Improving the thermal stability, electroactive β phase crystallization and dielectric constant of NiO nanoparticle/C–NiO nanocomposite embedded flexible poly(vinylidene fluoride) thin films

Abstract

Using a simple chemical precipitation process followed by sintering at 400 °C, NiO nanoparticles (NPs) and C–NiO nanocomposites (NCs) have been synthesized and characterized by X-ray diffraction, UV-visible spectroscopy, zeta potential measurement, field emission scanning electron microscopy and energy dispersive X-ray spectroscopy. Then, NiO NP or C–NiO NC loaded PVDF thin films were fabricated via a simple solvent casting or solution casting method. Thermogravimetric analysis confirmed good thermal stability of the nanocomposite films. Strong ion–dipole interaction between the negative surfaces of NiO NPs or C–NiO NCs and –CH₂ dipoles of polymer chains leads to the formation of a long stabilized TTTT conformation, i.e. formation of large number of electroactive β polymorphs in the modified PVDF thin films. Detailed study of the dependency of the dielectric properties on filler content (NiO NPs/C–NiO NCs) and frequency illustrates significant increase in the dielectric constant in the three phase C–NiO NCs–PVDF system than in the two phase NiO NPs–PVDF system. The dielectric constant is found to be as large as 317.4 at 20 Hz with a relatively low tangent loss value, and good flexibility when 20 mass% C–NiO NCs is incorporated in the PVDF matrix. These results have been explained in terms of Maxwell–Wagner–Sillars interfacial polarization at the NiO NPs/C–NiO NCs and insulating polymer matrix interface, evolution of a conductive network and formation of a microcapacitive structure in the NiO NPs or C–NiO NCs modified PVDF thin films.

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1. Introduction

Recently, electroactive or ferroelectric polymer nanocomposite (NC) thin films with high dielectric constant have drawn great attraction as a suitable candidate for next generation energy harvesters, piezoelectric nanogenerators, capacitive sensors, energy storage devices, photovoltaic self-charging cells, non-volatile memory, as well as in biomedical applications such as artificial muscles and smart skins.^1–8 Low cost, biocompatibility, non-toxicity, light weight, simple solution-based processes and high flexibility of ferroelectric polymers such as poly(vinylidene fluoride) (PVDF) ([–CH₂–CF₂–]_n) and its copolymers such as poly(vinylidene fluoride-hexafluoropropylene) [P(VDF-HFP)], poly(vinylidene fluoride-trifluoroethylene) [P(VDF-TrFE)], poly(vinylidene fluoride-chloro-trifluoroethylene) (PVDF-CTFE) etc. have been actively explored to develop electroactive and high dielectric nanocomposite (NC) thin film for versatile application in the field of energy systems and healthcare.^7,9

PVDF is a semicrystalline plastic polymer with at least five crystallite polymorphs α, β, γ, δ and ε. The most popular and thermodynamically stable phase is nonpolar α-phase. The nonpolar α-phase has a monoclinic unit cell with TGTG' (T-trans, G-gauche⁺, G′-gauche) dihedral conformation. The polar crystallite polymorph β and γ possess a orthorhombic unit cell with TTTT i.e. all trans and TTTGTTG' conformation respectively. The β phase exhibits maximum dielectric, piezoelectric, ferroelectric and pyroelectric property than other existing polymorphs. The δ and ε phases are basically the polar and anti-polar version of α and γ phases respectively.^7,10

The large spontaneous polarization, piezoelectricity, pyro-electricity and dielectric properties of β phase of PVDF encourage researchers to convert the nonpolar α phase to electroactive β phase in PVDF and to develop high dielectric PVDF thin films. Generally, melt crystallization of PVDF results in nonpolar α polymorphs and electroactive β phase is achieved by stretching the α-PVDF films.^11–13 Mechanical stretching is not preferable for fabrication of microelectronic devices because this process is not suitable for development of thin films directly on substrates.¹⁴ Alternatively, polar β phase can also be attained by solution evaporation method using appropriate solvents (e.g. from dimethyl formamide) at low temperature (below 70 °C) which smoothens the thin film deposition on desired substrate. But the film obtained in this process shows porosity which leads to a non-transparent appearance and a reduction of the electro-mechanical properties.¹⁵ Recently, improvement of electroactive β phase nucleation as well as good dielectric properties have been observed by the impregnation of different micro or nano-fillers such as clays,^16–19 ceramic nanoparticles (NPs),²⁰ transition metal oxides NPs,^9,21 metal NPs,^22,23 metal hydroxide NPs,²⁴ hydrated salts,^25–27 carbon nanotube,^28,29 carbon nanofiber,³⁰ ferrites,^31,32 hydrated metal oxide nanoparticle³³ etc. The incorporation these types of modifiers into the PVDF matrix effectively modify not only charge transport or storage and distribution in the NC dielectric films but also breakdown strength of the materials due to the interfacial effects.^34,35 Although many studies have been performed to improve electroactive β phase nucleation and dielectric constant by incorporating different metal oxides but there are no such report on the effect of NiO NPs and amorphous carbon doped NiO NCs to enhance polar β phase nucleation and dielectric properties of PVDF thin films. Low cost and large theoretical capacitance (2584 F g⁻¹ within 0.5 V) of nickel oxide (NiO) makes the material one of the most promising electrochemical capacitor material.³⁶

