# Coupling various in-situ methods to reveal the peculiar semi-crystalline morphology of melt-processed fluorinated copolymers

Sara Zanchi<sup>a</sup>, Régis Cipriani<sup>b</sup>, Fabrice Domingues Dos Santos<sup>c</sup>, Sylvie Tencé-Girault<sup>a, b</sup>, Sébastien Roland<sup>a, *</sup>

<sup>a</sup> Laboratoire PIMM, Arts et Metiers Institute of Technology, CNRS, Cnam, HESAM Université, 151 Boulevard de l'Hopital, 75013, Paris, France

<sup>b</sup> Arkema, CERDATO, Route du Rilsan, 27470, Serquigny, France

<sup>c</sup> Piezotech S.A.S., Arkema - CRRA, rue Henri Moissan, 69493, Pierre-Bénite Cedex, France

## ARTICLE INFO

Keywords:

Morphology

VDF-based terpolymer

AFM/SAXS-WAXS

## ABSTRACT

Poly(vinylidene fluoride-co-trifluoroethylene) (P(VDF-co-TrFE)) copolymer and poly(vinylidene fluoride-ter-trifluoroethylene-ter-chlorotrifluoroethylene) P(VDF-ter-TrFE-ter-CTFE) terpolymer thin films are usually produced by solvent-casting followed by annealing to enhance their electroactive properties in electronic devices. In this study, we investigate the effect of a different thermal treatment, namely crystallization from the melt, on the crystalline structure and morphology of P(VDF-ter-TrFE-ter-CTFE) terpolymers with varying CTFE content (ranging from 0 to 8.3 mol %) using a combination of different in-situ techniques. First, in-situ Wide-Angle X-Ray Scattering (WAXS) shows the different crystal-crystal transitions during cooling for all the samples. For the terpolymers, two sharp Bragg peaks ((110) and (200) family of planes) of the low temperature crystalline phase can be detected. Their sharpness reflects a high lateral extension of the crystalline domains. Moreover, coupled morphological studies by in-situ Small-Angle X-Ray Scattering (SAXS) and Atomic Force Microscopy (AFM) show that the crystallization starts with the formation of very extended semi-crystalline domains with a sub-micrometric thickness that contributes to the SAXS signal. For significant amounts of CTFE, crystallization is spread out, leading to the appearance of increasingly finer lamellar stacks during cooling, filling the spaces left unoccupied during crystallization at higher temperatures. This contrasts with the very rapid crystallization observed in the polymer with 0 mol % of CTFE. Finally, the Curie transition is clearly visible in the AFM mechanical images, providing an original method for detecting the structural transitions occurring in these polymers.

## 1. Introduction

Belonging to the fluorinated electroactive polymers, poly(vinylidene fluoride-co-trifluoroethylene) (P(VDF-co-TrFE) copolymers and poly(vinylidene fluoride-ter-trifluoroethylene-ter-chlorotrifluoroethylene) P (VDF-ter-TrFE-ter-CTFE) terpolymers thin films have been extensively studied and have found many applications in the field of flexible organic electronics in the last twenty years [1–6]. The electroactive behavior of these semi-crystalline polymers is mainly due to the strong electric dipoles originated by the C–F bonds. These dipoles are oriented perpendicularly to the polymer chain in the crystalline domains and can rotate more or less cooperatively, depending on the crystalline phase. The chemical composition (monomers ratios) and the thermal treatment are two key parameters in order to set both the nature of the crystalline

phases and the morphology, and thus the electroactive behavior of these materials.

This category of polymers belongs to the family of semi-crystalline polymers, which are characterized by the presence of an amorphous and at least one crystalline phase. When a polymer is able to crystallize into more than one crystalline structure, these phases are called polymorphs. These different crystalline phases can be distinguished by different chain conformations or/and the nature of inter-chain interactions. Thus, polymorphism can be revealed starting from different crystallization conditions (temperature, pressure, solvent). Several examples of semi-crystalline polymers showing polymorphism associated with specific and various macroscopic properties are described in the literature. The crystalline phases of polyamide 11 (PA11), driven by hydrogen bonding between chains, and their complex behavior with

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: sebastien.roland@ensam.eu (S. Roland).</footnote>
respect to the temperature have been widely studied [7–11]. For example, the α and α' triclinic phases of PA11 are obtained by solvent casting and after crystallization from the melt, respectively; only the α' undergoes the crystal-crystal Brill transition [8]. Poly(ether ketone ke-tone) (PEKK) is another example of semi-crystalline polymer which exhibits a polymorphism, which arises from different factors such as the chemical structure, specifically the amount of *para* and *meta* phenyl links, and the crystallization conditions [12–15]. Finally, the rich polymorphism of poly(vinylidene difluoride) (PVDF), linked to various *trans* (T) and *gauche* (G) conformations sequences of the CH₂-CF₂ units, must be mentioned [16].

In PVDF, the α phase (TG-TG conformations) is the most thermodynamically stable. Nevertheless, after the crystallization, it is possible to induce phase transitions by applying different stimuli (thermal, mechanical, electrical, etc.). For example, the polar β phase (all-*trans*, TTTT), characterized by an interesting electroactive behavior, is commonly obtained by stretching a film of non-polar α-PVDF [16]. A similar polar phase was observed for both P(VDF-*co*-TrFE) copolymers and P(VDF-*ter*-TrFE-*ter*-CTFE) terpolymers, after solvent-casting without the need for additional stretching. Tashiro et al. [17] and Bargain et al. [18] published the phase diagrams describing the nature of crystalline phases of these polymers as a function of temperature and chemical composition for P(VDF-*co*-TrFE) copolymers and P(VDF-*ter*-TrFE-*ter*-CTFE) terpolymers, respectively. These two articles are considered as benchmarks for understanding the correlation between chemical composition, crystalline structure, and the resulting electroactive behavior of these materials. Among the key results of their articles, it should be highlighted that the VDF and TrFE units are isomorphs and co-crystallize in the range 50–80 mol % of VDF [17]. The CTFE units are included in the crystalline phase, up to 10–12 mol % [18]. However, most of the existing literature does not focus on the study of thermal treatment above the melting point, but use the same well-established thermal annealing protocol above the Curie transition. This protocol, applied after the evaporation of the solvent, is known to be sufficient to roughly double the polymer crystallinity and to enhance the ferroelectric performances of these polymers [1, 6, 19–21].

