The Effect of PVP on Thermal, Mechanical, and Dielectric Properties in PVDF-HFP/PVP Thin Film

Abstract

In this research, the influences of the addition of PVP to PVDF-HFP polymers and the preparation of thin films using a solvent casting method were studied. The PVDF-HFP and polymer blend PVDF-HFP/PVP thin films with a nanostructured surface were investigated using scanning electron microscopy, differential scanning calorimetry, nanoindentation, and dielectric spectroscopy. The results showed that the PVP formed a dispersed phase (the poorer conductive islands) in the PVDF-HFP polymer matrix, which reduced its mechanical properties. The crystallinity of PVDF-HFP polymer decreased with the addition of PVP by 7.4%, but the PVP induced the formation of the polar β-phase of PVDF-HFP. Therefore, an improved dielectric response is expected, but it was not significantly improved even though the polar β-phase was detected. The contrasting effect was attributed to less conductive PVP islands on the surface of the PVDF-HFP/PVP polymer blend, which decreased its conductivity.

1. Introduction

Polymer composites, based on poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP) polymer, can be employed in many electronics, i.e., sensor systems [1,2], actuators [3], nanogenerators [4], microelectronic components, and medical applications [5], due to their cost-effective, mechanically flexible, and high-temperature stability. The chemically inert polymer PVDF-HFP has been used in numerous scientific fields, e.g., as membranes [6], polymer electrolyte for batteries [7,8], electrospinning produced fibres [9,10], as supercapacitors [11], and so on. The wide range of usage of PVDF-HFP polymers is a consequence of its ferroelectric, piezoelectric, and pyroelectric properties [12,13], good mechanical strength [14], thermal stability of up to 143 °C [6,15,16], and low degree of crystallinity (about 50%) [17]. For a thin film of PVDF-HFP/PVP, it has already been proven that it can show antimicrobial activities against Gram-positive bacteria L. monocytogenes after 3 and 6 h of exposure, but a reduction in colony forming units (CFU) was observed for P. anomala and A. flavus after 6 h [18].

The semi-crystalline properties of PVDF-HFP polymer are strongly dependent on its preparation conditions, solvents, and additives, which affects the degree of crystallinity and the presence of individual phases (α, β, γ, δ, and ε). The polar γ-phase occurrence increases with crystallization temperatures and time [19]. Still, the non-polar α-phase is the most common and is usually present after melt crystallization at temperatures below 160 °C. The most polar β-phase of PVDF-HFP is formed by mechanical deformation of α-phase PVDF-HFP film or by doping the polymer with fillers, most commonly with organic nanoparticles [20] or clay minerals [21].

As we already showed in our previous research [18], the presence of polar β-phase was induced by the addition of amorphous water-soluble polyvinylpyrrolidone (PVP) polymers. PVP has a weakly basic nature; therefore, it is capable of forming complexes not only with anionic reagents but also with ionic polymers [22]. For polymer blend PVDF/PMMA, it has been shown that 25% of β-phase was present after in situ blend crystallizations from ethanol with the addition of 1 wt.% of PMMA. In this study, it has been indicated that a good mixing ability of polymers results from the presence of hydrogen bonding, which increases the movement resistance of PVDF chain segments to the crystal front [23]. Therefore, the PVDF chain is forced to form its extended all-trans conformation that leads to crystallization in the β-phase. In another study, the polyaniline nanorods were dispersed in PVP and then incorporated into PVDF and, therefore, enabled the formation of dPANI@PVP/PVDF nanocomposites with a relatively low dielectric loss while maintaining a sufficiently high dielectric constant [24]. Moreover, the percolative nanocomposite Ag@PVP/PVDF, which is based on PVDF with the addition of silver core–shell nanoparticles, coated with the 5–10 nm layer of PVP for better dispersion into the matrix, and possessing low dielectric loss and high permittivity, has been prepared [25]. In the literature, some studies also described that the surface modification of ferroelectric particles (e.g., BaTiO₃) with PVP polymers, exhibiting fine core–shell structure, which was homogeneously dispersed in the PVDF matrix; in this way, new materials for energy storage have been prepared. With 40 vol% of BaTiO₃ loading, the prepared materials exhibited the largest dielectric constant of 65 at 25 °C and 1 kHz [26].

In recent years, many studies have been performed on polymers and their composites using the nanoindentation technique [27]. Nanoindentation measurements proved that the incorporation of various organic, inorganic, or organic-inorganic fillers help improve or modulate the mechanical properties of PVDF [28,29,30].