In our present study, NiO NPs and amorphous carbon doped NiO or C–NiO NCs are taken as the nanomodifiers for achieving large electroactive β phase nucleation and dielectric PVDF NC thin films. First, NiO NPs and C–NiO NCs have been prepared via soft chemical approach and then PVDF NC thin films with different amount of filler contents (1–20 mass%) have been synthesized via simple solution casting method. Thereafter the effect of the two nanomodifiers on the improvement of polar β phase crystallization and dielectric properties have been investigated on the basis of size, nature of fillers, charge on the surface of the nanofillers, filler contents and interfacial interaction between the nanofillers and PVDF chains. It has been found that with the same filler loading, the dielectric constant of the three-phase C–NiO–PVDF NC thin film is much higher than the dielectric constant of two-phase NiO–PVDF NC thin film. The origin of large dielectric constant of the as-synthesized PVDF thin films modified C–NiO NCs than NiO NPs loaded thin film is attributed to formation of large number of micro-capacitors in the three-phase NC thin films and large amount of interfacial polarization at the interfaces of the C–NiO NCs and PVDF matrix than in the two-phase NiO–PVDF system.

2. Experimental

2.1. Materials

The chemicals used in present work are nickel(II) acetate tetrahydrate (Ni(CH₃COO)₂·4H₂O) (Merck, India), lithium hydroxide (LiOH) (Merck, India), carbon powder (procured from CGCRI), poly(vinylidene fluoride) (PVDF) pellets (Aldrich, Germany. M_w: 180 000 GPC, M_n: 71 000) and dimethyl sulfoxide (DMSO) (Merck, India).

2.2. Preparation of NiO NPs and C–NiO NCs

Initially, 4.9 g nickel(II) acetate tetrahydrate was dissolved in 20 ml double distilled water followed by addition of LiOH (1 M) in double distilled water under vigorous stirring for 12 hours. The pH of the solution was adjusted to 9 for complete transformation of the salt to nickel hydroxide (Ni(OH)₂) precipitate. Similarly, carbon doped Ni(OH)₂ precipitate was prepared by dissolving 600 mg carbon powder and 4.9 g Nickel(II) acetate tetrahydrate in 20 ml water under vigorous magnetic stirring for 2 hours followed by precipitation as described above. The Ni(OH)₂ precipitate and C–Ni(OH)₂ precipitate were collected by centrifuging the mixtures and then washed with ethanol and double distilled water until the pH reached 7. The precipitates were then dried in a dust free oven at 80 °C for 24 hours. Finally, the NiO NPs and C–NiO NCs were obtained by sintering the Ni(OH)₂ precipitate and C–Ni(OH)₂ precipitate at 400 °C for 4 hours in air. The obtained samples were reserved in vacuum desiccators.

2.3. Fabrication of NiO NPs and C–NiO NCs loaded PVDF thin films

400 mg of PVDF pellets and different amounts of NiO NPs or C–NiO NCs (1–20 mass%) were dissolved in 10 ml DMSO. Thereafter the whole mixture was stirred for 12 hours at 60 °C followed by ultrasonic treatment for 30 minutes for uniform dispersion of the nanofillers in PVDF matrix. The NC thin films were obtained by casting the mixture in clean Petri dishes and evaporating the solvent (DMSO) at 80 °C for 24 hours in a dust free oven. Finally, the NC films with thickness about 60–70 μm were collected from the Petri dishes and stored in vacuum desiccator for further characterization. Pure PVDF films were also prepared via similar process for comparison. Details of compositions of thin films were designated in Table 1.

