Enhanced electroactive β-phase nucleation and dielectric properties of PVdF-HFP thin films influenced by montmorillonite and Ni(OH)₂ nanoparticle modified montmorillonite

Abstract

The present work lays emphasis on the synthesis of electroactive and high dielectric poly(vinylidenefluoride-co-hexafluoropropylene) (PVdF-HFP) composite films incorporated with montmorillonite (MMT) and Ni(OH)₂ nanoparticle (NP) modified MMT using a facile and low cost solution casting technique. Formation of Ni(OH)₂ NPs, different polymorphs and morphology of the composite samples have been investigated using UV-visible spectroscopy, X-ray diffraction, Fourier transform infrared spectroscopy, differential scanning calorimetry and field emission scanning electron microscopy. A remarkable improvement in β-phase nucleation (∼85.22% for 1 mass% MMT doped at and 82.10% for 7.5 mass% NiMMT doped polymer composites films) and promising dielectric values are the standpoints of the work. Strong ion–dipole interaction between the negatively charged surfaces of MMT and –CH₂ dipoles of the polymer chains in the MMT loaded PVdF-HFP samples and formation of hydrogen bonds between the –OH group of Ni(OH)₂ NPs and F atom of the polymer chains in the Ni(OH)₂ NP-MMT/PVdF-HFP system play a dynamic role in enhancing the electroactive β-phase nucleation and the dielectric properties of the samples. A large accumulation of charges at the interfaces of the dopant and polymer matrix via Maxwell–Wagner–Sillars interfacial polarization effect occurred due to the disparity in the conductivity of the dopants and the polymer matrix is liable for the significant enrichment of dielectric properties of the composite films. Excellent thermal stability in dielectric properties has also been observed in these composite samples over a wide temperature spectrum. These electroactive and high dielectric polymer composite thin films may be potential candidates for application in piezoelectric nanogenerators, energy storage devices, thin film transistors and capacitive sensors etc.

---

1. Introduction

The modern era has recognized dielectric spectroscopy as a crucial tool in the exploration of the dielectric relaxation process in different polymeric compounds when treated with an electric field. Notable reformation of physical, chemical, electrical and piezoelectric properties of dielectric polymers such as poly(vinylidene fluoride) (PVdF) and its copolymers are the noteworthy consequences of this process.¹

PVdF is one of the rarest, flexible, high molecular weight, complex electroactive hydrofluorocarbon polymers that has well-established piezoelectric, pyroelectric and dielectric properties. Added improvement of the piezoelectricity and mechanical behavior properties can be acknowledged when other monomer units like hexafluoropropylene (HFP) are attached to the main PVdF chain. Henceforth, PVdF copolymers like poly(vinylidenefluoride-co-hexafluoropropylene) (PVdF-HFP) turn out to be unique prospects with similar crystalline structures to that of PVdF but enriched flexibility and chemical resistance. Owing to properties like nontoxicity, high stability, shape and size tailor ability and recycling aptitude, these copolymers are gaining impetus in widespread energy harvesting technologies.²

Piezoelectric and capacitive sensors, thin film transistors, batteries, chemical warfare protection, electroacoustic transducers, electrochemical convertors, infrared detectors, actuators, ferroelectric memory devices and artificial muscles are a few expanses of utilization of the polymer.

Five distinguished polymorphs α, β, γ, δ and ε describe PVdF-HFP. The non-polar α-polymorph, characterized by a TGTG′ conformation of polymeric chains and obtained directly from the melt at moderate or high supercoolings or polymerization, is the most common and dominant phase amongst the five phases. However, the β-polymorph, having an all trans (TTTT) zigzag conformation, can be thought to be superior among the phases owing to its largest polarization per unit cell (8 × 10⁻³⁰ C m) in comparison to the other two polar phases i.e. γ-phase and δ-phase having a T₃GT₃G′ and TGTG′ conformation, respectively.^3,4 Hence, cultivation of β-phase nucleation in PVdF-HFP has received great interest in recent times for its superior performance in piezoelectric, pyroelectric, ferroelectric and dielectric extents.

Various studies have been undertaken to enhance the β-phase activity with the incorporation of nano modifiers. Priya and Jog were the pioneers who worked on the development of the polar phase in PVdF in the presence of organically modified bentonite clay Cloisite®6A and Cloisite®20A and results confirmed an increase in the melting temperature and crystallization rate of the clay polymer composites.^5–7 A complete β-phase transformation of PVdF on the addition of Cloisite®30B was attained by Shah et al.⁹ Thakur et al.⁸ also concluded an enhancement of β-phase and dielectric properties of PVdF on addition of clays like Kaol and Hal.

Cost-effectivity, good compatibility, good cationic exchange and absorbance capacity of MMT make this clay an attractive material as a dopant.^10,11 Repetitive octahedrons of alumina stuck in between two tetrahedral silica sheets defines the structure of MMT. Intercalation of MMT flakes in the polymer matrix favors the advancement in the mechanical strength, thermal stability and the dielectric constant which arises from the assimilation of two materials with complementary dielectric properties.¹²

Zhang et al. studied the use of MMT and PVdF and crystallinity was enhanced in their study.¹³ The crystallization behavior of PVdF/MMT composites were also investigated by Yu et al.¹⁴ Similar studies were also carried out by Dillon et al. using nanoclays (Cloisite®15A and Cloisite®25A) and PVdF.¹⁵

Kelarakis et al. were the first to study the structural and dielectric properties of a series of organo-MMT/PVdF-HFP composites.¹⁶ Very few studies have been done by researchers on the phase inversion mechanism and dielectric properties of PVdF-HFP/MMT composites. Most studies revolve around their applications in membranes, separators for the lithium-ion batteries^17–20 solid state dye sensitized solar cells¹² and fuel cells.²¹

Numerous works regarding organically modified MMT doped polymer have been reported.^5–7 However, there are a very few studies on inorganically modified MMT polymer composites. Ni(OH)₂ NPs can be synthesized in a simplistic and cost effective way. Also, they are environmentally benign, have a high theoretical specific capacitance and remarkable redox behavior which makes them a potential material of interest.²²

The present work deals with a comparative study on the increase of the electroactive β-phase nucleation, its transformation mechanism and dielectric properties of a PVdF-HFP thin film due to incorporation of negatively charged MMT layers and Ni(OH)₂ NP modified MMT in various percentages (1–10 mass%) in the polymer matrix via a simplistic solution casting method. Furthermore, a detailed dielectric study is also carried out over a temperature range from 30 °C to 140 °C. Promotion of electroactive β-phase nucleation and enrichment of the dielectric properties of the composite samples have been suitably discussed in terms of strong electrostatic interaction and large accumulation of charges between the interfaces of the dopants and polymer chains.

