Significant enhancement of the electroactive β-phase of PVDF by incorporating hydrothermally synthesized copper oxide nanoparticles

Abstract

The influence of copper oxide nanoparticles (CONPs) on the polymorphism of poly(vinylidene fluoride) (PVDF) was systematically investigated in this work. Copper oxide nanoparticles having an average diameter of 60 nm were synthesized by a simple, cost effective, environmentally benign modified hydrothermal method. Then a series of copper oxide nanoparticles incorporating flexible, self-standing PVDF films were prepared by the solution casting technique. The impact of the CONP loading on the structural and morphological properties of PVDF were studied by X-ray diffraction, scanning electron microscopy and Fourier transform infrared spectroscopy techniques. The thermal properties of the sample were investigated by differential thermal analysis, thermogravimetric analysis and differential scanning calorimetry techniques. The incorporation of CONPs leads to faster crystallization and thereby promotes the formation of electroactive β-phase enriched PVDF films. Strong interfacial interactions between the negatively charged nanoparticle surface and positively charged –CH₂ dipoles of the PVDF lead to a significant enhancement of the electroactive β-phase. The 5 wt% CONP–PVDF composite film exhibits a maximum β-phase fraction of 90% due to having the highest interfacial area between the well-dispersed nanoparticles’ surfaces and the polymer.

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1. Introduction

Recently, electroactive polymers have attracted much attention due to their various potential applications in energy harvesting, sensors, actuators and in the biomedical field.^1–4 Of the few polymers showing piezoelectric, pyroelectric and ferroelectric properties, such as Nylon-11, poly(lactic-co-glycolic acid) (PLGA) and polylactic acid (PLLA),^5–9 poly(vinylidene fluoride) (PVDF) and its copolymers have the most versatile electroactive properties. Therefore PVDF and its copolymers such as poly(vinylidene fluoride–trifluoroethylene) [P(VDF–TrFE)] and poly(vinylidene fluoride–hexafluoropropylene) [P(VDF–HFP)] are emerging as the popular polymers of choice for a growing number of potential applications.^10–13 Poly(vinylidene fluoride) ([–CH₂–CF₂–]_n), a flexible, cost-effective fluoropolymer, has received much interest due to its wide range of significant applications in piezoelectric nanogenerators, high charge storage capacitors, electro-striction for artificial muscles, magneto-striction, non-volatile memories in microelectronics, thin film transistors and pulsed lasers.^1,10–13 PVDF is a semicrystalline polymer having five different crystalline phases, i.e., α, β, γ, δ and ε.¹⁴ Among all of these phases, the non-polar α-phase is the most common and thermodynamically stable state at ambient temperature and pressure. On the other hand the β- and γ-phases are the polar phases of PVDF. The α-phase possesses a monoclinic unit cell structure with a TGTG′ (T – trans, G – gauche+, G′ – gauche) dihedral conformation, whereas the β- and γ-phases contain orthorhombic unit cell structures with TTT (all trans) and 3TG3TG′ conformations respectively.^14–17 However the β-phase has become more important than the other phases due to its better piezoelectric, ferroelectric and pyroelectric properties.^8,9,11,24 Thus, many research works regarding the enhancement of the electroactive β-phase content in PVDF have been reported for various applications of this particular polymer.^8,9,14,18,19 Different techniques, including the uniaxial or biaxial drawing of α-PVDF films, simultaneous poling, stretching and quenching the α-PVDF films, melting crystallization of α-PVDF under high pressure, the supercritical carbon dioxide technique and the electro-spinning method, are used to achieve a higher β-phase content in PVDF.⁹ The enhancement of the electroactive β-phase fraction in the PVDF matrix can also be achieved by the incorporation of various sub-micron or nanosized filler materials including metals, metal oxides, clays, ceramics, semiconductor oxides, carbon nanotubes, graphene, inorganic salts and ferrites.^{8,9,11,18–26} Since the sizes of fillers are in the nano-dimension, the interfacial interactions between the nanofillers and the polymer matrices increase significantly due to the high surface-to-volume ratio of the nanofillers. Thus the addition of a small amount of nanofiller improves the physicochemical properties of the polymer without compromising the polymer flexibility. Gold,¹⁸ silver,¹⁹ iron oxide and cobalt oxide²⁴ nanomaterials are used to enhance the electroactive β-phase content in PVDF. It is well-known that the use of metal and metal oxide nanoparticles can also significantly improve the mechanical, electrical and magnetic properties of polymeric composite films without losing the flexibility of the composite films.^18,20,21,25 Of the different metal oxide nanoparticle fillers, copper oxide nanoparticles are of paramount importance due to their key role in the enhancement of the electrical properties of PVDF.^20,25 A significant enhancement of the dielectric constant (ε′ ∼ 10³) of copper oxide–PVDF nanocomposite films has been previously reported.²⁰ It will be interesting to study the effect of the incorporation of comparatively cost-effective copper oxide nanoparticles (CONPs) on the nucleation of the electroactive β-phase in PVDF, as well as to study the structural and thermal properties of the copper oxide–PVDF nanocomposite films, which have not been reported till date.