Table 1 Sample designations showing respective compositions

Sample name	Amount of PVDF (g)	Percentage of NPs/NCs (mass%)	Amount of NPs/NCs (g)	Name of the filler added to PVDF
PVDF	0.4	0	0	NiO NPs
PNiO1	0.4	1	0.004
PNiO5	0.4	5	0.020
PNiO10	0.4	10	0.040
PNiO15	0.4	15	0.060
PNiO20	0.4	20	0.080
PCNiO1	0.4	1	0.004	C–NiO NCs
PCNiO5	0.4	5	0.020
PCNiO10	0.4	10	0.040
PCNiO15	0.4	15	0.060
PCNiO20	0.4	20	0.080

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2.4. Characterization

The crystalline structure of the fillers and different polymorphs of NC thin films were studied using X-ray diffractometer (Model-D8, Bruker AXS Inc., Madison, WI) at ambient condition under 40 kV operating voltage using Cu-K_α radiation with a scan speed of 0.3 s per step and scanning range from (2θ) 20 to 80° for fillers and from 15 to 65° for thin films respectively.

UV-visible absorption spectra of the films were measured by a UV-vis spectrophotometer (USA and UV-3101PC, Shimadzu) in the wavelength range 200–800 nm.

The zeta potential measurements of the fillers were performed with Zeta-sizer-5000 (Malvern Instruments, UK) to determine the surface charge.

The morphologies, microstructures, chemical analysis i.e. energy dispersive X-ray (EDX) spectroscopy and mapping of the films were investigated using field emission scanning electron microscope (FESEM) (INSPECT F50, Netherlands).

Thermal gravimetric analysis (TGA) was carried out using a TGA/SDTA851e (Mettler Toledo AG) at a temperature range from 50 to 600 °C for thin films with heating rate 10 °C min⁻¹ under nitrogen atmosphere with a constant flow rate of 40 ml min⁻¹. Ceramic crucibles having 70 μl capacity was used as sample holder.

The formation of different phases in the NC thin films were further investigated using Fourier transform infrared spectroscopy (FTIR-8400S, Shimadzu) with a resolution of 4 cm⁻¹ in a wavenumber range from 400 cm⁻¹ to 1100 cm⁻¹. 50 scans were carried out for each sample. The fraction of β-phase (F(β)) in the NC thin films were quantified using Lambert–Beer law using FTIR spectra which is,

where, A_α = the absorbance at 764 cm⁻¹ and A_β = the absorbance at 840 cm⁻¹. K_β (7.7 × 10⁴ cm² mol⁻¹) and K_α (6.1 × 10⁴ cm² mol⁻¹) are the absorption coefficients at 840 cm⁻¹ and 764 cm⁻¹ wave number respectively.^24,31

A differential scanning calorimeter (DSC-60, Shimadzu (Asia Pacific) Pte. Ltd., Singapore) was used to characterize the crystallization and melting behaviour of the pure PVDF and NC thin films. The thin films were heated from 80 °C to 200 °C at a heating rate of 10 °C min⁻¹ under N₂ gas atmosphere. The degree of crystallinity (X_c) of the samples was investigated using the following equation from DSC thermographs:

X_c = ΔH_m/ΔH_100% (2)

where, ΔH_m = the enthalpy of fusion of the thin films and ΔH_100% = the enthalpy of fusion of 100% crystalline PVDF (104.6 J g⁻¹).^24,33

The variation of capacitance (C) and tangent loss (tan δ) with increasing frequency of the NC thin films were collected using a digital LCR meter (Agilent, E4980A) using a sample holder containing circular Ag electrodes at ambient condition. Data were recorded in the range of 100 Hz to 2 MHz under applied voltage of 1 V. Dielectric constant (ε) and ac conductivity (σ_ac) of NC thin films were calculated using the following eqn (3) and (4) respectively,

ε = Cd/ε₀A (3)

and

σ_ac = 2πfε₀ε tan δ (4)

where, d = thickness, A = area of the thin films, f = the frequency in Hz applied across the NC thin films and ε₀ = permittivity of free space with value 8.854 × 10⁻¹² F m⁻¹.^16,25

3. Results and discussions

3.1. Characterization of NiO NPs and C–NiO NCs

Fig. 1a shows the XRD pattern of pure NiO NPs and amorphous carbon doped NiO NCs. The formation of NiO NPs in both samples is confirmed by the diffraction peaks at 2θ values 37.33° (111), 43.1° (200), 63.1° (220), 75.53° (311) and 79.6° (222) which corresponds to (111), (200), (220), (311) and (222) planes of cubic phase of NiO NPs (JCPDS 4-0835). The diffraction spectrum of C–NiO NCs shows more prominent and intense characteristic peaks of NiO NPs than the pure NiO NPs. It may be due to the presence of amorphous carbon in C–NiO that results in the formation of well crystalline NiO NPs. The absence of broader characteristic peaks of amorphous carbon may be due to impregnation of carbon in NiO nanostructure.

