# Physico-Chemical Understanding of Plasticizers Interaction with P(VDF-TrFE-CFE) Electroactive Polymer

Colin Lesenne, David Audigier, Jean-fabien Capsal

► To cite this version:

Colin Lesenne, David Audigier, Jean-fabien Capsal. Physico-Chemical Understanding of Plasticizers Interaction with P(VDF-TrFE-CFE) Electroactive Polymer. Macromolecular Materials and Engineering, In press, 10.1002/mame.202500275. [https://doi.org/10.1002/mame.202500275](https://doi.org/10.1002/mame.202500275)

HAL Id: hal-05329221

https://hal.science/hal-05329221v1

Submitted on 24 Oct 2025

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

<figure><img src="image_1.png" /></figure>

Distributed under a Creative Commons Attribution 4.0 International License
RESEARCH ARTICLE

OPEN ACCESS

# Physico-Chemical Understanding of Plasticizers Interaction with P(VDF-TrFE-CFE) Electroactive Polymer

Colin Lesenne

David Audigier

Jean-Fabien Capsal

INSA Lyon, LGEF, UR682, Villeurbanne, France

**Correspondence:** Jean-Fabien Capsal (jean-fabien.capsal@insa-lyon.fr)

**Received:** 7 July 2025 |

**Revised:** 7 October 2025

**Funding:** This work was supported by the European Union-funded Live-Mirror project (grant no. 101099220) from the EIC Horizon Pathfinder program.

**Keywords:** actuators | broadband dielectric spectroscopy | di-n-octyl phthalate | DNOP | FTIR | halogen bonds | plasticizer retention | P(VDF-TrFE-CFE)

## ABSTRACT

The introduction of plasticizers in P(VDF-TrFE-CFE), a relaxor ferroelectric polymer, has shown major improvements over the past few years. The drastic gains in softness and dielectric permittivity at low frequency are undeniably relevant for actuation, but the interactions involved in such mechanism remain unclear. A parametric study of the terpolymer composition on its interaction with di-n-octyl phthalate (DNOP) is investigated here. In combination with a room-temperature Curie-like transition obtained for VDF/(VDF+TrFE) ratio below two-thirds, a large proportion of TrFE, contrary to CTFE-based terpolymers, leads to plasticized films of Young's moduli below 50 MPa. It was disclosed by FTIR that the amount of termonomer triggers the retention of DNOP, here theorized by electrostatic interactions, in accordance with the saturation of the mechanical properties for low-CFE content terpolymers. For the first time, this study unveils a close synergy between chemical affinity and physical properties of plasticized terpolymers. As a result, an extremely soft, crack-free terpolymer film (Y = 17.6 MPa) was successfully synthesized with 20 wt.% of DNOP. Although its electrical breakdown strength is undoubtedly affected, a combination of high molecular weight and purity unlocks great potential for substantial, low-powered displacements of 4D-printed multilayer actuators.

## 1 | Introduction

The approach undertaken by Live-Mirror [1], a collaboration regrouping multiple research parties, is to establish the additive manufacturing process of an electroactive polymer (EAP)-based multilayer structure capable of supporting and locally curving a few-mm-thick, 5 m-wide mirror [2, 3]. EAPs have gain in interest during the past decades due to their wide range of applications, from artificial muscles and micro-pumps to sensors and actuators. Among them, polyvinylidene fluoride (PVDF)-based copolymers and terpolymers have shown promising properties. While P(VDF-TrFE) exhibits a ferroelectric behavior, the addition of a third monomer inside the matrix, such as chlorotrifluoroethylene (CTFE) [4, 5] or chlorofluoroethylene (CFE) [6], has demonstrated conformation changes into a relaxor-ferroelectric

with a minimized hysteresis. In combination with a high dipole moment of the VDF units (2.1 D [7]), the addition of CFE defects (2.1–1.8 D [8]) result in a greater electromechanical coupling ($k_{33}$) than with CTFE (0.64 D [9]). This indeed originates from the presence of two hydrogen atoms at the tail of the termonomer, contrary to CTFE with two fluorine atoms weakening the difference in local dipolar moment with its head. A 3D representation shown in Figure 1 allows the visualization of two termonomer units with ideal head-tail configurations, showcasing the major difference in overall dipolar moment.

It is worth noting that the proportion of head–tail (H–T) and tail–tail (T–T) configurations, among two terpolymers with alleged similar compositions, strongly varies depending on the polymerization process, leading to non-negligible results in terms of

This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly
cited.

© 2025 The Author(s). Macromolecular Materials and Engineering published by Wiley-VCH GmbH
<figure><img src="image_2.png" /><figcaption>FIGURE 1 | 3D representations of H-T-H-T-H-T P(VDF-TrFE-CTFE) and P(VDF-TrFE-CFE) with molecular electrostatic potential and overall dipolar moment under MolView. Values were extracted from IQmol software.</figcaption></figure>

thermal and dielectric properties [10]. In general, it is admitted that the suspension polymerization of 8% of CFE to a 60/40 copolymer delivers the best energy storage capabilities [11] and electrostrictive performances [12]. However, the tuning of the intrinsic dielectric properties largely depends on the polymerization process, and two similar compositions on the paper likely deliver varying outcomes. In 2014, Capsal et al. reported major improvements in the electromechanical response of terpolymers thanks to the incorporation of phthalate plasticizers [13]. They measured that the addition of 15 wt.% of bis(2-ethylhexyl) phthalate (DEHP) increased by up to 215 times the mechanical energy, and the resulting electrostriction coefficient by 28 times. Later, the addition of DINP through various concentrations also unveiled a 25% enhancement in strain for a same content as DEHP [14]. In combination with a high dielectric permittivity at low frequencies and low Young's modulus, this led to unprecedented strains related to Maxwell forces (Equation 1), expressed as:

$$S_{Maxwell} = -\epsilon_0 \epsilon_r (1 + 2\nu) E^2 / 2Y \quad (1)$$

with $\epsilon_0$ the vacuum permittivity, $\epsilon_r$ the relative permittivity of the dielectric material, E the applied electric field [13]. It should be noted that those stresses have a quadratic relationship with the electric field. As compared to a neat terpolymer, while it has been demonstrated a 20x enhancement of the mechanical energy density for CTFE-based terpolymers, CFE-based ones underwent a 100x increase [15], leading this study to pursue with P(VDF-TrFE-CFE). Nonetheless, an important amount of plasticizer is known to be detrimental for the EAP stability with an increase in electrical breakdown probability [16]. Its mitigation requires the use of high-molecular-weight terpolymers dissolved in pure solvents, followed by a filtration for the removal of high-density gels [17]. However, the influence of the terpolymer composition on the plasticizer benefits has only been investigated through few CTFE and CFE contents [15]. In summary, it was observed that a gain in softness accompanied by an enhanced molecular mobility of the amorphous phase (low-frequency dielectric permittivity) sky-rocketed the electrostrictive coefficient. However, while CTFE-based terpolymers are indeed cheaper and more commercially available, its non-negligible contrast with CFE in terms of molecular constitution and dipolar moment makes it unlikely to deliver similar outcomes. Therefore, the scope of this study is to analyze different CFE-based terpolymers, with various VDF, TrFE, and CFE contents, to validate or renew our vision on the plasticizer impact on the electromechanical performances. The selection of a candidate of high softness, dielectric permittivity, and electromechanical stability for low-frequency actuation will be key for the Live-Mirror project. Therefore, the analysis of

the thermal, mechanical, structural, and dielectric properties of plasticized P(VDF-TrFE-CFE) is undertaken here.

## 2 | Materials and Methods

### 2.1 | Materials

Poly(vinylidene-trifluoroethylene-chlorofluoroethylene), P(VDF-TrFE-CFE) powders were provided by Piezotech (Arkema group, France) with various VDF, TrFE, and CFE contents. Each terpolymer will be named after their approximate VDF and CFE contents as T(VDF:CFE), for instance T(79:13) refers to a terpolymer with 79% of VDF, 8% of TrFE, and 13% of CFE.

They were synthesized by suspension polymerization [18] and their precise compositions, determined by ¹⁹F RMN spectroscopy, are displayed in Table 1. Methyl-ethyl-ketone (MEK, Analab NORMAPUR Reag. Ph.Eur., ACS, 100.0%) was supplied by VWR Chemicals, and Cyclopentanone (>99%) by Carl-Roth. Di-n-octyl phthalate (DNOP, >98%, GC) plasticizer was supplied by Sigma-Aldrich, and the Durapore 5 µm PVDF membrane filters by Merck (Millipore).