This present work deals with the synthesis of copper oxide nanoparticles by a modified hydrothermal method. Hydrothermally synthesized copper oxide nanoparticles are incorporated into the PVDF matrix. Detailed analyses of the structural and thermal properties of the CONP loaded PVDF films have been included in this study. The effect of the copper oxide nanoparticles on the nucleation of the electroactive β-phase in PVDF, along with the reason behind the significant improvement of the β-phase fraction, has also been discussed from the physicochemical point of view.

2. Experimental

2.1 Materials

Poly(vinylidene fluoride) (PVDF) pellets (M_w = 27 500 g mol⁻¹) were obtained from Sigma Aldrich, USA. Dry N,N-dimethylformamide (DMF; Merck, India), copper(II) sulphate pentahydrate (CuSO₄·5H₂O, M_w = 249.685 g mol⁻¹; Merck, India), sodium hydroxide (NaOH; Merck, India) and ethyl alcohol (C₂H₆O; Merck, India) were used in this work. All the materials were used in the experiments without further purification.

2.2 Synthesis of copper oxide nanoparticles

The copper oxide nanoparticles were synthesized by a modified hydrothermal method.²⁷ In this typical procedure, initially a 0.1 M solution of CuSO₄·5H₂O was prepared in distilled water. A 1 M NaOH solution was then added drop-by-drop to the solution of CuSO₄·5H₂O with vigorous stirring until the pH of the resultant solution reached 13. After that, the solution was transferred to a stainless steel autoclave and fully sealed. The autoclave was heated in a furnace at a temperature of 110 °C for 18 hours. The autoclave was then set for natural cooling. After cooling to room temperature (30 °C), the obtained solution was centrifuged (4000 rpm) for 30 minutes. The as-obtained product was thoroughly washed several times using distilled water and ethyl alcohol to remove the residual impurities. Finally, the solution was dried to afford a powder and the powder sample was then kept in a vacuum desiccator for the complete removal of the residual solvent contained in the powder.

2.3 Synthesis of copper oxide loaded PVDF nanocomposite films

The copper oxide nanoparticle loaded PVDF films were synthesized by the simple solution casting method. In this typical synthesis procedure, 500 mg of PVDF was dissolved in 20 ml DMF with vigorous stirring at 60 °C and the complete dissolution of PVDF in DMF was achieved. Then a certain weight percent (1–5 wt%) of the as-synthesized copper oxide nanoparticles were added to the solution of PVDF and vigorously stirred for 16 hours followed by 30 min sonication to obtain a homogeneous mixture. The nanocomposite films were prepared by casting the mixture in a thoroughly cleaned and dried Petri dish, and the solvent was evaporated at 90 °C. Neat PVDF films were also prepared following the same procedure along with the loaded PVDF films. 1 wt%, 3 wt% and 5 wt% copper oxide nanoparticle loaded PVDF films were prepared in this work and are named PCO1, PCO3 and PCO5 respectively. The prepared neat PVDF is named P0.

2.4 Characterization methods

The structural properties of the as-synthesized CONPs and composite films were studied by X-ray diffractometry (Bruker-D8) with Cu-Kα radiation (wavelength 1.541 Å) using Bragg–Brentano goniometer geometry and the θ–2θ mechanism. The XRD patterns of all the samples are recorded with a scan speed of 0.3 s per step and under an operating voltage of 40 kV with 2θ varying from 10 to 50°.

Absorption spectra of the as-synthesized CONPs and CONP loaded PVDF films were recorded by UV-visible spectrophotometry (Jasco V-630 Spectrophotometer) in the wavelength range 200–1100 nm.

The surface morphology of the nanoparticles and the as-synthesized copper oxide nanoparticle doped PVDF films were studied by field emission scanning electron microscopy (FESEM) (Quanta, FEG 250). Energy dispersive X-ray spectroscopy (EDAX) of the samples was also carried out in the same FESEM instrument. The structure, shape and size distribution of the as-synthesized copper oxide nanoparticles were studied by transmission electron microscopy (TEM) using an FEI.TECHNAI.T-20 G2 SUPERTWIN (200 kV) instrument.

Vibrational spectra for all the samples were recorded at room temperature by Fourier transform infrared (FTIR) spectroscopy (Jasco FT/IR-460 PLUS) with a resolution of 1 cm⁻¹. The fraction of β-phase F(β) in the as-synthesized nanocomposite films was calculated from the FTIR spectra using the Beer–Lambert Law as given by,¹⁴

where A_α and A_β are the absorbance at 764 cm⁻¹ and 840 cm⁻¹ respectively. K_α (6.1 × 10⁴ cm² mol⁻¹) and K_β (7.7 × 10⁴ cm² mol⁻¹) are the absorption coefficients at 764 cm⁻¹ and 840 cm⁻¹ respectively.¹⁴

Differential thermal analysis (DTA) and thermogravimetric analysis (TGA) were performed using TGA/SDTA851 Mettler Toledo apparatus in an air atmosphere at a heating rate of 10 °C min⁻¹. The crystallization and melting behavior of the pure and copper oxide nanoparticle loaded PVDF films were analyzed using a differential scanning calorimeter (DSC-60, Shimadzu (Asia Pacific) Pte. Ltd, Singapore). All the samples were heated from 30 °C to 200 °C at a rate of 10 °C min⁻¹ in a nitrogen gas atmosphere. The degree of crystallinity (X_c) of the samples was calculated using the following equation:²⁸

X_c = ΔH_c/ΔH_100% (2)

where ΔH_c is the enthalpy of crystallization and ΔH_100% is the melting enthalpy of 100% crystalline PVDF with a value 104.6 J g⁻¹.

