# Nanometric and Macroscopic Electroactive Response of P(VDF-co-TrFE) Copolymers and P(VDF-ter-TrFE-ter-CTFE) Terpolymers

Sara Zanchi, Antoine Pascaud, Fabrice Domingues Dos Santos, Sylvie

Tencé-Girault, Sébastien Roland

► To cite this version:

Sara Zanchi, Antoine Pascaud, Fabrice Domingues Dos Santos, Sylvie Tencé-Girault, Sébastien Roland. Nanometric and Macroscopic Electroactive Response of P(VDF-co-TrFE) Copolymers and P(VDF-ter-TrFE-ter-CTFE) Terpolymers. ACS Applied Polymer Materials, 2023, 5 (7), pp.4966-4976. 10.1021/acsapm.3c00534. hal-05488019

HAL Id: hal-05488019

https://hal.science/hal-05488019v1

Submitted on 1 Feb 2026

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# Nanometric and Macroscopic Electroactive Response of P(VDF-co-TrFE) Copolymers and P(VDF-ter-TrFE-ter-CTFE) Terpolymers

Sara Zanchi, Antoine Pascaud, Fabrice Domingues Dos Santos, Sylvie Tencé-Girault,*
and Sébastien Roland*

**ABSTRACT:** The knowledge and the improvement of the electric responsiveness of P(VDF-co-TrFE) copolymer and P(VDF-ter-TrFE-ter-CTFE) terpolymer films are essential for their application in electronic devices. Many factors affect the performances of these electroactive films, such as the chemical composition and the thermal treatment. In this article, the response to an external electric field of copolymer and terpolymer films at both the nanometric and macroscopic scales is studied. The response at the nanometric scale is analyzed by piezoresponse force microscopy on thin films (~100 nm thick), while that at the macroscopic scale is analyzed by measuring the electric displacement-electric field loops on thick films (~20 µm thick). The effect of the chemical composition (CTFE content) is found to be crucial, both at the nanometric and macroscopic scales, in defining the electroactive behavior of the samples, while the thermal treatment seems more exacerbated at the nanometric scale. The orientation, shape, and size of the crystalline lamellae affect the orientation of the electric dipoles, which are thus more or less sensitive to the applied electric field. The controlled growth of edge-on lamellar stacks, obtained after crystallization from the melt at low cooling rates, provides a better piezoresponse of the film.

**KEYWORDS:** piezoelectric force microscopy, PVDF, electric displacement, CTFE

<figure><img src="image_2.png" /></figure>

## 1. INTRODUCTION

Fluorinated electroactive polymers, such as poly(vinylidene fluoride-co-trifluoroethylene) P(VDF-co-TrFE) and poly(vinylidene fluoride-ter-trifluoroethylene-ter-chlorotrifluoroethylene) P(VDF-ter-TrFE-ter-CTFE), have attracted increasing interest in the past few decades, because of their various potential applications in the field of printed organic electronics. The piezoelectric and ferroelectric (FE) behavior of P(VDF-co-TrFE) copolymers is mainly exploited for sensors, memories, and energy harvesting devices, while the electrostrictive and relaxor ferroelectric (RFE) behavior of P(VDF-ter-TrFE-ter-CTFE) terpolymers is essentially used for actuators and transistors.¹⁻⁴ Both FE and RFE behaviors are related to the nature of the crystalline phases, which, among others, is closely linked to the chemical composition of the polymer.

In particular, as discussed recently, the fraction of CTFE in P(VDF-ter-TrFE-ter-CTFE) terpolymers plays a key role in defining their macroscopic electromechanical properties.⁵ The inclusion of 6–10 mol % of the bulky termonomer containing the chlorine atom induces the formation of a single crystalline phase (RFE phase) with increased cell parameters perpendicular to the chain direction and reduces the size of the ferroelectric dipolar domains within the crystalline phase. These ferroelectric nanodomains contain few dipoles, which

align themselves under an external electric field but lose their orientation when the electric field is removed. This is due to the lack of cooperativity of dipoles on a large scale, given that CTFE interrupts the alignment of the dipoles along the chain and the inter-planar distances associated to the RFE phase are significantly higher than those of the ferroelectric and defective ferroelectric phases.⁶

In addition to the chemical composition, the morphology of these semi-crystalline polymers must be considered to understand their electroactive behavior. Several studies for P(VDF-co-TrFE) copolymers demonstrated that thermal annealing or melting/crystallization (MC) treatments produced different crystalline morphologies in thin films that correspond to different stacking of lamellae (rice-like grains after annealing and needle-like grains after MC).⁷⁻¹¹ These stacks of lamellae can present various orientations with respect to the substrate. In a previous study, we showed that the geometry of these semi-crystalline domains also depends on
both the cooling rate (for the MC samples) and the CTFE content.¹² The orientation of the crystalline lamellae is crucial for the electric responsiveness, since electric dipoles have a specific orientation with respect to the polymer chain (along the *b* direction of the crystalline lattice).¹³,¹⁴

The electric response of different morphologies at the nanometric scale can be measured by piezoresponse force microscopy (PFM).¹⁵,¹⁶ Up to now, this technique has been used on few polymers in the shape of films or fibers, namely PLA,¹⁷⁻²⁰ polyamide-11,²¹,²² and PVDF-based copolymers and terpolymers. Concerning the latter, some literature exists about the ferroelectric properties of different crystalline morphologies and orientations at the surface of thin films.²³⁻³⁵
The first attempt to measure the local polarization of a P(VDF-co-TrFE) copolymer thin film dates back to 1992 and

led to the first image of polarization switching at the microscopic scale (1 µm) for these systems.²⁶ PFM analyses were performed again on P(VDF-co-TrFE) copolymers only in 2008 and concerned the polarization switching kinetics of Langmuir–Blodgett films.²⁷ A pulsed electric field was locally applied to a pre-polarized surface (with the opposite sign). PFM phase images were used to measure the spontaneous polarization of the film, and hysteresis loops of the piezoresponse with respect to the applied voltage were plotted. PFM has also been used to observe the thermal evolution of the ferroelectric behavior at the nanoscale: the piezoresponse of the film is stable until the Curie transition and then disappears abruptly.²⁸ Some authors have also linearly coupled the piezoresponse at the local scale given by PFM with the macroscopic polarization state of these thin films under an external electric field.²⁹