## 3 | Methods

### 3.1 | Formulation

Each terpolymer was dissolved into a 60:40 MEK:Cyclopentanone mixture with a polymer-to-solvent mass fraction of 16 wt.%, followed by the addition of 10 wt.% (with respect to the terpolymer content) of DNOP plasticizer. This concentration has previously been determined as optimal to preserve a high-enough dielectric strength combined with a high electromechanical response. Once formulated, a filtering step through a 5 µm PVDF membrane allowed the removal of any residual gel from in situ polymerization, improving the breakdown strength of the resulting EAP solution [17]. A Doctor Blade (elcometer 3700) was used to cast 1 mm-thick films onto glass substrates which were let for 2 h at room temperature to allow the solvents evaporation (Figure 2). The films were introduced into an oven at 60°C for 2 h and annealed at 20°C below $T_{onset}$ for another 2 hours to ensure a homogeneous crystallization process for increased breakdown strength and dielectric permittivity, compliant with a median crystal fraction [19]. Determined by DSC, $T_{onset}$ is usually 10°C below $T_m$.
<table><caption>TABLE 1 | Terpolymers compositions.</caption><thead><tr><th>P(VDF-TrFE-CFE) terpolymers</th><th>VDF (mol.%)</th><th>TrFE (mol.%)</th><th>CFE (mol.%)</th><th>MFI</th></tr></thead><tbody><tr><td>T(62:7)</td><td>62.3</td><td>30.7</td><td>7.0</td><td>3.54</td></tr><tr><td>T(61:8)</td><td>60.7</td><td>31.4</td><td>7.9</td><td>0.05</td></tr><tr><td>T(58:8)</td><td>57.7</td><td>34.2</td><td>8.1</td><td>10</td></tr><tr><td>T(79:8)</td><td>79.3</td><td>12.5</td><td>8.2</td><td>4</td></tr><tr><td>T(54:9)</td><td>53.6</td><td>37.2</td><td>9.2</td><td>5.36</td></tr><tr><td>T(49:13)</td><td>49.5</td><td>37.5</td><td>12.9</td><td>1.8</td></tr><tr><td>T(79:13)</td><td>79.5</td><td>7.1</td><td>13.4</td><td>2.2</td></tr><tr><td>T(52:13)</td><td>51.7</td><td>34.8</td><td>13.5</td><td>2.8</td></tr></tbody></table>

<figure><img src="image_3.png" /><figcaption>FIGURE 2 | Illustration of the film-making process.</figcaption></figure>

## 3.2 | Thermal Characterization

DSC measurements were performed with a DSC 131 EVO calorimeter from *Seteram Instrumentation* under nitrogen flow. Polymeric powders were introduced into closed aluminum crucibles and subjected to a first cooling ramp down to -50°C, followed by a slow heating ramp at 10 K min⁻¹ up to 160°C. A second slow cooling at 10 K min⁻¹ down to room temperature allowed for a steady crystallization process. The routine was applied twice but the reveal of relaxor ferroelectric (RFE) to paraelectric (PE) transitions (Curie), determined with a ± 1°C uncertainty, was mostly observable during the first cycle due to its thermal history. Results were analyzed with *Calisto Processing* using integrated baselines.

## 3.3 | Mechanical Characterization

The obtention of the mechanical properties of each film required 50 µm-thick, 6 cm-long, and 1 cm-wide samples. One end was clamped to a 10 N load cell and the other one to motorized linear stages applying a fixed displacement of 1% at a given frequency. The monitoring and collecting of the data were done under *LabView* at room temperature with a ± 2 MPa uncertainty.

## 3.4 | Structural Characterization

The structural characterization of purified terpolymers was carried out with an ATR module, equipped with a crystal diamond, on a *ThermoFischer FTIR Nicolet iS50*. For each sample, the represented spectrum is the average of three spectra randomly

collected across both surfaces. Each one is obtained as a mean of 32 spectra collected between 525 and 4000 cm⁻¹ using KBr optics at room temperature.

## 3.5 | Dielectric Characterization

Broadband dielectric spectroscopy (BDS) was performed using a *Novoncontrol* impedance analyzer. The real and imaginary parts of the dielectric permittivity (respectively ε' and ε'', tan(δ) = ε''/ε') were collected over a frequency range of 1 MHz to 100 mHz with eight measurements per decade. Isothermal spectra, assumed to remain stable with a ± 0.2°C precision, were carried out every 3°C and over a temperature range of -60°C to 60°C to cover the glass and Curie transitions. Dielectric breakdown tests, illustrated in Figure 18a, were performed using 30 samples of 50-70 µm in thickness. Submerged in a silicon oil bath, a 1 kV s⁻¹ linear ramp was applied thanks to an Agilent 33220A generator, connected to a Trek 20/20C Amplifier and until electrical breakdown. Values of the maximum input voltage (peak-peak) were obtained using a Keysight EDUX1002G oscilloscope. The DC breakdown probability P(E) was fitted to a Weibull model with the following equation:

$$P(E) = 1 - \exp[-E/\lambda]k \quad (2)$$

where E is the breakdown electric field of each sample, λ is the breakdown electric field for P(E) = 63.2%, and k is the standard deviation of the experiment. The measured probabilities are calculated as the ratios between the cumulated number of samples to the total number of samples.
a)

<figure><img src="image_4.png" /></figure>

b)

<figure><img src="image_5.png" /><figcaption>**FIGURE 3** | (a) Thermograms of terpolymer powders, (b) evolution of Curie and melting temperatures with %TrFE, and (c) evolution of their enthalpies of fusion and crystallinity with %CFE.</figcaption></figure>

<table><caption>**TABLE 2** | Thermal properties of terpolymer powders.</caption><thead><tr><th>P(VDF-TrFE-CFE) terpolymers</th><th>Copolymer base</th><th>T<sub>m</sub> (°C)</th><th>T<sub>c</sub> (°C)</th><th>ΔH<sub>m</sub> (J.g<sup>-1</sup>)</th><th>χ<sub>c</sub> (%)</th></tr></thead><tbody><tr><td>T(62:7)</td><td>70/30</td><td>130.9</td><td>42.6</td><td>18.71</td><td>43.9</td></tr><tr><td>T(61:8)</td><td>70/30</td><td>128.1</td><td>41.1</td><td>16.74</td><td>39.3</td></tr><tr><td>T(58:8)</td><td>60/40</td><td>124.6</td><td>—</td><td>9.55</td><td>22.4</td></tr><tr><td>T(79:8)</td><td>85/15</td><td>121.5</td><td>81.4</td><td>15.15</td><td>35.6</td></tr><tr><td>T(54:9)</td><td>60/40</td><td>129.9</td><td>24.7</td><td>15.69</td><td>36.8</td></tr><tr><td>T(49:13)</td><td>60/40</td><td>113.7</td><td>-1.9</td><td>11.99</td><td>28.1</td></tr><tr><td>T(79:13)</td><td>85/15</td><td>105.2</td><td>77.8</td><td>13.37</td><td>31.4</td></tr><tr><td>T(52:13)</td><td>60/40</td><td>112.0</td><td>-3.4</td><td>10.94</td><td>25.7</td></tr></tbody></table>

# 4 | Results and Discussion

## 4.1 | DSC Results

Figure 3a displays the thermograms associated to each terpolymer powder. Terpolymers with low CFE/TrFE ratios, such as T(62:7), T(61:8), and T(54:9), all have melting temperatures ($T_m$) close to 130°C. For samples with CFE contents neighboring 13%, the larger proportion of TrFE confers T(52:13) higher melting points than T(79:13), which flirts with 100°C. Thermal data is summarized in Table 2. Regarding the Curie temperatures ($T_c$), they appeared when subjecting the powders to negative temperatures, facilitating the apparition of an associated endothermic peak ($\Delta H_{Curie} \sim 1$ J.g$^{-1}$, coherent with literature [20]). For instance, the first endothermic peak of T(62:7), neighboring 43°C, is relatively sharp compared to T(79:8)'s and is associated to a narrow Curie transition. The latter is broader due to an inhomogeneous chain distribution from the polymerization process which is underlined by the doubled melting peak. Therefore, it would be expected a broader and lower dielectric permittivity around its $T_c$.

Similarly to what has previously been observed for P(VDF-TrFE) [7, 21], an increase in TrFE content (for a similar %CFE) is accompanied by a lower $T_c$. This originates from the bulky TrFE units expanding the interchain distance for a reduced intermolecular friction [8]. In favor of the dipole rotation and reorientation, this results in a lower activation energy of the RFE/PE phase transition. In fact, while T(52:13) has a $T_c$ below 0°C, T(79:13) founds his almost hidden by its $T_m$ and $T_{cr}$ (similar to PVDF). Such behavior outlines major differences with CTFE-based terpolymers which, for a similar composition as T(79:13), has shown outstanding electrostrictive performances at room temperature [22]. Also, it is commonly stated for terpolymers with similar VDF/TrFE bases that the proportion of termonomer decreases the Curie temperature. In combination with large CFE units further expanding the lattice parameters, the introduction of defects (gauche bond) in the crystalline phase leads to smaller, nanodomains. For instance, the $T_c$ of a terpolymer with a VDF/TrFE base of 60/40 decreases from 24.7°C for T(54:9) to -3.4°C for T(52:13). Discrepancies are also noticeable for different molecular weight terpolymers, the high $M_w$ T(61:8) and the low $M_w$ T(58:8). While the first one displays two endothermic
peaks at 41.1°C and 70.7°C, with associated enthalpies of 0.31 and 0.26 J g⁻¹ respectively, the latter only displays one of 0.32 J g⁻¹ at 74.2°C. While the first peak can be associated to a Curie transition, the high-temperature peaks are most likely related to the transition of well-ordered ferroelectric phases of different sizes to paraelectric phases [23]. The evolution of T_c with %VDF, for several P(VDF-TrFE-CFE), is represented in Figure 3b.

