also can influence the nucleation and polymer chain mobility and thus the crystallization process [44]. Percentage of crystallinity was calculated for all the composites using the DSC analysis.
$$ CI=\frac{\left({\Delta H}_f\right)}{\left(1-\varnothing \right)\Delta {H}_f^0}100 $$
Fig. 4

DSC a second melting and b cooling curves for the PVDF-HFP nanocomposites

where CI represents the crystallinity index, ΔHf the difference between the measured heat of fusion and the measured heat of crystallization, Ø the weight percentage of the filler in the composite and ΔHf0 the enthalpy of crystallization for pure PVDF-HFP (105 J/g). The crystallinity index for the pure sample is 40% which decreases upon the addition of nanofillers. This substantiates the changes or phase transformations observed for the β-phase and the γ-phase as evident from the FTIR and XRD studies.

The onset temperature and enthalpy values provided in Table 2 also follow the same trend of crystallinity and substantiate the influence of tri phasic filler combination in regulating the polymer properties.
Table 2

Melting and crystallization data for the nanocomposites from DSC




Tc onset

Tm onset














































3.3 Dielectric property

A composite system is supposed to be useful in energy storage applications, if it has a high dielectric constant and low dielectric loss [44]. Since energy storage capacity is also important for those materials used in energy generating devices, analyzing the dielectric properties of current polymer nanocomposite systems has utmost significance [40]. The frequency-dependent dielectric constant, dielectric loss, and conductivity values are provided in Fig. 5. Figure 5a, b shows higher values for the dielectric constant and lower values for dielectric loss for all the nanocomposites studied. At 1 kHz, the neat polymer had a dielectric constant of 8, which changes to 39 in the case of PVDF-HFP/GS11 nanocomposite. At the same time the dielectric loss value for the neat PVDF-HFP was 0.24 and for the nanocomposite with maximum filler loading the value became 1.32. The higher frequencies witness a decreasing tendency of the dielectric constant since relaxation of polymer chains happens [45]. A nanocomposite system consists of an insulating polymer matrix and nanofillers with relatively higher conductivity values. This charge difference induces surface polarization and charge accumulation at the filler-polymer interfaces. This interfacial polarization happens at low frequency region where the conducting nanofillers act as small capacitors and this phenomenon is referred to as Maxwell-Wagner-Sillars effect [46, 47]. This is also evidenced from the AC conductivity values shown in Fig. 5c. It is clear that the conductivity values show a gradual increase depending on the concentration of the filler particles added to PVDF-HFP. The conductivity also was enhanced with frequency due to the interfacial polarization effect [44]. At low frequency region, the conductivity depends on angular frequency and dielectric loss factor and the dipolar relaxation causes the interfacial polarization to dominate. However, at higher frequency region, the dipoles relax slowly and the applied electric field reduces the space charge accumulation [44].
Fig. 5

Frequency-dependent dielectric constant (a) dielectric loss (b), and conductivity (c) for PVDF-HFP composites; variation of dielectric constant and loss at various frequencies (d) and relaxation times (e) for all nanocomposites

The dielectric constant and loss values for all samples are compared at different frequencies in Fig. 5d. For the neat polymer, dielectric loss and dielectric constant values are close at lower frequencies. As expected, all the samples showed high dielectric constant values compared to the loss values, which is the essential criteria for typical energy storage materials. The dielectric relaxation can also be seen at the lower frequency side. The relaxation times for the samples were also calculated by taking the reciprocal of frequency at which the dielectric loss is minimum [44]. The decreasing trend observed for the samples with higher filler concentration indicates the dipolar polarization occurring in the nanocomposites.