2. Experimental

2.1. Materials

Poly(vinylidene fluoride-co-hexafluoropropylene) pellets (Aldrich, Germany, M_w: 455 000 GPC; M_n: 110 000), montmorillonite clay (MMT) (<1 μm) (NanoCorUSA), dimethyl sulfoxide (DMSO) (Merck India), and nickel acetate tetrahydrate extra pure [(CH₃·COO)₂Ni·4H₂O] (LOBA Chemie) are the precursors used in the composite polymer film synthesis.

2.2. Synthesis of Ni²⁺ ion adsorbed MMT (NiMMT)

Briefly, 200 mg of MMT and 400 mg of nickel acetate was added to 20 ml of deionized water in a beaker and the mixture was stirred overnight. Then the Ni²⁺ ion adsorbed MMT composite was collected by centrifugation (16 000 rpm for 15 min) and was dried in a hot-air oven at 80 °C for further synthesis.

2.3. Synthesis of MMT and Ni(OH)₂ NP modified MMT (NiMMT) doped PVdF-HFP thin films

4% PVdF-HFP solution was prepared at 60 °C by dissolving 0.2 g of PVdF-HFP in 5 ml DMSO and formerly synthesized purified clay (MMT and NiMMT) was added to the PVdF-HFP solution in definite amounts (1–10 mass%, Table 1). The mixture was then stirred overnight using a magnetic stirrer followed by 30 min ultra-sonication for complete homogeneous mixing.

Table 1 Sample designations and calculated amount of clay mineral components added to PVdF-HFP during the synthesis of pure and composite polymer films

Sample name	Amount of PVdF-HFP (g)	Percentage of clay minerals (mass%)	Amount of clay minerals (g)	Name of the material doped
PVdF-HFP	0.2	0	0
PMT 1%	0.2	1.0	0.002	Montmorillonite (MMT)
PMT 2.5%	0.2	2.5	0.005
PMT 5%	0.2	5.0	0.010
PMT 7.5%	0.2	7.5	0.015
PMT 10%	0.2	10.0	0.020
PNiMT 1%	0.2	1.0	0.002	Ni(OH)2 Modified Montmorillonite (NiMMT)
PNiMT 2.5%	0.2	2.5	0.005
PNiMT 5%	0.2	5.0	0.010
PNiMT 7.5%	0.2	7.5	0.015
PNiMT 10%	0.2	10.0	0.020

---

A clean Petri dish was taken in which the mixture was poured to cast the film (50–60 μm) and placed at 80 °C in a dust free hot-air oven for 12 h for complete evaporation of the solvent. Pure PVdF-HFP films were also prepared without adding the dopants under similar conditions as a control in this study. The synthesized MMT and Ni(OH)₂ modified MMT PVdF-HFP films were then stored in a vacuum desiccator for further characterizations (Fig. 1).

Fig. 1 Schematic of the synthesis procedure of the pure and composite polymer films.

2.4. Material characterization

2.4.1. X-ray diffraction (XRD). The synthesized pure PVdF-HFP thin film, MMT and NiMMT doped PVdF-HFP composite thin films, pure MMT and NiMMT powder samples were all examined using an X-ray diffractometer (Model-D8, Bruker AXS Inc., Madison, WI) from 3° to 70° 2θ values operating under a voltage of 40 kV with a scan speed of 0.3 s per step using nickel filtered Cu-Kα radiation.

2.4.2. Fourier transform infrared spectroscopy (FTIR). The IR spectra of the prepared composite films were studied with the aid of a Fourier transform infrared spectrometer (FTIR-8400S, Shimadzu) to examine the influence of the doped MMT and NiMMT clay on the phase nucleation of the polymer PVdF-HFP. The absorbance data of the films (50–60 μm in thickness) were noted in the wavenumber range from 400 cm⁻¹ to 1000 cm⁻¹ with a resolution of 4 cm⁻¹.

Each sample was scanned 50 times while collecting the FTIR data. The fraction of β-phase, F(β) of the composite thin films was calculated using the Lambert–Beer law which is given as:

where A_β = the absorbance at 840 cm⁻¹, A_α = the absorbance at 764 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 respectively.

2.4.3. UV-vis spectroscopy. A UV-visible spectrophotometer (Lambda 25, Perkin Elmer, USA and UV-3101PC, Shimadzu) was used to obtain the absorption spectra of the pure PVdF-HFP film and MMT and NiMMT doped PVdF-HFP composite films in the wavelength range 200–800 nm.

2.4.4. Field emission scanning electron microscopy (FESEM). FESEM (INSPECT, F50, Netherlands) was implemented to capture the morphological and structural changes in the clay minerals (MMT and NiMMT) doped PVdF-HFP films in comparison to the pure polymer film. All the samples were gold coated via plasma spraying at 0.1 mbar pressure. A pressure of about 5 × 10⁻³ Pa at a distance of 12 mm away from the detector with spot size 3 mm was introduced inside the vacuum chamber and the samples were kept in it. An emission current of 170 μA and an operating voltage of 20 kV were applied while capturing the FESEM images.