Fig. 1 (a) XRD spectra and (b) UV-visible absorption spectra of pure NiO NPs and C–NiO NCs.

Fig. 1b presents the UV-visible absorption spectra of the sample. Absorption peak at 330 nm (∼3.75 eV) also confirmed the formation of NiO NPs. The absorption peak exists due to the interionic 3d⁸–3d⁸ transition in Ni²⁺ ions.^37,38 Lower absorbance value and blue shifting (∼13 nm) of the characteristic absorption peak of NPs in C–NiO are due to the presence of amorphous carbon and its strong interaction with the NiO NPs.

Fig. 2a and b represent the zeta potential distribution of NiO NPs and C–NiO NCs respectively. The data have been recorded at pH 5.8 for NiO NPs and pH 6 for C–NiO NCs in DMSO solvent, the interaction pH of the NPs or NCs–PVDF mixture in DMSO. Both the samples show negative zeta potential which indicates the surface of the NPs or NCs are negatively charged. Pure NiO NPs shows little high negative value in zeta potential than amorphous carbon doped NiO NPs.

Fig. 2 Zeta potential distribution of (a) NiO NPs and (b) C–NiO NCs; FESEM images of (c) NiO NPs and (d) C–NiO NCs; (e) EDX spectrum and (f) mapping of C–NiO NCs.

The morphology and microstructure of the NPs and NCs have been investigated using field emission scanning electron microscopy (FESEM) and represented in Fig. 2c and d. Formation of well-defined and uniformly distributed spherical NiO particles with average diameter ∼ 15–25 nm for pure NiO NPs (Fig. 2c) and ∼10–15 nm for amorphous carbon doped NiO NPs (Fig. 2d) have been observed respectively. Presence and distribution of carbon in C–NiO NCs sample has been verified using EDX spectroscopy and mapping. Fig. 2e and f show the EDX spectrum and mapping of the C–NiO NCs. EDX spectrum and mapping of the C–NiO NCs clearly indicate the presence of carbon and its uniform distribution in NiO matrix.

3.2. NiO NPs/C–NiO NCs–PVDF thin films

3.2.1. Thermal stability. Investigation of thermal stability of the pure PVDF and NC thin films have been carried out using TGA studies. Fig. 3 represents the TGA thermographs of pure PVDF and fillers modified PVDF thin films.

Fig. 3 TGA thermographs of pure PVDF and fillers modified PVDF thin films: (a) NiO NPs modified PVDF thin films and (b) C–NiO NCs modified PVDF thin films.

For pure PVDF film one single weight loss occurred at 430 °C. Increase in thermo-degradation temperature of about ∼18 to 30 °C for NiO NPs doped PVDF samples (Fig. 3a) and about ∼20 to 30 °C for C–NiO NCs loaded PVDF samples (Fig. 3b) is observed. Thus, good thermal stability is observed for NC thin films which may be due to the strong interfacial interaction between the polymer chains and the negatively charged fillers near the filler–polymer interface.

3.2.2. Digital images, morphology and microstructure of PVDF–NC thin films. Fig. 4a represents digital images of the films. Decrease in optical transparency with increasing loading of the NiO NPs or C–NiO NCs have been observed (Fig. 4a1) without any loss in flexibility. Image of flexible PVDF films with highest doping (20 mass%) of NiO NPs and C–NiO NCs are shown in Fig. 4a2 and a3 respectively. Morphology, microstructure and dispersion as well as interaction of the nanofillers in PVDF matrix have been investigated by FESEM. Fig. 4b shows the FESEM image of pure PVDF films and Fig. 4c–f represents the FESEM images of the fractured surfaces of PNiO10, PCNiO10, PNiO20 and PCNiO20 respectively. Both the NiO NPs or C–NiO NCs incorporated films are composed uniformly dispersed particles in the polymer matrix implying close interaction with the polymer chains.

Fig. 4 (a) Digital images (a1) and flexible PNiO20 (a2) and PCNiO20 (a3) samples; FESEM images of (b) pure PVDF, (c) PNiO10, (d) PCNiO10, (e) PNiO20 and (f) PCNiO20 (red circle showing NiO NPs or amorphous carbon doped NiO NPs).