For increasing VDF contents, the Curie temperatures close with their crystallization and melting points acting as barriers for pragmatic applications. For ground-based telescopes, usually placed at high altitude for a lower impact of atmospheric seeing, they operate under colder temperatures and would require a T_c closer to 0°C, such as a T(52:13). As for the degree of crystallinity (χ_c), it is calculated using the following expression:

$$ \chi_c(\%) = 100 \times \frac{\Delta H_f}{\Delta H_{f,100}} \quad (3) $$

with ΔH₁₀₀ the enthalpy of fusion of a fully crystalline terpolymer, which value of 42.5 J g⁻¹ has been determined by extrapolation of data acquired by Klein et al. [24]. The melting and Curie transition enthalpies are coherent with the literature for low [12] and high [25] termonomer contents. The decreasing crystallinity with the CFE content outlines the disrupting behavior of the termonomer in the crystal, also widely observed in the literature.

## 4.2 | Tensile Test Results

For a similar dielectric permittivity, while tough terpolymers will deliver the greatest blocking force for strong actuation, a limited strain (proportional to ε·Y⁻¹) would prevent the obtention of large displacements. In order to concede tradeoffs between decoupled electrostriction and blocking force, the investigation of the plasticizer influence on the mechanical properties in necessary. In fact, few studies have reported the mechanical properties of P(VDF-TrFE-CFE). Contrary to what is often stated with P(VDF-TrFE-CTFE) [10], the decrease in Young's modulus is not necessarily linked to an increase in termonomer content, even though it is indeed responsible for the decohesion of the crystalline phase. Indeed, while T(79:13) showcases a high termonomer content of 13.4%, it appears as one of the toughest materials tested here with a measured Y₁₀₀mHz of 279 MPa. A clearer distinction emerges when plotting Y as a function of TrFE (Figure 4a) or the average van der Waals (vdW) radius for a chain of 100 monomer units (Figure 4b). An explanation is the difference in termonomer microscopic composition associated with their respective van der Waals radii. In addition to one chlorine atom (r_vdW,Cl = 1.75 [26]), while CTFE possesses three fluorine ones (r_vdW,F = 1.4) capable of bringing more steric hindrance, CFE only possesses one. Hence, the addition of CFE units, while drastically increasing the dipolar moment of the polymer chain, does not (or less) participate in the overall reduction in Young's modulus. This is further confirmed when comparing two terpolymers of similar crystallinity, T(79:8) and T(54:9), with the latter being 3x softer. Henceforth, the major reductive toughness agent is TrFE, explaining why the softer terpolymers all dispose of at least 35% of TrFE.

The second representation has the benefit of showcasing the saturation in Young's modulus with the average vdW radius. While

the role of the DNOP indeed softens most terpolymers, T(79:13) witnesses a slightly higher module after its introduction. The plasticizing effect is known to provide extra chain mobility to the polymer backbone due to weaker intermolecular interactions. In semicrystalline polymers, such as PLA [27] and more relevantly P(VDF-TrFE-CTFE) [14], it ultimately leads to a facilitated activation energy for crystallization. Although all tested terpolymers indeed gain in crystallinity, T(79:13) sees its drastic increase from 22.0% to 34.1% (Figure S1) resulting in the associated toughening. It is important to note that contrary to T(62:7), which softness saturates at 100 MPa no matter how much DNOP is added, T(52:13) achieves 17.6 MPa for a 20 wt.% concentration (Figure S2). It is nonetheless likely to reach electrical breakdown at lower electric fields (as discussed later) but could still find potential applications for ultra-soft EAP-based devices.

## 4.3 | Infrared Spectroscopy (FTIR) Results

### 4.3.1 | P(VDF-TrFE-CFE) Conformation and Plasticizer Quantification

Figure 5 shows the spectrograms of pure and plasticized (10 wt.% of DNOP) terpolymers, alongside liquid DNOP (dotted line), between 525 and 1800 cm⁻¹. Although the coexistence of α, β and γ phases usually causes the attribution of each peak to be disputed [28], the composition discrepancies between selected terpolymers can provide a qualitative comparison of each phase prominence. One the one hand, the distinct peaks or bends located at 614, 766, 796, and 975 cm⁻¹ are attributed to TG conformations from the α-phase [29, 30]. On the other hand, as the strong peak at 840 cm⁻¹ is argued to correspond both to β and γ phases, only the bend at 1284 cm⁻¹ will be considered to correspond to the all-trans conformation of the β-phase. Finally, the T₃G conformation from the γ-phase can be distinguished at 772, 811, and 1234 cm⁻¹.

Several studies manage to determine the proportion of electroactive phase through the absorbance of defined peaks at 766 and 840 cm⁻¹ [31, 32]. However, for all terpolymers except the high VDF content T(79:13) and T(79:8), the 766 cm⁻¹ bend is hidden by the T₃G conformation at 772 cm⁻¹. A proper deconvolution considering all reported, distinguishable peaks and bends could be undertaken but appears too unprecise to give credible results in this study. Those two particular terpolymers indeed exhibit similar α-phase signatures as some P(VDF-CTFE) reported in the literature [33, 34], with distinct peaks at 766 and 796 cm⁻¹. Also attributed to the β-phase, their low TrFE contents are identified by the lack of bend at 1091 cm⁻¹, contrary to T(54:9) and T(52:13). As for the CFE, T(62:7) terpolymer stands with the highest absorbance for all-trans configurations due to its low termonomer content (7%), as opposed to T(49:13) and T(52:13) terpolymers. Its inclusion is also known to favor the formation of TG conformation and is confirmed by the increasing absorbance at 614 and 975 cm⁻¹. As DNOP is introduced into the films, T(79:13) witnesses a 12.1% gain in crystallinity with an overall increase in absorbance over the whole spectra. More precisely, a noticeable crystallization in α-phase with sharper peaks at 614, 766, and 1209 cm⁻¹ takes over the T₃G conformation of the γ-phase with lower absorbencies at 811 and 833–840 cm⁻¹. The analysis of its influence on the conformation remains limited as its own absorbance (displayed as yellow dotted line) overlaps
a)

<figure><img src="image_6.png" /><figcaption>FIGURE 4 | Evolution of the Young's modulus with (a) %TrFE or (b) van der Waals radii for pure and plasticized terpolymers at 100 mHz.</figcaption></figure>

b)

<figure><img src="image_7.png" /></figure>

<figure><img src="image_8.png" /><figcaption>FIGURE 5 | FTIR spectra of pure and plasticized terpolymers.</figcaption></figure>

many conformation-related peaks. Nonetheless, this is a mean of verifying the overall miscibility with the amorphous phase. By collecting the spectra of pure terpolymers and the spectrum of DNOP, reconstructions following a Beer Lambert-like law allows for the determination of their plasticizer retention.

Their quantification is calculated using the sharp peak at 1723 cm⁻¹, unaffected by the terpolymer, representing the stretching of the C=O ester group. After annealing and before any characterization, the surface wiping with ethanol partially removes unconfined DNOP, testifying a reduced plasticizer content compared to its initial weight fraction. A close-up view of the C=O group, displayed in Figure 6a, allows to visualize discrepancies in DNOP retention among terpolymers.

Their quantification, by mass as measured in a solid, can be obtained through reconstructed spectra as:

$$ \text{abs(reconst.)} = (1 - \chi_{\text{DNOP}}) \times \text{abs(terpopur)} + \chi_{\text{DNOP}} \times \text{abs(DNOP)} \quad (4) $$

with $\chi_{\text{DNOP}}$ as their true weight fractions, defined as:

$$ \chi_{\text{DNOP}} = \frac{\text{abs(terpoplast)}_{1723 \text{ cm}^{-1}}}{\text{abs(DNOP)}_{1723 \text{ cm}^{-1}}} \quad (5) $$

It should be mentioned that all tested samples display no significant discrepancies across their surfaces, and between the upper and bottom surfaces, as highlighted in Figure S3 with low standard deviations (Table S1). The homogeneity of the films will be considered sufficient to approximate the data collected by ATR, usually penetrating through 0.5–5 µm, as being representative of the whole sample through its thickness. Figure 7 displays the measured and reconstructed spectrum of pure and plasticized T(54:9). When manually selecting a plasticizer fraction of $\chi_{\text{DNOP}}$ = 0.095, the reconstructed spectrum showcases an overall very accurate fit with the true, measured spectrum. Small discrepancies are however noticeable and will be discussed later. Table 3 displays the coefficients used to recreate the plasticized spectra from the pure terpolymers and DNOP separately alongside their calculated loss (0.1–$\chi_{\text{DNOP}}$), and in Figure 6b, the remaining %DNOP with CFE. The plasticizer retention was normalized to that of T(79:13). As visually observed after the casting and annealing of T(62:7), it stands as the greatest plasticized terpolymer with a remaining content of 5.8% only. For terpolymers with higher %CFE, the true plasticizer content is closer to 10 wt.%, with almost no loss for T(79:13) and T(52:13), testifying to a great retention. Near 8% of CFE and similar VDF/TrFE ratios, the high molecular weight T(61:8) terpolymer endures a stronger loss in plasticizer compared to the low Mw T(58:8), suggesting a non-negligible impact of the macromolecule dimensions with regards to their interactions with DNOP.