3. Results and discussion

3.1 Copper oxide nanoparticles

X-ray diffraction analysis (XRD). Fig. 1 illustrates the X-ray diffraction (XRD) pattern of the as-synthesized copper oxide nanoparticles, and reveals the highly crystalline nature of the sample. The diffraction peaks positioned at 2θ values 35.3°, 38.5°, 52.3°, 61.2°, 66.0°, 67.9° and 71.7° correspond to the (1̄11), (111), (020), (1̄13), (3̄11), (113) and (311) planes of the crystalline phase of monoclinic cupric oxide (CuO) in accordance with JCPDS no. 41-0254 (a = 4.685 Å, b = 3.423 Å, c = 5.132 Å, space group: C12/c1). The peaks positioned at 2θ values 13.7°, 16.8°, 21.3°, 27.6°, 32.3° are assigned to the (110), (111), (220) and (311) planes of the crystalline phase of cubic cuprous oxide (Cu₂O) (JCPDS no. 34-1354, a = 4.217 Å, space group: Pn3̄m). The XRD analysis indicates that the as-synthesized copper oxide nanoparticles are in the mixed crystalline phases of CuO and Cu₂O with a dominating fraction of the CuO phase.

Fig. 1 X-ray diffraction pattern of the hydrothermally synthesized copper oxide nanoparticles.

Electron microscopy analyses. The field effect scanning electron micrograph (FESEM) of the as-synthesized copper oxide nanoparticles is shown in Fig. 2a. The inset of the figure shows the particle size distribution curve. The figure shows that the platelet-like particles with an average size of ∼170 nm are stacked together. Fig. 2b shows the transmission electron microscopic (TEM) micrograph of the as-synthesized copper oxide nanoparticles. The size distribution curve obtained from the TEM micrograph is shown in the inset of Fig. 2b. The figure shows that the copper oxide nanoparticles possess a hexagonal like structure with well-defined boundaries. As seen from the size distribution curve the average diagonal length of these nanoparticles is ∼60 nm. Fig. 2c shows the selected area electron diffraction (SAED) pattern of the copper oxide nanoparticles. The single crystalline nature of the nanoparticles is evident from the ‘dot’ patterned SAED result. The measured d-spacing values from the SAED pattern are 0.116 nm, 0.150 nm and 0.148 nm, which confirms that the synthesized copper oxide sample is in the mixed crystalline phase of CuO and Cu₂O as the d-spacing values are in good agreement with those of CuO and Cu₂O crystals (JCPDS no. 34-1354 and 35-1091). Fig. 2d exhibits the EDAX result of the hydrothermally synthesized sample, which shows the presence of only the characteristic peaks corresponding to O and Cu elements, confirming the high purity of the synthesized copper oxide nanoparticles.

Fig. 2 (a) FESEM image of the synthesized copper oxide nanoparticles. Inset shows the size distribution curve of the copper oxide nanoparticles; (b) TEM image of the synthesized copper oxide nanoparticles. Inset shows the size distribution curve of the copper oxide nanoparticles; (c) SAED pattern of the as-synthesized copper oxide nanoparticles; (d) EDAX result of the as-synthesized copper oxide nanoparticles.

UV-visible spectroscopy. The UV-visible absorption spectrum of the as-synthesized copper oxide nanoparticles as shown in Fig. 3 shows a strong absorption peak positioned at a wavelength of 258 nm. The presence of this absorption peak can be attributed to the existence of the cuprous oxide (Cu₂O) phase of the CONPs. Similarly, the existence of an absorption peak for cuprous oxide has already been reported by Abboud et al.²⁹ Another broad absorption peak centered at ∼650 nm can also be seen in Fig. 3. The presence of the cupric oxide (CuO) phase in the synthesized CONPs is responsible for this broad absorption peak.²⁹ Thus, the UV-visible absorption spectrum of the copper oxide nanoparticles also confirms the co-existence of both the CuO and Cu₂O phases in the CONPs.

Fig. 3 UV-visible absorption spectrum of the as-synthesized copper oxide nanoparticles.

