Tunable photoluminescence emissions and large dielectric constant of the electroactive poly(vinylidene fluoride–hexafluoropropylene) thin films modified with SnO₂ nanoparticles

Abstract

In the present study, SnO₂ nanoparticles (NPs) have been synthesized hydrothermally at 100 °C and 150 °C. The NPs have been characterized by X-ray diffraction (XRD), UV-visible spectroscopy and field emission scanning electron microscopy (FESEM). Thereafter, poly(vinylidene fluoride–hexafluoropropylene) (PVDF–HFP) with different contents (1–15 wt%) of SnO₂ NPs has been synthesized by a simple solution-casting method. An electroactive β-phase formation mechanism in the NP-incorporated PVDF–HFP thin films is confirmed by XRD, Fourier transform infrared spectroscopy and differential scanning calorimetry. FESEM images show interfacial interaction in the thin films between the polymer matrix and the NPs. Maximum β-phase nucleation of ∼74% has been reached by 10 wt% addition of SnO₂ NPs synthesized hydrothermally at 100 °C. The photoluminescence emission spectra show two UV and visible emissions at two excitation wavelengths. Uniform distribution of the NPs in the polymer matrix leads to large dielectric constant ∼2578.44 at 20 Hz due to addition of 10 wt% SnO₂ NPs (100SO) in PVDF–HFP matrix. The large increase in dielectric constant has been successfully explained in terms of MWS polarization effect and percolation theory.

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1. Introduction

Over the last decade, electroactive polymer nanocomposites (NCs), with high dielectric constant, excellent piezoelectric, pyroelectric, ferroelectric and electro-optic properties, have gained interest due to their variety of potential applications. These include high energy density arrays of capacitors, acoustic transducers, actuators, infrared detectors, piezoelectric sensors and nanogenerator, rail runs, pulsed lasers, information storage memory devices, developed flexible and stretchable energy harvesting systems in both biomedical and electronic fields.^1–10 Among these, energy harvesting technologies, which are produced with piezoelectric materials have attracted great attention because these allow the conversion of ambient energy, such as sunlight, wind, vibration, rain drops and sea waves, to electric energy.¹¹

Among the most important electroactive semicrystalline thermoplastic polymers, the poly(vinylidene fluoride) (PVDF) homopolymer and its copolymers (such as poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF–HFP), poly(vinylidene fluoride-co-trifluoroethylene) (P(VDF–TrFE)), and poly(vinylidene fluoride–trifluoroethylene–chlorofluoroethylene) (P(VDF–TrFE–CFE))), with good peizo-electric, pyro-electric, ferro-electric coefficients, excellent thermal stability and chemical resistance, have been considered as prospective materials for the abovementioned applications.^12–18 Nowadays, there has been a great interest to enhance the piezoelectricity and electrochemical coupling factor of PVDF, which is sweeping real applications. The semicrystalline polymer of PVDF occurs in at least five forms, i.e. α, β, γ, δ and ε, but the most common crystalline phases of PVDF are α, β and γ; moreover, the β form assumes paramount importance among all other forms due to its excellent piezoelectricity and good elastic strength and therefore study on improvement of the β-phase nucleation in PVDF has currently received a lot of attention.^19–22 The thermodynamically stable polymorph of PVDF and its co-polymers is the nonpolar α-phase with TGTG′ (T-trans, G-gauche⁺, G′-gauche) dihedral conformation.²³ However, the all trans (TTTT) conformation, i.e. the electroactive β-phase, leads to highest spontaneous polarization per unit cell.²⁴ The polar γ-phase with TTTGTTG′ conformational sequence shows average piezoelectricity.²⁶

Thus, there is a great challenge to develop cost-effective and reliable methods to improve the fraction of electroactive β-phase in PVDF and its co-polymers. The polar β-phase nucleation in PVDF and its co-polymers can be obtained from the nonpolar α-phase by, for example, controlled annealing or melting crystallization under high pressure²⁷ of α PVDF and its co-polymer, the electro-spinning method,²⁸ low-temperature polar solvent crystallization,²⁹ addition of nucleation agents or surface-charged nucleation agents,³⁰ applying a strong electric field³¹ and, most recently, using supercritical carbon dioxide (SC CO₂).³² Recent studies are suggesting that the addition of nanoparticle (NPs),³³ some fillers, such as metal oxide NPs, metal NPs, carbon black, carbon nanotubes, graphene, carbon nanofibers, inorganic salts, clays, ceramic nanofillers, transition metal oxide NPs, nickel hydroxide, and some other polymer, into the polymer matrix not only effectively enhances the polar β-phase in PVDF and its co-polymers but also improves the dielectric properties of the polymer NCs thin films.^33–40 Several efforts have been made⁴¹ to improve the nucleation of β crystallites in PVDF, dielectric properties and optical properties, but rarely the study deal with use of the cost-effective electroactive PVDF–HFP to develop optically active and high dielectric electroactive PVDF–HFP NCs. Some studies on PVDF–HFP/SnO₂ composite films used as anode materials in lithium ion batteries have been reported earlier.⁴² However, studies on the effect of SnO₂ NPs on the electroactive β-phase and its dielectric properties are very rare. Tin oxide (SnO₂) is an important metal oxide and n-type wide band gap semiconductor and has outstanding optical, electrical and electrochemical properties.⁴³