### 4.3.2 | Hypothetical Electrostatic Interactions: Van de Waals and Halogen Bonds

A first theory could be that pending CFE groups are responsible for an interaction with plasticizers. Halogen bonds, characterized as an electrostatic interaction between an R-X group (X =
<figure><img src="image_9.png" /><figcaption>b)</figcaption></figure>

<figure><img src="image_10.png" /></figure>

**FIGURE 6** | (a) FTIR spectrum of plasticized terpolymers after annealing around 1723 cm⁻¹ and (b) retention of DNOP with CFE.

<figure><img src="image_11.png" /></figure>

<figure><img src="image_12.png" /></figure>

<figure><img src="image_13.png" /></figure>

<figure><img src="image_14.png" /></figure>

**FIGURE 7** | Spectrum reconstitution close-up of T(54:9) around (a) 1723 cm⁻¹ (C=O stretching), 1091 cm⁻¹ (β-phase), 614 cm⁻¹ (α-phase) and 1271 cm⁻¹ (C—O stretching).
<table><caption>TABLE 3 | Table of plasticizer loss.</caption><thead><tr><th>Terpolymer</th><th>100 × χ<sub>DNOP</sub> (%)</th><th>% loss</th></tr></thead><tbody><tr><td>T(62:7)</td><td>58%</td><td>42%</td></tr><tr><td>T(61:8)</td><td>54.5%</td><td>45.5%</td></tr><tr><td>T(58:8)</td><td>75%</td><td>25%</td></tr><tr><td>T(79:8)</td><td>83.2%</td><td>16.8%</td></tr><tr><td>T(54:9)</td><td>87.5%</td><td>12.5%</td></tr><tr><td>T(49:13)</td><td>90.7%</td><td>9.3%</td></tr><tr><td>T(79:13)</td><td>100%</td><td>0%</td></tr><tr><td>T(52:13)</td><td>94.4%</td><td>5.6%</td></tr></tbody></table>

halogen) and a Lewis base (entity with a lone pair of electrons), originate from the presence of a σ-hole, a positive electrostatic potential region at the outer shell of the halogen surface [35]. In this case, C—O...Cl—C, C=O...Cl—C or π...Cl—C bonds (and possibly a combination of some), represented in Figure 8, could be responsible for such affinity between phthalates and chloro(tri)fluoroethylene.

Although fluorine atoms are rarely considered to participate in such phenomenon [36], they strengthen the electrostatic potential ($V_{s,max}$) of Cl. In fact, a $H_3C—Cl$ unit does not exhibit such feature ($V_{s,max}$ of -1.5) while a $F_3C—Cl$ unit shows one of 16.3 [37]. This has been further confirmed by Liu et al. [38] which confirmed the preeminence of hydrogen bonds, C—H...O=C, rather than halogen bonds, with the formation of three-center hydrogen-bonded complexes for plasticized PVC. Using FTIR, they identified the great shift in the C=O stretching frequency towards lower wavenumbers, with increasing plasticizer content, as being unrelated to an unstable C=O...Cl—C halogen bond. One must however bear in mind that in a terpolymer (whether CTFE or CFE), a σ-hole could be related to a CFC—Cl unit, which electrostatic potential can be compared to that of a $H_2FC—Cl$ with a value of 3.5. Unlike PVC, no hydrogen atom shares the same carbon as a chlorine atom, expect if CFE acts as an end-group. In TrFE however, fluorine can weaken the C—H bond to potentially generate a hydrogen bond with another polar entity of DNOP (Figure 9). Another way of representing the interactions generated by the incorporation of plasticizer is by accordingly subtracting the spectrum of pure DNOP, relative to its concentration, to the spectra of the plasticized terpolymers. The

alternative to Equation (4) is expressed below with Equation (6).

$$ \text{abs (reconst.)} = (1 + \chi \text{DNOP}) \times \text{abs (terpoplast.)} \\ - \chi \text{DNOP} \times \text{abs (DNOP)} \qquad (6) $$

The presence of halogen or hydrogen bonds (HB) under FTIR can be identified as red (normal or “proper” HB) or blue (weak or “improper” [39] HB) shifts. Usually, an electron-rich Lewis base (Y) attracts the electron from a Lewis acid hydrogen atom (H) and weakens, thus forces the elongation of the X—H bond from which arises a red-shift. For a poor X—H bond however, its contraction has been observed for a variety of Xs and Ys, including Cl and F-containing complexes [40], leading to the observation of blue-shifts. This phenomenon, in plasticized terpolymers, can be illustrated in Figure 9.

Figure 10 shows both the measured and reconstructed spectra (using Equation 6) of T(52:13) with 5, 10, and 20 wt.% of DNOP. With its C=O signature located at 1723.6 cm⁻¹ for pure DNOP, a shift comprised between 0.9 and 2.4 cm⁻¹ is observed when increasing the plasticizer content from 5 to 20 wt.%. Such phenomenon is however not noticed for T(79:13) and will be considered insignificant to receive further investigation. Three other discrepancies, thus potential interactions, are however reported.

The first one involves the out-of-plane bending of the δ(= C—H) from the benzene [41], with a significant shift in wavenumber from 741 cm⁻¹ for neat DNOP to 746.8 cm⁻¹ for a 20 wt.% concentration. Such change in vibration frequency could find its origin in a CFC—Cl...π interaction which has been recently underlined by Ibrahim et al. [42] for a $X_3C—Cl...π$-system. Near 1171 cm⁻¹, corresponding to the stretching of symmetric $CF_2$ from PVDF, the subtracted spectra systematically display greater absorbances, accompanied by a slight blue shift by up to 4 cm⁻¹. A possible explanation comes from the increase in free volume due to the plasticizer long carbon chain, following both ester groups, enhancing the distance between fluor groups and liberating more numerous vibrations. A third discrepancy is located near the all-trans-related peak surrounding 1284 cm⁻¹. The plasticizer's spectrum itself possesses a broad peak at 1270 cm⁻¹, corresponding to the C—O stretching of an aromatic ester. Figure 11a represents its deconvolution, displaying mainly three different signals, that can correspond to its three rotamers (i-o, i-i and o-o [38] in Figure 11b).

<figure><img src="image_15.png" /><figcaption>FIGURE 8 | Possible halogen bonds formed between the terpolymer and plasticizer.</figcaption></figure>
Prominently, two broad peaks have their maxima located at 1259 and 1283 cm⁻¹ and one smaller peak has its at 1269 cm⁻¹. With an increasing plasticizer content, one can notice a stronger response at 1293 cm⁻¹ as opposed to a weaker band gradually shifting from 1275 to 1279 cm⁻¹. Such equilibrium could testify a constrained movement from a particular rotamer leading to a switch in conformation, possibly due to its implication in an electrostatic interaction with the terpolymer matrix. It is further supported by the greater deviation (ΔA₁₂₇₁cm⁻¹) observed between the measured and reconstructed spectra (from Equation 4) with an increasing %CFE. Represented in Figure 13a, it indeed suggests a stronger impediment of the chlorine entities to the C—O stretching. The interpretation of multicomponent infrared spectra being utterly complex here makes it tricky to fully confirm and further investigations, through NMR spectroscopy [43] (also undeniably valuable for terpolymer structural characterization) and DFT simulations under Gaussian09 software [38, 42], should be conducted to support this theory.

### 4.3.3 | Quantitative Approach for 10 wt.% of Plasticizer

A quantitative approach of the retention by atom enumeration is assessed. Considering that M<sub>VDF</sub> = 64 g mol⁻¹, M<sub>TrFE</sub> = 82 g mol⁻¹ and M<sub>CFE</sub> = 80.5 g mol⁻¹, an averaged monomer unit of a terpolymer which composition is %VDF / %TrFE / %CFE accordingly weighs:

$$ \text{Mavg} = 0.01 \times (M_{\text{VDF}} \times \%\text{VDF} + M_{\text{TrFE}} \times \%\text{TrFE} + M_{\text{CFE}} \times \%\text{CFE}) \text{ [g.mol} - 1] \quad (7) $$

The terpolymer formulations tested here contained approximately 11.4 g of terpolymer, which gives n<sub>avg</sub> = 11.42/M<sub>avg</sub> mol per formulation. The amount of CFE (n<sub>CFE</sub>) can therefore be expressed as n<sub>avg</sub> × %CFE mol per formulation. The aforementioned crystallinity being calculated based on enthalpies, it can be used to estimate the number of CFE units inside (N<sub>CFE,c</sub>) and outside (N<sub>CFE,a</sub>) the crystallites with and N<sub>CFE,c</sub> = N<sub>a</sub> × χ<sub>c</sub> × n<sub>CFE</sub> and N<sub>CFE,a</sub> = N<sub>a</sub> × (1-χ<sub>c</sub>) × n<sub>CFE</sub>. In the meantime, each formulation received 1.27 g of DNOP which, for M<sub>DNOP</sub> = 350.96 g mol⁻¹, is equivalent to n<sub>DNOP</sub> = 3.25 × 10⁻³ mol or N<sub>DNOP</sub> = 1.96 × 10²¹ molecules.