2.4.5. Differential scanning calorimetry (DSC). The phase nucleation and melting nature of the pure PVdF-HFP and MMT and NiMMT incorporated PVdF-HFP films were explored by means of a differential scanning calorimeter (DSC-60, Shimadzu (Asia Pacific) Pte. Ltd., Singapore). All the films were heated from 30 °C to 600 °C at a heating rate of 10 °C min⁻¹ under N₂ gas atmosphere using ceramic crucibles of capacity 70 μl.

The DSC curves as obtained helped in the evaluation of the enthalpies (ΔH_m) of fusion of the pure and composite films. The degree of crystallinity (X_c) of the same was assessed from the equation given below:

X_c = ΔH_m/ΔH_100% (2)

where ΔH_m = heat of melting or enthalpy of fusion and ΔH_100% = melting enthalpy of 100% crystalline PVdF with a value of 104.6 J g⁻¹.

2.4.6. Dielectric measurements. The electrical properties specifically dielectric properties of the pure film of co-polymer PVdF-HFP, and MMT and NiMMT embedded polymer matrix were studied in the frequency range of 100 Hz to 2 MHz at room temperature and also by variation of temperature at the step of 30 °C till 140 °C using a digital LCR meter (Agilent, E4980A) and a PID temperature controller. Across the two opposite surfaces of all the film samples, 1 V ac voltage was applied. The dielectric constant (ε) and the ac conductivity (σ_ac) of the samples were evaluated using the following eqn (3) and (4) respectively,

ε = Cd/ε₀A (3)

σ_ac = 2πfε₀ε tan δ (4)

where, C, d, A and tan δ are the capacitance, thickness, area and tangent loss of the samples respectively, f is the applied frequency in Hz and ε₀ is permittivity of free space with value 8.854 × 10⁻¹² F m⁻¹.

3. Results and discussions

3.1. UV-vis absorption properties

The optical absorbance spectra of the pure polymer film and MMT or NiMMT doped PVdF-HFP composite films, examined under the ultraviolet-visible electromagnetic spectrum are exhibited in Fig. 2. No characteristic absorption peak is detected for the pure PVdF-HFP film in the UV-vis spectrum.

Fig. 2 UV-visible absorption spectra of pure PVDF-HFP and composite films; (a) MMT doped PVdF-HFP films and (b) Ni(OH)₂ NP modified MMT doped PVdF-HFP films.

A small and weak absorption peak is observed at 243 nm in Fig. 2(a) due to the metal cation charge transfer transition which increases in intensity with the increment of the doping concentrations of MMT in the polymer films signifying the homogeneous embedment of the clay in the PVdF-HFP matrix.²³

In the NiMMT doped PVdF-HFP films (Fig. 2(b)), an absorption peak is noticed at 214.67 nm (∼5.80 eV) which established the formation of Ni(OH)₂ NPs in the polymer matrix.²⁴ Electronic transfer of Ni²⁺ in the oxygen octahedral sites from the ground state ³T_1g (P) to the ³T_1g (F) state results in the appearance of the characteristic absorption peak of the Ni(OH)₂ NPs. A slight shift of this characteristic peak is noticed with an increase in the concentrations of NiMMT in the PVdF-HFP. The peak shift is clearly denoted in Fig. 2(b) which may be due to the strong electrostatic interactions of Ni(OH)₂ NPs with the PVdF-HFP chains or because of the slight variation in size and morphology of the Ni(OH)₂ NPs formed in the polymer matrix, upsetting the electronic transition in the conduction or valence bands and resulting in the prominent interaction of the Ni(OH)₂ NPs in MMT clay with the polymer. The characteristic absorption peak of the Ni(OH)₂ NPs strengthens gradually with the increase in doping concentration which suggests the complete formation of Ni(OH)₂ NPs on the surface of MMT in the PVdF-HFP matrix.

3.2. X-ray diffraction analysis

XRD patterns of pure MMT, Ni²⁺ ion adsorbed MMT, pure PVdF-HFP and MMT/NiMMT doped PVdF-HFP composite films with various weight percentages of MMT and NiMMT (1–10 mass%) are displayed in Fig. 3.

Fig. 3 X-ray diffractograms of (a) MMT, (b) Ni²⁺ ion adsorbed MMT, (c) MMT doped PVdF-HFP composite films and (d) NiMMT doped PVdF-HFP composite films.

Characteristic reflections of the crystalline phases of the pure PVdF-HFP membrane appear in the XRD spectrum at 2θ = 17.9° (020), 26.8° (021) and 38.8° (131) which relates to the bigger spherulites of the non-polar α-phase crystals while 2θ = 18.5° (100) and 20.1° (110) corresponds to the smaller spherulites of the γ-phase crystal which are found to co-exist with the α-phases.^12,25,26

The characteristic peak at 2θ = 20.3° corresponds to the superposition of the (110) and (200) β-phase diffraction in the spectrum.^27,28

The diffraction peak at 2θ = 7.2° in the XRD spectrum reveals the presence of MMT²⁹ whereas in the Ni²⁺ adsorbed MMT powder, the clay peak is shifted to 2θ = 6.2° as shown in Fig. 3(a) and (b). On the other hand, the spectra of MMT or NiMMT doped PVdF-HFP composite films reflects the shift of this characteristic clay peak towards the left side of the spectrum resulting in a diffused peak around 2θ = 6.44° for all the composite films. The observed results demonstrate the strong intercalation of the clay particles in the polymer PVdF-HFP matrix. This type of structure is formed due to the interaction between the clay layers and PVdF-HFP or because of shear induced interaction. This reflection goes intense with the increase in the doping concentration of MMT and NiMMT in the PVdF-HFP matrix (Fig. 3(c) and (d)).^30,31