Our present study deals with fabrication of low cost, high dielectric, electroactive PVDF–HFP NC thin films by incorporating the optically active SnO₂ NPs. First, SnO₂ NPs have been prepared via a hydrothermal process, treating for 12 hours at two different temperatures, 100 °C and 150 °C. Therefore, PVDF–HFP thin films with different percentages of SnO₂ NPs have been synthesized by a solution-casting method. The effect of the NPs on electroactive β-phase crystallization, optical and dielectric properties in PVDF films has been investigated thoroughly in terms of surface charge, size and the interfacial polarization, which occurs between the interfaces of the NPs and the polymer chains near the percolation threshold.

2. Experimental

2.1. Materials

The materials used in our present studies are stannic chloride (SnCl₄·5H₂O) (Merck, India), lithium hydroxide (Merck, India), poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF–HFP) pellets (Aldrich, Germany. M_w: 180 000 GPC, M_n: 71 000) and dimethyl sulfoxide (DMSO) (Merck, India).

2.2. Preparation of SnO₂ NPs and PVDF–HFP composites

0.5 M stannic chloride (SnCl₄·5H₂O) was dissolved in 50 mL de-ionized water with vigorous stirring. Then, 1.0 M LiOH was added after dissolving SnCl₄·5H₂O. The mixture was vigorously stirred for 1 hour and then transferred into two 25 mL Teflon-lined stainless-steel autoclaves. The two autoclaves were heated, one at 100 °C the other at 150 °C, for 12 hours and subsequently cooled to room temperature. The products were collected by centrifugation and rinsed with deionized water several times to adjust the pH to 7. Finally, the obtained samples were dried at 80 °C overnight and collected in vacuum for further characterizations.

PVDF–HFP/SnO₂ NPs NC thin films were prepared via simple solution-casting method. PVDF–HFP (0.2 g) was dissolved in 5 mL DMSO and the desired percentage of SnO₂ NPs added to this solution. The mixture was continuously stirred at 60 °C overnight. The mixture was then cast onto clean Petri dishes and dried at 80 °C in an oven until complete evaporation of the solvent. The films were collected and stored in a vacuum. Pure PVDF–HFP film was also prepared by the same procedure for comparison. The thickness of the synthesized films was in the range of 50–75 μm. Table 1 presents details of the different amounts of NPs added to the PVDF–HFP matrix.

Table 1 Conditions and the different amounts of SnO₂ NPs added in PVDF–HFP

Sample name	Amount of PVDF–HFP (g)	Hydrothermal temperature of SnO2 NPs (°C)	Percentage SnO2 NPs (mass%)	Amount of SnO2 NPs (g)
Pure PVDF–HFP	0.2	NA	0	0
100SO1	0.2	100	1	0.002
100SO2.5	0.2	100	2.5	0.005
100SO5	0.2	100	5	0.01
100SO7.5	0.2	100	7.5	0.015
100SO10	0.2	100	10	0.02
100SO15	0.2	100	15	0.03
150SO1	0.2	150	1	0.002
150SO2.5	0.2	150	2.5	0.005
150SO5	0.2	150	5	0.01
150SO7.5	0.2	150	7.5	0.015
150SO10	0.2	150	10	0.02
150SO15	0.2	150	15	0.03

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2.3. Characterization

2.3.1. X-ray diffraction (XRD). The XRD patterns of all the samples were recorded by X-ray diffractometer (Model-D8, Bruker AXS Inc., Madison, WI) at atmospheric pressure and room temperature using nickel-filtered Cu-Kα radiation with a scan speed of 0.3 s per step under an operating voltage of 40 kV with 2θ varying from 15° to 50°.