Although the standard process to print organic electronic devices is solvent-casting, recently the melting of a polymer film followed by its crystallization was investigated. This crystallization method has been widely used by the research group of Lei Zhu for structural studies of P (VDF-*ter*-TrFE-*ter*-CTFE) terpolymers [22,23]. But, in their studies, a crystal orientation was sought by applying a multistep process including quenching from the melt, stretching, and annealing. This process complicates the understanding of the crystalline phases obtained after crystallization from the molten state.

In a recent publication [24], we observed that controlled cooling from the melt at different rates, without further treatment, resulted into a significantly different morphology with respect to the one obtained after annealing. A subsequent paper highlighted the promising electro-active response of this melt-crystallized (MC) morphology, consisting of very extended semi-crystalline structures with a possible preferential orientation [25]. Although it is known that the MC treatment induces the crystallization of very extended semi-crystalline domains, a key question regarding the possible effect of this crystallization method on the crystalline phases of P(VDF-*co*-TrFE) copolymers and P(VDF-*ter*-TrFE-*ter*-CTFE) terpolymers remains unanswered.

In this work, we combined Wide and Small-Angle X-ray Scattering (WAXS and SAXS) and Atomic Force Microscopy (AFM) to perform an in-situ structural and morphological study on thin films of P(VDF-*co*-TrFE) copolymers and P(VDF-*ter*-TrFE-*ter*-CTFE) terpolymer crystallized from the melt. Different chemical compositions were selected (from 0 to 8.3 mol% of CTFE), in order to assess the effect of this thermal treatment on CTFE-poor and CTFE-rich terpolymers.

## 2. Experimental part

### 2.1. Materials and samples preparation

A P(VDF-*co*-TrFE) copolymer (Terpo65-0) and three P(VDF-*ter*-TrFE-*ter*-CTFE) terpolymers (Terpo66-4.7, Terpo67-6.4, 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 SI-1). The VDF/(TrFE + VDF) ratio was kept constant (around 66 %) for the four 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 bar-coating and spin-coating.

#### 2.1.1. 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. The melt-crystallized treatment was performed on these 20 µm thick bar-coated films, in an oven, prior to subsequent analysis. After an isotherm at T_m + 25 °C for 15 min (T_m is the melting temperature), the oven was turned off and the door was left partially open, in order to slowly cool the polymer films to room temperature.

These films were used to perform DSC, in-situ WAXS and SAXS, and DMA analyses.

#### 2.1.2. Spin-coating

A silicon wafer, 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. These thin samples were prepared in order to perform in-situ AFM analyses.

### 2.2. Differential Scanning Calorimetry (DSC)

DSC measurements were performed using a Q1000 series TA Instruments. Disks of polymer films were cut using a 6 mm cylindrical shape punch and superimposed in a DSC aluminum pan to reach 9–10 mg. Heating and cooling ramps were done at 10 °C/min under nitrogen flow (50 mL/min). After melting each polymer at T_m + 25 °C for 5 min, the thermograms of the cooling ramps down to 0 °C and the following heating ramps up to T_m + 25 °C were collected.

### 2.3. Wide-Angle X-ray Scattering (WAXS)

Wide-Angle X-rays Scattering (WAXS) experiments were performed on the Xenocs Nano-inXider SW system in transmission mode using Cu Kα radiation (λ = 1.54 Å) from an X-ray microsource (GeniX3D) operating at 50 kV-0.6 mA (30 W). Patterns were collected on a detector Pilatus3 (Dectris) allowing a 2θ range between 5° and 62° (q range between 0.35 and 4.2 Å⁻¹), the collimation was chosen at 800 µm to maximize intensity. For the in-situ thermal experiments, the film sample (thickness ca. 20 µm), was wrapped in aluminum foil and placed in a Linkam heating stage (HFSX350) and melted at T_m + 25 °C for 5 min, then submitted to a cooling ramp (rate - 10 °C/min) having plateaus every 10 °C to allow the WAXS data acquisitions. After stabilizing the temperature for 30 s, the WAXS patterns were recorded during 60 s at each defined temperature. The reduction of 2D datasets from the Pilatus3 detectors to 1D was achieved by radially averaging the raw data detector counts using Foxtrot data reduction software [26]. WAXS intensities were corrected for the flat field and for the Ewald sphere solid angle. WAXS spectra, I(2θ), were fitted using the Fityk 0.9.8 software in the full 2θ range. With this software, each WAXS spectrum is decomposed into crystalline and amorphous contributions, in a wide angular
range $2\theta \in [10^{\circ}, 60^{\circ}]$. The crystalline peaks are associated with sharp peaks while amorphous signal is fitted by broad peaks. The $2\theta_{hkl}$, or $q_{hkl}$, of each crystalline peak gives the inter-planar distance $d_{hkl}$ using the Bragg's law (Equation 1):

$$2d_{hkl} \sin \theta_{hkl} = \lambda \qquad (1)$$

The scattering vector $q_{hkl}$ is defined by Equation 2:

$$q_{hkl} = 4\pi \frac{\sin \theta_{hkl}}{\lambda} = \frac{2\pi}{d_{hkl}} \qquad (2)$$

For the fluorinated copolymers considered in this study, the crystalline structure is orthorhombic with cell parameters obeying the rule, a ≠ b ≠ c and α = β = γ = 90°. The inter-planar distance $d_{hkl}$ for an orthorhombic cell is given by Equation 3:

$$d_{hkl} = \frac{1}{\sqrt{\frac{h^2}{a^2} + \frac{k^2}{b^2} + \frac{l^2}{c^2}}} \qquad (3)$$

The main Bragg peaks are associated with the family of planes (110) and (200), with the respective inter-planar distances, $d_{110}$ (Equation 4) and $d_{200}$ (Equation 5):

$$d_{110} = \frac{1}{\sqrt{\frac{1}{a^2} + \frac{1}{b^2}}} = \frac{a}{\sqrt{1 + \left(\frac{a}{b}\right)^2}} \qquad (4)$$

The hexagonal structure is obtained when $d_{110} = d_{200}$, which gives $\frac{a}{b} = \sqrt{3}$.

$$d_{200} = \frac{a}{2} \qquad (5)$$

The weight crystallinity, $\chi_c$, is calculated, for $2\theta \in [10^{\circ}, 60^{\circ}]$, using Equation 6:

$$\chi_c = \frac{A_c}{A_c + A_a} \times 100 \% \qquad (6)$$

in which $A_c$ and $A_a$ are the total integrated intensity of crystalline peaks and of amorphous shape, respectively (see Fig. SI-1).