3.2.3. UV-visible spectroscopy. UV-visible spectroscopy also suggests intimate interaction between the nanofillers and the PVDF matrix. Fig. 5 represents the UV-visible absorption spectra of pure PVDF and NC thin films. No such characteristic absorbance peak has been observed for pure PVDF. The absorbance spectra of NiO NPs loaded PVDF thin films show a characteristic absorption peak at ∼305 nm (∼4.07 eV) (Fig. 5a) and the absorbance spectra of C–NiO NCs loaded PVDF thin films show a characteristic absorption peak at ∼312.3 nm (∼3.97 eV) (Fig. 5b). In comparison to the position of pure nanofillers characteristic absorbance peaks (shown in Fig. 1b), a blue shifting in the absorbance peak ∼25 nm in NiO NPs loaded PVDF thin films and ∼4.7 nm in C–NiO NCs doped PVDF thin films have been observed. Strong electrostatic interaction between the negatively charged surfaces of nanofillers and polymer chains may result in slight change in the morphology, size and surface properties of the nanofillers i.e. NiO NPs which may influence 3d⁸–3d⁸ transition in Ni²⁺ ions in NiO NPs resulting blue shifting in absorption spectra of NC thin films. For both nanofillers loaded samples the intensities of the absorbance peaks have been increased with increasing loading concentration of the nanofillers implying fine homogeneous distribution and dispersion of the nanofillers in polymer matrix.

Fig. 5 UV-visible spectra of pure PVDF and nanofillers modified PVDF thin films: (a) NiO NPs modified PVDF thin films and (b) C–NiO NCs modified PVDF thin films.

3.2.4. Investigation of electroactive β phase. The crystalline phase and structure of the films have been evaluated using X-ray diffraction (XRD) analysis. Fig. 6 presents the XRD pattern of pure PVDF and nanofiller modified PVDF thin films. The XRD spectrum of pure PVDF shows diffraction peaks at 2θ values of 17.46° (100), 18.2° (020), 19.85° (021) and 26.44° ((201), (310)) corresponding to nonpolar α phase. But upon the incorporation of the nanofillers (NiO NPs or C–NiO NCs) in PVDF matrix all diffraction peaks due to nonpolar α phase vanishes, only one characteristic diffraction peak at 2θ = 20.5° ((110), (200)) emerged prominently implying electroactive β phase nucleation in the NC samples.^7,16,24

Fig. 6 XRD pattern of pure PVDF and nanofillers modified PVDF thin films: (a) NiO NPs modified PVDF thin films and (b) C–NiO NCs modified PVDF thin films.

The characteristic peaks due to NiO NPs have also been observed in the diffraction patterns of NC samples and the intensities of the peaks increased with loading of NiO NPs or C–NiO NCs indicating uniform distribution of the nanofillers inside the polymer matrix. The diffraction peak value at 20.5° illustrates that the amount of electroactive β phase has increased with the doping of both nanofillers inside the PVDF films.

Identification of different polymorphs and quantification of the phase content in PVDF has been suitably carried out using Fourier transform infrared spectroscopy (FTIR). In the present work, basically specific absorbance bands like 766 cm⁻¹ and 840 cm⁻¹ corresponding to nonpolar α phase and polar β phase respectively are recognized and have been utilized to calculate the electroactive β phase content using the eqn (1).³¹ Fig. 7a and b illustrate the FTIR spectra of pure PVDF and NiO NPs or C–NiO NCs modified PVDF thin films. The characteristic absorbance bands corresponding to nonpolar α phase present in the FTIR spectrum of pure PVDF at 490 cm⁻¹ (CF₂ waging) 532 cm⁻¹ (CF₂ bending), 615 and 764 cm⁻¹ (CF₂ bending and skeletal bending), 796 and 975 cm⁻¹ (CH₂ rocking) are totally diminished for both NiO NPs or C–NiO NCs modified PVDF thin films and characteristic absorbance bands due to electroactive β phase at 479 cm⁻¹ (CF₂ deformation), 510 cm⁻¹ (CF₂ stretching), 600 cm⁻¹ (CF₂ wag) and 840 cm⁻¹ (CH₂ rocking, CF₂ stretching and skeletal C–C stretching) appear prominently. A small shoulder band at 812 cm⁻¹ (CF₂ asymmetric stretching) assigned to polar γ phase has also been detected in the NC samples.^7,24,39 Significant improvement in the absorbance intensity of main β phase characteristic band i.e. 840 cm⁻¹ has been observed with increasing loading of both nanofillers.