2.3.2. UV-visible spectroscopy. Absorption spectra of SnO₂ NPs were obtained by UV-visible spectrophotometer (UV-3101PC, Shimadzu) in the wavelength range 200–800 nm.

2.3.3. Surface charge analysis. The surface charge of the SnO₂ NPs was investigated using zeta potential measurement on a Zeta-sizer-5000 (Malvern Instruments, UK).

2.3.4. Field emission scanning electron microscopy (FESEM). The morphology and shape of the SnO₂ NPs and their interaction with the PVD–HFP matrix in the NP-modified polymer matrix were investigated using field emission scanning electron microscopy (FESEM) (INSPECT F50, the Netherlands).

2.3.5. Fourier transform infrared spectroscopy (FTIR). The effect of SnO₂ NPs on phase crystallization of the PVDF–HFP matrix was further characterized using FTIR (FTIR-8400S, Shimadzu). The absorbance data of the films were analyzed by performing 100 scans of each sample in the wavenumber range from 400 cm⁻¹ to 1100 cm⁻¹.

The fraction of β-phase (F(β)) in the films was evaluated using the Lambert–Beer law:

where A_α is the absorbance at 764 cm⁻¹, A_β is the absorbance at 840 cm⁻¹, and 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.¹⁹

2.3.6. Differential scanning calorimetry (DSC). The phase crystallization and melting behavior of pure PVDF–HFP and the SnO₂ NP-doped PVDF–HFP thin films (100SO and 150SO) were investigated using a differential scanning calorimeter (DSC-60, Shimadzu (Asia Pacific) Pte. Ltd., Singapore). All samples were heated from 80 °C to 200 °C at a heating rate of 10 °C min⁻¹ under an N₂ gas atmosphere.

The heat of melting or enthalpies of fusion (ΔH_m) of pure PVDF–HFP and SnO₂ NP-doped PVDF–HFP thin films (100SO and 150SO) were calculated from DSC curves. The degree of crystallinity (X_c) of the samples was evaluated using the equation as follows:

X_c = ΔH_m/ΔH_100% (2)

where ΔH_m is the heat of melting or enthalpy of fusion and ΔH_100% is the melting enthalpy of 100% crystalline PVDF with a value 104.6 J g⁻¹.²⁵

2.3.7. Dielectric properties measurement. The capacitance (C) and tangent loss (tan δ) were collected in the frequency range of 20 Hz to 2 MHz at room temperature using a digital LCR meter (Agilent, E4980A). 1 V AC voltage was applied across the two opposite surfaces of the samples. The dielectric constant (ε) and the AC conductivity (σ_ac) of the samples were evaluated using eqn (3) and (4), respectively, as follows:

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 the permittivity of free space (8.854 × 10⁻¹² F m⁻¹).⁴⁴

3. Results and discussion

3.1. Characterization of SnO₂ particles

We prepared SnO₂ NPs for two different hydrothermal temperatures, i.e. 100 °C and 150 °C. The XRD patterns of the as-synthesized SnO₂ NPs (synthesized at 100 and 150 °C) are show in Fig. 1a. All the diffraction peaks from X-ray reflections confirm the formation of tetragonal phase SnO₂ NPs (JCPDS no. 41-1445); there are no other peaks, suggesting high purity of the synthesized SnO₂ NPs. The UV-Vis spectroscopy of both sample 100SO and 150SO also have been shown in Fig. 1b. Absorption peaks for both samples arise at ∼255 nm.⁴⁵ Fig. 1c and d show the FESEM images of SnO₂ NPs synthesized hydrothermally at 100 °C and 150 °C, respectively. Formation of comparatively large particles with diameter ∼800 nm has been observed at 100 °C and well-defined, uniformly-distributed SnO₂ NPs with diameter ∼30–50 nm are formed at 150 °C.

Fig. 1 (a) XRD spectra of SnO₂ NPs synthesized at 100 °C and 150 °C; (b) UV-visible absorption spectra in DMSO medium and FESEM images of SnO₂ NPs: (c) 100SO and (d) 150SO.