If a Bragg peak is associated with only one (hkl) family of planes, its corrected full width at half maximum, $\Delta 2\theta_{hkl}^{cor}$, gives the extension of crystalline domains, perpendicular to the (hkl) planes, using the Scherrer formula (Equation 7) [27]:

$$D_{hkl} = \frac{0.9 \times \lambda}{\Delta 2\theta_{hkl}^{cor} \times \cos \theta_{hkl}} \qquad (7)$$

Where $\Delta 2\theta_{hkl}^{cor}$ is the full width corrected for the instrumental resolution [20]. Since we will focus only on (110) and (200) plans, this equation will refer to a lateral extension of the crystalline lamellae.

## 2.4. Small-Angle X-ray Scattering (SAXS)

As reported in the literature [24], these fluorinated terpolymers crystallized from the melt show a SAXS signal at very low q values ($q_{min}$ around 0.003 Å⁻¹). This q value is accessible for the Nano-inXider with a low collimation (200 µm), but the acquisition time required for this collimation is not compatible with in-situ measurements. For this reason, in-situ SAXS thermal experiments were carried out on the SWING beamline at the Soleil synchrotron facility (project n°: 20210035). The beam energy was set at 10 keV, leading to a wavelength of 1.24 Å. SAXS patterns were recorded using an EigerX4M detector (Dectris) at 4.29 m from the sample, leading to a q range between 0.002 Å⁻¹ and 0.24 Å⁻¹. SAXS patterns were recorded every 30 s, during 200 ms, along a cooling ramp at - 10 °C/min (Linkam heating stage THMS 600) from T_m + 25 °C (after an isotherm of 5 min) and down to room temperature.

The reduction of 2D datasets from the EigerX4M detector to 1D was

achieved by radially averaging the raw data detector counts using Foxtrot data reduction software [26], then normalized to the number of transmitted photons. After subtraction of blank (spectra acquired with the same environment, but without sample), the SAXS intensity profiles $I(q)$ and $q^2I(q)$ (Lorentz-corrected intensity) were then plotted and analyzed.

## 2.5. Atomic Force Microscopy (AFM)

A Multimode 8 (Bruker) microscope equipped with a JHV scanner (High Temperature up to 250 °C) controlled by a Nanoscope V controller and operating in the PeakForceQNM mode was used. The temperature was controlled by a Thermal Applications Controller (TAC) and cooled by circulating water. High temperature in-situ AFM analyses were performed by using ScanAsyst-Air-HR cantilevers with a nominal frequency, spring constant, and tip radius of 70 kHz, 0.4 N/m, and 2 nm, respectively. The spin-coated samples were heated from room temperature up to T_m + 25 °C, then this temperature was kept for 5 min before the cooling step. The sample was cooled down to room temperature with plateaus (ca. 5 min) at intermediate temperatures to allow the in-situ AFM imaging. All the thermal ramps were manually kept at 10 °C/min.

## 2.6. Dynamical Mechanical Analysis (DMA)

DMA analysis was performed on two melted-crystallized samples, the Terpo65-0 and the Terpo67-8.3. Samples for DMA experiments were cut from these films by using a rectangular shape punch. DMA tests were performed using a Q800 TA Instruments Dynamic Mechanical Analyzer. Storage modulus, E', and loss factor, tan δ = E''/E', were measured in tension mode following a heating ramp at 2 °C/min from - 80 °C to a temperature close to the melting temperature (120 °C and 95 °C, for Terpo65-0 and the Terpo67-8.3, respectively). Then, a cooling ramp was applied to analyze the reversibility of the behavior. A sinusoidal tensile strain, of approximately 0.15%, was applied at a frequency of 1 Hz.

# 3. Results

The aim of this article is to study the crystalline structure of thin films made of P(VDF-ter-TrFE-ter-CTFE) terpolymers after their crystallization from the melt. Four different samples are included in this study, having an increasing content of CTFE (Terpo65-0, Terpo66-4.7, Terpo67-6.4, and Terpo67-8.3).

A first overview of the crystallization state obtained from the molten state can be done through Differential Scanning Calorimetry (DSC). Fig. 1a-b shows the DSC thermograms during the first cooling from the melt and the successive heating ramp, respectively. The characteristic temperatures (crystallization, Curie transition, and melting) are reported in Table 1. As a reminder, at the Curie temperature ferroelectric polymers undergo a thermally-driven phase transition towards a paraelectric state. This transition is reversible at cooling, with a more or less marked thermal hysteresis. For these polymers, this transition is a crystal-crystal transition from an orthorhombic to a hexagonal structure. During the cooling ramp (Fig. 1a), three intense exothermic peaks are detected for Terpo65-0 (black curve), the crystallization peak at 133 °C and two peaks at 62 °C and 58 °C, which are associated with the Curie transition. For Terpo66-4.7 (red curve), the crystallization peak is located at 115.5 °C and there is a unique Curie peak at 37 °C. The enthalpy of both peaks is lower than that for Terpo65-0 ones, but the Curie peak is still clearly visible and characteristic of a first order transition. The thermograms at cooling for Terpo67-6.4 and Terpo67-8.3 (blue and green curves, respectively) are very similar. The crystallization takes place at 106 °C for Terpo67-6.4 and at 103 °C for Terpo67-8.3, while their crystal-crystal low temperature transition (pseudo-Curie) takes place at ~25 °C and ~22 °C, respectively. It is
<figure><img src="image_1.png" /><figcaption>**Fig. 1.** DSC thermograms for Terpo65-0, Terpo66-4.7, Terpo67-6.4, and Terpo67-8.3. a) First cooling from the molten state down to 0 °C and b) second heating from 0 °C up to the molten state.</figcaption></figure>

<caption>**Table 1**</caption>

Characteristic temperatures extracted from DSC thermograms (cooling: T<sub>crystallization</sub> and T<sub>Curie</sub>, heating: T<sub>Curie</sub> and T<sub>melting</sub>) for Terpo65-0, Terpo66-4.7, Terpo67-6.4, and Terpo67-8.3. The * symbol refers to the pseudo-Curie transition, which occurs for CTFE-rich terpolymers.