3.4 Piezoelectric properties

Energy harvesting performance for all samples is tested by means of assembled experimental set up, which is represented in Fig. 6a. The PVDF-based test samples were excited by using an electrodynamic shaker system [48]. The setup consists of the test sample sandwiched between a conductive seismic mass on top and a conductive foil layer below. The seismic mass and the foil layer form the electrodes with wires connecting them to the shunt resistance box (not shown in the figure) voltage across the load resistance is measured using a data acquisition system. Beneath the foil layer and flat plate is an insulating layer that isolates the sample from the metal shaker platform, which oscillates harmonically in the vertical direction. An accelerometer is used to measure the motion of the shaker table through the data acquisition system. The sample electroded with silver on both sides is shown in Fig. 6b. When the sample is given a compressive force with the help of specific weight and the shaker, an electric potential difference is generated as shown in Fig. 6c. The piezoelectric response of the present quaternary phasic sample is a collective influence from the polymer as well as the filler.
Fig. 6

a Schematic illustration showing the experimental set up. b Photograph of the PVDF-HFP nanocomposite film. c Force responsive signal formation happening within the sample

The governing equation for the first order dynamics is
$$ C\ v(t)+\frac{v(t)}{R}={d}_{33}F(t)\ \mathrm{where}\ F(t)= ma(t) $$
where F is the force on the sample in thickness direction (due to base acceleration a and seismic mass m), C is the static capacitance, R is the load resistance, v is the voltage across the load resistance, d33 is the equivalent piezoelectric strain constant, and an over-dot represents the derivative with respect to time.
For harmonic excitation at frequency ω with base acceleration amplitude |A|, the magnitude of the complex voltage V across the load resistance per unit base acceleration is obtained from Eq. (4) as
$$ \left|\frac{V}{A}\right|=\frac{Rm{d}_{33}\omega }{\sqrt{{\left( RC\omega \right)}^2+1}} $$
and therefore the power output per base acceleration squared is
$$ \left|\frac{P}{A^2}\right|=\frac{R{\left(m{d}_{33}\omega \right)}^2}{{\left( RC\omega \right)}^2+1} $$

The frequency response function (FRF) that relates the voltage output to base acceleration is very useful since it can be used for parameter identification (i.e. d33 identification) at a given frequency, load resistance, and sample capacitance through Eq. (5) as well as performance comparison in power generation using Eq. (6).

It is also possible to define V/F (volts/force) FRF since |F| = |mA| in Eq. (4), yielding
$$ \left|\frac{V}{F}\right|=\frac{R{d}_{33}\omega }{\sqrt{{\left( RC\omega \right)}^2+1}} $$
and the power output FRF for force input becomes
$$ \left|\frac{P}{A^2}\right|=\frac{R{\left({d}_{33}\omega \right)}^2}{{\left( RC\omega \right)}^2+1} $$
We can further proceed to identify the figure of merit in power generation. At a given frequency, the harvested power output is maximized when the load impedance matches the output impedance of the device. For this case, the optimal load resistance is
$$ {R}_{opt}=\frac{1}{\omega C} $$
Substituting this into the power expression yields
$$ {\left(\frac{P}{A^2}\right)}_{max}=\frac{\omega {m}^2{d_{33}}^2}{2C} $$
which means that the figure of merit is
$$ {\left(\frac{P}{A^2}\right)}_{max}\infty \frac{{d_{33}}^2}{\varepsilon_{33}} $$
where ε33 is the overall permittivity of the film (since the capacitance is \( C={\varepsilon}_{33}{A}_e/h \) where Ae is the surface electrode area and h is the film thickness, i.e. the electrode spacing).
The capacitance values of the samples and the d33 constant values calculated are provided in Table 3. Maximum capacitance of 1050 pF and good piezoelectric response are observed for the sample containing the G-T/SrTiO3 filler combination in 1:2 ratio. This can be because of the uniform dispersion and the good crystalline behavior of the composite at this particular composition. While increasing the filler concentration, the piezoelectric power generation is expected to decrease. The proposed mechanism of piezo response of the samples can be the combined effect of PVDF dipoles and the surface charges present on the filler particles. The oppositely charged polar surfaces of the G-T and SrTiO3 interact with the PVDF dipoles and develops positive and negative charge densities on the nanocomposite and causes charge induced polarization. This increases the β-phase formation and in addition to this effect a stress induced polarization is also imparted by the external mechanical force [49]. This self-polarization of semi crystalline polymer nanocomposite is responsible for the observed piezoelectric power generation and this material can be applied to fabricate self-powered devices such as nanogenerators [32].
Table 3