Table 4 summarizes the molar ratio of chlorine atoms, in the amorphous (r<sub>a</sub> = N<sub>CFE,a</sub> / N<sub>DNOP</sub>) and crystalline (r<sub>c</sub> = N<sub>CFE,c</sub> / N<sub>DNOP</sub>) phases, to available oxygen atoms from DNOP based on

<figure><img src="image_16.png" /><figcaption>FIGURE 9 | Illustration of possible red and blue shifts in plasticized terpolymers.</figcaption></figure>

Equation (7). Their evolution with %CFE is displayed in Figure 12. At 8% of CFE, the formulation contains an equal amount of chlorine and oxygen atoms, considered here as a probability requisite for at least one interaction to take place. While it has been established the absence of DNOP inside the crystal, pending Cl are seen as potential anchor sites. The crystal morphology (rice-like ellipsoid) elongating when increasing the termonomer content optimizes the specific area, providing more pending linkages and rigid amorphous phase for high CFE contents. Indeed, while a shape factor (length/width) of 4 was found for a 10.3% CTFE content, it dropped to 1.7 for the copolymer [44], losing about 22% of possible linkage groups if a similar volume is considered. This contribution, although more pronounced for high CFE contents, could be considered as relatively small and r<sub>c</sub> would mostly represent the fraction of Cl atoms that do not participate in an interaction with DNOP.

Nonetheless, a strong discrepancy was observed between low and high-molecular-weight (M<sub>w</sub>) terpolymers. The latter is indeed supposed to dispose of less chain ends, acting as pending linkages, capable of expanding the number of interactions with the plasticizer. In fact, it was measured a 54.5% DNOP retention for the high M<sub>w</sub> T(61:8), while its counterpart, a T(58:8) with an MFI of 10 (low M<sub>w</sub>), showcases a significantly greater retention of 75%. Note that the latter remain drastically below the retention expectation, raising the limitations of such reasoning for low M<sub>w</sub> terpolymers. Until further assessment with GPC characterization [45] to precisely determine the influence of M<sub>w</sub> on said parameters, T(58:8) is removed from Figure 12. It should also be born in mind that low M<sub>w</sub>, despite being beneficial for the plasticizer retention, are determinantal to the terpolymer electrical breakdown strength [17]. Therefore, a fraction of the red line, characterizing the number of Cl atoms in the crystalline phase, can still get increasingly involved in halogen bonds with CFE by the medium of pending linkages. The available Cl entities capable of participating in electrostatic interactions in the amorphous and constrained amorphous phases is represented with the blue line. One can observe that before 8% of CFE, the retention is only partial due to an overall lack of Cl atoms (r<sub>tot</sub> < 1). Combined with the low specific area of crystallites, DNOP is overpopulated and gets freed from the amorphous phase. Between 8% and 11% of CFE, there are enough chlorine atoms to possibly interact with every oxygen atom from DNOP, but the crystallinity is still too important to imply them all in electrostatic interactions. At 11% of CFE, the crystallinity is sufficiently low so that there is an equal amount of chlorine atoms in the amorphous phase to interact with every oxygen atom from DNOP, synonym of an expected full retention. The crystal morphology increasingly involved in the formation of O...Cl bonds, a 13% CFE content should sustain even more DNOP without triggering seepage. Such saturation has been tested on T(62:7), T(54:9), and T(79:13) terpolymers. Synthesized with concentrations of 5, 10, and 20 wt.%, the remaining plasticizer fraction, χ<sub>DNOP</sub>, is calculated with the same procedure as Equation (4) and is represented in Figure 13b.

In agreement with the grippy surface of T(79:13) conserved before and after casting/annealing for all plasticizer concentrations, FTIR reveals a linear increase in absorbance with %DNOP, testifying a great retention due to its high termonomer content. By assuming a 100% retention, other values were normalized to
<figure><img src="image_17.png" /></figure>

<figure><img src="image_18.png" /></figure>

<figure><img src="image_19.png" /><figcaption>**FIGURE 10** | (a) Normalized spectra of plasticized T(52:13) near the C=O bond, (b) un-normalized subtracted spectra near the C—O bond, and (c) un-normalized subtracted spectra near the aromatic, out-of-plane, 1,2 disubstituted δ( = C—H) bending.</figcaption></figure>

that of T(79:13). For T(62:7), an early saturation appears after
the addition of 5 wt.% of DNOP. In fact, adding 10–20 wt.%
resulted in a remaining plasticizer fraction of respectively 5.24
and 6.38 wt.%, illustrating the rejection of a surplus that cannot
be contained by the amorphous phase and interfaces. This
observation supports previous studies reporting a limited gain
in softness when incorporating larger quantities of DEHP [46],
where a significant plasticizer migration towards the surfaces
(seepage) and ion mobility led to high dielectric losses at low
frequency.

**4.3.4** | Impact of Plasticizer Retention on the Microstructure

Coming back to its potential application, the integrity of the
4D-printed, annealed films on low surface roughness glass

is essential. Previous printing attempts of various plasticized
terpolymers have revealed a non-negligeable probability of
crack formation after annealing, compromising the stability of
actuators even before its use. In fact, when overpassing the
melting temperature of terpolymers (previous results implied an
annealing below Tm), plasticizer unequally migrate and create
high-stress regions, initiating crack formation and propagation.
The printing of DNOP-doped films already led to more stable
recrystallization than DINP, but regarding previous results, it is
unlikely for different terpolymer compositions and affinities to
produce similar outcomes. EAPs doped with 10 and 20 wt.% have
been subjected to thermal annealing at 160°C for 2 h, on glass,
followed by a slow cooling rate (1 h), enabling a full reorganization
of the crystallites and plasticizer homogenization. At 10 wt.%,
once cooled down to room temperature, cracks in the mm-wide
range are observable, in Figure 14, for T(62:7) and T(54:9) but not
for T(52:13).
a)

<figure><img src="image_20.png" /></figure>

b)

<figure><img src="image_21.png" /><figcaption>FIGURE 11 | (a) Deconvolution of the C—O stretching band via Origin and (b) DNOP rotamers.</figcaption></figure>

<table><caption>TABLE 4 | Retention parameters based on Equation (7). M<sub>avg</sub> corresponds to the average molecular mass of a terpolymer unit, r<sub>tot</sub> to the overall ratio of Cl to O atoms in a 10 wt.% DNOP formulation, and r<sub>c</sub> and r<sub>a</sub> to this ratio in the crystalline and amorphous phases.</caption><thead><tr><th>Terpolymer</th><th>M<sub>avg</sub> (g mol<sup>-1</sup>)</th><th>r<sub>tot</sub> (%)</th><th>r<sub>c</sub> (%)</th><th>r<sub>a</sub> (%)</th></tr></thead><tbody><tr><td>T(62:7)</td><td>70.68</td><td>87</td><td>37</td><td>50</td></tr><tr><td>T(61:8)</td><td>70.96</td><td>98</td><td>38</td><td>59</td></tr><tr><td>T(58:8)</td><td>71.49</td><td>275</td><td>61</td><td>214</td></tr><tr><td>T(79:8)</td><td>67.60</td><td>106</td><td>38</td><td>69</td></tr><tr><td>T(54:9)</td><td>72.21</td><td>112</td><td>41</td><td>71</td></tr><tr><td>T(49:13)</td><td>72.81</td><td>156</td><td>44</td><td>112</td></tr><tr><td>T(79:13)</td><td>67.49</td><td>174</td><td>55</td><td>120</td></tr><tr><td>T(52:13)</td><td>72.49</td><td>164</td><td>42</td><td>122</td></tr></tbody></table>

Such phenomenon can originate from various causes. The surface roughness of the glass can act as defects and serve as crack-ignition sites. The films being however casted on similar glasses and the observed cracks being homogeneously spread over the surface, it is considered as not responsible. The substrate's Young's modulus is also known to improve the critical cracking thickness [47], but here again, all films were casted on glass. Also, the thicker the film, the more likely it is to generate cracks during the drying process. Here, all films were deposited with the same method using a BEVS 1806A/50 adjustable applicator set for a 1 mm-thickness before evaporation. The terpolymer weight fraction equaling 16 wt.% for all formulations, the post-drying thickness also appears similar and does not influence one film more than another. Lastly, thermal treatments causing residual internal stresses at crystalline-amorphous interfaces of polymers also stand as potential sources of cracks. Here, the annealing temperature of 160°C, largely over their respective melting points,