3.2 Copper oxide loaded PVDF nanocomposite films

X-ray diffraction analysis (XRD). X-ray diffraction (XRD) spectroscopy was used to analyze the polymorphism of the neat PVDF and copper oxide nanoparticle loaded PVDF films. Fig. 4a shows the XRD patterns of the neat PVDF (P0) and PVDF films loaded with different concentrations of CONPs, i.e., PCO1, PCO3 and PCO5. As seen from the figure, the diffraction pattern of neat PVDF reveals the semi-crystalline nature of the polymer. The peaks positioned at 2θ = 17.5° (100), 18.2° (020), 19.8° (021), and 26.6° ((201), (310)) of the neat PVDF can be assigned to the non-polar α-phase of PVDF.^14,23 A well-defined peak for the loaded PVDF films positioned at 2θ value 35.3° (1̄11) can be readily indexed to the monoclinic CuO phase (JCPDS-41-0254) of the CONPs, whereas peaks at 2θ values 13.57° (110), 22.54° (200) and 33.17° (311) can be assigned to the cubic Cu₂O phase (JCPDS-34-1354) of CONPs. The occurrence of these peaks attributed to the different phases of CONPs indicates the successful incorporation of CONPs in the PVDF matrix. The XRD patterns also show that with the increasing loading fraction of the nanoparticles, the intensity of the peaks corresponding to the α-phase of PVDF (positioned at 17.5°, 18.2°, 19.8° and 26.6°) decreases gradually. However, a new peak positioned around 2θ = 20.6° ((020), (101)), which is the characteristic peak of β-phase PVDF, appears in the copper oxide nanoparticle loaded PVDF films.^14,23 A closer observation of the XRD patterns reveals that for sample PCO5, the intensity of the peak positioned at 26.6° corresponding to the non-polar α-phase of PVDF is almost completely diminished, whereas the relative intensity of the peak positioned at 2θ = 20.6° (characteristic peak of β-phase PVDF) becomes maximized. It can be also seen from the XRD pattern that a peak positioned at 2θ = 38.9° (211) corresponding to the γ-phase of PVDF^14,23 is present for the neat as well as the loaded PVDF films. The ratio I_20.6°/I_18.2° is a measurement of the α- and β-phase content in the neat PVDF and CONP incorporating nanocomposite films, where I_20.6° and I_18.2° are the intensities of the peaks positioned at 20.6° (characteristic peak for the β-phase) and 18.2° (characteristic peak for the α-phase) respectively. Fig. 4b represents the content dependence of the ratio (I_20.6°/I_18.2°). For the neat PVDF film the ratio is about 0.93 and this ratio increased with the increasing weight fraction of the CONPs in the PVDF matrix. The maximum value of the ratio is ca. 12.05, obtained for sample PCO5 (5 wt%). Thus, as a whole, the XRD patterns of the loaded PVDF films confirm that the incorporation of the CONPs results in a phase transformation from the α-phase to the electroactive β-phase of PVDF.

Fig. 4 (a) X-ray diffraction patterns of the neat PVDF (P0) and CONP loaded PVDF films (PCO1, PCO3, PCO5); (b) ratio of intensities at 20.6° and 18.2° of samples P0, PCO1, PCO3 and PCO5.

Field emission scanning electron microscopy (FESEM). Fig. 5 exhibits the FESEM micrograph of the neat PVDF. The surface morphology of the neat PVDF shows that the matrix possesses a highly compact and uniform morphology. In addition to that, a well-defined spherulite (marked with a red circle) with a diameter of ∼8 μm is evident in the FESEM micrograph. In Fig. 6a, the low magnification FESEM micrograph of sample PCO5 shows the surface morphology of the film, which reveals a uniform distribution of the spherulites (marked as red circles in the figure) with diameters of 2–3 μm throughout the surface of the composite film.

Fig. 5 FESEM micrograph of the neat PVDF film. Red circle shows the spherulite.

Fig. 6 (a) FESEM micrograph of the CONP loaded PVDF film (PCO5) at low magnification. Red circles show the spherulites; (b) FESEM micrograph of sample PCO5 at high magnification; (c) EDAX analysis of sample PCO5.

The uniform distribution of the spherulites confirms the good, homogeneous dispersion of the copper oxide nanoparticles in the PVDF matrix. The smaller size of the spherulites for sample PCO5 in comparison to that of the neat PVDF represents the faster nucleation kinetics in the CONP–PVDF composite films²⁸ due to the presence of additional CONP nucleation centers, and the enhancement of the β-phase fraction may be achieved in the composite film due to the existence of the large number of favorable nucleation centers.²⁸ Again the presence of the added CONPs in the PVDF film can be clearly seen from the high magnification FESEM micrograph of sample PCO5 in Fig. 6b. It is evident from the figure that the surface of the PVDF film is completely modified with the addition of CONPs. Impinging spherulites (marked by straight red lines) are observed in the figure due to the faster kinetics in the presence of CONPs. Thus the FESEM images for both the neat and composite films support the enhancement of the electroactive β-phase due to the incorporation of the CONPs which is consistent with the result obtained from the X-ray diffraction patterns of the samples. Fig. 6c represents the EDAX result of the CONP loaded PVDF film (PCO5). The figure shows the presence of characteristic peaks corresponding to Cu, O, C and F elements where C and F are the polymeric components of the PVDF matrix, and Cu and O are the components of the CONPs as previously shown in Fig. 2d. Hence the presence of copper oxide in PVDF is further confirmed by the EDAX result.