It is quite vibrantly reflected from the spectra of the composite films that peaks corresponding to the γ-phase and α-phase crystals (2θ = 17.9°, 18.5°, 20.1°, 26.8°, 38.8°) disappeared with the doping of MMT and NiMMT content in the PVdF-HFP matrix and only one distinct characteristic peak at 2θ = 20.3° ((110), (200)) appeared which confirmed the nucleation of the electroactive β-phase nucleation in the synthesized composite films.²⁵ The nucleation of the β-phase is substantial in the 1 mass% doping concentration of MMT in the case of MMT–PVdF-HFP composite films (Fig. 3(c) and (d)). Surprisingly two weaker XRD peaks have been observed at 9° and 29° in the 5% MMT doped PVdF composite (Fig. 3(c)). This may be due to the presence of some unknown phases in the composite at that particular concentration. On the other hand, the intensity of the β-phase increased gradually in the NiMMT–PVdF-HFP composite films and was found to be maximum for 7.5 mass% doping of NiMMT clay. The variation of the nucleating action of the two dopants may be attributed to the disparity of morphology and surface charges of the MMT/NiMMT of the doped films. The catalytic activity of MMT excels over NiMMT in β-phase crystallization in the PVdF-HFP matrix which is due to the significant stabilizing effect of negatively charged flat surfaces of MMT layers via ion–dipole interaction (between negatively charged flat surfaces of clay and –CH₂ dipoles of polymer chains) than positively charged Ni(OH)₂ NPs modified MMT.

3.3. Fourier transform infrared spectroscopy

Investigation has been executed to define the various crystalline polymorphs present in unmodified PVdF-HFP, MMT and NiMMT modified PVdF-HFP composite thin films using one of the techniques recognized as Fourier transform infrared (FTIR) spectroscopy and hereafter to determine the α, β and γ phases of PVdF-HFP as shown in Fig. 4.

Fig. 4 IR spectra of (a) pure PVdF-HFP and MMT doped PVdF-HFP films, (b) evaluation of β-phase content [F(β)] for MMT–PVdF-HFP films, (c) pure PVdF-HFP and NiMMT doped PVdF-HFP films and (d) evaluation of β-phase content F(β) for NiMMT–PVdF-HFP films.

In the untreated PVdF-HFP films, the absorbance bands are detected at 490 cm⁻¹ (–CF₂ wagging), 530 cm⁻¹ (–CF₂ bending), 615 cm⁻¹ (skeletal bending), 764 cm⁻¹ (–CF₂ bending), 795 cm⁻¹ (–CH₂ rocking) and 975 cm⁻¹ (twisting) in the IR spectra which are related to the presence of the non-polar α-phase in PVdF-HFP. Small absorbance peaks of the β-phase are also found at 509 cm⁻¹ (–CF₂ stretching) and 840 cm⁻¹ (–CH₂ rocking, –CF₂ stretching and skeletal C–C stretching) in the spectrum³² (Fig. 4(a) and (c)).

In both MMT and NiMMT treated PVdF-HFP films, all the absorbance bands due to the α-phase ceased to exist and the absorbance peaks at 509 cm⁻¹ (–CF₂ stretching) and 840 cm⁻¹ (–CH₂ rocking, –CF₂ stretching and skeletal C–C stretching) are prominent, signifying the strong emergence of the electroactive β-phase^33,34(Fig. 4(a) and (c)).

The main absorbance peak of the β-phase at 840 cm⁻¹ is found to have the strongest intensity for 1 mass% of MMT embedded PVdF-HFP film while in the NiMMT–PVdF-HFP composite films the absorbance peak at 840 cm⁻¹ gradually has been strengthened with increasing concentrations of NiMMT and reached its maximum value at 7.5 mass% of NiMMT content which indicates intimate interactions in the composite films for that particular concentration.

Evaluation of the relative fraction of the β-phase content F(β) in pure PVdF-HFP, MMT and NiMMT loaded polymer composite films have been executed from the IR spectra using the Lambert–Beer law stated in eqn (1). Graphical representation of the percentage variation of F(β) (%) with the MMT and NiMMT content is shown in Fig. 4(b) and (d) respectively. About 45.31% F(β) is obtained for the pure PVdF-HFP film. With the increase in concentration (mass%) of both MMT and NiMMT in the PVdF-HFP matrix, F(β) increases and almost saturates at higher doping concentrations. Maximum F(β) values of ∼85.22% for 1 mass% loading of MMT and ∼82.10% at 7.5 mass% incorporation of NiMMT in PVdF-HFP composite film were achieved. Henceforth, stabilization and nucleation of the electroactive β-phase, enhanced with the addition of MMT and NiMMT content are well approved from FTIR spectroscopic measurements.

3.4. Differential scanning calorimetry

Different phase, thermal property, enthalpy and degree of crystallinity of the pure and composite polymer film samples have been investigated using differential scanning calorimetry thermograms shown in Fig. 5. The melting temperature (T_m) is observed around 149 °C for the pure PVdF-HFP film^12,35 while for the MMT–PVdF-HFP films, the melting temperature shows a prominent shifting to higher temperature by about 2–4 °C (Fig. 5(a)). A slight variation of melting temperature with concentration is also perceived for Ni(OH)₂ modified MMT doped polymer films with respect to the pristine PVdF-HFP film (Fig. 5(d)). Enhancement in the β-phase in the doped films is responsible for this escalation of melting temperatures.^7,9,13

Fig. 5 DSC thermograms of (a) pristine PVdF-HFP and MMT doped PVdF-HFP films (b) assessment of melting enthalpies and (c) degree of crystallinity of pure and MMT doped polymer films. (d) Pristine PVdF-HFP and NiMMT doped PVdF-HFP films, (e) assessment of melting enthalpies and (f) degree of crystallinity of pure and NiMMT doped polymer films.