<figure><img src="image_22.png" /><figcaption>FIGURE 12 | Plasticizer retention related to the chlorine repartition in 10 wt.% plasticized terpolymers.</figcaption></figure>
<figure><img src="image_23.png" /><figcaption>FIGURE 13 | (a) Evolution of $ΔA_{1271cm}^{-1}$ (C–O) with %CFE and (b) retention capability of T(62:7), T(54:9), and T(79:13) plasticized (5, 10, and 20 wt.%) terpolymers with their respective $r_a$ at 10 wt.%.</figcaption></figure>

<figure><img src="image_24.png" /></figure>

<figure><img src="image_25.png" /><figcaption>FIGURE 14 | Annealed-on-glass (160°C) 10 wt.% T(62:7), T(54:9), and T(52:13).</figcaption></figure>

erases all thermal history and place them under equal conditions. Therefore, as both high-TrFE content terpolymers (of similar softness), T(52:13) and T(54:9), behave differently, the crack formation can be assimilated to the lower %CFE of the latter. The inhomogeneous recrystallization of T(62:7) with 20 wt.% led to the formation of cm-wide cracks and of unconfined DNOP aggregates visible in Figure 15. Those are trapped by a thin layer of terpolymer, highlighted with greater absorbance measured by FTIR. Its migration towards the surface in contact with the substrate being forbidden (contrary to Figure 13b) and its retention only partial, islets homogeneously spread over the film are created and demonstrate a phase segregation. In average, they appear circular (minimizing the internal energy) at the center of the film with a diameter of 1.73 mm and slightly elongated with a 2:1 ratio at its periphery. FTIR also reveals a remaining concentration of 3.14% for the uniform surfaces. As for the crack formation, the capability of each terpolymer to contain or release such stress (i.e., toughness) would be assimilated

to the energy absorbed per unit volume and quantified by tensile testing. However, even under non-standardized samples, their impressive softness and elongation led to the impossibility of straining until fracture with the available characterization methods.

### 4.3.5 | Broadband Dielectric Spectroscopy Measurements of T(52:13) Terpolymer

The evolution of the dielectric permittivity $ε_r$ and tan(δ) with temperature and frequency of pure and 20 wt.% plasticized T(52:13) is shown in Figure 16. The pure terpolymer shows at least three distinguishable relaxation modes, α, α' and Curie, respectively associated to the manifestation of the glass transition [48], the constrained amorphous phase [49], and the broad relaxor ferroelectric to paraelectric transition [20]. The latter being isochrone, it merges with the α mode
<figure><img src="image_26.png" /><figcaption>FIGURE 15 | Annealed-on-glass (160°C) 20 wt.% T(62:7) and T(52:13) with their FTIR spectra near 1723 cm⁻¹.</figcaption></figure>

<figure><img src="image_27.png" /><figcaption>FIGURE 16 | Dielectric permittivity εr and tan(δ) of respectively a,b) pure and c,d) 20 wt.% plasticized T(52:13) terpolymer.</figcaption></figure>

near 1 kHz and its temperature, comprised between -6°C and -3°C, is consistent with the endothermic peak measured by DSC. The addition of DNOP triggers no significant change in the α mode as being related to the segmental motion of the mobile amorphous phase (Figure 17). The apparition of a sub-glass β mode is related to the freezing temperature of DNOP approaching -25°C. The presence of the α' mode is uncovered by resorting to the Kramers-Kronig transformation partially cancelling the influence of DC conductivity, according

to:

$$ \epsilon'' = -\pi/2 \times [d\epsilon'/d \ln(2\pi f)] \quad (8) $$

As depicted in Figure S4, the addition of DNOP induces a gradual high-frequency shift of α' unveiling the freeing of the rigid amorphous phase. According to the Williams-Landel-Ferry (WLF) equations [50], the ratio of all electrical and mechanical
<figure><img src="image_28.png" /><figcaption>FIGURE 17 | Fits of the α and α' relaxation modes for pure, 10 and 20 wt.% plasticized T(52:13).</figcaption></figure>

relaxation phenomena can be expressed as:

$$ \ln(\tau/\tau_0) = -C1 \times (T - T_0) / (C2 + T - T_0), \quad T < T_g + 100 \text{ K} \quad (9) $$

with τ the characteristic relaxation time at T, τ₀ the relaxation time at a reference temperature T₀, and C₁ and C₂ are constants. It is noteworthy that this equation is only considered valid in a short range of 100 K after the glass transition. The evolution of ln(τ) with temperature for pure, 5, 10, and 20 wt.% plasticized T(52:13) is represented in Figure 17. The data relative to the α mode is represented with open triangles, while the data relative to the α' mode is represented with open circles. Fits were estimated using WinFit software with Havriliak–Negami equations:

$$ \varepsilon' = \varepsilon_{\infty} + \Delta\varepsilon/(1 + (i\omega\tau)^{\alpha})^{\beta} \quad (10) $$

where ε∞ is the dielectric permittivity at the high frequency limit, Δε is the difference between the static and high-frequency permittivity, τ is the characteristic relaxation time, and α and β are coefficients assimilated to the asymmetry and broadness of the relaxation. As a non-Arrhenius behavior, a drastic shift in ln(τ) implies a faster relaxation as implicitly connected to the amount of RAF released (or being impeded to form from the melt) by the additive. As for the α mode, it can be considered unaffected by the addition of plasticizer as WLF fits almost perfectly overlap each other.

The apparition of another relaxation mode at higher temperatures characterizes a combination of ionic conductivity of the free charge carriers and a polarization of the electrodes. While the distinction between both phenomenon remains uncertain, such interfacial phenomenon have already been observed and are mostly associated to Maxwell–Wagner–Sillars polarization. The latter is characterized by a maximum in tan(δ) discernibly shifting from 133 mHz at 3°C to 5.63 Hz at 60°C. The greater free volume conferred further from the glass transition indeed provides a facilitated molecular mobility of the plasticizer, thereby requiring less time (τ) to complete full polarization. Such shift is also

fostered by the plasticizer concentration, as a 10 wt.% inclusion only provides a maximum in tan(δ), at 60°C, at 177 mHz (Figure S4). A deeper investigation in quasi-static conditions should be conducted to precisely determine the role of each polarization in the dielectric behavior [51] of those formulations. In the meantime, a drastic increase in dielectric permittivity strikes with a value of 8.3 × 10³ at 100 mHz and 60°C. It is noteworthy that when an electric field is applied, the saturation of dipoles gets increasingly involved in the reduction of εr, narrowing the effective electrical input to low values. Moreover, the apparition of high dielectric losses at low frequency due to ionic conductivity is detrimental for the integrity of the terpolymer as responsible for lower DC breakdown strength.

### 4.3.6 | Dielectric Breakdown Strength Results

It is stated that the electrical breakdown of plasticized terpolymers, related to an exceeded current flow, decreases with the addition of plasticizers. In fact, for a terpolymer with 7.5% of CFE, it was observed a drastic shift in breakdown strength (for a probability of 63.2%) from 100.4 to 34.1 V µm⁻¹ respectively for pure and 12 wt.% DINP-doped terpolymers [16]. For pragmatic applications, this narrows the safe operating electric field. Associated to the presence of defects, another possible explanation can be the uncontrolled seepage leading to extremely high dielectric losses from ionic conductivity.

As seen in Figure 18b, the use of high-purity solvents and a filtration process already delivered enhanced breakdown strength for the neat terpolymer with a λ coefficient of 214 V µm⁻¹. The successful formulation of a 20 wt.% doped terpolymer, such as T(52:13), implies a significant presence of defects and a drop in breakdown strength with λ = 107 V µm⁻¹. The experimental procedure implied a narrower breakdown field distribution (k of 10.97 compared to 5.56 for the pristine) but the strong associated Weibull fit (R² ~ 0.98) extends the safe operating electric field well above previous works [16] and saturation fields.

The complete process of plasticizer retention and migration through the amorphous is yet to be fully understood. Evidence shows some chemical affinity related to the formation of electrostatic interactions as potential halogen bonds in the form Cl...O and Cl...π complexes. Those naturally originate from the presence of termonomer but are also favorited by the numerous pending linkages of low molecular weight terpolymers. Originally stated as being responsible for the freeing of a constrained amorphous, a combination of both phenomena is likely to confer plasticized terpolymer their boost in dielectric properties at low frequency. Those options will be investigated in further research. For now, it is undoubtedly stated that any quantitative comparison of two plasticized terpolymers properties must beforehand verify their true plasticizer concentrations via FTIR. For instance, the dielectric characterization of plasticized T(62:7) and any other terpolymer can only be conducted DNOP contents below its saturation, here reported to approach 6%.