UV-visible spectroscopy. Fig. 7 shows the UV-visible absorbance spectra of the neat PVDF along with the CONP loaded PVDF films (PCO1, PCO3 and PCO5) within the range 200–1100 nm. The UV-visible spectrum of the neat PVDF infers that it is transparent in the higher wavelength region starting from the visible region. However, the presence of a well-defined absorption peak positioned at a wavelength of ∼270 nm for the CONP loaded PVDF films is evident from the figure. It is worthwhile to mention here that the absorption peak for the CONPs is positioned at a wavelength of ∼258 nm as seen from Fig. 3. Thus, the UV-visible spectra for the loaded PVDF films reveal that the absorption peak for the CONPs is being shifted in the presence of the host polymer matrix. This observation may be explained by the presence of the host polymer matrix surrounding the CONPs, which may affect the electronic transition in the valence band and conduction band in the copper oxide nanoparticles due to the strong electrostatic interactions of the CONP filler with the polymer matrix.⁹ The figure also shows that the intensity of the absorption peaks is increased with the increasing loading fraction of the copper oxide nanoparticles in the PVDF matrix which may be attributed to the homogeneous dispersion of the nanoparticles in the PVDF matrix.⁹

Fig. 7 UV-visible absorption results of neat PVDF (P0) and the CONP loaded PVDF films (PCO1, PCO3 and PCO5).

Fourier transform infrared spectroscopy. Fig. 8a shows the recorded Fourier transform infrared (FTIR) absorption spectra of the neat PVDF and CONP loaded PVDF films, i.e., PCO1, PCO3 and PCO5. The FTIR spectrum of the neat PVDF shows characteristic peaks at 488 cm⁻¹ (CF₂ wagging), 532 cm⁻¹ (CF₂ bending), 615 cm⁻¹ and 764 cm⁻¹ (CF₂ bending and skeletal bending), 796 cm⁻¹ and 976 cm⁻¹ (CH₂ rocking) corresponding to the IR bands of the non-polar α-phase of PVDF.⁸ Two small peaks corresponding to the β-phase of PVDF⁸ are also observed at 510 cm⁻¹ (CF₂ stretching) and 840 cm⁻¹ (CH₂ rocking, CF₂ stretching and skeletal C–C stretching) for the neat PVDF. It can be clearly seen from the figure that all the characteristic absorption bands corresponding to the non-polar α-phase of PVDF are gradually diminished with the increasing loading fraction of the CONPs in PVDF. Moreover, the relative intensities of the peaks corresponding to the β-phase increase with increasing CONP loading. A closer observation shows that for sample PCO5 all the peaks corresponding to the α-phase have almost completely disappeared and only the characteristic absorbance bands assigned to the polar β-phase are dominant, appearing at 510 cm⁻¹, 600 cm⁻¹ (CF₂ waging) and 840 cm⁻¹.⁸ Therefore, the FTIR results strongly indicate an α to β phase transformation of PVDF with the loading of the CONPs in the PVDF matrix. To quantitatively study the phase transformation for the CONP loaded PVDF films, the fraction of β-phase (F(β)) was calculated for the nanocomposite films using eqn (1). The variation of F(β) with CONP content is shown in Fig. 8b. A rapid increment in the F(β) value for sample PCO1 is evident from the figure, then F(β) increases almost linearly with the increasing filler content of the CONPs. The maximum β-phase nucleation occurs for the 5 wt% CONP loading, and is ∼90%. This enhancement of the β-phase fraction with the filler loading may be readily explained by the interaction between the filler and the polymer matrix. At low filler content the amount of interfacial area between the polymer and the CONPs is low, while the interfacial area increases with the increasing content of CONPs in PVDF as the CONPs are homogeneously dispersed in the polymer matrix. As a consequence, the number of aligned chains having the ‘all trans’ conformation increases, resulting in an increase in the β-phase fraction. The variation of F(β) also suggests that the most intimate interaction between the polymer and the CONPs occurs at 5 wt% loading of the filler.

Fig. 8 (a) FTIR spectra of the P0 and CONP loaded PVDF films (PCO1, PCO3 and PCO5). Dotted red and black lines correspond to the β-phase and α-phase of PVDF respectively; (b) variation of the β-phase fraction (F(β)) with CONP loading in the PVDF matrix.