3.2. Characterization of PVDF–HFP/SnO₂ NPs NCs

3.2.1. X-ray diffraction analysis. The X-ray diffraction (XRD) pattern of the pure PVDF–HFP and PVDF–HFP/SnO₂ NPs NC thin films has been shown in the Fig. 2. The XRD pattern previously reported in the literature reveals the presence of at least four crystal phases of PVDF (α, β, γ and δ). The diffraction spectrum of the pure PVDF–HFP film and its NCs can be analyzed by reference to the standard XRD pattern of PVDF because the HFP co-monomer does not significantly change the lattice symmetry when compared to PVDF.¹⁶

Fig. 2 XRD pattern of (a) pure PVDF–HFP and 100SO-doped PVDF–HFP thin films, (b) pure PVDF–HFP and 150SO loaded PVDF–HFP thin films. Ratio of I_20.6 and I_18.2 of pure PVDF–HFP and SnO₂ NPs-loaded PVDF–HFP thin films: (c) 100SO-loaded PVDF–HFP thin films and (d) 150SO-loaded PVDF–HFP thin films.

The standard XRD peaks of pure PVDF at 2θ = 17.7° (100), 18.4° (020), 19.9° (110), and 26.5° (021) all are assigned to the α-PVDF, whereas the 2θ = 20.8° peak, due to the superposition of the (110) and (200) plane, corresponds to β-PVDF and that at 38.9° (211) is due to the γ-phase. However, all the peaks for the pure PVDF–HFP are well matched with the standard peaks of α-PVDF.⁴⁶ However, for the PVDF–HFP/SnO₂ NPs films the intensities of all the characteristic diffraction peaks corresponding to α- and γ-phases become less significant but the diffraction peak at 20.4° ((110), (200)) appeared more strongly, indicating a conversion of the crystalline structure from the nonpolar α-phase to electroactive β-phase in the SnO₂ NPs-modified PVDF–HFP thin films. It may be suggested from the diffraction patterns that the highest β-phase conversion has occurred at 10 wt% addition of SnO₂ NPs prepared at 100 °C, whereas only 5 wt% loading of the SnO₂ NPs prepared at 150 °C in the polymer matrix obtains the maximum β-phase crystallization in NC thin films. Smaller sizes of the SnO₂ NPs prepared at 150 °C play a crucial role in developing and stabilizing more β polymorphs or crystallites in the NPs-modified PVDF–HFP thin films at low concentration. Further addition of NPs shows again the conversion of nonpolar α-crystallites.

More quantitatively, the α-phase and β-phase nucleation may be assessed through the ratio of the intensity of the peaks at 20.4° (200) and 18.2° (020) (I_20.4/I_18.2).⁴⁵ For pure PVDF–HFP thin film, this ratio is found to be 0.94. Results suggest that this ratio depends on wt% of NPs. 10 wt% loading of SnO₂ NPs prepared at 100 °C in the polymer matrix (Fig. 2c) shows the highest value ∼5.35798 and 4.61594 has been found for 5 wt% loading of SnO₂ NPs prepared at 150 °C (Fig. 2d).

3.2.2. Fourier transform infrared spectroscopy. Electroactive β-phase nucleation in SnO₂ NPs-modified PVDF–HFP thin films have been further investigated by Fourier transform infrared (FTIR) spectroscopy. FTIR is one of the best techniques to determine and quantify the different phases of PVDF–HFP.

Fig. 3 illustrates the IR spectra of the pure PVDF–HFP and NPs-loaded PVDF–HFP thin films. The absorbance bands due to nonpolar α-phase (TGTG′) centered at 488 cm⁻¹ (CF₂ wagging), 532 cm⁻¹ (CF₂ bending), 615 and 764 cm⁻¹ (CF₂ bending and skeletal bending, respectively), 796 and 976 cm⁻¹ (CH₂ rocking) are clearly observed in the spectrum of pure PVDF–HFP thin film. Two tiny bands also observed at 840 cm⁻¹ (CH₂ rocking, CF₂ stretching and skeletal C–C stretching) and 813 cm⁻¹ (CF₂ asymmetric stretching), which correspond to electroactive or polar β-phase and γ-phase, respectively.

Fig. 3 (a) FTIR spectra of pure PVDF–HFP and 100SO-doped NCs thin films, (b) evaluation of β-phase content with increasing SnO₂ NPs content from IR spectra, (c) FTIR spectra of pure PVDF–HFP and 150SO-doped NCs thin films and (d) evaluation of β-phase content with increasing SnO₂ NPs content from IR spectra.