Melting enthalpies are evaluated from the thermographs and thereafter the degree of crystallinity for all the samples was calculated using eqn (2). Variation of melting enthalpy and degree of crystallinity of the samples are graphically illustrated in Fig. 5(b), (e), (c) and (f). A detectable upsurge in the enthalpy of fusion and hence the degree of crystallinity for both the composite samples compared to the pure PVdF-HFP film are noticed with the increasing concentration of the additives.³⁶ Increase in the crystallinity of the composite samples may be due to the active interaction of the additives with polymer chains which induces the formation of polar β polymorphs from non-polar α spherulites.²⁷ The degree of crystallinity has been found highest at ∼58% and 50% for 1 mass% and 7.5 mass% of MMT and NiMMT doped polymer films respectively as both the concentrations exhibit maximum β-polymorphs. However, higher loading concentrations of the additives in the polymer matrix reduce the crystallinity due to confinement of the movement of the polymer chains.^9,13

3.5. Scanning electron microscopy

FESEM images clearly demonstrate the surface morphology of the MMT flakes, Ni²⁺ ion adsorbed MMT, pure PVdF-HFP and the composite film samples as shown in Fig. 6. The good uniform distribution and strong interaction of the MMT flakes in both MMT and NiMMT doped PVdF-HFP samples can be described from the encircled areas in the polymer matrix as depicted in the FESEM images (Fig. 6(d) and (e)).^37,38

Fig. 6 FESEM micrographs of (a) MMT flakes, (b) Ni²⁺ ion adsorbed MMT, (c) pristine PVdF-HFP surface, (d) 10 mass% MMT doped PVdF-HFP films and (e) 10 mass% NiMMT doped PVdF-HFP films.

3.6. Phase transformation mechanism

Schematic representation of the most anticipated electroactive β-phase interaction mechanism between the MMT surfaces or NiMMT surfaces and the PVdF-HFP chains has been illustrated in Fig. 7.

Fig. 7 Graphical illustration of the proposed β-phase transformation mechanism.

When MMT flakes are intercalated in the polymer matrix, the interaction between the negatively charged surface of the doped clay mineral and the positive –CH₂ dipoles from (CH₂–CF₂) monomer of the PVdF-HFP alters the polarity and gives rise to the nucleation of the β-phase in composite samples. However, impregnation of Ni(OH)₂ NP modified MMT clay in the polymer matrix, portrays a different picture of the interaction mechanism. Ni²⁺ ions strongly interact with the negatively charged surfaces of the MMT flakes and the H atoms of the Ni(OH)₂ NPs go through a strong interaction with the F atoms of the polymer chains via formation of hydrogen bonds establishing longer TTTT conformation in NiMMT modified PVdF-HFP samples (Fig. 7). The additives act as a substrate for β-phase nucleation when prepared films are dried at 80 °C. Thus, the stimulation of stabilized alignment of the polymer chains due to these substrates results in the formation of elongated electroactive all trans conformation i.e. β-phase nucleation.^9,13,39 Maximum β-phase is attained at 1 mass% of MMT doped polymer films while the same is obtained at 7.5 mass% in case of NiMMT–PVdF-HFP films. Results achieved throws light on the fact that negatively charged MMT is much efficient in enhancing the TTTT conformation compared to positively charged NiMMT. These results have also been justified by XRD, FTIR and DSC data.

3.7. Dielectric properties

3.7.1. Additive content (mass%) dependency of the dielectric properties. The disparity of dielectric constant (ε), tangent loss (tan δ) and ac conductivity (σ_ac) of pure PVdF-HFP film, MMT and Ni(OH)₂ NPs modified MMT incorporated PVdF-HFP composite films as a function of MMT and NiMMT content (mass%) were measured at 100 Hz, 1 kHz and 1 MHz at ambient condition have been displayed in Fig. 8 and 9.

Fig. 8 MMT content (mass%) dependence on dielectric constant and tangent loss property of MMT doped polymer films (PMT 1%, PMT 2.5%, PMT 5%, PMT 7.5% and PMT 10%) at (a) 100 Hz, (b) 1 kHz and (c) 1 MHz; NiMMT content (mass%) dependence on dielectric constant and tangent loss property of NiMMT doped polymer films (PNiMT 1%, PNiMT 2.5%, PNiMT 5%, PNiMT 7.5% and PNiMT 10%) at (d) 100 Hz, (e) 1 kHz and (f) 1 MHz.

Fig. 9 Clay content dependence of ac conductivity at 100 Hz, 1 kHz and 1 MHz of (a) MMT doped polymer films (PMT 1%, PMT 2.5%, PMT 5%, PMT 7.5% and PMT 10%) (b) NiMMT doped polymer films (PNiMT 1%, PNiMT 2.5%, PNiMT 5%, PNiMT 7.5% and PNiMT 10%).

A gradual increase in dielectric constant is observed for both MMT and NiMMT loaded PVdF-HFP films with the increase in their content (mass%) at the mentioned frequencies. Highest dielectric constants of 28, 26 and 19 are calculated for 7.5 mass% doping of MMT in PVdF-HFP matrix at 100 Hz, 1 kHz and 1 MHz frequencies respectively, while the dielectric constant values ∼41, 38 and 29 for 10 mass% doping of Ni(OH)₂ NPs modified MMT in PVdF-HFP composite films are obtained at the respective frequencies (Fig. 8) using eqn (3).

Maxwell–Wagner–Sillars (MWS) interfacial polarization which occurs due to the difference in conductivity of the dopants and polymer is accountable for the enrichment of dielectric constant in the doped films with increase in doping concentration compared to the pure polymer film.^40–42

The dielectric constant of the NiMMT–PVdF-HFP film is higher than the MMT–PVdF-HFP film. This disparity is mainly due to the formation of large interfaces in the three phase NiMMT–PVdF-HFP composite film than the two phase MMT–PVdF-HFP composite film. With increasing MMT content in the polymer matrix the interfacial area per unit volume associated with the interfaces of clay particles and the polymer chains crop up. Also the inter-particle distance reduces. All these, along with the strong interaction formed between the negatively charged surface of MMT and positive –CH₂ dipoles, influence the significant strengthening of the average localized polarization accompanied with the clay particles and the coupling between the adjoining grains. Hence the dielectric constant of the MMT modified polymer films improves with respect to the pure polymer film as well as with the increase in doping concentration.