However, only a few studies were performed on PVDF copolymer (PVDF-HFP) films when determining their mechanical properties. The research of Yuennan and Muensit [31] showed that the addition of magnesium chloride hexahydrate (MgCl₂·6H₂O) fillers to PVDF-HFP acted as a plasticizer and improved the flexibility of the film, which is advantageous in piezoelectric applications.

In this article, we study the effect of the addition of PVP polymers with polar pyrrolidinone substituents on the degree of interactions with the polar polymer matrix PVDF-HFP. Furthermore, we also investigated how the formation of polar β-phase affects mechanical and dielectric properties.

2. Materials and Methods

The PVDF-HFP, obtained from Sigma Aldrich (St. Louis, MO, USA), was dissolved in dimethylformamide (DMF), Carlo Erba Reagents (Cornaredo, Italy), by mixing for 4 h on a magnetic stirrer with 400 rpm at 80 °C. After achieving a homogenous solution, the thin film of PVDF-HFP was formed on a Teflon plate by solution casting and drying for 2 h at 80 °C.

Polymer solution PVDF-HFP/PVP was obtained by dissolving PVDF-HFP and PVP Sigma Aldrich (St. Louis, MO, USA) in DMF separately for 4 h on a magnetic stirrer with 400 rpm at 80 °C. The final PVDF-HFP/PVP solution was prepared by adding the PVP solution into the PVDF-HFP one in mass ratio 25:75 by mixing both components on a magnetic stirrer for additional 2 h. The thin film of PVDF-HFP/PVP was prepared in the same way and at the same conditions as explained above for the PVDF-HFP film.

The morphology of formed thin films was investigated with a scanning electron microscope using a field-emission gun Supra 36 VP, Carl Zeiss (Braunschweig, Germany). The accelerated voltage was 1.5 kV, and the working distance 4.3 mm and the figures were formed from the signal of secondary electrons. The thin films were placed on an adhesive carbon tape and sputtered with 10 nm of carbon to ensure better sample conductivity. Thermal properties of polymers and polymer blend were determined by using differential scanning calorimetry—DSC (Q2500, TA Instruments, New Castle, DE, USA). Heat–cool–reheat tests were performed according to standard ISO 11,357 in the temperature range from −80 °C to 250 °C for PVP, from −80 °C to 300 °C for PVDF-HFP, and from 40 °C to 300 °C for polymer blend PVDF-HFP/PVP with heating and cooling rates of 10 °C/min under an inert (nitrogen) atmosphere. To determine the values of the glass transition, melting and crystallization temperatures, and enthalpies, the TRIOS software was used. The degree of crystallinity of PVDF-HFP polymer was calculated using the following equation:

$${\chi }_c=\frac{\Delta H_f}{\Delta H_f^0} ·100 \%$$

(1)

where Χ_c is the degree of crystallinity, $\Delta H_f$ is the fusion enthalpy (calculated from DSC curve), and $\Delta H_f^0$ is the fusion enthalpy for 100% crystallinity (for PVDF-HFP polymer, the value of 104.7 J/g was used [6]). The degree of crystallinity for polymer blend PVDF-HFP/PVP was calculated with the same equation, where the values of crystallinity were multiplied with weight fraction coefficient 1/φ, where φ represents the weight fraction of the PVDF-HFP as the crystallized component (in our case 0.75).

The XRD diffractogram of the PVDF-HFP/PVP polymer blend was performed on a D4 Endeavor diffractometer (Bruker Corporation, Billerica, MA, USA) at room temperature using a quartz monochromator Cu Kα radiation source (λ = 0.1541 nm) with a Sol-X dispersive detector in the range of 2θ from 5° to 65° (step size 0.04°). The XRD diffractogram has been smoothed in the computer program Origin. Optical images of spherulites were taken with a polarized light optical Microscope Axioskop 2, Carl Zeiss (Oberkochen, Germany) at the magnification 200×.

Mechanical properties were characterized with nanoindentation techniques on the surface of the tested thin film samples. Nanoindentation is a useful non-destructive technique for determining mechanical properties such as elastic moduli and hardness at the nano-scale of many materials in bulk or as a coating deposited on the surface [32,33]. All tests were performed on a Nanoindenter G200 XP instrument manufactured by Agilent Technologies, Inc (Santa Clara, CA, USA). Continuous Stiffness Measurement (CSM) was performed using a standard three-sided pyramidal Berkovich probe with the tip oscillation frequency of 45 Hz and 2 nm harmonic amplitude. For the characterization of the samples, several sheets were cut and placed into a holder (diameter 8 mm, height 1 mm), melted at 140 °C for 10 min, and cooled down in air. Thirty-six indents were performed onto two samples for each coating with a 200 µm distance between adjacent indentations to exclude interaction effects. The highest depth of the indents was 2500 nm; however, for the calculation of the elastic modulus and hardness, values at the depths of 1000 and 2000 nm were used. For analyzing the results, the Poisson ratio of 0.33 was used for PVDF-HFP [34]. All measurements were conducted at room temperature.