Yet, the effect of parameters such as the thickness or the morphology of the film were not taken into account. Lee *et al.* studied the effect of the thermal treatment upon the polymer chain orientation, crystalline morphology, and ferroelectricity of P(VDF-co-TrFE) copolymer thin films for non-volatile memory devices.²³ As-cast, thermally annealed and melt/crystallized samples were first studied by atomic force microscopy (AFM). The morphology of the annealed sample is constituted by a high amount of edge-on-oriented crystalline lamellae, while the as-cast and melt/crystallized ones do not have or have only few edge-on-oriented lamellae. This was responsible for a more pronounced ferroelectric behavior of the annealed sample, which was the only one successively analyzed by PFM. The crystalline orientation was also probed by Wu *et al.*, who measured an important piezoresponse for their isothermally annealed sample (edge-on orientation), while no signal for the isothermally melt/crystallized one (flat-on orientation) was observed.³⁰ In this case, while the observed morphologies were very different from those observed after a non-isothermal cooling, the PFM analyses led to the same conclusion and were closely related to the orientation of the crystalline lamellae. The effects of thermal treatments above the melting temperature, this time with a non-isothermal crystallization, have been further studied on P(VDF-co-TrFE) copolymers.³¹ Phase images show that the polarization is confined within the needle-like structures (edge-on lamellar orientation) and follow their direction. Those dipoles are able to switch when the electric field is reversed, while the areas between the needles are not polarizable. The ferroelectric behavior of needle-like crystals was also confirmed by Kim and Khang, who highlighted the importance of the polymer chain and crystalline lamellae orientation to facilitate

the switching of dipoles.³² Nevertheless, those crystalline domains do not cover the whole film surface, thus making the ferroelectric response heterogeneous.

The effect of solvent annealing (SA) has also been studied by PFM, this time both for P(VDF-co-TrFE) copolymers and P(VDF-ter-TrFE-ter-CTFE) terpolymers.³³,³⁴ While for co-polymers, the SA treatment does not lead to a stronger ferroelectric behavior than the thermal annealing, for terpolymers, it is responsible for the existence of a ferroelectric crystalline fraction (instead of a single RFE phase) and, consequently, of a ferroelectric behavior. Thus, the existing studies lack a comprehensive PFM characterization of P(VDF-ter-TrFE-ter-CTFE) terpolymers through different thermal treatments.

In this paper, we investigated the effect of the chemical composition and of the processing conditions on the electric responsiveness of P(VDF-ter-TrFE-ter-CTFE) terpolymers at both the nanometric and macroscopic scales. The former was achieved by analyzing PFM phase images for 100 nm thick spin-coated films, while the macroscopic behavior under an electric field was probed by performing electric displacement *vs* electric field loops (*D*-*E*) on 20 µm thick bar-coated films. The morphologies of the thin and thick films are the same as those observed and explained in our recent publication.¹² The FE/RFE behaviors with respect to the CTFE fraction (0–10% CTFE) were first studied for a standard and well-known morphology, resulting from the annealing above the Curie temperature and below the melting temperature. As already mentioned, for thin films, this morphology is made of rice-like domains, consisting of parallel crystalline lamellae stacks. The domains may have many different orientations, and thus, the film is homogeneously polarizable. The second part of this article focuses on melt/crystallized samples and particularly on ways to induce changes in the orientation and the stacking of the lamellae, namely the CTFE content and the cooling rate from the melt. Once again, the PFM results and macroscopic *D*-*E* loops were compared, revealing the importance of the crystalline morphology for the electroactive behavior of these terpolymers, at both the nanometric and macroscopic scales. Understanding the behavior of PVDF copolymers and terpolymers under electrical fields is crucial for fully exploiting their potential in future applications. This knowledge can guide the design and optimization of devices based on these materials, such as sensors, actuators, and energy harvesters.

## 2. EXPERIMENTAL PART

2.1. Materials and Methods. A P(VDF-co-TrFE) copolymer (Terpo65-0) and two P(VDF-ter-TrFE-ter-CTFE) terpolymers (Terpo66-4.7 and Terpo67-8.3) powders were provided by Piezotech. Samples were named TerpoX-Y, where X and Y refer to the VDF/(TrFE + VDF) ratio and the amount of CTFE (mol %), respectively (see Table 1). The VDF/(TrFE + VDF) ratio was kept constant (around 66%) for the three chosen chemical compositions. Polymer powders were dissolved in methylethylketone at 14 wt % by using a reflux setup at 80 °C for 3 h to obtain a homogeneous polymer solution. Solutions were then filtered with a 0.6 µm mesh filter. Depending on the analysis, two different techniques were used for the preparation of the samples, namely spin-coating and bar-coating.

2.1.1. *Spin-Coating.* A conductive silicon wafer (100), P type, resistivity = 0.001–0.005 Ohm·cm⁻¹; BT Electronics), cleaned with piranha solution (H₂SO₄/H₂O₂, 3:1) for 1 h, was used as a substrate. The polymer solution was diluted to 2 wt % and spin-coated at 2000 rpm rotational speed for 60 s. The solvent evaporated during the spin-coating, leaving a 100 nm thick film. Two different thermal treatments were performed in the Linkam THMS600 heating stage. The
<table><caption>Table 1. List of the Samples</caption><thead><tr><th>name</th><th>VDF/(VDF + TrFE) mol %</th><th>CTFE mol %</th><th>T<sub>m</sub> (°C)</th><th>treatment for AFM/PFM</th><th>treatment for D-E loops</th></tr></thead><tbody><tr><td>Terpo65-0</td><td>65</td><td>0</td><td>150</td><td>-A10, -MC10, -MC100</td><td>-A, -MC</td></tr><tr><td>Terpo66-4.7</td><td>66</td><td>4.7</td><td>135</td><td>-A10, -MC10</td><td>-A, -MC</td></tr><tr><td>Terpo67-8.3</td><td>67</td><td>8.3</td><td>115</td><td>-A10, -MC10</td><td>-A, -MC</td></tr></tbody></table>

annealing (A) was done below the melting temperature ($T_m$), at $T_m - 15$ °C for 1 h, using 10 °C/min as heating and cooling ramps. The melting/crystallization (MC) treatment consisted in melting the film at $T_m + 25$ °C for 15 min, followed by cooling at 10 °C/min or 100 °C/min. The samples were named TerpoX-Y-A10, TerpoX-Y-MC10, and TerpoX-Y-MC100, respectively, and they were used to perform both AFM and PFM analyses.