<table><thead><tr><th rowspan="2"></th><th colspan="3">First cooling</th><th colspan="3">Second heating</th></tr><tr><th>T<sub>crystallization</sub> (°C)</th><th>T<sub>Curie</sub> (°C)</th><th>T<sub>melting</sub> (°C)</th><th>T<sub>Curie</sub> (°C)</th><th>T<sub>melting</sub> (°C)</th><th>T<sub>melting</sub> (°C)</th></tr></thead><tbody><tr><td>Terpo65-0</td><td>133.0</td><td>62.0</td><td>58.0</td><td>70.5</td><td>89.0</td><td>154.0</td></tr><tr><td>Terpo66-4.7</td><td>115.5</td><td>37.0</td><td></td><td>40.0</td><td></td><td>140.0</td></tr><tr><td>Terpo67-6.4</td><td>106.0</td><td>~25*</td><td></td><td>~25*</td><td></td><td>130.0</td></tr><tr><td>Terpo67-8.3</td><td>102.0</td><td>~22*</td><td></td><td>~22*</td><td></td><td>129.5</td></tr></tbody></table>

worth noting that for Terpo66-4.7, Terpo67-6.4, and Terpo67-8.3 (highlighted by a green arrow in Fig. 1a) there is an extended exothermic signal with low intensity, located in the 75-90 °C range. The following heating ramp, as depicted in Fig. 1b, shows the reversibility of the Curie or the pseudo-Curie transition for all the samples. While the temperature values for Terpo65-0 show a strong hysteresis in temperature (T<sub>C</sub> are 89 °C and 70.5 °C), those for Terpo66-4.7 (40 °C), Terpo67-6.4 (~25 °C), and Terpo67-8.3 (~22 °C) exhibit a very low hysteresis. The thermograms indicate that the higher the CTFE content, the lower the melting and Curie temperatures.

To understand the temperature-dependent evolution of the

structures and the morphologies of P(VDF-co-TrFE) copolymer and P(VDF-ter-TrFE-ter-CTFE) terpolymers, different techniques, including DSC, SAXS-WAXS, AFM, and DMA were used. These experiments were performed on samples with varying thicknesses, around 20 µm for DSC, SAXS-WAXS, and DMA and 100 nm for AFM. Nonetheless, while the temperature kinetics observed by the samples and the temperatures measured on our different instruments exhibit slight variations, it is through the interpretation and parallel analysis of these temperature-dependent evolutions (DSC, SAXS-WAXS, AFM) that a comprehensive understanding of this crystallization process can be established. This approach mirrors our previous work, which involved coupling SAXS with observations from various microscopies [24].

### 3.1. *In-situ* WAXS experiments: crystalline structure

WAXS analyses during the cooling ramp (10 °C/min) from the molten state (5 min at 200 °C, in order to erase the thermal history) were performed. This was done to ascertain the nature of the crystal-crystal transition that occurs in each sample. Fig. 2a-d shows the WAXS spectra collected with 10 °C steps, starting from the crystallization from the melt and down to room temperature (or 0 °C for Terpo66-4.7 and Terpo67-6.4). The 2θ interval (12°-23°) used for Fig. 2 is narrower than the range of the acquired data, in order to focus on the main signal of the spectra. For all the samples, these WAXS spectra were fitted with large amorphous peaks and sharp crystalline peaks, as illustrated more comprehensively in Fig. SI1. The fits led to quantitative information: the inter-planar distances were deduced from the position of the crystalline peaks and the crystallinity was deducted from the integrated areas of the amorphous and crystalline peaks (see Experimental Section and Fig. SI-1) [19,20].

Terpo65-0 (Fig. 2a) crystallizes in the paraelectric (PE) phase at 130 °C. In this high temperature hexagonal crystalline phase, the polymer chains are characterized by a conformational disorder and rotate violently around the chain axis [28]. The inter-planar distance of the PE phase (d<sub>PE</sub>) decreases linearly upon cooling, as shown in Fig. 2a'. At 70 °C, the onset of the Curie transition is observed on the WAXS spectra. At this temperature, an additional small crystalline peak (blue asterisk in Fig. 2a') appears, associated with the ferroelectric phase (FE). A significant modification of the main diffraction peak is also observed and distinguishing the contributions of the paraelectric phase and the defective ferroelectric phase (DFE) becomes challenging (see the black and red lines in Fig. 2a'). Starting from 60 °C, two distinguished crystalline peaks are observed, one much weaker and wider than the other. Based on previous studies on as-cast and annealed P(VDF-co-TrFE) copolymers [19], we associate these peaks to the DFE and FE phases, respectively. Both DFE and FE phases have an orthorhombic crystalline lattice, characterized by all-trans conformations for the FE phase and by the presence of gauche conformational defects for the DFE one. It is worth highlighting that the transition from the PE to the DFE phase is steeper during the cooling from the melt than during the cooling from the annealing temperature [19]. This is consistent with the width of the Curie peak, which is significantly narrower after crystallization from the melt (Fig. 1a) [19].

For Terpo66-4.7, the crystallization in the PE phase from the melt initiates at around 110 °C, as shown in Fig. 2b-b'. The d<sub>PE</sub> decreases until 50 °C. Below this temperature, two peaks are needed to fit the crystalline signal of this WAXS spectrum. These two peaks are narrow; they have the same width and very similar intensities. As the temperature decreases, these two peaks gradually spread apart. Fig. 2c-c' shows a similar behavior for Terpo67-6.4. The crystallization of the PE phase takes place at 100-110 °C, its inter-planar distance decreases monotonously and, below 50 °C, two peaks are detected. Similarly to the Terpo66-4.7, they continue to progressively separate down to 0 °C. We associate these peaks with the (200) and (110) families of crystalline planes of the relaxor ferroelectric phase (RFE), already observed on annealed solvent-cast films [18]. The difference with solvent-cast
<figure><img src="image_2.png" /><figcaption>**Fig. 2.** WAXS spectra and inter-planar distances evolutions during the cooling ramp from the melt for a), a') Terpo65-0, b), b') Terpo66-4.7, c), c') Terpo67-6.4, and d), d') Terpo67-8.3. In a'), the red and black lines show the impossibility to distinguish the contributions of the PE and DFE phases at the beginning of the Curie transition. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

samples is that, after crystallization from the melt, the diffraction peaks are much narrower and can therefore be separated, whereas they were not after solvent evaporation and annealing. This RFE crystalline phase is similar to the DFE one, but it has larger inter-planar distances and, from a practical standpoint, it is not polarizable under an external electric field. At 0 °C, the values for d₂₀₀ and d₁₁₀ are 4.8 Å (or 4.85 Å) and 4.65 Å (or 4.7 Å), for Terpo66-4.7 (or Terpo67-6.4). Finally, the WAXS spectra for Terpo67-8.3 have also an analogous behavior during the cooling ramp, as shown in Fig. 2d-d', the PE crystallizes at 100 °C and the (200) and (110) peaks of the RFE phase start to be distinct below 50 °C. At room temperature, their inter-planar distances are 4.9 Å and 4.8 Å, respectively.