Measured capacitance values and d33 values for the samples





803 pF

5.70 pC/N


930 pF

5.08 pC/N


987 pF

6.24 pC/N


1050 pF

7.52 pC/N


859 pF

5.94 pC/N

Figure 7 shows the generated voltages at different frequencies of vibration for all the polymer nanocomposite samples. The frequency was varied between 15 and 50 Hz and the piezoelectric output voltage was observed to increase towards higher frequencies of vibration. The corona poling done on the sample surface helped the dipoles to align so that good electrical signals are generated. Compared to the neat polymer (Fig. 7a), all nanocomposites show higher output voltages substantiating the contribution of crystallinity and reinforcement from the filler particles [32, 33]. The tri phasic filler combination showed high output voltages with the maximum of 2 V (Fig. 7d) in the case of PHP/GS12 sample. It is concluded that the SrTiO3 particles and their uniform distribution have special role in enhancing the mechanical energy harvesting property of PVDF-HFP.
Fig. 7

Piezo electrical output voltage generated at different frequencies of vibration a PHP, b PHP/GT1, c PHP/GS11, d PHP/GS12, e PHP/GS13, and f comparison of samples at 45 Hz frequency

Though the sample voltages were enhancing with the applied frequency of mechanical vinration, it is found that high dosage of vibrations are also not good. The output signals were actually decreasing when the frequencies crossed 50 Hz. Since substantial values of output voltages were obtained at 45 Hz of vibrational frequency, all samples were compared for their performances as given in Fig. 7f. The PHP/GS12 showed 10.5 times higher output voltage than the neat polymer (PVDF-HFP). Higher filler concentration ended up with non uniform dispersion and decreased output performance.

It is also significant that the piezoelectricity depends on the sample fabrication method and sample dimension. For instance, with the electrospun fibers of PVDF-HFP/cobalt-doped ZnO system, we have obtained 2.8 V output performance recently [32]; however, in this work, simple solution mixing is adopted, which is the reason for the comparatively lower voltage. In all the composite films studied here, around 0.1 g of the sample (circular with 2 cm diameter and 0.2 mm thickness) was analyzed.

The specific sample, PHP/GS12 was also checked for its mechanical stretchability and consistence of piezoelectric performance by subjecting to stretching, bending and twisting movements. The output voltages were in the range of 1–2 V as illustrated in Fig. 8. The pictures shown in Fig. 8a, b also prove the flexible nature of it. Since the sample showed consistent performance in piezoelectric studies, its application in flexible and wearable electronics can be proposed.
Fig. 8

a, b Foldable films of PHP/GS12 nanocomposite. c Output voltage obtained at bending, stretching, and twisting the PHP/GS12 sample

4 Conclusion

In conclusion, we have demonstrated the fabrication of PVDF-HFP hybrid nanocomposite applying the principle of filler synergy. A ternary filler combination of rGO, TiO2 and SrTiO3 is employed in PVDF-HFP following the simple method of solution casting. Dielectric properties reveal the energy storage capability of the nanocomposite and the dielectric constant enhanced with the concentration of the filler. The composites achieved good piezoelectric response and this was illustrated in terms of crystallinity derived from FTIR, XRD and DSC studies. A peak to peak output voltage of 2 V is obtained for the specific sample containing G-T/SrTiO3 filler combination at 1:2, which is 10.5 times higher than the neat PVDF-HFP. Finally, we believe that the designed flexible product at 1:2 filler composition will be very useful for manipulating functional materials particularly in the field of self-powering devices.


Funding information

This publication was made possible by NPRP grant 6-282-2-119 from the Qatar National Research Fund (a member of Qatar Foundation). The statements made herein are solely the responsibility of the authors.


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Copyright information

© Springer International Publishing AG, part of Springer Nature 2018

Authors and Affiliations

  1. 1.Center for Advanced MaterialsQatar UniversityDohaQatar
  2. 2.G. W. Woodruff School of Mechanical Engineering, Georgia Institute of TechnologyAtlantaUSA
  3. 3.Department of PhysicsB. S. Abdur Rahman Crescent Institute of Science and TechnologyChennaiIndia
  4. 4.Materials Science & Technology Program (MATS), College of Arts & SciencesQatar UniversityDohaQatar

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