### 2.1.2. Bar-Coating.
The polymer solution was uniformly spread on a glass substrate thanks to a bar-coater and left 30 min at room temperature, in order to let the solvent evaporate and to get 20 µm thick solvent-cast films. Two different thermal treatments were performed on these films in an oven: an annealing treatment (A) and a melting/crystallization treatment (MC). The former was performed at $T_m - 15$ °C for 1 h, and the latter was performed by cooling after an isotherm at $T_m + 25$ °C for 15 min. In both cases, the oven was turned off after the thermal treatment and the door was left partially open, to slowly cool the polymer films down to room temperature. The samples were named TerpoX-Y-A and TerpoX-Y-MC, and they were used to perform electric-displacement vs electric field (D-E) cycles, in order to obtain the hysteresis loops of the analyzed samples. Table 1 summarizes the samples used in this study and their melting temperatures.

### 2.2. Atomic Force Microscopy.
A Multimode 8 (Bruker) microscope controlled by a Nanoscope V controller and operating in the PeakForce QNM mode was used, in order to avoid the eventual surface damages due to the tip contact. ScanAsyst-Air-HR cantilevers with a nominal frequency of 130 kHz, a nominal spring constant of 0.4 N/m, and a nominal tip radius of 2 nm were used. Topography (contrast provided by the roughness of the sample) and adhesion images (contrast provided by a difference of adhesion forces between the tip and the sample) were analyzed, in order to characterize the morphologies of the different samples.³⁶ For thin films, average film thicknesses were estimated from AFM topography images of scratches made with a razor blade. Typical thickness was around 120 nm.

### 2.3. Piezoresponse Force Microscopy.
The same Multimode 8 (Bruker) microscope and Nanoscope V controller were used. The experiments were performed in Contact mode, with SCM-PIT-V2 platinum-iridium-coated conductive cantilevers (nominal frequency = 75 kHz, nominal spring constant = 3 N/m, and a nominal tip radius of 25 nm). The voltage of the cantilever was manually controlled, in order to apply the right signal to the sample. Two PFM protocols were used. The first one (Figure SI 1) was performed in order to probe the switching ability of the ferroelectric domains for the different samples and was divided in three steps:

1. Polarization at +10 V of a 5 µm × 5 µm surface

2. Polarization at -10 V of a 3 µm × 3 µm surface, inside the first one

3. Revealing of the two polarization steps, by the application of a lower AC voltage (amplitude 2 V; frequency ca. 60 kHz) to a larger area (7 µm × 7 µm)

This experiment gives information about the ferroelectric or relaxor ferroelectric behavior of the sample. The former is associated to a high remnant polarization and dipole reversal under electric field, while the latter is related to a very low remnant polarization and to an instantaneous loss of the orientation of the dipoles when the external electric field is removed.

The second protocol was performed in order to study the evolution of the piezoresponse with respect to the applied voltage and led to the plot of local hysteresis (Δφ-E) loops. This protocol is described in the Supporting Information (Figures SI 2 and 3).

### 2.4. Electric Displacement-Electric Field (D-E) Loops.
Films were placed between two gold electrodes and then the system was

kept in compression thanks to a hydraulic press (150 and 50 kN for the copolymer and the terpolymers films, respectively). Electrodes were connected to a high-voltage amplifier and an ammeter. The poling process was performed on 1 cm² surfaces of both annealed and melt/crystallized samples at room temperature. The charging-discharging electric field cycles were first applied to 25 V/µm followed by 25 V/µm steps until reaching 100 V/µm at 0.06 Hz frequency. The measured current was then integrated over time. This numerical treatment gave access to the D-E loops and to their characteristic parameters: the remnant polarization ($P_r$) and the coercive field ($E_c$). $P_r$ and $E_c$ are measured at the intersection of the D-E loop with the ordinate and abscissa axis, respectively.³⁷

## 3. RESULTS

The objective of this article is to conduct a multi-scale study of samples that are known to have three different electroactive behaviors: ferroelectric (FE), relaxor ferroelectric (RFE), and a mixed behavior (FE/RFE) between these two. As stated by Bargain *et al.*,⁵ the copolymer (without CTFE) is FE, while the terpolymer with 8 mol % CTFE is RFE. The FE/RFE behavior is located between 4 and 7 mol % CTFE. We have therefore selected the compositions according to these behaviors.

In the following section, the samples that are annealed and crystallized from the melt will be analyzed, compared, and discussed.

### 3.1. Annealed Samples.
The morphology of an annealed thin film of copolymer (Terpo65-0-A10) is shown in Figure 1.

<figure><img src="image_3.png" /><figcaption>Figure 1. AFM topography images for Terpo65-0-A10 at the (a) 5 µm × 5 µm and (b) 0.5 µm × 0.5 µm scales. Inset: adhesion image at the 0.25 µm × 0.25 µm scale.</figcaption></figure>

As reported in the literature, it is made of a multitude of thin rice-like semi-crystalline domains, which have many different orientations.¹⁰ Figure 1a shows a 5 × 5 µm² topographic AFM image in which the surface is rather smooth. It is worth noting that these domains are evenly distributed over the entire surface, yielding a very homogeneous film with low surface roughness (RMS roughness = 3.6 nm). The internal structure of the domains is shown in a zoomed-in AFM image (500 × 500 nm²) in Figure 1b. Each domain is approximately 100 nm long and 50 nm wide. The domains are made of the parallel stacking of some crystalline lamellae, which are separated by a layer of amorphous phase (darker in the topographic image and lighter in the adhesion image). Between each semi-crystalline domain, an amorphous phase can also be detected.
Both the topography and adhesion AFM images reveal that the rice-like domains have many different orientations: the crystalline lamellae can be parallel to the substrate (flat-on), perpendicular to it (edge-on), but they mainly have intermediate orientations between the two extremes.