However, the Ni(OH)₂ NP modified MMT doped polymer film introduces us with a three phase system having two interfaces, MMT/Ni(OH)₂ NPs and Ni(OH)₂ NPs/PVdF-HFP. Under the action of the electric field, the negatively charged surfaces of the MMT layers strongly interact with the Ni(OH)₂ NPs and the –OH groups of the NPs pursue the same with the F atoms of the polymer chains establishing large interfacial surfaces. Thus larger interfacial polarization and further enhancement of dielectric constant occurs here in comparison to the two phase MMT–PVdF-HFP composite films. Also, an increase in dopant content in the polymer matrix results in the intensification of fraction of electroactive β-phase nucleation. This further leads to the increase in the number of free amorphous chains in the film samples and also contributes to the increase in dielectric strength.^9,43 A non-ohmic trend is sensed for the tangent loss for the MMT and Ni(OH)₂ NP modified MMT–PVdF-HFP composite films. Ac conductivity of both the composites is found to increase in an almost linear fashion.

3.7.2. Frequency dependence of the dielectric properties. The frequency dependence of dielectric constant, tangent loss (tan δ) and the ac conductivity of the pure PVdF-HFP film, MMT and Ni(OH)₂ NPs modified MMT–PVdF-HFP composite films are illustrated in Fig. 10.

Fig. 10 Frequency dependency of he dielectric properties of pure PVdF-HFP, MMT and NiMMT doped PVdF-HFP films (a) and (d) dielectric constant; (b) and (e) tangent loss (c) and (f) ac conductivity.

The gradual decrease of the dielectric constant values at higher frequencies of all the synthesized samples and the increase of dielectric values with the increase in the doping concentration for both the MMT and Ni(OH)₂ NPs modified MMT doped PVdF-HFP films are visibly evident from the illustrations (Fig. 10(a) and (d)). Both of these observations can be enlightened by the well-known MWS interfacial polarization effect.⁹ According to the MWS effect, when the composite film samples are treated with an electric field, a huge accretion of charge carriers from the impurities (i.e. dopants) or the electrode of both the composite films propagate liberally and get sandwiched amidst the interfaces of the dopant particles and the polymer chains due to the dissimilarity in electrical conductivity between the dopant and polymer matrix. A large interfacial polarization and hence a substantial increase in dielectric constant are the two plausible upshots of the colossal deposition of the charge carriers at the interfaces of the dopants and polymer matrix. The charge carriers or dipoles get adequate time for their movement at lower frequencies, as because their movement and their gradual buildup in the interfacial region readily continues at lower frequencies and a large polarization and dielectric constant become the two significant outcomes of this charge accumulation. At higher frequencies the picture is quite different. Due to the presence of a lag between the changes in dipoles and the changes in electrical field, the dielectric constant becomes a complicated function of electrical field frequency. Increasing frequency restricts the movement of the charge carriers or dipoles due to the incongruity of frequency and the dielectric constant values initiates to decrease for all the samples. Again, the reduction of the effective number of dipoles may be the other reason for the diminishing dielectric constant values at higher frequencies.^9,13,44 Fig. 10(b) and (e) also provide a vivid representation of the non-linear increase of the tangent losses with the increase of frequency in both systems.

It is obvious from Fig. 10(c) and (f) that the ac conductivity graph linearly accelerates with the increase in frequency as well as with increase in doping concentration for both the MMT and NiMMT doped PVdF-HFP films. Lower frequency conductivity is mainly governed by dc conductivity while the same is dictated by ac conductivity when a higher frequency region comes into the scenario. The increase of dopant percentages notably contributes to increase in the ac conductivity. Hence, the MWS interfacial effect and dipolar relaxation modes become the unique roots for the linear rise of the ac conductivities of the composite films with rise in frequencies. Absence of a lower frequency plateau in the graphical representation can be the outcome of the rapid increase of conductivity at low frequencies which linearly rises with frequency.^9,45

3.7.3. Temperature dependence of the dielectric properties. Fig. 11 and 12 show the temperature dependency of the dielectric constant values, tangent loss and ac conductivity of the pristine PVdF-HFP and MMT or NiMMT doped PVdF-HFP films at three different frequencies (100 Hz, 1 kHz and 1 MHz).

Fig. 11 Variation of dielectric constant with temperature in pristine PVdF-HFP film and MMT doped PVdF-HFP film at (a) 100 Hz (b) 1 kHz (c) 1 MHz; variation of tangent loss with temperature in pristine PVdF-HFP film and MMT doped PVdF-HFP film at (d) 100 Hz (e) 1 kHz (f) 1 MHz; variation of ac conductivity with temperature in pristine PVdF-HFP film and MMT doped PVdF-HFP film at (g) 100 Hz (h) 1 kHz (i) 1 MHz.

Fig. 12 Variation of dielectric constant with temperature in the pristine PVdF-HFP film and NiMMT doped PVdF-HFP film at (a) 100 Hz (b) 1 kHz (c) 1 MHz; variation of tangent loss with temperature in the pristine PVdF-HFP film and NiMMT doped PVdF-HFP film at (d) 100 Hz (e) 1 kHz (f) 1 MHz; variation of ac conductivity with temperature in the pristine PVdF-HFP film and NiMMT doped PVdF-HFP film at (g) 100 Hz (h) 1 kHz (i) 1 MHz.