Fig. 7 FTIR spectra of pure PVDF and nanofillers loaded PVDF thin films: (a) NiO NPs modified PVDF thin films and (c) C–NiO NCs modified PVDF thin films and evaluation of β-phase content with increasing nanofillers content form IR spectra: (b) NiO NPs modified PVDF thin films and (d) C–NiO NCs modified PVDF thin films.

The variation of the relative fraction of the β phase content ((F(β)%)) with increasing doping concentration of NiO NPs or C–NiO NCs has been represented graphically in Fig. 7c and d. F(β) value is found to be ∼35% for pure PVDF. The electroactive β phase content attains ∼80% with nanofillers doping inside the PVDF matrix for all the cases. Maximum F(β) value ∼82% is achieved for 10 mass% loading of NiO NPs in PVDF and highest F(β) value ∼80% is reached by 5 mass% incorporation of C–NiO NCs in the PVDF due to maximum elongation and stabilization of the polymer chains in longer TTTT conformation. Further loading of the fillers restrict the movement and elongation of polymer chains in longer TTTT conformation which in turns reduces β phase content.^16,31

The thermo-analytical technique, DSC has also been used to complement XRD and FTIR results for identification of the different crystalline polymorphs of PVDF. Different melting peaks arise on the DSC thermograph due to different crystalline phases of PVDF.⁷ Fig. 8 shows the DSC thermographs of pure PVDF and NiO NPs/C–NiO NCs doped PVDF films. The thermograph of pure PVDF contains one strong melting peak at 163.4 °C assigned to α phase crystallization in pure PVDF. An increase in melting temperature (T_m) ∼ 4.3 °C for NiO NPs loading in PVDF and ∼4.9 °C for C–NiO NCs loading in PVDF imply electroactive β phase nucleation in the NC samples which are compatible with both XRD and FTIR analysis.^19,24,29

Fig. 8 (a) DSC thermographs of pure PVDF and NiO NPs modified PVDF thin films and (b) DSC thermographs of pure PVDF and C–NiO NCs modified PVDF thin films.

DSC is not only used to investigate the different crystalline polymorphs, but also to measure the enthalpy of fusion or melting enthalpy as well as degree of crystallinity of the polymer films. Melting enthalpy of the samples has been investigated first from the DSC thermographs and thereafter the degree of crystallinity of the samples calculated using eqn (2). Fig. 9 illustrates the variation of melting enthalpy and degree of crystallinity of pure PVDF and NiO NPs/C–NiO NCs doped PVDF films with nanofillers content. The melting enthalpy or degree of crystallinity for both nanofillers modified PVDF thin films show higher values compare to the melting enthalpy or degree of crystallinity of pure PVDF due to the nucleating or catalytic action of the nanofillers. The degree of crystallinity values have been increased with the nanofillers content and gained the maximum ∼43.4% and ∼38.8% for 10 mass% loading for both nanofillers. Higher loading of the nanofillers results in hindering the free movement of the polymer chains and this restricts formation of β phase spherulites i.e. grains in the NC thin films. Thus, decrease in degree of crystallinity values have been observed for higher loading of nanofillers in PVDF.^9,40

Fig. 9 Evaluation of melting enthalpy (a and c) and degree of crystallinity (b and d) of pure PVDF and NiO NPs/C–NiO NCs loaded PVDF thin films with increasing nanofillers content from DSC thermographs.

3.2.5. Formation mechanism of electroactive β phase. The formation of electroactive β phase in the NC samples has been investigated in details using XRD, FTIR and DSC. The percentage of β phase nucleation in the NiO NPs/C–NiO NCs loaded PVDF thin films have also been quantified from FTIR spectra of the samples. Based on the above observations, we try to explain the formation of electroactive β polymorph in term of ion–dipole or electrostatic interaction at the interface between the polymer chains and the dopants.^{9,16,24,31,41}

The zeta potential of the dopants have been investigated at the same pH of the mixture of NiO NPs/C–NiO NCs loaded PVDF in DMSO (at pH 5.8 for NiO NPs and pH 6 for C–NiO NCs) as the surface charge on the particle strongly depends on the pH of the solution.³¹ Both the dopants show negative zeta potential confirming negatively charged surface of the dopants. Thus, the nucleation of electroactive β phase in the NC thin films takes place at the negatively surface of the dopants. Upon the addition of the dopants in the PVDF–DMSO solution, the positive dipole –CH₂ of the polymer chains experiences strong electrostatic interaction with negative surfaces of nanofillers which in turn leads in the alignment of the polymer chains in longer stabilized TTTT or all trans conformation on the negatively charged surface of the dopants i.e. NiO NPs/C–NiO NCs resulting nucleation of electroactive β phase in the NC samples.^{9,16,31,41,42} The β phase formation mechanism i.e. interfacial interaction between NiO NPs/C–NiO NCs and polymer chains is schematically presented in Fig. 10.