A very similar morphology exists for P(VDF-ter-TrFE-ter-CTFE) terpolymer annealed thin films, as shown in Figure SI 4. When the CTFE content increases up to 8.3 mol %, the shape factor of the semi-crystalline domains, defined as the length/width ratio, increases from ca. 1.7 to 4 with a strong decrease of their width (down to 25–30 nm), as reported in our recent study.¹² The morphology shown in Figure 1a is the starting point for the PFM measurements in this study, performed on Terpo65-0, Terpo66-4.7, and Terpo67-8.3 thin films. Figure 2 depicts the phase images resulting from the first PFM protocol described in Section 2.3 (also illustrated in Figure SI 1) for Terpo65-0-A10, Terpo66-4.7-A10, and Terpo67-8.3-A10.

<figure><img src="image_4.png" /><figcaption>Figure 2. PFM phase images for (a) Terpo65-0-A10, (b) Terpo66-4.7-A10, and (c) Terpo67-8.3-A10. (d) Evolution of the mean phase along a straight 7 µm long profile, for all compositions.</figcaption></figure>

The phase signal of the piezoresponse in PFM is related to the displacement response and to the dipole orientation in the sample.²⁹ In the case of VDF-based copolymers and terpolymers showing ferroelectricity and piezoelectricity, the piezoelectric coefficient $d_{33}$ is negative.³⁸ This means that, when a positive tip bias is applied to a pure ferroelectric single-crystal domain with dipoles oriented upward (from the back electrode to the sample surface, $E = -10$ V), an out-of-plane expansion of the domains, which is in phase with the alternative driving voltage, is expected. The phase value must, therefore, be close to 0°. On the contrary, an out-of-plane phase value of ±180° corresponds to an out-of-phase response (contraction) of the domain when the dipoles are oriented downward (from the sample surface to the back electrode, $E = +10$ V).²⁹

The inner 3 µm × 3 µm square surface in Figure 2a–c is the result of the polarization switching from +10 to -10 V, while the surrounding dark frame only corresponds to the pre-polarization at +10 V. Since these ferroelectric polymers are

semi-crystalline and contain many ferroelectric domains with imperfect orientations, it is impossible to reach 0 and 180° phase values:²⁹ the +10 V polarized area corresponds to an out-of-phase surface, while the -10 V polarized area corresponds to an in-phase surface, even if the phase range is shorter than 180°. Due to the various orientations of the crystalline lamellae, the unpolarized outer frame should have an average phase value of ±90°. A value around -115° is measured, which is roughly the same for Terpo65-0-A10, Terpo66-4.7-A10, and Terpo67-8.3-A10 (Figure 2d).¹⁶,²⁹ The pre-polarization at +10 V of the 5 × 5 µm² surface leads to different electric responses for the three samples: Terpo65-0-A10 and Terpo66-4.7-A10 reach a -145 and -150° phase value, respectively, while the phase for Terpo67-8.3-A10 remains at a value closer to that of the unpolarized area (ca. -120°). Then, the application of a -10 V voltage to the inner 3 × 3 µm² area is responsible for the dipolar switching within the ferroelectric domains of the three samples. Terpo65-0-A10 has the strongest dipolar switching, since its phase value goes from -145 to -65°. This strong increase is related to high crystallinity and the pure ferroelectric behavior of Terpo65-0-A10.⁵ As shown in Figure 2b, Terpo66-4.7-A10 still shows a noticeable phase contrast between the in-phase and out-of-phase areas, but it is characterized by a lower dipolar switching (the phase goes from -150 to -90°). This is a consequence of the lower crystallinity and the mixed FE/RFE behavior of P(VDF-ter-TrFE-ter-CTFE) terpolymers with a CTFE amount lower than 6–7 mol %.⁵ Figure 2c shows that the dipolar switching for Terpo67-8.3-A10 is very weak, since the phase value only goes from -120 to -110°. The inefficiency of the polarization for this sample reflects both the low crystallinity and the pure relaxor ferroelectric behavior of P(VDF-ter-TrFE-ter-CTFE) terpolymers with a high amount of CTFE (>6–7 mol % CTFE).⁵

It must also be pointed out that the phase profiles are more or less symmetric with respect to the unpolarized area, as shown in Figure 2d. For Terpo65-0-A10, the out-of-phase state is achieved with a $\Delta\varphi$ of 30°, while the in-phase state is achieved with a $\Delta\varphi$ of 50°. For Terpo66-4.7-A10, the profile is symmetric (±30° to achieve both the out-of-phase state and the in-phase state), while for Terpo67-8.3-A10, the phase evolutions are very weak and stay close to the -110° initial value.

As reported in Section 2.3 and described precisely in Figures SI 2 and 3, another PFM protocol was used in order to highlight the evolution of the $\Delta\varphi$ between the unpolarized and polarized areas with respect to the applied voltage (from -10 to +10 V). Figure 3 shows the $\Delta\varphi-E$ loops for Terpo65-0-A10, Terpo66-4.7-A10, and Terpo67-8.3-A10. The electric field was calculated with respect to the thickness of each sample.