At low temperature (30–80 °C), the dielectric constant of pure PVdF-HFP doesn’t vary much with temperature and shows an almost parallel behavior pattern for all frequencies while it increases with the rise in temperature (above 80 °C), as depicted for the lower frequency range due to total polarization of PVdF-HFP.¹

At high frequencies, even with the rise in temperature there is a decrease in dielectric constant value. This behavior is associated with the glass transition temperature (T_g) of the polymer and can be explained by the α_a-relaxation process which arises from the thermal response of the molecular motions and Brownian motion of the main polymer chain segment. The dipoles respond to the applied heat energy and the polar group present in the polymer exhibits thermal motion.^1,16,46

At low temperature, the dipoles are very intimately attached to each other so limited orientation among themselves takes place and they influence each other. However, with the escalation in temperature, the orientation of the dipoles increases considerably because of the expansion of the bulk material. An increase in temperature reduces the ratio of the molecules per effective length of the dielectric material. So the interaction between the dipoles is reduced and their movement is improved. The dipoles can also get additional energy and time for orientation from the supplied heat. Thus in a two-fold way, the net increase in polarization and simultaneous increase in dielectric constant occurs at a low and medium frequency region over the temperature spectrum.^47,48

The dielectric constant of the electroactive polymer composites remains independent around 30–40 °C, thereafter, gradually rising with the temperature at lower frequencies with an increase in doping concentrations of the both MMT/NiMMT dopants. The dielectric constant value decreases and almost saturates even at higher temperatures when higher frequencies intervene (Fig. 11 and 12(a)–(c)). Absence of enough dipoles as well as not having sufficient time to catch up with the alteration in frequency of electrical field results in weak polarization and weak dielectric dependence on temperature at higher frequency.^1,48

Predominance of the interfacial MWS polarization effect where large charge gathering takes place due to the strong interactions between negatively charged MMT layers with positive –CH₂ dipoles for MMT doped polymer films and interaction between negatively charged surfaces of the MMT with the Ni²⁺ ions and the –OH groups with the F atoms of the polymer chains in Ni(OH)₂ NPs modified polymer films, increases the crosslinking density of the polymer chains and is responsible for the overall increase of dielectric values in the low frequency region. Dominant α_a relaxation in the low temperature area and the chain scissoring process which accelerates the movement of functional groups also can be attributed to the increase in dielectric constant in the low temperature region.^1,42,46,49

The dielectric constant value increases from 41 to 167 for Ni(OH)₂ NP modified polymer composites when the temperature rises from 30 °C to 140 °C which outstrips MMT doped polymer films where a dielectric value of 146 was obtained at the highest temperature (140 °C) establishing NiMMT doped polymer films to be a better dielectric material.

Ac conductivity and tangent losses of the pure and composite films increase with temperature at low frequencies but at higher frequencies, the reverse nature is observed in Fig. 11 and 12(d)–(i). The upswing of temperature highly amplifies the segmental flexibility by the faster internal bond rotation of the polymer chain which further expedites the intra- and inter-ionic hopping in the polymer chain and charge transportation is boosted.⁵⁰ Also, the MWS interfacial polarization phenomenon marks the production of a large number of induced dipoles, increased dielectric values and hence enhancement in α_c-relaxation. All these result in the increase of ac conductivity as compared to the pure polymer with temperature.

4. Conclusion

A simplistic solution casting process has been instigated for the synthesis of pure and MMT/NiMMT PVdF-HFP composite polymer films with various dopant percentages (1–10 mass%). UV-vis studies confirm the formation of Ni(OH)₂ NPs. Materialization of the polar β-phase from non-polar α-spherulites has been confirmed through the XRD, FTIR and DSC measurements. All trans (TTTT) conformation i.e. electroactive or polar β-polymorphs in the PVdF-HFP is enhanced up to 85.22% and 82.1% when 1 mass% MMT and 7.5 mass% Ni(OH)₂ NPs modified MMT have been incorporated in PVdF-HFP matrix respectively.

Dielectric studies have been performed at a temperature range of 30 °C to 140 °C for all films. Promising dielectric constant values are reached for the composite films. A dielectric constant of ∼28 and 41 were obtained at 100 Hz for the MMT and NiMMT doped PVdF-HFP films respectively which were enhanced to 146 and 167 respectively at the same frequency when the temperature increased to 140 °C. Though the β-phase nucleation occurs only at 1 mass% of MMT doped films compared to NiMMT doped polymer composites where the same occurs at much greater doping percentage, electrical studies reveal the latter as a much better candidate of dielectric potential. Henceforth, these composites can prove to be alluring in the electronic fields.

Acknowledgements

We are obliged to the University Grants Commission (UGC) Govt. of India for providing fellowship during the execution of this work.