## 5 | Conclusion

The analysis of different P(VDF-TrFE-CFE) compositions revealed significant (sometimes neglected) differences with
a)

<figure><img src="image_29.png" /><figcaption>**FIGURE 18** | (a) Electrical breakdown setup with (b) associated Weibull probability of pure and plasticized (20 wt.%) T(52:13) terpolymer. In green is the operating electric field range.</figcaption></figure>

b)

<figure><img src="image_30.png" /></figure>

CTFE-based terpolymers. While the tunability of each monomer fraction remains challenging depending on the polymerization process, terpolymers synthesized by suspension polymerization appear to deliver optimal (for ambient conditions) Curie-like temperatures for 52%-56% of VDF. In addition to closing the gap between its $T_c$ and $T_m$, a strong %VDF at the expense of TrFE leads to high Young's modulus, which further toughens with the addition of DNOP. A strong TrFE content favors free volume and chain flexibility, multiplying the electrostrictive coefficient. The obtention of amorphous-crystalline interfaces originating from the termonomer, it has been observed under FTIR an increasing retention of DNOP with %CFE possibly due to halogen bonds, justifying the struggle of some terpolymers to trap the plasticizer during the annealing process and overtime. The high plasticization of T(52:13) is confirmed through dielectric spectroscopy with a two-decade high-frequency shift of the $\alpha'$ mode while maintaining a breakdown strength compliant with low electric field operation. For the first time, a comprehensive study on the plasticizer's interaction with fluorinated terpolymers is proposed and led to the successful making of a 17.6 MPa ultra-soft film with promising applications. Further understanding of the plasticizer migration and accumulation at the interfaces combined with SAXS, NMR, and DFT calculations are suggested as valuable insights for future research.

## Acknowledgements

This work was supported by the European Union-funded Live-Mirror project (grant no. 101099220) from the EIC Horizon Pathfinder program. C.L. acknowledges the cooperation of Fabrice Domingues Dos Santos and Mickaël Pruvost from Arkema for their precious advice and the delivery of tailored terpolymer powders.

## Conflicts of Interest

The authors declare no conflicts of interest.

## Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.

## References

1. Live-Mirror, https://eic-live-mirror.eu.

2. G. Moretto, J. R. Kuhn, J.-F. Capsal, et al., "The ExoLife Finder (ELF) Telescope: New Adaptive Optics and Hybrid Dynamic Live-Optical Surfaces Strategies," in *Ground-Based and Airborne Telescopes VII*, ed. R. Gilmozzi, H. K. Marshall, and J. Spyromilio (SPIE, 2018), 158, https://doi.org/10.1117/12.2312599.

3. G. Moretto, J. R. Kuhn, J.-F. Capsal, et al., "Hybrid Dynamic Structures for Optical-Quality Surfaces Shape Control: Live-Mirror." in *Quantum Sensing and Nano Electronics and Photonics XVI*, ed. M. Razeghi, J. S. Lewis, and G. A. Khodaparast (SPIE, 2019), 66, https://doi.org/10.1117/12.2511625.

4. T. C. Chung, A. Petchsuk, and G. W. Taylor, "Ferroelectric Polymers with Large Electrostriction; Based on Semicrystalline VDF/TrFE/CTFE Terpolymers," *Ferroelectrics Letters Section 28* (2001): 135-143, https://doi.org/10.1080/07315170108202957.

5. H. Xu, Z.-Y. Cheng, D. Olson, T. Mai, Q. M. Zhang, and G. Kavarnos, "Ferroelectric and Electromechanical Properties of Poly(vinylidene-fluoride-trifluoroethylene-chlorofluoroethylene) Terpolymer," *Applied Physics Letters* 78, no. 16 (2001): 2360-2362, https://doi.org/10.1063/1.1358847.

6. F. Xia, Z.-Y. Cheng, H. S. Xu, et al., "High Electromechanical Responses in a Poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene) Terpolymer," *Advanced Materials* 14, no. 21 (2002): 1574-1577, https://doi.org/10.1002/1521-4095(20021104)14:21%3C1574::AID-ADMA1574%3E3.0.CO;2-%23.

7. T. Furukawa, "Structure and Functional Properties of Ferroelectric Polymers," *Advances in Colloid and Interface Science* 71-72 (1997): 183-208, https://doi.org/10.1016/S0001-8686(97)90017-8.
8. L. Yang, X. Li, E. Allahyarov, P. L. Taylor, Q. M. Zhang, and L. Zhu, "Novel Polymer Ferroelectric Behavior via Crystal Isomorphism and the Nanoconfinement Effect," *Polymer* 54, no. 7 (2013): 1709-1728, https://doi.org/10.1016/j.polymer.2013.01.035.

9. T. Soulestin, V. Lademiral, F. D. Dos Santos, and B. Améduri, "Vinylidene Fluoride- and Trifluoroethylene-containing Fluorinated Electroactive Copolymers. How Does Chemistry Impact Properties?," *Progress in Polymer Science* 72 (2017): 16-60, https://doi.org/10.1016/j.progpolymsci.2017.04.004.

10. Y. Lu, J. Claude, Q. Zhang, and Q. Wang, "Microstructures and Dielectric Properties of the Ferroelectric Fluoropolymers Synthesized via Reductive Dechlorination of Poly(vinylidene fluoride-co-chlorotrifluoroethylene)s," *Macromolecules* 39, no. 20 (2006): 6962-6968, https://doi.org/10.1021/ma061311i.

11. Y. Hambal, V. V. Shvartsman, D. Lewin, et al., "Effect of Composition on Polarization Hysteresis and Energy Storage Ability of P(VDF-TrFE-CFE) Relaxor Terpolymers," *Polymers* 13, no. 8 (2021): 1343, https://doi.org/10.3390/polym13081343.

12. R. J. Klein, F. Xia, Q. M. Zhang, and F. Bauer, "Influence of Composition on Relaxor Ferroelectric and Electromechanical Properties of Poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene)," *Journal of Applied Physics* 97, no. 9 (2005): 094105, https://doi.org/10.1063/1.1882769.

13. J.-F. Capsal, J. Galineau, M. Lallart, P.-J. Cottinet, and D. Guyomar, "Plasticized Relaxor Ferroelectric Terpolymer: Toward Giant Electrostriction, High Mechanical Energy and Low Electric Field Actuators," *Sensors and Actuators A: Physical* 207 (2014): 25-31, https://doi.org/10.1016/j.sna.2013.12.008.

14. N. Della Schiava, M. Le, J. Galineau, F. Domingues Dos Santos, P. Cottinet, and J. Capsal, "Influence of Plasticizers on the Electromechanical Behavior of a P(VDF-TrFE-CTFE) Terpolymer: Toward a High Performance of Electrostrictive Blends," *Journal of Polymer Science Part B: Polymer Physics* 55, no. 4 (2017): 355-369, https://doi.org/10.1002/polb.24280.

15. J.-F. Capsal, J. Galineau, M.-Q. Le, F. Domingues Dos Santos, and P.-J. Cottinet, "Enhanced Electrostriction Based on Plasticized Relaxor Ferroelectric P(VDF-TrFE-CFE/CTFE) Blends," *Journal of Polymer Science Part B: Polymer Physics* 53, no. 19 (2015): 1368-1379, https://doi.org/10.1002/polb.23776.

16. K. Thetraphpi, S. Chaipo, W. Kanlayakan, et al., "Advanced Plasticized Electroactive Polymers Actuators for Active Optical Applications: Live Mirror," *Advanced Engineering Materials* 22, no. 5 (2020): 1901540, https://doi.org/10.1002/adem.201901540.

17. F. Pedroli, A. Marrani, M. Le, C. Froidefond, P. Cottinet, and J. Capsal, "Processing Optimization: A Way to Improve the Ionic Conductivity and Dielectric Loss of Electroactive Polymers," *Journal of Polymer Science Part B: Polymer Physics* 56, no. 16 (2018): 1164-1173, https://doi.org/10.1002/polb.24636.

18. M. R. Gadinski, Q. Li, G. Zhang, X. Zhang, and Q. Wang, "Understanding of Relaxor Ferroelectric Behavior of Poly(vinylidene fluoride-trifluoroethylene-chlorotrifluoroethylene) Terpolymers," *Macromolecules* 48, no. 8 (2015): 2731-2739, https://doi.org/10.1021/acs.macromol.5b00185.

19. Q. Liu, C. Richard, and J.-F. Capsal, "Control of Crystal Morphology and its Effect on Electromechanical Performances of Electrostrictive P(VDF-TrFE-CTFE) Terpolymer," *European Polymer Journal* 91 (2017): 46-60, https://doi.org/10.1016/j.eurpolymj.2017.03.046.

20. T. C. Chung and A. Petchsuk, "Synthesis and Properties of Ferroelectric Fluoroterpolymers with Curie Transition at Ambient Temperature," *Macromolecules* 35, no. 20 (2002): 7678-7684, https://doi.org/10.1021/ma020504c.

21. T. Yagi, M. Tatemoto, and J. Sako, "Transition Behavior and Dielectric Properties in Trifluoroethylene and Vinylidene Fluoride Copolymers," *Polymer Journal* 12, no. 4 (1980): 209-223, https://doi.org/10.1295/polymj.12.209.

22. Y. Lu, J. Claude, B. Neese, Q. Zhang, and Q. Wang, "A Modular Approach to Ferroelectric Polymers with Chemically Tunable Curie Temperatures and Dielectric Constants," *Journal of the American Chemical Society* 128, no. 25 (2006): 8120-8121, https://doi.org/10.1021/ja062306x.