Since these measurements are made few minutes after the application of the external electric field, the phase signal is representative only of the remnant polarization, $P_r$, (polarization value when the electric field is removed) and thus strongly related to the crystalline structure and the chemical composition of the samples. The $\Delta\varphi$ values between the unpolarized external area and the inner polarized area are slightly different than those measured in Figure 2d. In particular, Terpo65-0-A10 has more symmetric $\Delta\varphi$ values (black curve, $\Delta\varphi \sim 55$ and $\sim -45$° in Figure 3 for the highest and lowest electric field, respectively) than in Figure 2d (black curve, $\Delta\varphi \sim 50$° and $\sim -30$° for -10 and +10 V respectively).
<figure><img src="image_5.png" /><figcaption>Figure 3. Δφ-E loops obtained with the PFM phase images for Terpo65-0-A10, Terpo66-4.7-A10, and Terpo67-8.3-A10. (see Section 2.3 and Figures SI 2 and 3 for the detailed procedure).</figcaption></figure>

For Terpo66-4.7-A10, the two protocols lead to a symmetric behavior with respect to Δφ, but the maximum values measured in Figure 3 (red curve, Δφ ~ 40°) are higher than those highlighted by the phase profile in Figure 2d (red curve, Δφ ~ 30°). Currently, we lack a definitive explanation for this discrepancy, which could potentially be attributed to the differences in the experimental conditions. Finally, the phase values for Terpo67-8.3-A are very similar and very low (around 10°) for the two protocols. In addition to the Δφ values, the second PFM protocol gives information about the electric field responsible for the extinction of the remnant polarization for each sample. These measured electric fields are not perfectly symmetric with respect to the 0 V/μm electric field: it seems that the extinction of the remnant polarization takes place at higher field values for positive applied electric fields than for the negative ones. Terpo65-0-A10 and Terpo66-4.7-A10 have very close values (+38 and -30 V/μm), while for Terpo67-8.3-A10, the values are 33 and -16 V/μm. The measurements for the latter are more difficult, since the phase contrasts are very low for Terpo67-8.3-A10 and the risk of error is higher. The asymmetry of the cycles is rather weak and can be associated with chemical and geometric differences of the two electrodes. On the one hand, the geometry (sharp tip vs flat back electrode) generates a non-homogenous distribution of the electric field.<sup>16,39,40</sup> On the other hand, the difference in surface chemistry of the two electrodes could induce asymmetric physical bonds between the two electrodes and the hydrogen and the fluorine atoms, and thus a difference in switching properties of the dipoles.<sup>41</sup>

Overall, the shapes of the loops presented in Figure 3 are characteristic of the different behaviors under an external electric field. Terpo65-0-A10 and Terpo66-4.7-A10 have a quite large and rectangular loop, with high Δφ values (which is representative of the remnant polarization), while Terpo67-8.3-A10 has a very thin loop, characterized by very low Δφ values. Moreover, a higher electric field is necessary for Terpo65-0-A10 and Terpo66-4.7-A10 compared to Terpo67-8.3-A10, to remove the remnant polarization. These curves measured at the nanometric scale may remind of the macroscopic loops of the electric displacement with respect to the applied electric field, even if the latter are recorded in situ (under electric field) and thus do not reflect only the

evolution of the remnant polarization but rather the instantaneous state of polarization.

The D-E loops at 100 V/μm for the 20 μm thick bar-coated and annealed films are shown in Figure 4.

<figure><img src="image_6.png" /><figcaption>Figure 4. D-E loops at 100 V/μm for Terpo65-0-A, Terpo66-4.7-A, and Terpo67-8.3-A.</figcaption></figure>

The cycle of Terpo65-0-A (black curve) is characterized by a high coercive field (41 V/μm) and high remnant polarization (47 mC/m²). For this sample, the hysteresis loop is large and its shape is similar to a rectangle, with an abrupt change of electric displacement at the coercive field. Terpo67-8.3-A (green curve) shows a very different behavior from that of Terpo65-0-A: its D-E loop is much thinner and sloped, with a coercive field of 15 V/μm and a remnant polarization of 9 mC/m². The cycle for Terpo66-4.7-A (red curve) has an intermediate behavior, with a $P_r$ of 16 mC/m², while its coercive field is the same as that of Terpo67-8.3-A (15 V/μm) and markedly lower than that of Terpo65-0-A. This trend is similar to the one presented by Bargain *et al.* with their D-E cycles at 150 V/μm for annealed terpolymer films with a CTFE fraction ranging from 0 to 8.5 mol %.⁵ It confirms that a 100 V/μm electric field is sufficient to study the electroactive behavior of these polymers, since their coercive field is quite lower than this value for all the analyzed samples. It can be assessed that the D-E cycle for Terpo65-0-A is characteristic of a ferroelectric behavior, while the one of Terpo67-8.3-A is characteristic of a relaxor ferroelectric behavior.

Our previous article focused on the morphology of P(VDF-co-TrFE) copolymers and P(VDF-ter-TrFE-ter-CTFE) terpolymers.<sup>12</sup> Thin and thick films were analyzed, and their morphology was found to be similar and to have similar characteristic dimensions, as recalled in Figure SI 5a,c for Terpo65-0-A and Terpo67-8.3-A thick films. Thus, it seems reasonable to compare the D-E cycles performed on these thin and thick films. While Figures 3 and 4 depict a similar behavior at the nanometric and macroscopic scale for Terpo65-0-A and Terpo67-8.3-A, the results for Terpo66-4.7-A are significantly different at the two scales. Its local Δφ-E loop, measured on thin film, shows a more ferroelectric behavior than that of the macroscopic D-E loop. This might be interpreted by a different amount of the crystalline phases (FE, RFE) in the thin and thick films. Crystallization process in thin films, during spin-coating, might induce the crystallization of a higher amount of ferroelectric phase.
<figure><img src="image_7.png" /><figcaption>Figure 5. AFM topography images for Terpo65-0-MC10 at the (a) 5 µm × 5 µm and (b) 0.5 µm × 0.5 µm scales. Inset: DMT modulus image at the 0.5 µm × 0.5 µm scale. The height profile below (a) is measured along the blue line.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>Figure 6. (a) PFM phase image for Terpo65-0-MC10 and (b) phase measurement along the edge-on and flat-on lamellae stacks. Red arrows indicate the propagation of polarization within the lamellae stacks and outside the in-phase polarized area.</figcaption></figure>