The contact modulus (CM) method was used to determine the dynamic properties (storage and loss modulus) of films using a 100 µm flat-punch tip. The viscoelastic properties were determined at various frequencies from 45 Hz to 1 Hz. A pre-compression of 5 µm was used to make sure that the tip is in contact with the surface. Fifteen tests were performed on each sample, and averaged values are provided as a result.

Dielectric properties of PVDF-HFP and PVDF-HFP/PVP thin films were measured with a Precision LCR Meter (HP 4284A) using the amplitude of the probing AC electric signal of 1 V. Two-layer electrodes were sputtered on the thin films: 10 nm of chromium for better adhesion and 100 nm of gold. The complex dielectric constant ε* = ε′ − iε″ was measured during cooling in the temperature range between 400 K and 150 K at a cooling rate of 1 K/min. The temperature of the samples was controlled by a lock-in bridge technique with a platinum resistor Pt100 and was stabilized within ±0.01 K. The real part of the complex AC conductivity σ* = σ′ + iσ″ was calculated via equation ${\sigma }^{\prime }=2\pi \nu {\epsilon }_0{\epsilon }^{″}$, where ${\epsilon }_0$ is the permittivity of free space.

3. Results

3.1. Morphology

The surface topography of the transparent thin films of PVDF-HFP and polymer blend PVDF-HFP/PVP is shown in Figure 1. On the surface of the PVDF-HFP thin film (Figure 1A), some random spherical structures are visible. The addition of the PVP polymer into the PVDF-HFP (PVDF-HFP/PVP polymer blend, Figure 1B) led to the formation of a nanostructured surface with spherical structures of an average diameter of 200 to 500 nm.

3.2. Thermal Properties

The DSC curves that indicate the thermal properties of PVP and both thin films are presented in the Figure 2. The values of the glass transition temperature (T_g), melting temperature (T_m), and melting enthalpy (ΔH) were determined from the re-heating curve (Figure 2A), while crystallization temperatures and enthalpies were determined from the cooling curve (Figure 2B). The obtained values of temperatures of phase transitions and fusion enthalpy are summarized in

Table 1. Parameters of the DSC curves.

Thin Films	Tg (°C)	Tm (°C)	ΔHm (J/g)	Tc (°C)	ΔHc (J/g)
PVP	164.8	-	-	-	-
PVDF-HFP	−33.8	139.2	31.2	99.2	30.3
PVDF-HFP/PVP	−20.3	115.7	16.9	69.6	12.9

. The PVP is an amorphous polymer with an average molecular weight of 40 kDa and exhibits only the glass transition temperature at 164.8 °C. For PVDF-HFP thin films, the crystallinity was 28.9%, while the addition of 25% of the PVP polymer (polymer blend PVDF-HFP/PVP) decreased the crystallinity to 21.6%.

PVP is a hygroscopic amorphous polymer and, therefore, a broad endothermic peak ranging from 80 to 120 °C was observed in the DSC curve (the first heating cycle is not shown). This peak was attributed to the presence of a residual moisture [35]. In polymers, water acts as a plasticizer [36], which can be observed in the DSC thermogram as a shift in phase transition temperatures to lower temperatures. Moreover, the literature reports that the presence of water in PVP polymer affects the increasing mobility of the polymer blend [37], which could also increase the dielectric response of such thin films.

3.3. Crystal Structures

The XRD diffractogram of the PVDF-HFP/PVP polymer blend is presented in Figure 3A. It can be clearly seen that the multi-phase crystal structure in polymer blend was observed. The XRD diffractogram of PVDF-HFP/PVP polymer blends, presented according to our previous publication, show a broad peak centered at 2θ = 20°, which was attributed to the presence of the β-phase, as a result of the sum of the diffraction in (110) and (200) planes [37]. The next, less intensive peak, centered at 26.6°, was attributed to the diffraction of the crystal plane (021) and corresponds to the presence of mixed phases: α and γ-phase [37]. The predominant crystalline phase of PVDF-HFP in the polymer blend was attributed to the γ phase, because the most intensive peak was observed at about 2θ = 40°. The spherulites, formatted during the crystallization, were observed by a polarizable light optical microscope, and they are shown in Figure 3B. A greater share of the small and medium crystal grains was attributed to the presence of α and β phase of PVDF, whereas larger crystal grains are present in the γ-phase [37].