It is worth noting that during the cooling from the melt, the four samples crystallize at their respective crystallization temperatures, in the paraelectric phase. For this PE phase, at a specific temperature, such as 100 °C, the inter-planar distance increases with the amount of CTFE, from 4.92 Å for Terpo65-0, to 4.96 Å for Terpo66-4.7, and to 4.99 Å for Terpo67-6.4 and Terpo67-8.3. This trend shows the incorporation of the bulkier CTFE unit in the PE crystalline cell. Below the Curie transition, two crystalline phases exist for Terpo65-0, the major FE phase and the minor DFE phase, as already reported for annealed solvent-cast samples [19]. On the other hand, below the pseudo-Curie transition, the two CTFE rich terpolymers appear to be characterized by a single crystalline phase, the RFE phase, whose inter-planar distances d₂₀₀ and d₁₁₀ spread apart and continuously with temperature. Terpo66-4.7 (Fig. 2b-b') evolves in a slightly different way, as it will be shown later. It is probably crystallized with a minor fraction of FE phase, which evolves discontinuously, and with a significant fraction of RFE phase, characterized by a continuous evolution. This contributes to the deceptively steeper evolution of d₁₁₀.

The crystallinity can be obtained from the fit of each WAXS spectrum (see Experimental Section and Fig. SI-1 and Fig. SI-2). Fig. SI2 shows that

for Terpo65-0 (in black), the initial stage of cooling maintains a constant crystallinity of ca. 15% at elevated temperature. Below 70 °C, it increases steeply and achieves 25% at room temperature. This increase is related to the Curie transition, where the sample regains the intra-chain order. For the three terpolymers, the crystallinity seems to have constant values, around 15% for Terpo66-4.7 and Terpo67-6.4, and around 12% for Terpo67-8.3. This behavior is related to the pseudo-Curie transition, which does not involve the restoring of well-defined intra-chain order for the three terpolymers.

### 3.1.1. Evolution of the crystal lattice parameters

More information about the crystalline lattice of these polymers was deduced from the inter-planar distances d₁₁₀ and d₂₀₀. The expression of d₁₁₀ and d₂₀₀, reported in Section 2.3 leads to the *a* and *b* parameters of the crystalline lattice for each sample. The dependence of these parameters with the amount of CTFE was studied at room temperature and at 100 °C, where all the samples are in the paraelectric phase (Fig. 3). For Terpo65-0, as also shown in Fig. SI-1a, two crystalline phases coexist at room temperature (FE and DFE). The respective (110) and (200) peaks are inseparable, preventing the calculation of lattice parameters. For this composition, the two values reported in Fig. 3a (empty symbols) are taken from the literature [29]. For the terpolymers, the (110) and (200) peaks can be distinguished at room temperature, as highlighted in Fig. SI-1b-d, and the lattice parameters can be calculated. For Terpo66-4.7, both *a* and *b* values are higher than those of a P (VDF-co-TrFE) copolymer. The lattice parameters are even larger for Terpo67-6.4, reaching a plateau for Terpo67-8.3. The increase of *a* and *b* with the amount of CTFE indicates that the bulkier CTFE unit is incorporated into the orthorhombic RFE crystal cell. The *a/b* ratio gives information about the symmetry of the orthorhombic crystalline lattice. As shown in Fig. 3a, at room temperature, the increase of CTFE leads to a symmetry progressively converging toward the hexagonal crystalline
<figure><img src="image_3.png" /><figcaption>Fig. 3. a) Evolution of the a and b lattice parameters (● and ▼ symbols, respectively) and of the a/b ratio (× symbol) with respect to the CTFE content. Empty symbols refer to values taken from the literature. The lattice parameters of a sample containing 10 mol % of CTFE are included for a comprehensive trend representation. b) Evolution of the a/b ratio (× symbol) during a cooling ramp from the molten state, for Terpo66-4.7, Terpo67-6.4 and Terpo67-8.3.</figcaption></figure>

lattice, where $a = \sqrt{3}b$. At 100 °C, all samples are in the paraelectric phase (Fig. SI-1a'-d'), which is characterized by a hexagonal crystalline lattice, where $d_{200} = d_{110}$ and $a = \sqrt{3}b$. The evolutions of a and b with the composition are less pronounced at 100 °C, compared to room temperature. They show a slight increase between 0 and 6 mol % of CTFE, followed by a plateau.

Fig. 3b shows the evolution of a/b for Terpo66-4.7, Terpo67-6.4, and Terpo67-8.3 with respect to the temperature, during the cooling from the melt. At the beginning, all the polymers crystallize in the hexagonal paraelectric phase, thus the a/b ratio is $\sqrt{3}$. When the Curie transition starts, a/b increases. The increase is steeper for Terpo66-4.7, which achieves a slightly higher a/b value at 0 °C. All the samples are characterized by an orthorhombic crystalline lattice below the Curie transition.

### 3.1.2. Lateral extension of the crystalline domains

The WAXS spectra contain also information about the lateral extension of the crystalline domains (so perpendicularly to the polymer chains), which can be calculated with the Scherrer equation, when the widths $\Delta2\theta_{110}$ and $\Delta2\theta_{200}$ and the positions $\theta_{110}$ and $\theta_{200}$ of the (110) and (200) peaks are known. At 100 °C, all the samples are in the paraelectric phase and the extension of their crystalline domains decreases with the CTFE content. It is ca. 36 nm for Terpo65-0, 25 nm for Terpo66-4.7, 22 nm for Terpo67-6.4, and 21 nm for Terpo67-8.3. At room temperature, the (110) and (200) peaks are discernible only for Terpo66-4.7, Terpo67-6.4, and Terpo67-8.3. Their width is very similar, giving extensions close to 15 nm for the three samples. The decrease in crystalline domain extension at low temperatures is linked to the crystal-crystal transition and the disappearance of the dynamic conformational disorder along the chains, which exists only in the PE phase [28]. It is worth to compare these values, obtained during the crystallization from the molten state, with those obtained after an annealing treatment on solvent-cast films, in the paraelectric phase. Bargain et al. [20] calculated the extension of the crystalline domains only in the PE phase (at 95 °C) for samples containing 0, 4.4, and 9.7 mol % of CTFE and found 24 nm, 17 nm, and 15 nm, respectively. These values are lower than those obtained for similar compositions after crystallization from the melt, meaning that the crystallization from the melt results in higher extensions of the crystalline domains. For solvent-cast samples at room temperature, this calculation could not be performed by Bargain et al. because the (110) and (200) peaks are not thin enough to be separated in the low-temperature orthorhombic phase.