All characteristic absorbance bands corresponding to nonpolar α-phase are present in the pure PVDF–HFP spectrum. For the 100SO/PVDF–HFP NC thin films, the IR spectra confirm the nucleation of polar β-phase by showing the strong appearance of the peaks at 444 cm⁻¹ (CF₂ rocking), 510 cm⁻¹ (CF₂ stretching), and 840 cm⁻¹ (CH₂ rocking, CF₂ stretching and skeletal C–C stretching). For the sample 100SO10, maximum enhancement of the characteristic peaks related to the electroactive β-phase (TTTT conformation) has been observed (Fig. 3a). However, for 5 wt% the doping of 150SO in PVDF–HFP matrix (Fig. 3c) obtains the maximum intensity of the main β-phase peak at 840 cm⁻¹ (CH₂ rocking, CF₂ stretching and skeletal C–C stretching).^13,14,19,25 The characteristic absorption peak related to the polar γ-phase at 812 cm⁻¹ (CF₂ asymmetric stretching)⁴⁶ is totally absent in the NC thin films.

The relative fraction of electroactive β-phase content (F(β)) has been calculated using eqn (1). The calculated value of (F(β)) for pure PVDF–HFP film has been found to be 32% and for the 100SO/PVDF–HFP thin films, a maximum ∼74% β-phase nucleation is achieved by loading 10 wt% of SnO₂ NPs synthesized at 100 °C (Fig. 3b). The relative fraction of β-phase (F(β)) attained a maximum value ∼70.9% for 5 wt% loading of SnO₂ NPs synthesized at 150 °C. The FTIR data are well consistent with the XRD results.

3.2.3. Differential scanning calorimetry. Differential scanning calorimetry (DSC) analysis has also been carried out to determine the thermal properties, such as melting enthalpy and degree of crystallinity, of the samples. DSC is the complementary technique of XRD and FTIR for the verification of different crystalline polymorphs of PVDF–HFP. Fig. 4 shows the DSC thermographs of pure PVDF–HFP and SnO₂ NPs-modified PVDF–HFP NC thin films.

Fig. 4 DSC thermographs of (a) pure PVDF–HFP 100SO-loaded PVDF–HFP thin films and (b) pure PVDF–HFP 150SO-loaded PVDF–HFP thin films.

The appearance of a broader endothermic peak at 147.86 °C in the thermograph of pure PVDF–HFP thin film suggests crystallization of the nonpolar α polymorph.⁴⁹ From Fig. 4a it can be found that the melting temperature increases to higher temperature up to 10 wt% loading of 100SO in PVDF–HFP matrix and decreases for higher loading of the NPs. A similar type of shifting of melting temperature to higher temperature in 150SO/PVDF–HFP NCs has also been observed (Fig. 4b). Maximum shifting in melting temperature to higher value is observed up to 5 wt% loading of SnO₂ NPs prepared at 150 °C in PVDF–HFP matrix. This increase in melting temperatures up to a certain concentration of NPs (10 wt% for 100SO and 5 wt% for 1500SO) may be due to the formation of the more crystalline polar β-phase in polymer thin films by catalytic action of the NPs.⁴⁸ These results are also consistent with XRD and FTIR data.

The melting enthalpies (ΔH_m) can be calculated from the DSC thermographs and then the degree of crystallinity (X_c) of the samples may also be obtained using eqn (2). Fig. 5 represents the change of ΔH_m and X_c with the NPs content. Melting enthalpy (ΔH_m) and the degree of crystallinity (X_c) increase up to 10 wt% for 100SO-loaded PVDF–HFP thin films and then the values decrease. Herein, also maximum melting enthalpy (ΔH_m) and the degree of crystallinity (X_c) value have been found for 5 wt% loading of 150SO-loaded PVDF–HFP thin films.

Fig. 5 Representation of (a) and (c) melting enthalpy and (b) and (d) degree of crystallinity of pure PVDF–HFP for the SnO₂ NPs/PVDF–HFP thin films with increasing NPs content evaluated from DSC thermographs.

3.2.4. Field emission scanning electron microscopy (FESEM). Surface morphology, NPs distribution and interaction with the polymer matrix have been investigated by field emission scanning electron microscopy (FESEM). Fig. 6a and 7a show the flat surface of pure PVDF–HFP thin film. Homogeneous and uniform distribution of SnO₂ nanostructures has been clearly observed in the SnO₂ NPs-loaded PVDF–HFP thin films up to 10 wt% loading of SnO₂ NPs synthesized at 100 °C (Fig. 6) and 5 wt% doping of SnO₂ NPs prepared at 150 °C in polymer matrix (Fig. 7). Further loading of both types of NPs leads to agglomeration of the NPs in the PVDF–HFP matrix.

Fig. 6 FESEM micrograph of (a) pure PVDF–HFP, (b) 100SO2.5, (c) 100SO5, (d) 100SO7.5, (e) 100SO10 and (f) 100SO15 samples.