Thermal properties. The thermal properties of neat PVDF and CONP loaded PVDF films have been investigated by TG-DTA techniques. Fig. 9a shows the DTA results of the neat PVDF and the PVDF nanocomposite films in the temperature range 30 °C to 600 °C. An enlarged view of the DTA results in the temperature range 140 °C to 170 °C is shown in the inset of the same figure. The figure reveals the presence of an endothermic peak positioned around 155.8 °C for the neat PVDF, which is attributed to the melting temperature of the neat PVDF.^8,21 It can also be seen from the figure that, with the increasing loading fraction (1–5 wt%) of the nanofillers (CONPs) the melting temperature of the composite films is increased in comparison to that of the neat PVDF film. The maximum shift (3 °C) in the melting temperature is observed for sample PCO5. A similar increase in melting temperature (5–7 °C) with the loading of the nickel particles (0 to 2.0 wt%) in the PVDF matrix was previously reported by Levedev et al.²¹ This increase in the melting temperature may be attributed to the change in the degree of crystallinity and homogeneity in the sub-molecular structure of the nanocomposite film compared to the neat films.^8,21 The presence of copper oxide nanoparticles in the PVDF matrix is responsible for such changes in the melting temperature of the nanocomposite films. The increment in melting temperature of the nanocomposite film also confirms the formation of the electroactive β-phase in the nanocomposite films.^8,21 An approximate 2 °C shift of the main exothermic peak for sample PCO5 in comparison to the neat PVDF film can be observed from the DTA curve (Fig. 9a), which infers the good thermal stability of the nanocomposite films. Fig. 9b shows the TGA analysis of the neat, PCO1 and PCO5 films. The figure shows a significant, step-like weight loss, which starts around 450 °C and continues up to 500 °C for all the samples. This weight loss is related to the decomposition of the films by a stripping mechanism, during which hydrogen fluoride (HF) gas is released from the samples. During this degradation process, about 40% of hydrogen fluoride gas was found to be released in this temperature range from the neat film, as calculated from the TGA curve.⁸ The TGA curves in Fig. 9b also confirm a fair enhancement in the decomposition temperature of the nanocomposite films, leading to the improved thermal stability of the films. The inset of the Fig. 9b shows the TGA curve of the synthesized neat copper oxide nanoparticles. The copper oxide nanoparticle powder undergoes almost no weight loss in the temperature region 30–600 °C except for the loss of water due to the formation of only metal oxide.³⁰

Fig. 9 (a) DTA of the samples P0, PCO1 and PCO5. Inset shows an enlarged view of the DTA curve in the temperature range 140–175 °C; (b) TGA thermographs of the samples P0, PCO1 and PCO5. Inset shows the TGA thermographs of the synthesized copper oxide nanoparticles.

To determine the kinetic parameters for the thermal degradation of the nanocomposite films the Coats–Redfern equation³¹ is usually used, which is as follows:

where E_d is the activation energy, R is the universal gas constant (8.814 J mol⁻¹ K), A is the pre-exponential factor, B is the heating rate (10 °C min⁻¹), T is the degradation temperature (K) and α_d is the degree of conversion. The degree of conversion (α_d) is denoted as the kinetics of the conversion of the polymer to the volatile decomposition product. At a given temperature α_d is defined by,³¹where m_i, m_T, and m_f are the initial sample weight, weight of the sample at temperature T and weight of the final sample respectively. In this work the kinetic parameters of thermal decomposition for samples P0, PCO1 and PCO5 are calculated around the maximum degradation temperature (where the maximum weight loss has occurred). In Fig. 10a–c, the plots of ln[−ln(1 − α_d)/T²] against 1/T (K⁻¹) for samples P0, PCO1 and PCO5 give linear curves with a negative slope. From the slope of the curves, the value of the decomposition activation energy (E_d) of the neat and nanocomposite films has been calculated. The calculated values of the decomposition activation energy and the corresponding values of the regression coefficient (R²) for the neat and loaded films are given in Table 1. In this context it should be noted that the value of the regression factor (R²) is close to unity for all of the samples. The activation energy of the loaded films is higher than that of the neat PVDF films, which infers that the loading of the CONPs in the PVDF matrix increases the thermal stability of the nanocomposite films in comparison to the neat PVDF. This enhancement in thermal stability for the nanocomposite films is also in fair agreement with the DTA-TGA results. The maximum activation energy value is obtained for sample PCO5, and is 654.22 kJ mol⁻¹. A similar increase in thermal stability by incorporatingWO₃·H₂O nanoparticles in the PVDF matrix has been reported by Thakur et al.⁹

Fig. 10 Coats–Redfern plot of (a) P0; (b) PCO1; and (c) PCO5.

Table 1 Activation energy at the major degradation region (around T_max) for the as-synthesized samples

Name of the sample	Maximum degradation temperature (Tmax) (°C)	Activation energy at the major degradation region (Ed) (kJ mol^−1)	Regression factor (R^2)
P0	478.7	545.64	0.9998
PCO1	480.9	628.41	0.9997
PCO5	477.7	654.22	0.9995

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The interface between the polymer and nanofillers plays a key role in determining the final properties of the nanocomposite films. From the TGA results, the nanofiller–polymer interface region (m_I) of the CONP–PVDF composite was obtained by using the following equation,³⁷

where m_I0 is the mass of the pristine PVDF at the temperature at which the mass loss rate is maximum and m(x)_I0 is the mass of the composite containing a given wt% of nanoparticles that has not degraded at the same temperature. Fig. 11 shows the variation of the mass fraction (m_I) of the polymer located at the interface with the loading fraction (1–5%) of the CONP nanofiller in PVDF. The figure shows a rapid increment in the m_I value for sample PCO1, then m_I increases almost linearly with the increasing loading fraction of CONPs. The maximum value of m_I (20%) is obtained for sample PCO5 (5 wt% CONPs loaded PVDF). As a result the number of particles interacting with the polymer matrix is highest for PCO5. A higher m_I value reflects the increment in the interaction between the partially positive CH₂ bonds of the PVDF chains and the electrostatically negatively charged CONPs, which again promotes β-phase nucleation. This fact is also supported by the FTIR and XRD results which show the highest β-phase fraction is obtained for sample PCO5. It is worthwhile to mention here that the CONPs are not agglomerated in the polymer matrix up to a filler loading of 5 wt%. Rather, they are discretely distributed in the polymer matrix leading to a higher m_I value at a higher filler loading fraction.