**3.2. Melt/Crystallized Samples.** In addition to the annealed samples, melt/crystallized samples were analyzed. Figure 5 shows the characteristic morphology for a Terpo65-0-MC10 thin film: both flat-on and edge-on crystalline lamellar stacks exist (Figure 5a). The observed flat-on lamellae are similar to those already observed by Wu *et al*. in their study about the morphology of P(VDF-co-TrFE) thin films after an isothermal crystallization from the molten state.30 The edge-on stacks are described in the literature as needle-like structures, but in Figure 5a, the lateral extension (several micrometers) of these rectilinear domains is higher than the ones reported.23,31 Recently, their dimensions were found to depend on the chemical composition and thermal treatment.12 Figure 5b shows the internal structure of such stacking, which is made of many crystalline lamellae with edge-on orientation and alternated with a layer of inter-lamellar amorphous phase. The sample shows a non-uniform thickness, due to the presence of the needle-like semi-crystalline structures: in Figure 5a, the edge-on stack is *ca*. 50 nm higher than the flat-on stack in the background of the image, whose thickness is *ca*. 100 nm.

The electric responsiveness of the melt/crystallized morphology was first studied on a Terpo65-0-MC10 thin film. Similarly to the results shown in Figure 2 for the annealed thin films, Figure 6a shows the PFM phase image for Terpo65-0-MC10 and Figure 6b shows the phase profiles along the dotted lines of Figure 6a. The external perimeter of the out-of-phase frame (+10 V) is less distinguishable than for Terpo65-0-A10, but a phase contrast is still present between the in-phase (−10 V) and the out-of-phase (+10 V) areas. This is highlighted in Figure 6b: the phase step between these two regions, within an edge-on lamella stack, is *ca*. 100°, while it is *ca*. 50° along the flat-on ones. It must be pointed out that the phase values are related to the polarization state of the thin film and then to the orientation of dipoles, which are perpendicular to the polymer chain and can rotate in a plane perpendicular to the polymer chain axis. For the flat-on lamellae, the chain axes are expected perpendicular to the substrate; thus, the dipoles and their rotation plane are parallel to the substrate, and therefore insensitive to the application of the electric field applied by the AFM probe. Therefore, the electro-responsiveness, observed in Figure 6 for the flat-on lamellae stacks, are unexpected. To explain the observed response, it is necessary
<figure><img src="image_9.png" /><figcaption>Figure 7. Topography and PFM phase images for (a,a') Terpo65-0-MC100 and (b,b') Terpo67-8.3-MC10.</figcaption></figure>

to imagine that the dipoles have a component parallel to the electric field and can rotate around the polymer chain. This dipole orientation could originate from a tilt angle of the chains with respect to the lamellae stacks surface, which would result in a nonzero dipole moment component parallel to the electric field. We do not have more data to support this hypothesis, but the possible existence of a tilt has already been published. Indeed, a tilt angle of 8° has already been reported for a P(VDF-co-TrFE) copolymer with 35 mol % of TrFE in the literature through *ab initio* calculations.⁴² This result is also consistent with the AFM observation of tilted crystalline lamellae within these rectilinear semi-crystalline domains.¹² It can also be related to a misorientation of lamellae stacks, which thus are not perfectly parallel to the substrate, as reported in the height profile of Figure 5a, in the range between 1.2 and 2.2 µm of horizontal distance. A combination of both contributions is also not excluded.

In addition, the in-phase polarization propagates beyond the -10 V polarized area along the edge-on lamellar stacks: the maximum phase value (ca. -45°) can be measured even ca. 100 nm far from the border of the out-of-phase area, as highlighted by the two red arrows in Figure 6a. This happens also at the border between the out-of-phase and the unpolarized regions, since the phase value of the edge-on stacks can be measured also beyond this boundary. This can be interpreted as a preferential propagation of the remnant polarization of the ferroelectric domains within the edge-on stacks beyond the polarized area. This preferential propagation of the polarization, probably favored by the switching cooperativity of the dipoles, is in agreement with the literature.³¹ Outside the +10 V polarized area, the phase value is -110° for both crystalline structures, meaning that there is no preferential dipolar direction if the sample does not

undergo a polarization, as already observed for the annealed thin films (Figure 2).

The results presented in Figure 6 are specific to a chemical composition and a single processing condition (crystallization from the melt at 10 °C/min). As explained in a previous study, for high amounts of CTFE and for high cooling rates from the molten state, the extension, shape, and orientation of the semi-crystalline lamellae stacks change dramatically with the amount of CTFE or/and with cooling rates from the molten state. For high amounts of CTFE or for high cooling rates, the lateral extension of the semi-crystalline domains decreases, they stack over a shorter distance, and they have a less defined orientation (petal-like).¹² It is thus possible to have equivalent morphologies for different chemical compositions of the terpolymer and thus for different crystal structures (FE or RFE). We selected a sample without CTFE cooled at high cooling rate (Terpo65-0-MC100) and a sample with high CTFE content slowly cooled (Terpo67-8.3-MC10), to compare their response under electric field.

As shown in Figure 7, the same PFM protocol was applied on Terpo65-0-MC100 and Terpo67-8.3-MC10, in order to study the effect of the morphology on the electric responsiveness for high cooling rates and high CTFE contents, respectively.

Figure 7a shows the topography image of a thin film of Terpo65-0-MC100, obtained in contact mode during the PFM analysis. The edge-on lamellae stacks still exist after a crystallization from the melt at 100 °C/min, but they stack over a shorter distance and they are less straight and less thick than for Terpo65-0-MC10. In addition to these edge-on stacks, a large number of crystalline lamellae stacks with ill-defined orientation (petal-like) exists. Figure 7a' gives the switching behavior of this morphology. The in-phase and out-of-phase areas are once again blurred in comparison to those of
Terpo65-0-A10 (Figure 2a), but a strong phase contrast can still be observed in some parts of the image. The thick edge-on stacks in both in-phase and out-of-phase regions are more visible than the rest of the morphologies (petal-like and flat-on lamellae stacks). The phase difference (Δφ) along these thicker edge-on stacks is ca. 40° (from -90 to -130°), and the polarization still propagates preferentially along these semi-crystalline domains. In the surroundings, the phase contrast is weaker and depends on the lamellar stacking orientation: the thin petal-like lamellar stacks still show a noticeable phase contrast (between 15 and 20°) at the boundary between the out-of-phase and the in-phase areas, while no phase contrast is visible in correspondence of the flat-on lamellae stacks (its phase is the same as that of the external unpolarized area, i.e. -110°).