References

1. D. S. Rana, D. K. Chaturvedi and J. K. Quamara, J. Optoelectron. Adv. Mater., 2010, 4(6), 838–844.
2. L. Wu, W. Yuan, N. Hu, Z. Wang, C. Chen, J. Qiu, J. Ying and Y. Li, J. Phys. D: Appl. Phys., 2014, 47, 135302.
3. V. Tomer, E. Manias and C. A. Randall, J. Appl. Phys., 2011, 110, 044107.
4. P. Martins, A. C. Lopes and S. Lanceros-Mendez, Prog. Polym. Sci., 2014, 39, 683–706.
5. L. Priya and J. P. Jog, J. Polym. Sci., Part B: Polym. Phys., 2002, 40(15), 1682–1689.
6. L. Priya and J. P. Jog, J. Polym. Sci., Part B: Polym. Phys., 2003, 41(1), 31–38.
7. L. Priya and J. P. Jog, J. Appl. Polym. Sci., 2003, 89(8), 2036–2040.
8. P. Thakur, A. Kool, B. Bagchi, S. Das and P. Nandy, Appl. Clay Sci., 2014, 99, 149–159.
9. D. Shah, P. Maiti, E. Gunn, D. Schmidt, D. D. Jiang and C. A. Batt, Adv. Mater., 2004, 16(14), 1173–1177.
10. G. B. B. Varadwaj and K. M. Parida, RSC Adv., 2013, 3, 13583–13593.
11. J. Kabilaphat, N. Khaorapapong, K. Saito and M. Ogawa, Appl. Clay Sci., 2015, 115, 248–253.
12. K. Prabakaran, S. Mohanty and S. K. Nayak, Int. J. Plast. Technol., 2014, 18(3), 349–361.
13. Y. Y. Zhang, S. L. Jiang, Y. Yu, G. Xiong, Q. F. Zhang and G. Z. Guang, J. Appl. Polym. Sci., 2012, 123, 2595–2600.
14. W. Yu, Z. Zhao, W. Zheng, B. Long, Q. Jiang, G. Li and X. Ji, Polym. Eng. Sci., 2009, 49, 491–498.
15. D. R. Dillon, K. K. Tenneti, C. Y. Li, F. K. Ko, I. Sics and B. S. Hsiao, Polymer, 2006, 47, 1678–1688.
16. A. Kelarakis, S. Hayrapetyan, S. Ansari, J. Fang, L. Estevez and E. P. Giannelis, Polymer, 2010, 51, 469–474.
17. A. Atanassov, G. Kostov, D. Kiryakova and L. Borisova-Koleva, J. Thermoplast. Compos. Mater., 2014, 27(1), 126–141.
18. V. Aravindan, P. Vickraman, S. Madhavi, A. Sivashanmugam, R. Thirunakaran and S. Gopukumar, Solid State Sci., 2011, 13(5), 1047–1051.
19. H. S. Jeong, S. C. Hong and S. Y. Lee, J. Membr. Sci., 2010, 364(1–2), 177–182.
20. M. Wang, F. Zhao, Z. Guo and S. Dong, Electrochim. Acta, 2004, 49, 3595–3602.
21. G. Gnana kumar, J. Mater. Chem., 2011, 21, 17382.
22. H. Du, L. Jiao, K. Cao, Y. Wang and H. Yuan, ACS Appl. Mater. Interfaces, 2013, 5, 6643–6648.
23. C. Zhang, W. W. Tjiu, W. Fan, Z. Yang, S. Huang and T. Liu, J. Mater. Chem., 2011, 21, 18011.
24. Y. Qi, H. Qi, J. Li and C. Lu, J. Cryst. Growth, 2008, 310, 4221–4225.
25. K. M. Kim, N. G. Park, K. S. Ryu and S. H. Chang, Electrochim. Acta, 2006, 51, 5636–5644.
26. J. Cao, B. Zhu, Y. Xu, J. Li and C. Chen, J. Macromol. Sci., Part A: Pure Appl. Chem., 2011, 45(6), 449–455.
27. T. U. Patro, M. V. Mhalgi, D. V. Khakhar and A. Misra, Polymer, 2008, 49, 3486–3499.
28. P. Thakur, A. Kool, B. Bagchi, S. Das and P. Nandy, Phys. Chem. Chem. Phys., 2015, 17, 1368.
29. S. Kakuta, T. Okayama, M. Kato, A. Oda and T. Abe, Catal. Sci. Technol., 2011, 1, 1671–1676.
30. Q. Feng, C. Honbu, K. Yanagisawa, N. Yamasaki and S. Komarneni, J. Mater. Chem., 2000, 10, 483–488.
31. S. Miao, Z. Liu, B. Han, J. Zhang, X. Yu, J. Du and Z. Sun, J. Mater. Chem., 2006, 16, 579–584.
32. C. H. Du, B. K. Zhu and Y. Y. Xu, J. Mater. Sci., 2006, 41, 417–421.
33. L. Wu, W. Yuan, N. Hu, Z. Wang, C. Chen, J. Qiu, J. Ying and Y. Li, J. Phys. D: Appl. Phys., 2014, 47, 135302.
34. L. Wu, G. Huang, N. Hu, S. Fu, J. Qiu, Z. Wang, J. Ying, Z. Chen, W. Lia and S. Tang, RSC Adv., 2014, 4, 35896.
35. Shalu, V. K. Singh and R. K. Singh, J. Mater. Chem. C, 2015, 3, 7305–7318.
36. Y. Xu, G. Ray and B. Abdel-Magid, Composites, Part A, 2006, 37, 114–121.
37. B. Bagchi, S. Kar, S. K. Dey, S. Bhandary, D. Roy, T. K. Mukhopadhyay, S. Das and P. Nandy, Colloids Surf., B, 2013, 108, 358–365.
38. Y. Feng, W. L. Li, Y. F. Hou, Y. Yu, W. P. Cao, T. D. Zhang and W. D. Fei, J. Mater. Chem. C, 2015, 3, 1250.
39. P. Martins, C. M. Costa, M. Benelmekki, G. Botelho and S. Lanceros-Mendez, CrystEngComm, 2012, 14, 2807–2811.
40. L. A. Ramajo, M. A. Ramírez, P. R. Bueno, M. M. Reboredo and M. S. Castro, Mater. Res., 2008, 11(1), 85–88.
41. J. Liu, C. G. Duan, W. G. Yin, W. N. Mei, R. W. Smith and J. R. Hardy, Phys. Rev. B: Condens. Matter Mater. Phys., 2004, 70, 144106.
42. A. C. Lopes, C. M. Costa, R. S. Serra, I. C. Neves, J. L. G. Ribelles and S. Lanceros-Méndez, Solid State Ionics, 2013, 235, 42–50.
43. S. A. M. Nabirqudri, J. Mater. Res., 2014, 29(24), 2957–2964.
44. N. Chilaka and S. Ghosh, Electrochim. Acta, 2014, 134, 232–241.
45. P. Thakur, A. Kool, B. Bagchi, N. A. Hoque, S. Das and P. Nandy, RSC Adv., 2015, 5, 62819.
46. S. El-Sayed, T. A. Abdel-Baset and A. Hassen, AIP Adv., 2014, 4, 037114.
47. V. S. Yadav, D. K. Sahu, Y. Singh and D. C. Dhubkarya, Proceedings of IMECS, 2010, III, ISBN 978–988-18210-5-8.
48. Z. M. Dang, J. P. Wu, H. P. Xu, S. H. Yao and M. J. Jiang, Appl. Phys. Lett., 2007, 91, 072912.
49. A. Hassen, T. Hanafy, S. El-Sayed and A. Himanshu, J. Appl. Phys., 2011, 110, 114119.
50. C. W. Liew, Y. S. Ong, J. Y. Lim, C. S. Lim, K. H. Teoh and S. Ramesh, Int. J. Electrochem. Sci., 2013, 8, 7779–7794.

---

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.

---