Fig. 10 Proposed schematic of electroactive β phase formation mechanism.

Easier local internal chain rotation results some gauge effect in some samples. Thus small presence of γ phase in some samples are due to the easier local internal chain rotation.^24,33 Slightly high zeta potential value of NiO NPs results maximum 82% β phase nucleation in the PCNiO10. Similar types of electroactive β phase crystallization via ion–dipole or electrostatic interaction have been previously shown in our studies on Fe₂O₃/Co₃O₄ NPs-PVDF,⁹ clay minerals-PVDF¹⁶ and WO₃·H₂O NPs-PVDF NCs.³³

3.2.6. Dielectric properties.
3.2.6.1. Fillers content dependency of the dielectric constant. Dependency of dielectric constant (ε) on doping concentration of the nanofillers (NiO NPs or C–NiO NCs) has been represented in Fig. 11. Slow and almost linear increase in dielectric constant up to 15 mass% doping of NiO NPs in PVDF is observed clearly from Fig. 11a followed by a steep rise for 20 mass% loading. Highest dielectric constant (ε) ∼ 35.4 (at 20 Hz) is obtained which is 4 times larger than that of pure PVDF (ε ∼ 9). More than 20 mass% addition of NiO NPs in polymer matrix degrades material's structural integrity and films become porous. Thus we cannot further increase the NPs loading for achieving larger dielectric constant. Subsequently, to achieve large dielectric value three phase C–NiO–PVDF NC thin films have employed. The dielectric constant of C–NiO NCs modified PVDF thin films increases continuously in a nonlinear fashion with increasing C–NiO NPs and reached the maximum dielectric value ∼317.4 at 20 Hz for 20 mass% addition of C–NiO NCs that is ∼9 times larger than that of NiO NPs–PVDF matrix (Fig. 11). For NiO NPs–PVDF system the increase in dielectric constant is not so remarkable at 20 mass% loading, but addition of 20 mass% C–NiO NCs in PVDF increases the dielectric constant to a very high value.

Fig. 11 Fillers content dependency of dielectric constants at 20 Hz, 1 kHz, 10 kHz and 1 MHz: (a) NiO NPs loaded PVDF thin films and (b) C–NiO NPs loaded PVDF thin films.

Increase in dielectric value for both systems may be attributed to both Maxwell–Wagner–Sillars (MWS) interfacial polarization effect and microcapacitors structure formation in nanocomposite films (shown in Fig. 12).^43–45 The addition of NiO NPs in the polymer matrix increases the interfaces and improves MWS interfacial polarization resulting in enhancement of dielectric constant. However, in case of C–NiO–PVDF system, a large number of interfaces and microcapacitors are formed due to the presence of the conducting carbon particle which are embedded in semiconducting NiO particles and insulating polymer chains which results in large accumulation of space charges at the C–NiO–PVDF interface. Consequently, PCNiO20 shows remarkable enhancement in dielectric value. Higher frequency dielectric constant is mainly dominated by dipolar polarization.^25,45–49 Thus for both system, dielectric constant shows almost similar weak increment with loading of the fillers because the number of dipoles do not vary much for those NC thin films as the F(β) value is almost fixed for all samples (Fig. 7b and d).

Fig. 12 Schematic diagram of interfacial polarization and microcapacitor model: (a) NiO NPs loaded PVDF thin films and (b) C–NiO NPs loaded PVDF thin films.

3.2.6.2. Frequency dependency of the dielectric properties. Frequency dependence dielectric constant, tangent loss and ac conductivity of pure PVDF and NiO NPs or C–NiO NCs loaded PVDF have been graphically presented in Fig. 13. It is observed from Fig. 12a and d that the dielectric constants decreased with increasing frequency for both systems after incorporation of NiO NPs or C–NiO NCs in polymer matrix. Rapid decrease in dielectric value have been observed from 20 Hz to 10³ Hz (35.4 to 22 for PNiO20 and 317.4 to 66.8 for PCNiO20) in the NC thin film with 20 mass% NiO NPs or C–NiO NCs. The dependency or variation of dielectric constants of other samples are not so remarkable.