Fig. 7 FESEM micrograph of (a) pure PVDF–HFP (b) 150SO2.5, (c) 150SO5, (d) 150SO7.5, (e) 150SO10 and (f) 150SO15 samples.

3.2.5. Mechanism of electroactive β-phase nucleation. Enhancement of electroactive or polar β-phase in the as-synthesized samples was investigated form XRD, FTIR and DSC data. FESEM micrographs confirm the strong interaction between the NPs' surfaces and the polymer matrix. However, the original β-phase transformation mechanism is still not been discussed clearly. It is essential to understand the mechanism of the transformations of the electroactive β-phase. Our previous study of clay mineral–PVDF composites, Fe₂O₃/Co₃O₄ NPs–PVDF NCs, and WO₃·H₂O NPs–PVDF NCs suggest that the strong interaction between positive –CH₂ dipoles and the negatively-charged surface of NPs or clay mineral induced the longer TTTT conformation on the NPs clay mineral surfaces.^19,25,41,45 To describe the phase transformation mechanism properly, we investigated the surface charge, i.e., zeta potential, of the NPs at the same pH 6 of the interaction mixture, i.e., the NPs/PVDF–HFP solution in DMSO. Zeta potential studies show negative values suggesting that the surfaces of the SnO₂ NPs are negatively-charged for both samples. When NC solutions are cast on Petri dishes and dried at 80 °C, strong interaction between the positive dipole –CH₂ and the negatively-charged surface of the NPs improves the polymer chains in longer all trans (TTTT) conformations in the NPs-loaded PVDF–HFP thin films. The possible electroactive β-phase formation mechanism and structure of PVDF–HFP are schematically shown in Fig. 8.

Fig. 8 (a) Possible mechanism of electroactive β-phase formation and (b) chemical structure of PVDF–HFP.

3.2.6. Photoluminescence spectroscopy. Fig. 9 shows photoluminescence emission spectra at room temperature of the SnO₂ NPs (100SO and 150SO)/PVDF–HFP NC thin films. In Fig. 9a and c, the strong emission peak around wavelength 430 nm has been observed at the excitation value, λ_ex = 256 nm. The peak intensity gradually increases with the increasing loading concentration of the NPs for both samples. At the same excitation value, another small emission peak centered at 355 nm is also noticed in Fig. 10a and c. Interestingly, when the excitation wavelength has been increased to λ_ex = 303 nm, the NC samples show two emission peaks, one centered at 355 nm, which appears strongly and another at ∼495 nm. Thus the NPs-doped PVDF–HFP films pose very interesting tunable photoluminescence emission properties both in the UV and visible regions making this material suitable for optoelectronics applications. Earlier reports showed that the SnO₂ thin films exhibit a board emission peak near 391 nm (E_g ∼ 3.1 eV).⁵⁰ As all the emission maxima centered at 430 nm (2.9 eV), 355 nm (3.49 eV) and 495 nm (2.50 eV) are lower than the band gap of the SnO₂ nanocrystal (E_g ∼ 4.1 eV), we cannot assign the emission bands to the direct recombination of a conduction electron in the Sn 4p band and a hole in the O 2p valence band. Therefore, all these broad peaks may originate from different luminescence centers such as dangling, tin interstitials or various defects. Oxygen vacancies, the intrinsic defects known as the common defect for nanocrystalline oxides, usually act as the radiative center in the luminescence process. The oxygen vacancies may be present in the oxides in the three different states such as , and states. The oxygen vacancy is an intrinsic donor in SnO₂.⁵¹ Because is a very shallow donor, it is expected that most oxygen vacancies will be in their paramagnetic states under the flat band condition and thus the emission of SnO₂NPs/PVDF–HFP thin films in the UV region is directly attributed to the formation of oxygen vacancies in paramagnetic states. The visible broad emission in our samples around 400–500 nm may be assigned to the formation of a luminescent center. The recombination of a conduction band electron with the center results in visible emission in the samples when the surface-trapped hole recombines with an electron in a deeply trapped to form a center.⁵²

Fig. 9 Photo luminescence spectra of pure PVDF–HFP and SnO₂ NPs-PVDF–HFP thin films: (a) and (c) λ_ex = 256 nm and (b) and (d) λ_ex = 303 nm at ambient conditions.