3.4. Mechanical Properties

The mechanical properties of PVDF-HFP and PVDF-HFP/PVP were determined using a nanoindentation technique, and the results are presented in the Figure 4. The results show that the addition of PVP to PVDF-HFP decreased the elastic modulus and hardness of the thin film. The elastic modulus of the PVDF-HFP/PVP film was 1.61 ± 0.01 GPa and 1.99 ± 0.02 GPa of the film without PVP. A larger difference between both samples was observed by the determination of surface hardness, which was for PVDF-HFP/PVP almost twice lower compared to PVDF-HFP.

Figure 5A shows a variation in the storage modulus (E′), loss modulus (E″), and loss factor (tan δ) with the modulation of frequency at a depth of 5 µm for PVDF-HFP and PVDF-HFP/PVP, respectively. For both thin films, the values of storage modulus were found to be two orders of magnitude higher compared to the loss modulus. However, the addition of PVP to PVDF-HFP (PVDF-HFP/PVP thin film) decreased both the E′ and E′′ modulus. The E′ for both thin films decreased for about 0.2 GPa with the decreasing frequency, while the E″ was almost independent upon frequency changes. The loss factor, which represents the ratio of E′/E″, exhibited the same trend for both thin films, with lower values for PVDF-HFP/PVP (Figure 5B).

3.5. Dielectric Properties

The results of dielectric measurements are presented in Figure 6. The real and imaginary parts of the complex dielectric constant ε* are shown along with the real part of the complex AC conductivity at four frequencies (1, 10, 100 kHz, and 1 MHz) in the temperature range between −123 °C and 127 °C (150 and 400 K). Imaginary part ε″ represents dielectric losses or, in another words, the electrical conductivity of the system. The dielectric response of the PVDF-HFP thin film shows dielectric relaxation in the temperature range between −73 °C and 27 °C (200 and 300 K) due to the dynamic transition from the solid glass state to the soft rubber phase occurring in the amorphous part of the PVDF-HFP polymer [38]. The results show that, at room temperature, the dielectric constant of PVDF-HFP was about 10 and 9 for the polymer blend of PVDF-HFP/PVP, respectively. The characteristic dynamic peaks in the graph of dielectric losses (ε”) on the temperature scale occured at the same temperature (24 °C or 297 K) in both thin films. The addition of PVP polymer did not cause any significant changes in the values of ε′, ε″, and σ′ or in the structure of PVDF-HFP polymer. We observed that the electrical conductivity slightly decreased after the addition of the amorphous PVP polymer. The addition of PVP polymer simultaneously reduced the proportion of crystalline phases in the sample (which consequently reduces the relaxation intensity [38], i.e., the relaxation peaks in ε″ and the relaxation strength in ε′), but on the other hand, it triggered the crystallization of PVDF-HFP polymers into the ferroelectric polar phase and stabilized it; no significant change in the dielectric response was observed.

4. Discussion

Some studies in the literature have already reported that the PVP polymer enhances thermal stability, crosslinking, and mechanical strength if added to the composite because pyrrolidone groups can enable a homogeneous incorporation of organic and inorganic salts [39]. However, according to our knowledge, no research studies on thin films involving the PVDF-HFP polymer have been conducted. The goal of our research was, therefore, to study such films and characterize their physical properties, i.e., optical, thermal, mechanical, and dielectric properties. Physical properties are influenced also by the crystalline phase in the polymorphous matrix polymer PVDF-HFP. Technologically, the most interesting phase is the polar β-phase, which exhibits pyro, piezoelectric, and dielectric properties. It has already been shown that the predominance of the β-phase could be determined from the temperature of the solvent evaporation; the majority of β-phase was formed at temperatures below 70 °C. Between 70 and 110 °C, a mixture of α- and β-phase was formed, and above 110 °C, the γ-phase became dominant [19]. It has already been proven in our previous research that although the temperature of the crystallization process of PVDF-HFP thin films was 80 °C, Raman spectroscopy confirmed the presence of the β-phase in the PVDF-HFP/PVP polymer blend. On the other hand, the XRD diffractogram revealed that the γ-phase was dominant as the most intensive peak was observed at about 40°. In the XRD diffractogram (Figure 3A), the peaks of β-phase were also present, but the intensity was lower compared to the γ-phase. However, the β-phase was not detected in pure PVDF-HFP thin films prepared by crystallization at the same conditions [18]. It has been reported that the interphases between the crystalline and amorphous phases maintain the conformational characteristics of the crystalline regions [19]; therefore, we expect the improvement of the mechanical and dielectric response of the PVDF-HFP/PVP polymer blend. In particular, this expectation is because Li et al. expected an interaction between the polar pyrrolidinone substituents of PVP and the polar polymer matrix (PVDF-HFP) and, thus, an improvement in dielectric properties [25].