Fig. 11 Variation of nanoparticle–polymer interface (m_I) with CONP content.

The thermal properties and polymorphism of the composites were also studied using DSC. Fig. 12a–c show the DSC heating cycles of the samples P0, PCO1 and PCO5 respectively. To get a deeper insight into the coexistence of different PVDF phases in the composite films, the DSC curves were fitted with multiple Gaussian functions corresponding to α-, β- and γ-phase PVDF. The multi-peak fitted DSC heating curves reveal the presence of three distinguishable peaks, indicating the coexistence of α-, β- and γ-phases in the neat as well as the PVDF nanocomposite films. The melting temperature for all the samples corresponding to the three different phases are shown in Table 2. The figures show that the relative intensity and the area under the peak related to the melting temperature of the β-phase increase with the increasing filler concentration of the CONPs. On the other hand, the intensity and the area under the curve assigned to α-PVDF diminishes significantly with the increasing loading fraction of the CONPs in the PVDF matrix. This observation can be explained by the increment in the β-PVDF nucleation with the expense of a reduction in the α-PVDF fraction in the composite films. The transformation from the α to β polymorph occurs due to the incorporation of CONPs. It can be seen from Fig. 12c that for sample PCO5 the relative intensity of the peak related to the β-phase of PVDF is greatest, while the peak related to α-phase PVDF is almost completely diminished. It is worthwhile to mention here that the XRD and FTIR analyses of the neat and PVDF nanocomposite films also validate the presence of the maximum β-phase fraction in sample PCO5. The DSC cooling cycles for neat PVDF and copper oxide nanoparticle loaded PVDF films are shown in Fig. 13. Though multiple endothermic peaks can be seen from the heating cycle of the neat and CONP loaded PVDF films, the cooling cycles of the samples show a single exothermic peak. This indicates the existence of polymorphs in the films. However, a single exothermic peak is obtained for the cooling cycle as, after the first heating cycle in DSC, the processing history is erased. The single exothermic peak present in the DSC cooling cycle is attributed to the crystallization temperature (T_c) of the neat and loaded PVDF films. The crystallization temperature for the CONP incorporated PVDF films is higher than that of neat PVDF as observed from Fig. 13 and it increases with the increasing content of CONPs in the PVDF matrix. This observation is expected as the incorporated copper oxide nanoparticles act as the nucleating agent in the PVDF matrix, thereby influencing the crystallization temperature. For sample PCO5 the crystallization temperature is increased by ∼2 °C in comparison to that of the neat PVDF. This indicates that the addition of well-dispersed CONPs as the nucleating agent lowers the free energy barrier of nucleation, thus accelerating the crystallization of PVDF, which manifests in a higher crystallization temperature for the nanocomposite films.²⁸ The faster nucleation kinetics is also evident from the FESEM micrograph of PCO5 (Fig. 6a) where smaller spherulites are observed in comparison to neat PVDF. DSC thermographs were also used to obtain the enthalpy of crystallization (ΔH_c J g⁻¹) and the degree of crystallinity (X_c) of the neat PVDF as well as the CONP–PVDF composite films. The degree of crystallinity of the samples was calculated using eqn (2). The crystallization temperatures (T_c), enthalpy of crystallization (ΔH_c), and the degree of crystallinity (X_c) for samples P0, PCO1 and PCO5 are reported in Table 3. The crystallinity decreases with the addition of CONPs to the PVDF matrix and PCO5 shows the minimum value of X_c. As is evident from the XRD and FTIR results, neat PVDF contains three phases, α, β and γ, but for sample PCO5 the α-phase of PVDF is almost completely diminished and the only present dominant crystalline phase is β-PVDF, resulting in a reduction in the overall crystallinity of the composite films. This indicates the intimate interaction of the CONPs with PVDF, and thus the strong influence of the nanoparticles on the structural modification of PVDF matrix.

Fig. 12 DSC heating cycles for samples (a) P0; (b) PCO1; and (c) PCO5. The DSC curves are fitted with multiple Gaussian functions corresponding to α-, β- and γ-phase PVDF.

Table 2 Melting temperature of the neat and CONP loaded PVDF films containing three different phases, α, β and γ

Sample name	Melting temperature of α-PVDF (Tm)α (°C)	Melting temperature of β-PVDF (Tm)β (°C)	Melting temperature of γ-PVDF (Tm)γ (°C)
P0	156.0	161.5	164.6
PCO1	158.3	162.8	166.7
PCO5	153.5	162.4	166.9

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Fig. 13 DSC cooling cycle for samples P0, PCO1 and PCO5.