Fig. 10 NPs content-dependent dielectric constants and tangent losses at 20 Hz and 1 MHz for the samples 100SO0, 100SO1, 100SO2.5, 100SO5, 100SO7.5, 100SO10 and 100SO15 are shown in (a) and (b), respectively and for samples 150SO0, 150SO1, 150SO2.5, 150SO5, 150SO7.5, 150SO10 and 150SO15 are show in (c) and (d), respectively. (Insets represent the best fitted curve in accordance with percolation theory).

3.2.7. Dielectric properties.
3.2.7.1. SnO₂ NPs content dependence. The dependence of dielectric constant and tangent loss of SnO₂ NPs-loaded PVDF–HFP thin films on NPs content at two different frequencies, 20 Hz and 1 MHz, have been shown in Fig. 10. Remarkable improvement in dielectric constant has been observed due to incorporation of both NPs in the PVDF–HFP matrix. The dielectric constant has been observed to increase up to 10 wt% loading of the SnO₂ NPs synthesized at 100 °C in PVDF–HFP in both frequency and dielectric value but has decreased for further loading of the NPs (Fig. 10a and b). The maximum values of the dielectric constant, ∼2578.45 at 20 Hz and 400.6 at 1 MHz, have been achieved by embedding 10 wt% loading of SnO₂ NPs synthesized at 100 °C. Only 5 wt% incorporation of SnO₂ NPs prepared at 150 °C into the polymer matrix shows a fast improvement in dielectric constant, with values ∼865.55 at 20 Hz and 200.4 at 1 MHz, and then a decrease in dielectric constant has been noted for further loading of 150SO. Tangent losses for both the NPs-loaded samples (100SO and 150SO) have increased non-linearly due to the occurrence of different Debye-like relaxation.^16,25

The improvement of dielectric constant is very slow up to 7.5 wt% for 100SO-doped polymer thin films and 2.5 wt% for 150SO-loaded samples, but the dielectric value increase very fast at 10 and 5 wt% addition of 100SO and 150SO in PVDF–HFP matrix, respectively, which may be due to reaching the percolation threshold value.^16,25 The remarkable improvement of dielectric constants may be explained according to percolation theory as follows:

ε = ε₀(1 − f/f_c)^−q (5)

where ε₀ is the dielectric constant of pure PVDF–HFP, f is the weight fraction of the NPs, f_c is the percolation threshold value, and q is the critical exponent.⁵³ The experimental dielectric values are in good agreement with the percolation eqn (5). Best fitted curves according to the percolation equation are shown in the inset of Fig. 10a and c. The dielectric constant upon reaching the percolation threshold of the composite system shows very large values in dielectric constant, which are mainly due to the formation of long-range connectivity of the SnO₂ NPs in the polymer matrix.⁵⁴ Herein, the semiconductive SnO₂ nanoparticles act as a weak conductive phase upon reaching the percolation threshold value. Upon reaching the percolation threshold, the distance between SnO₂ nanoparticles in polymer matrix is small enough for electron transport in the matrix by the tunneling effect.⁵⁴ Second, the enhancement of dielectric constant may also be explained easily with the help of Maxwell–Wagner–Sillars (MWS) interfacial polarization at low frequency.^25,47 The SnO₂ NPs are uniformly distributed and well-separated from each other in the polymer matrix at low loading concentration (up to 10 wt% of 100SO and 5 wt% of 150SO), which makes large interfaces and between the NPs and polymer chains.⁴⁶ Thus, large accumulation of space charge between the interfaces of SnO₂ NPs and the polymer matrix at the respective concentration (10 wt% of 100SO and 5 wt% of 150SO) of the NPs results in a large improvement in dielectric value due to MWS effect.^1,12 Further loading leads to agglomeration of the NPs in the polymer matrix resulting in reduction in the effective interaction interfacial interfaces between the SnO₂ NPs and PVDF–HFP chains. Thus, reduction in dielectric constant has been observed for both NPs-loaded samples after reaching their respective percolation concentrations.

Fig. 11a and b display variation of the AC conductivity of the samples with different concentration of NPs at 20 Hz, 10 kHz, 100 kHz and 1 MHz. Almost sharp increase in the AC conductivity of the NPs-doped samples has been noticed upon the addition of 100SO up to 10 wt% and 150SO up to 5 wt% at the high frequency of 1 MHz, but for lower frequencies (100 kHz, 10 kHz and 20 Hz), the variation of AC conductivity of the NC thin films with NPs content shows a slow increase without altering the nature of variation at 1 MHz. With increase of the frequency the AC conductivity also increase on the basis of an apparent power law: σ(ω) = ωⁿ (0.6 ≤ n ≤ 1).¹⁶

Fig. 11 SnO₂ NPs content dependence of AC conductivities at 20 Hz, 10 kHz, 100 kHz and 1 MHz (a) 100SO-loaded samples and (b) 150SO-loaded samples.