Fig. 13 Frequency dependence of dielectric properties of pure PVDF and NiO NPs or C–NiO NCs loaded PVDF thin films; (a and d) dielectric constants, (b and e) tangent losses and (c and f) ac conductivities.

The relaxation of dielectric constant at lower frequencies (below 10³ Hz) and high value of dielectric constant are due the MWS relaxation occurring at the interface of NiO–PVDF, C–NiO–PVDF and C–NiO NCs. The behaviour of dielectric constant at lower frequencies suggests that the synthesized NC system is a typical percolation system. Similar type of frequency dependency was observed by Yao et al. in BaTiO₃–carbon nanotube–PVDF three-phase system.⁴³ The dielectric response are mostly dominated by space charge accumulation at the interfaces of NiO–PVDF, C–NiO–PVDF and C–NiO NCs. Achievement of high dielectric values are significantly controlled by MWS interfacial polarization as more interfaces are obtained upon the addition of NiO NPs or C–NiO NCs in polymer matrix especially near percolation threshold concentration of the fillers.^43–48 Fig. 13b and e represent the variation of tangent loss value with increasing frequency. For NC samples the tangent loss decreases with frequency due to MWS effect. The decrease in tangent loss values for the NCs with low content of NiO NPs filler comparing with neat PVDF has been observed in Fig. 13b. It may be interpreted as the phase inversion in PVDF matrix, the phenomenon where an electroactive β-phase rich continuous phase has been formed due to strong interfacial interaction with the NPs which in turns results in a lower leak current and thus a decrease in tangent loss values.¹⁶ And, still high dielectric constant has been achieved with quite low loss value (maximum tangent loss ∼ 1.7 for PCNiO20). This lower loss value may be attractive for versatile application in the fields of electric energy storage devices and the arrays of thin film supercapacitors of the synthesized flexible percolative NC thin films. The dependency of ac conductivity of the NiO NPs and C–NiO NCs modified PVDF thin films have been illustrated in Fig. 13c and f. The low frequency plateau is not observed for both nanofiller modified samples. A fast and linear increment of the conductivity for all doped samples have been observed with increasing frequency due to the presence of NiO NPs or C–NiO NCs in the polymer matrix. The increase of ac conductivity with doping concentration of the nanofillers are mainly due to the formation of more conductive networks, MWS interfacial polarization contribution and dipolar relaxation modes in both NC systems.^43–49 AC conductivity is found to be ∼1.7 × 10⁻⁸ S cm⁻¹ for PNiO20 and ∼5.9 × 10⁻⁷ S cm⁻¹ PCNiO20 at 20 Hz.

4. Conclusions

In summary, the chemically synthesized NPs and NCs have been doped in PVDF matrix via simple solution casting method to develop electroactive and high dielectric flexible polymer thin films. Electroactive β phase transformation has been observed for both nanofiller modified PVDF thin films. About 82% electroactive β phase nucleation is achieved by incorporating 10 mass% NiO NPs and 80% is obtained by adding 5 mass% of C–NiO NCs. The β phase nucleation occurs in the NC samples due to strong ion–dipole interaction i.e. electrostatic attraction between the negatively charged fillers surfaces and –CH₂ dipoles of the polymer chains. The uniform dispersion or distribution of both nanofillers in PVDF matrix results in large interfacial polarization and formation of large number of microcapacitance structure in the NC samples mainly in three phase C–NiO–PVDF thin films. Highest dielectric constant ∼ 317.4 at 20 Hz is obtained by adding 20 mass% C–NiO NCs in PVDF matrix which is 35 times and 9 times larger than that of pure PVDF and two phase NiO NPs–PVDF thin film (PNiO20) respectively. Thus, the well fabricated electroactive, low loss and high dielectric flexible polymer thin films may be the attractive candidate for developing capacitive sensors, actuators, UV protectors or detectors, energy storage devices, piezoelectric nanogenerator and energy harvesters.

Acknowledgements

Authors are grateful to University Grants Commission (UGC), Government of India (F. 17-76/2008 (SA-1)) for the financial assistance.

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Footnotes

† Present address: Fuel Cell and Battery Division, Central Glass and Ceramic Research Institute, Kolkata-700032, India.

‡ Present address: Department of Physics, IIEST, Howrah, West Bengal – 711103, India.

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