As already mentioned, the crystallization of PVDF-HFP thin films was performed at 80 °C. The results, already published in [21], showed that after crystalization, the film exhibited a rich multi-phase structure, including the presence of the polar β-phase. As already mentioned, the crystallization of PVDF-HFP thin film was performed at 80 °C. The results, already published in [21], showed that after crystallization, the film exhibited a rich multiphase structure, including the presence of the polar β-phase. However, the β-phase in the composite, prepared by “drawing at low temperatures”, did not exhibited piezoelectric properties due to the random orientation of dipoles [40]. It turned out that in decreased electric fields, the α-phase was the most polarizable, slightly less β-phase, while due to the lowest remanent polarization, the γ-phase was the least polarizable 37. The loss of polarisation is attributed to the greater mobility of polymer molecular chains [41]. However, at the γ-phase of the PVDF, the crystal grains grow side by side, and a little block space between the adjacent crystal grains can be observed [37]. In our case, where the multiphase structure was observed, the crystal grains had different sizes. The larger light spherulites represented the γ-phase (Figure 3B).

After the addition of PVP, the degree of crystallinity in the formed PVDF-HFP/PVP film decreased, but we detected that the addition of a PVP polymer affected the PVDF-HFP crystal structure, while it crystallized in polar β- and γ-phases [18]. For the presence of the β-phase, of the formation of hydrogen bonds between hydroxyl groups (–OH) and carbonfluorine (–CF) is desirable, but the critical role for the polar molecular conformation of PVDF-HFP is to possess -OH groups [42]. In the following, the more intensive peak at 839 cm⁻¹ of the PVDF-HFP/PVP thin film, detected with Raman spectroscopy, confirmed that the addition of PVP resulted in the stabilization of the polar β-phase of the PVDF-HFP polymer [43].

However, the addition of PVP polymer deteriorates the mechanical properties of the PVDF-HFP/PVP thin film compared to the PVDF-HFP thin film. The elastic modulus of the PVDF-HFP/PVP polymer blend decreased by around 20%. At higher contents of PVP polymer, stronger linkages may occur between the components of the polymer blends and the polymeric chains, which can reduce the motion of molecules [44]. A higher amount of PVP accumulated in the dispersed phase (PVP islands on the surface, Figure 1B) in the PVDF-HFP/PVP thin film, which decreases the mechanical properties of films [44]. The hardness of the polymer blend reduced for around 50%.

Due to the confirmed presence of the polar β-phase [18], we expected stronger dielectric responses from the films, prepared by the addition of PVP. However, the micrographs showed that nanostructured surface of the PVDF-HFP/PVP thin film contained PVP island structures (diameter between 200 and 500 nm, height 150 nm) with poorer conductivity. It turned out that these islands overshadowed the increase in dielectric responses and decreased the conductivity of thin PVDF-HFP/PVP films.

5. Conclusions

A polymer blend based on inert PVDF-HFP with the addition of PVP was successfully prepared and characterized by its thermal, mechanical, and dielectric properties. The addition of 25% of PVP polymer slightly decreased the elastic modulus of the polymer blend and reduced the thin film’s hardness by half. Despite the detection of the ferroelectric β-phase of PVDF-HFP polymer, the polymer blend PVDF-HFP/PVP did not exhibit an improved dielectric response, because the PVP was present in the dispersed phase, which led to the formation of less conductive islands. However, the dielectric properties of thin films improved in a moisture environment, since their conductivity increased. With the presented results, we can conclude that the PVDF-HFP/PVP polymer blend can be further used in conductive biocompatible components in the field of biomedicine, sensors, and smart scaffolds. For future studies, crystallization parameters will be optimized in a manner in which large amounts of β phase will be detected. After that, the polarity of the thin film will be checked and the nanogenerators will be possibly formed.