3.2.7.2. Frequency dependence of dielectric properties. The frequency dependence of dielectric constant, tangent loss (tan δ) and AC conductivity of pure PVDF–HFP and PVDF–HFP/SnO₂ NPs NC thin films at ambient conditions have been shown in Fig. 12. The dielectric constant of the samples decreases with frequency increasing when the NPs concentration increases up to 15 wt% for 100SO- and 150SO-doped polymer nanocomposite films. At low frequency, space charge, i.e., dipoles, may follow the direction of the applied electric field, but for higher frequencies the dipoles are unable to follow the electric field and lag behind the electric field resulting in the decrease of dielectric constant.^12,16 At the maximum value of frequency, ∼1/t, the dielectric constant shows a relaxation (t is the relaxation time).⁴⁶ This reduction of dielectric constant can also be explained by the MWS polarization effect. Upon the application of the electric field, the charges move freely and accumulate at the interfaces between the polymer matrix and the NPs due to differences in conductivity and dielectric properties of insulating PVDF–HFP and semiconducting SnO₂ NPs.²⁵ At low frequencies, the movement and transfer of charge carriers occurs effortlessly due to there being enough time to move to the electrode, but at higher frequencies, the interfacial movement of the dipoles is restricted resulting in reduction of the dielectric constant. Similarly, tan δ for all samples decreases with increasing frequency in the same manner. Only the tan δ of the 150SO5 sample shows a Debye-like dipolar relaxation peak (Fig. 12e). The highest dielectric constants observed have been ε ∼ 2578.45 for 100SO10 and ε ∼ 865.55 for 150SO5 at 20 Hz.

Fig. 12 Frequency dependence of dielectric properties of pure PVDF–HFP and 100SO- and 150SO-modified PVDF–HFP thin films: (a) and (d) dielectric constant, (b) and (e) tangent loss and (c) and (f) AC conductivity.

Fig. 12c and f represent the variation of the AC conductivity of pure PVDF–HFP and NC thin films with increasing frequency. The AC conductivities of the samples increase almost linearly with increasing frequency. The conductivity of the sample can be divided into two regions, one is for the lower frequency range dominated by DC conductivity, and the other region is a higher frequency region characterized by AC conductivity, which mainly depends on frequency. In our samples, a lower frequency plateau is totally absent. Thus, our samples pose only AC conductivities. Therefore, the increase in AC conductivity of the NC samples with increasing frequency and up to a certain percentage of the fillers (10 wt% for 100SO and 5 wt% for 150SO), is mainly due to the MWS interfacial contribution and dipolar relaxation modes.⁴⁶

4. Conclusions

The effect of SnO₂ NPs (of two different sizes) on the electroactive β-phase crystallization, photoluminescence emissions and dielectric properties of PVDF–HFP thin films has been studied in detail in the present study. Incorporation of SnO₂ NPs promotes ∼74% electroactive β-phase nucleation for 10 wt% loading of NPs synthesized at 100 °C and ∼70.9% for 5 wt% addition of NPs prepared at 150 °C. Strong ion-dipole or electrostatic interaction between the –CH₂ dipoles and negatively-charged NPs' surfaces transform the polymer chains from nonpolar TGTG′ conformation to polar all trans (TTTT) conformation or electroactive β polymorphs. Tunable photoluminescence emissions, both in the UV and visible regions, have been observed for both NPs-loaded samples. Remarkable improvement in dielectric constant due to incorporation of the SnO₂ NPs in the PVDF–HFP matrix has also been noticed. The very large increase of dielectric constant at certain doping concentrations is due to reaching the percolation threshold value and the large accumulation of charge due to the MWS effect. The highest dielectric constants achieved were ∼2578.45 for 10 wt% loading of 100SO in PVDF–HFP matrix and ∼865.55 for 5 wt% loading of 150SO, at 20 Hz. Therefore, these successfully fabricated electroactive, optically active and high dielectric polymer NC thin films may be used suitably to develop the energy harvesters, sensors, actuators, energy storage devices, thin film capacitors, solar cells, optical detectors, optoelectronic devices and piezoelectric nanogenerators.

Acknowledgements

Authors acknowledge University Grand Commission (UGC), Government of India for the financial support to this study.

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Footnote

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

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