23. H.-M. Bao, J.-F. Song, J. Zhang, Q.-D. Shen, C.-Z. Yang, and Q. M. Zhang, "Phase Transitions and Ferroelectric Relaxor Behavior in P(VDF-TrFE-CFE) Terpolymers," *Macromolecules* 40, no. 7 (2007): 2371-2379, https://doi.org/10.1021/ma062800l.

24. R. J. Klein, J. Runt, and Q. M. Zhang, "Influence of Crystallization Conditions on the Microstructure and Electromechanical Properties of Poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene) Terpolymers," *Macromolecules* 36, no. 19 (2003): 7220-7226, https://doi.org/10.1021/ma034745b.

25. Z. Zhang, Q. Meng, and T. C. M. Chung, "Energy Storage Study of Ferroelectric Poly(vinylidene fluoride-trifluoroethylene-chlorotrifluoroethylene) Terpolymers," *Polymer* 50, no. 2 (2009): 707-715, https://doi.org/10.1016/j.polymer.2008.11.005.

26. A. Bondi, "van der Waals Volumes and Radii," *The Journal of Physical Chemistry* 68, no. 3 (1964): 441-451, https://doi.org/10.1021/j100785a001.

27. L. Benkraled, A. Zennaki, L. Zair, et al., "Effect of Plasticization/Annealing on Thermal, Dynamic Mechanical, and Rheological Properties of Poly(Lactic Acid)," *Polymers* 16, no. 7 (2024): 974, https://doi.org/10.3390/polym16070974.

28. X. Cai, T. Lei, D. Sun, and L. Lin, "A Critical Analysis of the α,β and γ Phases in Poly(vinylidene fluoride) Using FTIR," *RSC Advances* 7, no. 25 (2017): 15382-15389, https://doi.org/10.1039/C7RA01267E.

29. M. Kobayashi, K. Tashiro, and H. Tadokoro, "Molecular Vibrations of Three Crystal Forms of Poly(vinylidene fluoride)," *Macromolecules* 8, no. 2 (1975): 158-171, https://doi.org/10.1021/ma60044a013.

30. M. A. Bachmann, W. L. Gordon, J. L. Koenig, and J. B. Lando, "An Infrared Study of Phase-III Poly(vinylidene fluoride)," *Journal of Applied Physics* 50, no. 10 (1979): 6106-6112, https://doi.org/10.1063/1.325780.

31. P. Martins, A. C. Lopes, and S. Lanceros-Mendez, "Electroactive Phases of Poly(vinylidene fluoride): Determination, Processing and Applications," *Progress in Polymer Science* 39, no. 4 (2014): 683-706, https://doi.org/10.1016/j.progpolymsci.2013.07.006.

32. N. Pereira, S. Gonçalves, J. C. Barbosa, et al., "High Dielectric Constant Poly(vinylidene fluoride-trifluoroethylene-chlorofluoroethylene) for Capacitive Pressure and Bending Sensors," *Polymer* 214 (2021): 123349, https://doi.org/10.1016/j.polymer.2020.123349.

33. Y. Huang, J.-Z. Xu, T. Soulestin, et al., "Can Relaxor Ferroelectric Behavior be Realized for Poly(vinylidene fluoride-co-chlorotrifluoroethylene) [P(VDF-CTFE)] Random Copolymers by Inclusion of CTFE Units in PVDF Crystals?," *Macromolecules* 51, no. 14 (2018): 5460-5472, https://doi.org/10.1021/acs.macromol.8b01155.

34. H. Xie, X. Yuan, H. Luo, and D. Zhang, "The Effects of Chlorotrifluoroethylene (CTFE) Units on the Dielectric and Energy Storage Properties of P(VDF-CTFE)," *IET Nanodielectrics* 7, no. 1 (2024): 18-25, https://doi.org/10.1049/nde2.12075.

35. T. Clark, M. Hennemann, J. S. Murray, and P. Politzer, "Halogen Bonding: The σ-Hole," *Journal of Molecular Modeling* 13, no. 2 (2007): 291-296, https://doi.org/10.1007/s00894-006-0130-2.

36. P. Politzer, J. S. Murray, and M. C. Concha, "Halogen Bonding and the Design of New Materials: Organic Bromides, Chlorides and Perhaps Even Fluorides as Donors," *Journal of Molecular Modeling* 13 (2007): 643-650, https://doi.org/10.1007/s00894-007-0176-9.

37. P. Politzer, P. Lane, M. C. Concha, Y. Ma, and J. S. Murray, "An Overview of Halogen Bonding," *Journal of Molecular Modeling* 13, no. 2 (2007): 305-311, https://doi.org/10.1007/s00894-006-0154-7.

38. Y. Liu, R. Zhang, X. Wang, et al., "The Strong Interaction Between Poly(vinyl chloride) and a New Eco-Friendly Plasticizer: A Combined Experiment and Calculation Study," *Polymer* 55, no. 12 (2014): 2831-2840, https://doi.org/10.1016/j.polymer.2014.04.037.
39. P. Hobza and Z. Havlas, "Blue-Shifting Hydrogen Bonds," *Chemical Reviews* 100, no. 11 (2000): 4253-4264, https://doi.org/10.1021/cr990050q.

40. J. Joseph and E. D. Jemmis, "Red-, Blue-, or No-Shift in Hydrogen Bonds: A Unified Explanation," *Journal of the American Chemical Society* 129, no. 15 (2007): 4620-4632, https://doi.org/10.1021/ja067545z.

41. P. Gassmann, C. Bohlmann, and V. Pintus, "Towards the Understanding of the Aging Behavior of p-PVC in Close Contact with Minced Meat in the Artwork POEMETRIE by Dieter Roth," *Polymers* 15, no. 23 (2023): 4558, https://doi.org/10.3390/polym15234558.

42. M. A. A. Ibrahim, O. A. M. Ahmed, N. A. M. Moussa, S. El-Taher, and H. Moustafa, "Comparative Investigation of Interactions of Hydrogen, Halogen and Tetrel Bond Donors with Electron-Rich and Electron-Deficient π-Systems," *RSC Advances* 9, no. 56 (2019): 32811-32820, https://doi.org/10.1039/C9RA08007D.

43. D. Von Der Heiden, A. Vanderkooy, and M. Erdélyi, "Halogen Bonding in Solution: NMR Spectroscopic Approaches," *Coordination Chemistry Reviews* 407 (2020): 213147, https://doi.org/10.1016/j.ccr.2019.213147.

44. S. Zanchi, M. Engel, A. Pascaud, et al., "Unraveling the Morphological Diversity of P(VDF-ter-TrFE-ter-CTFE) Semi-Crystalline Terpolymers via Combined AFM and SAXS Experiments," *Polymer Testing* 120 (2023): 107973, https://doi.org/10.1016/j.poly testing.2023.107973.

45. Z. Wang, Z. Zhang, and T. C. M. Chung, "High Dielectric VDF/TrFE/CTFE Terpolymers Prepared by Hydrogenation of VDF/CTFE Copolymers: Synthesis and Characterization," *Macromolecules* 39, no. 13 (2006): 4268-4271, https://doi.org/10.1021/ma060738m.

46. M. Q. Le, J.-F. Capsal, J. Galineau, et al., "All-Organic Electrostrictive Polymer Composites with Low Driving Electrical Voltages for Micro-Fluidic Pump Applications," *Scientific Reports* 5, no. 1 (2015): 11814, https://doi.org/10.1038/srep11814.

47. B. S. Tomar, A. Shahin, and M. S. Tirumkudulu, "Cracking in Drying Films of Polymer Solutions," *Soft Matter* 16, no. 14 (2020): 3476-3484, https://doi.org/10.1039/C9SM02294E.

48. J.-F. Capsal, E. Dantras, and C. Lacabanne, "Molecular Mobility Interpretation of the Dielectric Relaxor Behavior in Fluorinated Copolymers and Terpolymers," *Journal of Non-Crystalline Solids* 363 (2013): 20-25, https://doi.org/10.1016/j.jnoncrys ol.2012.12.008.

49. Q. Liu, "Development of Electrostrictive P(VDF-TrFE-CTFE) Terpolymer for Inkjet Printed Electromechanical Devices" (Institut National des Sciences Appliquées de Lyon).

50. M. L. Williams, R. F. Landel, and J. D. Ferry, "The Temperature Dependence of Relaxation Mechanisms in Amorphous Polymers and Other Glass-Forming Liquids," *Journal of the American Chemical Society* 77, no. 14 (1955): 3701-3707, https://doi.org/10.1021/ja01619a008.

51. M. Samet, V. Levchenko, G. Boiteux, G. Seytre, A. Kallel, and A. Serghei, "Electrode Polarization vs. Maxwell-Wagner-Sillars Interfacial Polarization in Dielectric Spectra of Materials: Characteristic Frequencies and Scaling Laws," *The Journal of Chemical Physics* 142, no. 19 (2015): 194703, https://doi.org/10.1063/1.4919877.

**Supporting Information**

Additional supporting information can be found online in the Supporting Information section.

**Supporting file:** mame70119-sup-0001-SuppMat.docx