Table 3 Crystallization temperature (T_c), enthalpy of crystallization (ΔH_c) and degree of crystallinity (X_c) of the neat and CONP loaded PVDF films

Sample name	Crystallization temperature (Tc) (°C)	Enthalpy (ΔHc)	Crystallinity (Xc) (%)
P0	134.6	36.9	35
PCO1	136.0	34.8	33
PCO5	136.8	26.7	26

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3.3 Mechanism of electroactive β-phase nucleation

The results obtained from the XRD, FTIR, and thermal studies of the nanocomposite films confirm that the incorporation of copper oxide nanoparticles in the PVDF matrix leads to an increase in electroactive β-phase nucleation. Neat PVDF films primarily contain α-phase PVDF and the transformation from the α to the electroactive β polymorph in PVDF is possible due to the incorporation of CONPs as well as the homogeneous dispersion of these CONPs in the PVDF matrix. Therefore, in this context it is meaningful to analyze the interaction between the copper oxide nanoparticles and the polymer, which promotes the formation of the electroactive β-phase. The improvement of the electroactive β-phase fraction may occur due to the strong interactions between the PVDF matrix and embedded nanoparticles.^8,9,24 This type of ion–dipole interaction between the PVDF matrix and the loaded ferrite nanoparticles to form the electroactive β-phase has also been reported by Martins et al.³² During the synthesis of the copper oxide nanoparticles we maintained the pH of the copper precursor solution at 13 by the use of NaOH solution. The variation of the zeta potential of the copper oxide nanoparticles with the pH of the respective precursor solution has been reported by Hassan et al.³³ This variation of zeta potential confirms that the surface of the as-synthesized copper oxide nanoparticles are negatively charged. When the negatively charged nanoparticles are added to the PVDF matrix in the solution phase, the nanoparticles will act as substrates for the formation of the electroactive β-phase. The partially positive –CH₂ dipoles of the PVDF chains experience strong electrostatic interactions with the negatively charged nanoparticle surfaces. The interaction leads to the alignment of the stabilized PVDF chains on the surface of the nanoparticles in a longer ‘all-trans’ (TTT) conformation resulting in the formation of the electroactive β-phase. Thus the surface of the nanoparticles acts as the nucleation center for the formation of the electroactive β-phase. Fig. 14 depicts the schematic of the ion–dipole interaction mechanism between the nanoparticles and the polymer chains leading to the formation of the electroactive β-phase. It is worthwhile to mention here that our experimental observations show the presence of a small fraction of the electroactive γ-phase of PVDF for the loaded PVDF films. Similar results have also been reported by Lopes et al. for aluminosilicate–PVDF composites.^34–36 This γ-phase formation may occur due to the gauge effect which may develop due to the easier local internal chain rotation.

Fig. 14 Proposed schematic of the β-phase formation mechanism.

4. Conclusion

In this work, a series of flexible CONP–PVDF composite films were synthesized by a simple solution casting technique and the effect of the CONPs on the polymorphism of PVDF was rigorously studied. The hexagonal copper oxide nanoparticles with an average size of 60 nm were synthesized by the modified hydrothermal method. Our study confirms the successful inclusion of well-dispersed copper oxide nanoparticles in the PVDF matrix. XRD, FESEM, FTIR and thermal analyses of the composite films confirm that the loading of copper oxide nanoparticles in PVDF leads to the successful transformation from the α to the electroactive β polymorph in PVDF. The β-phase fraction in PVDF increases with the increasing CONP loading in the polymer matrix. PCO5 shows the highest β-phase fraction of ca. 90%. The enhancement in the β-phase fraction of PVDF is attributed to the interactions of the negatively charged nanoparticle surfaces and the positive –CH₂ groups of the PVDF polymer chains, which leads to the alignment of the stabilized, longer ‘all-trans’ (TTT) conformation of the PVDF chains on the surface of the nanoparticles, resulting in the nucleation of the electroactive β-phase. The calculated nanoparticle–polymer interface is also maximal for PCO5. The highest rate of enhancement (with nanofiller loading) in the β-phase fraction and nanoparticle–polymer interface values is observed for the 1 wt% CONP loaded PVDF composite film. However, the values are further enhanced with increasing CONP filler content in the PVDF matrix. This indicates that the CONPs are nicely dispersed in the PVDF matrix up to 5 wt% loading. The FESEM micrographs show that the size of the spherulites is decreased for the CONP loaded PVDF films in comparison to the neat PVDF, which is facilitated due to the faster nucleation kinetics in PVDF due the presence of the homogeneously dispersed CONP nucleating agent. This phenomenon is also supported by the DSC results. The faster nucleation kinetics help to enhance the β polymorph in the nanoparticle loaded PVDF. The thermal stability of the composite films is not degraded with the CONP loading. Rather, the melting temperature is increased by ∼3 °C for PCO5 in comparison to neat PVDF. In this context it should be noted that the copper oxide nanoparticles are synthesized by a simple, cost effective and environmentally benign hydrothermal method. Thus the as-synthesized copper oxide nanoparticle loaded PVDF films with a significantly enhanced β-phase fraction may be used in future cost effective, efficient piezoelectric devices.

Acknowledgements

The authors wants to acknowledge DST INSPIRE, Government of India (IF140204, IF140209) for the financial support and Dr Sukhen Das for providing different characterization facilities.

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