Based on the WAXS analysis, it can be concluded that, regardless of the crystallization process (either annealing after solvent-casting or crystallization from the melt), a combination of FE and DFE phases

crystallizes for the copolymer, whereas the RFE phase crystallizes for the terpolymers. The CTFE unit is included in the different crystals (RFE and PE). The main difference between these two crystallization methods lies in the crystal extension, which is notably larger after crystallization from the melt. This is consistent with previous AFM studies [24].

### 3.2. In-situ SAXS and AFM experiments: morphologies from the melt

In order to understand the evolution of the morphology during the cooling from the melt, in-situ SAXS analyses were performed on these samples at Soleil Synchrotron (see Experimental Section for more details). Each sample was kept at $T_m + 25$ °C (175 °C, 160 °C, 155 °C, and 150 °C for Terpo65-0, Terpo66-4.7, Terpo67-6.4, and Terpo67-8.3, respectively) for 5 min and, then, the SAXS data were collected during the cooling at 10 °C/min down to room temperature. The results are shown in Fig. 4.

Fig. 4 shows two sets of data for Terpo65-0, Terpo66-4.7, Terpo67-6.4, and Terpo-8.3. First, it presents a 3D representation of the $q^2I(q)$ SAXS spectra recorded during the cooling process from the melt. At the same time, AFM images acquired while cooling from the melt are shown and reveal the presence of sharp and long objects. In SAXS analyses, it is noteworthy that no scattering is observed in the molten state as the intensity is very weak. During cooling, a signal emerges at temperatures that vary depending on the compositions. These temperatures are illustrated on the 3D representations of the SAXS spectra: 130 °C for Terpo65-0, 115 °C for Terpo66-4.7, 105 °C for Terpo67-6.4, and 100 °C for Terpo67-8.3. They are in good agreement with the crystallization temperature of the different compositions, deduced from DSC analyses and reported in Table 1.

In parallel, in-situ temperature AFM analyses were carried out to complement the SAXS experiment and to elucidate the crystallization process during cooling from the melt. For obvious experimental reasons, the cooling rates were slightly different. Nonetheless, AFM images (mechanical contrast DMT Modulus obtained in PeakForce QNM mode) recorded at temperatures close enough to the crystallization temperatures unveil the appearance of long rectilinear objects, as already reported in previous studies [24]. Notably, the temperatures at which these objects appear follow a similar evolution of the crystallization temperature with the CTFE content. At the crystallization temperature, initial large objects appear, exhibiting flat-on and edge-on orientations. For Terpo65-0 and Terpo66-4.7, both orientations manifest simultaneously at the chosen first isothermal temperature. In the case of the other two samples, it seems that only the edge-on orientations appear and coexist with the molten state. These objects induce scattering in SAXS, with intensity decreasing as the CTFE content increases (please refer to the intensity scales in Fig. 4). The observed SAXS scattering
<figure><img src="image_4.png" /><figcaption>Fig. 4. 3D representation of SAXS spectra (Lorentz corrected, $q^2I(q)$) acquired during the crystallization from the melt and AFM images (mechanical contrast DMT Modulus in Peak Force QNM mode) of selected regions at the onset of crystallization during cooling, for a) Terpo65-0, b) Terpo66-4.7, c) Terpo67-6.4, and d) Terpo67-8.3. Each AFM image is labeled with the corresponding temperature at the time of capture. In the AFM insets, the long and bright objects correspond to edge-on lamellar stacks.</figcaption></figure>

shows a peak at low $q$, around $0.01 \AA^{-1}$ for low CTFE contents, which shifts towards $0.02 \AA^{-1}$ for higher CTFE contents.

### 3.2.1. Discussion on crystallization studied by in-situ SAXS experiments

Before going further into the correlation between SAXS and AFM data, we will describe a few, albeit sparse, studies in the literature on the investigation of semi-crystalline polymer crystallization using in-situ SAXS.

For example, Song et al. described crystallization from the melt in two stages [30]. First, the growth of spherulites fills the space, then, it is followed by the creation of new crystallites and the growth of existing ones. These two stages may occur simultaneously. The first stage is characterized by a shoulder in the SAXS signal, without a real peak, which would correspond to disordered lamellar structure. Similarly, Tashiro, notably in the Handbook of Polymer Crystallization [31], describes crystallization as the emergence of denser, uncorrelated domains in the melt, characterized by a specific size that he equates to a radius of gyration $R_g$. As crystallization progresses, a correlation length within these domains appears, becoming the lamellar period when crystallization is well established. However, according to these authors, this model does not seem applicable to fluorinated copolymers due to a «too strong parasitic scattering» in their SAXS spectra [20]. Finally, Camara et al. modeled the scattering function of the crystalline lamellae during crystallization. In their model, crystalline lamellae evolve gradually from an uncorrelated state to the final state where they are stacked periodically [32]. These authors also caution against excessive and hasty interpretation of the correlation peak appearing in the $q^2I(q)$ representation [32]. Indeed, the Lorentz representation of the scattered intensity may show a correlation peak even when the original scattered intensity does not show a distinct correlation peak, due to the different power-laws slopes in the lowest and highest $q$-range.

Therefore, we will use both representations, $I(q)$ and $q^2I(q)$, for a correct interpretation, analyzing the beginning of the crystallization in the context of uncorrelated dense domains. These representations are displayed, for selected significant temperatures in Fig. 5 for Terpo65-0 and Terpo67-8.3, and in Fig. SI-3 for the remaining two compositions.

In the melt, the scattered intensity is very low and shows no characteristic size, regardless of the composition. Then, during cooling, a broad bump appears at low $q$ values, corresponding to a length $(2\pi/q)$ of about 77 nm for Terpo65-0 (70 nm for Terpo67-8.3). This signal is more intense for Terpo65-0 than for Terpo67-8.3. It increases in intensity while shifting towards larger $q$ values, with a small shift for Terpo65-0, ca. 20 nm from 77 to 57 nm, and a larger one for Terpo67-8.3, ca. 35 nm from 70 to 35 nm. We associate these broad signals to the sub-micron dimension of the long semi-crystalline domains (i.e. their thickness) observed by AFM, as shown in Fig. 4. These domains appear uncorrelated at the onset of crystallization.

In parallel, extensive analyses by AFM were carried out for Terpo65-0 and Terpo67-8.3 to monitor the evolution of microstructures with further cooling. Isotherms were captured at 10 °C intervals during the acquisition of AFM images, as shown in Fig. 6. Notably, for Terpo65-0, Fig. 6a, it is evident that there is no evolution of the aggregates, which show no perceptible change in shape, size, or number during the cooling process from 120 °C and below.