Figure 7b shows the topography image (in contact mode) of a Terpo67-8.3-MC10 thin film. As for Terpo65-0-MC100, the edge-on lamellar stacks are thinner than for Terpo65-0-MC10 (Figure 6a) and the rest of the image is made of lamellae stacks that have many different orientations. The phase image for this sample is presented in Figure 7b': no phase contrast exists between the in-phase and out-of-phase polarized areas, which means that the dipoles within the edge-on lamellae do not have a better switching behavior compared to the flat-on lamellae under an electric field. The PFM image for Terpo66-4.7-MC10 (Figure SI 6) completes the series: for this sample, with a similar morphology, the phase contrast is strong within the edge-on lamellar stacks (90° phase shift between the out-of-phase and the in-phase areas) and almost negligible (around 10°) outside the edge-on stacks, i.e. within the short and ill-oriented petal-like stacks. The phase contrast within its edge-on lamellae stacks is comparable to that of Terpo65-0-MC10, meaning that their response to an external electric field is similar. The short extension and the petal-like orientation of the other semi-crystalline domains of Terpo66-4.7-MC10 is responsible for the very weak response under electric field (10° in Figure SI 6), while the extended lamellae of Terpo65-0-MC10 still show a 50° phase difference (Figure 6). This may be due to the different size of these lamellae: even if in both cases, the orientation is not the most favorable to induce a strong piezoresponse, the large extension helps for the propagation of the polarization and leads to a higher and more uniform phase value (and polarization strength).

Figures 7 and SI 6 show that the similar morphologies, obtained by increasing the cooling rate (Terpo65-0-MC100) or by increasing the CTFE content (Terpo66-4.7-MC10 and Terpo67-8.3-MC10), lead to different responses to the applied electric field. The different electro-responsiveness is related to the nature of the existing crystalline phases, which determines the FE, RFE, or mixed FE/RFE behaviors.

The macroscopic polarization behavior of Terpo65-0-MC, Terpo66-4.7-MC, and Terpo67-8.3-MC under an external electric field is shown in Figure 8. The shape of the D-E loops is very similar to that of the respective annealed samples (Figure 4), but some differences should be discussed. For Terpo65-0-MC (black curve), an increase of 30% is observed for the remnant polarization, while the coercive field slightly increases from 41 to 43 V/μm. This translates into a more rectangular D-E loop, which is related to a more pronounced ferroelectric behavior. The Pr and Ec values for Terpo66-4.7-MC (red curve) do not show a clear difference with respect to those of Terpo66-4.7-A, but the shape of its D-E loop displays a weak double loop hysteresis (DLH). For Terpo67-8.3-MC

<figure><img src="image_10.png" /><figcaption>Figure 8. D-E loops at 100 V/μm for Terpo65-0-MC, Terpo66-4.7-MC, and Terpo67-8.3-MC.</figcaption></figure>

(green curve), the Pr and Ec values (7 mC/m² and 10 V/μm, respectively) are globally lower than those of Terpo67-8.3-A and a double loop hysteresis (DLH) is well visible.

## 4. DISCUSSION

The results presented in Sections 3.1 and 3.2 are mainly related to the effect of the chemical composition and of the processing conditions on P(VDF-co-TrFE) and P(VDF-ter-TrFE-ter-CTFE) films (0, 4.7, and 8.3 mol % CTFE), at both the nanometric and macroscopic scales.

These results highlighted that the processing conditions (annealing or melting/crystallization at different cooling rates) may also have an impact on the electric responsiveness of the analyzed samples. The PFM images showed that, for the same chemical composition, the phase contrasts are more homogeneous for the annealed samples, but the highest values are obtained within the edge-on lamellae stacks of the melt/crystallized samples that show a ferroelectric behavior. For example, Terpo65-0-MC10 (Figure 6) and Terpo66-4.7-MC10 (Figure SI 6) edge-on stacks reach 100 and 90° phase difference values (Δφ between the out-of-phase and the in-phase areas), respectively. These values are higher than those of the respective annealed samples (Δφ ~ 80° for Terpo65-0-A and Δφ ~ 60° for Terpo66-4.7-A). This might be related to the different orientations and lateral extensions of the semi-crystalline domains. In edge-on lamellae stacks, the polymer chains are parallel to the substrate, and dipoles are perpendicular to the polymer chain and can rotate in a plane perpendicular to the chain axis, thus oriented parallel to the electric field. The annealed samples are made of small rice-like domains with many different orientations with respect to the external electric field, meaning that dipoles have many different orientations and their overall response is surely lower than that of a pure edge-on lamellae stacks. The phase difference measured outside the edge-on stacks confirm this trend, since these lamellae stacks are almost flat-on (even if with a more or less pronounced tilt): the phase difference (Δφ) in these regions are 50 and 10° for Terpo65-0-MC10 and Terpo66-4.7-MC10, respectively. These values are lower than those for the respective annealed samples, meaning that this orientation is less favorable to have a strong polarization under electric field. The thickness of the samples and its heterogeneity must also be discussed, since it has a strong impact on the piezoresponse of the samples. While the annealed ones have a very low
roughness and a thin morphology, the MC ones are characterized by a heterogeneous thickness that is given by the formation of large edge-on stacks. Since the maximum applied voltage is the same for the A and MC samples, due to technical constraints, the thicker edge-on lamellar stacks of the melt/crystallized samples are indeed subjected to a lower electric field. Within a melt/crystallized sample, the edge-on stacks are again subjected to lower electric fields than the rest of the sample. Nonetheless, the highest response is measured for the edge-on lamellae, meaning that their favorable orientation counteracts and wins the eventual effect of the thickness. The large amount of petal-like semi-crystalline domains for Terpo65-0-MC100 and Terpo66-4.7-MC10 (Figures 7a' and SI 6) is responsible for a very heterogeneous response to the external electric field, since their orientation is less favorable for the dipolar switching. It should be pointed out that the effect of the thermal treatment on the electric responsiveness at the nanoscale is much less important for Terpo67-8.3: after annealing (Figure 2c), a weak phase contrast exists (Δφ ~ 10°), while after melting/crystallization, no contrast can be detected (Figure 7b').