In contrast, for Terpo67-8.3, Fig. 6b-a distinct pattern emerges, revealing the presence of increasingly finer lamellar aggregates as the temperature decreases, starting notably from 90 °C down to 70 °C. These fine aggregates adapt by filling the melted spaces that were left unoccupied during crystallization at high temperatures. Their number and size remain constant during each isotherm, and crystallization resumes after further cooling at the next temperature level. This phenomenon is consistently observed at all isothermal temperatures down to 70 °C, beyond which no new aggregates appear. The observed shifts in the SAXS signals of Terpo65-0 and Terpo67-8.3 would therefore be associated with the evolution in the size distribution of semi-crystalline objects that appear during crystallization. This evolution shows little change for Terpo65-0 (20 nm), while a significant shift is detected for Terpo67-8.3 (35 nm).

These SAXS and AFM observations of crystallization during the cooling of Terpo65-0 and Terpo67-8.3 can be related to the shape of the crystallization exotherms peaks measured in DSC (Fig. 1). In particular,
<figure><img src="image_5.png" /><figcaption>Fig. 5. SAXS spectra *I*(*q*) in Log-Log representation and *q*²*I*(*q*) for representative temperatures during cooling from the melt for a) Terpo65-0 and b) Terpo67-8.3.</figcaption></figure>

the crystallization peak of Terpo65-0 is very intense and sharp, while the crystallization peak of Terpo67-8.3 is broader and appears to extend to lower temperatures (broad exothermic signal highlighted with a green arrow in Fig. 1a). SAXS and AFM observations for both samples seem to confirm the more or less extensive crystallization occurring during cooling from the melt.

### 3.2.2. *In-situ* SAXS study of the Curie and pseudo-Curie transition

At lower temperatures, below 90 °C for Terpo65-0 and 80 °C for Terpo66-4.7, a new signal becomes clearly visible in SAXS, around $q = 0.03 \AA^{-1}$, identified by an asterisk on the $I(q)$ representation and by a black arrow on the $q^2I(q)$ representation (see Fig. 5a and SI-3a). This new signal appears as a discontinuity during cooling in the 3D representation for these two compositions (Fig. 4a and b). Therefore, we associate it with the Curie transition and the emergence of a strong contrast between the crystalline lamellae and the amorphous phase, when the less dense PE phase disappears in favor of the dense FE phase, as previously discussed in the literature [19,20]. This SAXS signal corresponds to the periodic stacking of crystalline lamellae and the amorphous phase. At the Curie transition, the associated contrast, and thus the intensity, significantly increases with the appearance of the FE phase. For terpolymers with higher CTFE contents, 6.4 and 8.3, an increase in intensity around $q = 0.03 \AA^{-1}$ is observed, but without a well-defined peak, probably partly due to the nearby signal around $q = 0.02 \AA^{-1}$. This intensity increase occurs progressively and does not show a discontinuity in the 3D representations. This evolution corresponds to

the PE to RFE transition, which is a continuous transition [20].

At low temperatures, and in the $q^2I(q)$ representation, SAXS signals are found, with a similar shape to those published by Zanchi et al. [24] This present study confirms that the main signal observed around $q = 0.01 \AA^{-1}$ for Terpo65-0 and around $q = 0.02 \AA^{-1}$ for Terpo67-8.3 is associated with the size and shape of the semi-crystalline domains. Meanwhile, the signal around $q = 0.03 \AA^{-1}$, with its intensity increasing below the PE to FE or RFE transition, is associated with the periodic stacking of crystalline lamellae inside these semi-crystalline domains.

Given the complexity of this part, some key results of the SAXS analyses should be highlighted and summarized. During cooling, two contributions to the SAXS spectra exist. The first appears at the onset of the crystallization and is related to the shape and size of semi-crystalline domains. The second signal is associated with the periodic stacking of the crystal/amorphous phase within the domains. The intensity of the later signal increases at lower temperatures, especially below the crystal-crystal transition.

To conclude, the correlation between in-situ SAXS and in-situ AFM experiments suggests that, from the initial stages of crystallization, semi-crystalline objects with sub-micron thickness emerge and contribute to the scattering observed in the SAXS spectra. This signal evolves slightly as crystallization progresses, concurrent with the growth of new objects. The contrast originating from the crystal/amorphous interface remains relatively low and undiscernible at the early stages of crystallization. However, as the temperature decreases and the Curie transition is reached, the crystal/amorphous contrast increases significantly, leading
<figure><img src="image_6.png" /><figcaption>**Fig. 6.** In-situ AFM images (mechanical contrast DMT Modulus in Peak Force QNM mode) of select regions during cooling after crystallization has started for (a) Terpo65-0 (scale bar = 600 nm) and (b) Terpo67–8.3 (scale bar = 400 nm). Each image is labeled with the corresponding plateau temperature at the time of capture.</figcaption></figure>

to an increase in the intensity of the lamellar signal, which becomes visible. This behavior differs from that of a typical semi-crystalline polymer, which would form spherulites, as they are much larger and do not induce scattering in the explored SAXS $q$ range.

### 3.3. Complementary study of the low temperature mechanical transition by AFM/DMA

Unexpectedly, this crystal-crystal transition from the high-temperature PE phase to the FE or RFE phases is also observed by AFM. A significant evolution of the contrast in the DMT Modulus images appear to occur during further cooling at temperatures close to the Curie transition temperature. Upon a second heating, this transition is reversible and can be associated with the presence or not of a hysteresis

for Terpo65-0 and Terpo67–8.3, respectively. These transitions are indicated by red and blue arrows in Fig. 7b and Fig. 8b. In PeakForce mode, the mechanical properties of the materials can be probed. This contrast variation in the AFM images reflects the mechanical change close to the Curie transition as already evidenced through DMA by Bargain et al. [18,20] DMA analyses on films were therefore conducted to highlight the evolution of the mechanical modulus associated with the Curie transition from the less dense and dynamical disordered PE phase to the denser all-trans FE phase and vice versa, depending on the direction of temperature change. Fig. 7b and Fig. 8b show the results obtained for samples Terpo65-0 and Terpo67–8.3, respectively. As the temperature of Terpo65-0 decreases, the transition associated with the Curie temperature is evidenced by a rather abrupt change in the storage modulus, from a value of 80 MPa to a value of around 400 MPa. The transition temperature is around 55 °C. One of the characteristics of the Curie transition is the hysteresis between heating and cooling, also observed for this sample. The reverse change in mechanical properties occurs at around 90 °C during heating. This phenomenon seems to be detected by AFM microscopy in PeakForce mode. Indeed, the same temperature difference that causes hysteresis in DMA (around 40 °C) is clearly shown in the AFM images (Fig. 7).

On the contrary, the hysteresis is not visible for Terpo67–8.3, as shown in Fig. 8b. For this composition, the transition seems to occur at a temperature between 20 and 50 °C, depending on the technique used. By AFM, this mechanical shift is again reflected in the loss of contrast in DMT Modulus images during cooling, leading to the disappearance of lamellar stacks (between 60 °C and 50 °C)—an outcome likely attributed to the detected change around 20–30 °C in mechanical properties by DMA. Upon raising the temperature beyond the transition point, the lamellar stacks become visible again by AFM. The temperature shift observed between DMA and AFM may be partly due to the different applied frequencies, 1 Hz in DMA and 2000 Hz in AFM. Indeed, the temperature of the pseudo-Curie transition, associated with the maximum of the dielectric constant, increases with frequency for these relaxor ferroelectric polymers [20,33].

In the literature, the Curie transition of P(VDF-co-TrFE) copolymers has already been easily detected by various techniques, such as dielectric spectroscopy, DMA, X-ray diffraction, and DSC. In contrast, the detection of the pseudo-Curie transition of P(VDF-ter-TrFE-ter-CTFE) terpolymers is more challenging. On the one hand, it was first associated to a maximum of the $\epsilon'$ dielectric [20], which shifts with the frequency [33]. On the other hand, this transition does not exhibit a well-defined

<figure><img src="image_7.png" /><figcaption>**Fig. 7.** AFM images (DMT Modulus in PeakForce mode) during (a) cooling followed by heating of a Terpo65-0 thin film; (b) evolution of the storage modulus $E'$ of a Terpo65-0 film measured by DMA (tensile) during heating followed by cooling. The red and blue arrows show the temperatures at which changes in mechanical properties are detected by AFM, while the black arrow represents hysteresis in DMA. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>**Fig. 8.** AFM images (DMTModulus in PeakForce mode) during (a) heating followed by cooling of a Terpo67–8.3 thin film; (c) evolution of the storage modulus E' of a Terpo67–8.3 film measured by DMA (tensile) during heating followed by cooling. The red and blue arrows show the temperatures at which changes in mechanical properties are detected by AFM, while the small black arrow represents the absence of hysteresis in DMA. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

endothermic signal in DSC thermograms, it shows a continuous evolution of SAXS-WAXS signals and a small drop of the storage modulus in DMA curves, as shown in Fig. 8b. In conclusion, highlighting this transition through mechanical AFM finding suggests that it is a particularly sensitive technique for detecting mechanical transition.

## 4. Conclusions

This study aims to understand the relationship between crystalline structure and morphology in terpolymers with varying CTFE content. Here, we explored films crystallized directly from the melt, and compared with previous studies on films crystallized from a solvent followed by annealing. Regardless of the crystallization process used, the dominant crystalline phase is the FE phase, in the absence of CTFE. In contrast, for CTFE contents above 6 mol %, a single crystalline form, identified as the RFE phase, emerges. In both crystallization methods, compositions around 4–6 mol % of CTFE exhibit a coexistence of RFE and FE phases.

The key difference between these two methods lies in the size of the resulting semi-crystalline objects, characterized by small rice-like domains (ca. 30 nm) for annealed solvent-prepared films and elongated rectilinear domains, with a thickness around 100 nm, for films crystallized from the melt. The estimated crystal extension, as deduced from WAXS spectra, aligns with this observed morphological difference, indicating larger crystal extension after melt crystallization.

In addition, a detailed analysis of WAXS spectra consistently reveals the inclusion of CTFE groups into the crystalline phases at any given temperature. The pseudo-Curie transition from the orthorhombic RFE phase to the hexagonal PE phase is clearly highlighted through the evolution of lattice parameters.

The combination of the analysis of the in-situ SAXS and the in-situ AFM experiments help identify two contributions in the scattering SAXS profiles. One appears at the onset of crystallization and is associated with the shape and dimension of semi-crystalline domains, while the other arises from the periodic stacking of the crystal/amorphous phase within these domains. The later signal increases in intensity at lower temperatures, especially below the Curie transition temperature. The chemical composition strongly affects the crystallization kinetics. For terpolymers containing a significant amount of CTFE, the crystallization takes place in a wide temperature range: increasingly finer lamellar stacks are formed during cooling, filling the empty spaces existing after the crystallization at higher temperatures. In contrast, the

crystallization is very fast for the copolymer. The overall crystallization behavior of these VDF-based copolymers and terpolymers contrasts with the typical spherulitic crystallization commonly observed in semi-crystalline polymers.

Finally, the high sensitivity of AFM analysis, in mechanical mode, enabled us to observe the change in mechanical properties coinciding with the Curie transition, already observed only by DMA analyses. To the best of our knowledge, this is the first time such a result is reported in the literature for these semi-crystalline fluorinated polymers.

## CRediT authorship contribution statement

**Sara Zanchi:** Writing – review & editing, Writing – original draft, Methodology, Investigation, Formal analysis, Data curation, Conceptualization. **Régis Cipriani:** Investigation. **Fabrice Domingues Dos Santos:** Writing – review & editing, Resources. **Sylvie Tencé-Girault:** Writing – review & editing, Supervision, Methodology, Investigation, Funding acquisition, Formal analysis, Data curation, Conceptualization. **Sébastien Roland:** Writing – review & editing, Supervision, Methodology, Investigation, Funding acquisition, Formal analysis, Data curation, Conceptualization.

## Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Data availability

Data will be made available on request.

## Acknowledgements

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), authors concerned: all. This work was partially funded by the French National Research Agency (ANR FETA project, ANR-18-CE06-0024), authors concerned: S. Z., F. D. S., S. T.-G. and S. R. We thank the SOLEIL synchrotron facility and the SWING beamline for access to the instrumentation (project n° 20210035). The beamline manager, Javier Perez, is thanked for his advice before and his assistance
during the in-situ experiments. 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. LEM laboratory (Arkema CERDATO, Serquigny, France) and in particular Sylvie Lebreton and François Bargain are warmly thanked for the kind welcome and access to X-ray diffraction facility. Finally, Ilias Iliopoulos and François Bargain are thanked for the enlightening discussions as well as for reading and clarifying this manuscript.

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