Thus, the processing conditions have a strong effect on the crystalline morphology independently from the chemical composition, but the electric response of the induced morphologies deeply depends on the nature and amount of the crystalline phases. Terpo65-0 exhibits strong FE behavior, while Terpo67-8.3 exhibits RFE behavior: the edge-on lamellae stacks of the former lead to a pronounced dipolar switching for both thermal treatments (Figures 2a, 6, and 7a'), while the edge-on lamellae stacks of the latter cannot enhance the responsiveness to the external electric field (Figures 2c and 7b').

At the macroscopic scale, the D-E loops for the annealed and melt/crystallized samples (Figures 4 and 8, respectively) do not show significantly different behavior. Terpo65-0-MC has a more pronounced FE loop, with higher polarization values than those of Terpo65-0-A, while the evolution of the loops for the two terpolymers consists in a slight thinning of the hysteresis cycle and in the appearance of double loop hysteresis (DLH) for both Terpo66-4.7-MC and Terpo67-8.3-MC. Yet, their polarization and coercive field values do not undergo a pronounced evolution. Thus, the effect of the processing conditions is much more important at the local scale since the orientation of the lamellar stacks plays a key role, whereas at the macroscopic scale, these specific orientations disappear as they are averaged across the thickness of the sample.¹²

## 5. CONCLUSIONS

This article examines the impact of chemical composition and processing conditions on the electric responsiveness of P(VDF-ter-TrFE-ter-CTFE) terpolymers. The analysis was conducted at both nanometric and macroscopic scales, using PFM for thin spin-coated films and D-E loops for 20 µm thick films. Three chemical compositions were tested with CTFE contents of 0, 4.7, and 8.3 mol %. Two thermal treatments were studied, namely annealing in the paraelectric phase and melting followed by crystallization at a chosen cooling rate.

For thick films, the electrical response properties are influenced by the composition of the film, showing an evolution of behavior from ferroelectric toward relaxor ferroelectric with an increase in the CTFE content. This evolution is relatively unaffected by the thermal processing

despite a noticeable difference in the crystal morphology and probably due to the relatively large thickness of the films. Within the film thickness, the two types of morphology are distributed isotropically giving a similar electrical response, due to the isotropic distribution of the dipoles in the material.

For thin films, in addition to the composition of CTFE, the size and orientation of the semi-crystalline domains play a very important role in the orientation of the dipoles and, therefore, in the behavior under electric field. When the samples are annealed, the lamellar stacks are small, resulting in a homogeneous film thickness. The electrical response of these films is thus homogeneous at the nanometer scale and corresponds to the CTFE composition. However, when the stacks are much larger, the samples exhibit different electrical field responses related to the orientation of the dipoles within the crystalline lamellae: edge-on-oriented lamellae are much more responsive than flat-on-oriented lamellae. For the sample with 4.7 mol % CTFE, this melt/crystallized morphology appears to highlight a mixture of ferroelectric/relaxor ferroelectric behavior.

For sensor and actuator applications, controlling the growth and orientation of these lamellae stacks in macroscopic films would further improve their electric responsiveness.

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

## ASSOCIATED CONTENT

### *Supporting Information

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsapm.3c00534.

Additional information and scheme of the first PFM protocol, detailed procedure of the second PFM protocol, AFM topography images of Terpo66-4.7-A10 and Terpo67-8.3-A10, SEM images on the surface of Terpo65-0-A, Terpo65-0-MC, Terpo67-8.3-A, and Terpo67-8.3-MC thick films, and PFM topography, phase images, and phase measurements along two sections of Terpo66-4.7-MC10 (PDF)

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

## AUTHOR INFORMATION

### Corresponding Authors

Sylvie Tencé-Girault – Laboratoire PIMM, Arts et Métiers
Institute of Technology, CNRS, Cnam, HESAM Université,
75013 Paris, France; Arkema, CERDATO, 27470 Serquigny,
France; ● orcid.org/0000-0001-5090-2334;
Email: sylvie.girault@ensam.eu

Sébastien Roland – Laboratoire PIMM, Arts et Métiers
Institute of Technology, CNRS, Cnam, HESAM Université,
75013 Paris, France; ● orcid.org/0000-0001-5104-2698;
Email: sebastien.roland@ensam.eu

### Authors

Sara Zanchi – Laboratoire PIMM, Arts et Métiers Institute of
Technology, CNRS, Cnam, HESAM Université, 75013 Paris,
France

Antoine Pascaud – Laboratoire PIMM, Arts et Métiers
Institute of Technology, CNRS, Cnam, HESAM Université,
75013 Paris, France

Fabrice Domingues Dos Santos – Piezotech S.A.S., Arkema—
CRRA, 69493 Pierre-Bénite Cedex, France
Notes

The authors declare no competing financial interest.

■ ACKNOWLEDGMENTS

This research was carried out within the framework of the
Industrial Chair Arkema (Arkema/CNRS-ENSAM-Cnam,
Arkema N° AC-2018-413, CNRS N° 183697). This work
was partially funded by the French National Research Agency
(ANR FETA project, ANR-18-CE06-0024). The authors
gratefully acknowledge the Institut Carnot ARTS for
participating in the funding of the PeakForce QNM upgrade
of the Multimode AFM microscope used in this study. Sylvie
Lebreton, from Arkema CERDATO, is warmly thanked for the
beautiful SEM images, which appear in SI and support this
study. We gratefully acknowledge Estelle Nguyen for the
contribution of cover artwork (https://shorturl.at/betFW).
Finally, Ilias Iliopoulos is thanked for the enlightening
discussions as well as for reading and clarifying this manuscript.

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