# Crystallization behavior of co-polyesters based on hydroxy fatty acids extracted from tomato peel agro-wastes

Kylian Hallavant, Antonella Esposito, Bénédicte Bakan, Denis Lourdin, Allisson Saiter-Fourcin

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Kylian Hallavant, Antonella Esposito, Bénédicte Bakan, Denis Lourdin, Allisson Saiter-Fourcin. Crystallization behavior of co-polyesters based on hydroxy fatty acids extracted from tomato peel agro-wastes. Polymer, 2026, 342, pp.129314. 10.1016/j.polymer.2025.129314. hal-05366298v2

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Submitted on 21 Nov 2025

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# Crystallization behavior of co-polyesters based on hydroxy fatty acids extracted from tomato peel agro-wastes

Kylian Hallavant<sup>a</sup>, Antonella Esposito<sup>a ID</sup>, Bénédicte Bakan<sup>b ID</sup>, Denis Lourdin<sup>b,*</sup>, Allisson Saiter-Fourcin<sup>a,** ID</sup>

## ARTICLE INFO

**Keywords:**

Bio-based materials

Thermal analysis

Melting enthalpy

Microstructure

## ABSTRACT

This work investigates the crystallization behavior of bio-based slightly crosslinked polyester networks synthesized from long-chain hydroxy fatty acids. The crystallization kinetics and the melting behavior were assessed using a combination of modulated-temperature differential scanning calorimetry (MT-DSC) and fast scanning calorimetry (FSC). The characterization of the crystalline phases was performed by X-ray diffraction (XRD) and polarized-light optical (POM) microscopy. A methodology based on calorimetric investigations was used, allowing to estimate the theoretical value of the melting enthalpy $\Delta h_m^0$ of the initial cutin monomers as well as of the derived co-polyester networks. The results obtained by this method allowed to estimate the crystalline content $X_c$. The combination of XRD and calorimetric analyses evidenced the existence of polymorphs with different stabilities over time at room temperature, characterized by a monotropic transition from the metastable crystal phase ($\alpha$-phase) to more stable crystal phases ($\beta$'and $\beta$-phases) depending on the crystallization time accorded to the monomers and the crosslinked co-polyesters. This work delves deeper into the complex crystallization behavior of biobased polyesters, which could eventually allow a better control over the properties of these materials.

## 1. Introduction

A more efficient biomass valorization and the quest for more sustainable materials are two important research focus in both the industrial and research communities. Polyesters have attracted a lot of attention because of their easiness of production and their tunable properties, which makes them appealing for a wide range of applications including high-performance polymers [1,2]. Polyesters can be synthesized using solvent-free and catalyst-free procedures, from synths extracted from the biomass derived from crop production. In this regard, plant cutin is a source of long-chain hydroxy fatty acids (HFAs); these building blocks can be extracted from industrial pomaces and used for the synthesis of bio-based and/or bioinspired polyesters [3,4]. Recently, by tuning the biorefinery process of these HFAs, polyesters highlighting peculiar barrier, mechanical and shape memory properties have been produced [4]. These properties seem to be associated with high crystallinities, and different mechanical behaviors are observed depending

on their crosslinking state. Indeed, the polymerization of cutin monomers (C16 HFAs) produces semi-crystalline weakly crosslinked macromolecular networks, whose thermal properties depend on their crosslinking density [3,5].

Performing a fine structural characterization of these polyesters, as well as understanding the mechanisms by which they are produced, are crucial to develop suitable applications. Indeed, the microstructure of semi-crystalline polymers can be fine-tuned to improve barrier and mechanical properties. Assessing the crystallization ability of a polymer is essential, since it affects macroscopic properties such as stiffness, strength, chemical resistance, thermal stability, morphology and gas permeability [6,7], which has to be considered for a proper evaluation of the processability, applicability, and biodegradability of the material [8]. In particular, high values of $X_c$ result in harder, stiffer and less ductile behavior [9]. To understand the relationship between properties and microstructure, a precise knowledge of the main thermodynamical parameters (specific heat capacity $c_p$, glass transition temperature $T_g$,

<footnote><sup>a</sup> Univ Rouen Normandie, INSA Rouen Normandie, CNRS, Normandie Univ, Groupe de Physique des Matériaux UMR 6634, F-76000, Rouen, France</footnote>

<footnote><sup>b</sup> INRAE, UR 1268, Biopolymers, Interactions & Assemblages (BIA), F-44316, Nantes, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>** Corresponding author.</footnote>

<footnote>E-mail addresses: denis.lourdin@inrae.fr (D. Lourdin), allison.saiter@univ-rouen.fr (A. Saiter-Fourcin).</footnote>
melting enthalpy $\Delta h_m$) and of the crystalline content $X_c$ is then required. The estimation of $X_c$ can be done through calorimetry, using the ratio of the measured melting enthalpy $\Delta h_m$ to the theoretical melting enthalpy $\Delta h_m^0$ of an ideally fully crystallized polymer (which is never achieved due to chain entanglements). For common polymers the value of $\Delta h_m^0$ may be found in handbooks, but for relatively new materials this parameter is not always available. Most of the studies involving such new materials treat crystallinity in a relative way, or estimate $\Delta h_m^0$ based on analogous saturated polymers, eventually extrapolating the values obtained for different crystallinities [10]. Recently, Cebe et al. [11] and Fosse et al. [12] proposed a calorimetric method to estimate $\Delta h_m^0$ relying on Fast Scanning Calorimetry (FSC) for polymers that crystallize rapidly and cannot be fully amorphized using the cooling rates associated with conventional Differential Scanning Calorimetry (DSC). In this work, the crystallization kinetics and the melting behavior of co-polyesters based on hydroxy fatty acids were assessed using a combination of modulated temperature differential scanning calorimetry and fast scanning calorimetry. The characterization of the crystalline phases was performed by X-ray diffraction and polarized-light optical microscopy. Finally, the methodology proposed by Cebe et al. [11] and Fosse et al. [12] was used to estimate the melting enthalpy $\Delta h_m^0$ of the initial cutin monomers as well as of the derived co-polyester networks, thus enabling to assess the crystalline content $X_c$ of these new functional biobased polyesters.

## 2. Experimental section

### 2.1. Materials

Bio-sourced co-polyesters of hydroxy fatty acids were obtained from tomato peel agro-wastes using a solvent-free and catalyst-free polycondensation process, as described by Marc et al. [3]. The industrial tomato pomaces were provided by Conserveries de Bergerac (UNIPROLEDI, Bergerac, France). The crude monomer extract consists of a 9 (10),16-dihydroxyhexadecanoic acid (whose chemical structure is presented in Fig. 1), 5 % dicarboxylic fatty acids, and a minor fraction of co-extracted phenolics. Following the purification process as described by Marc et al. [3], only a low content of phenolic compounds remained in the purified monomers.

Two materials were studied, (1) the purified cutin monomers (referred to "purified CM" in the following) and (2) the polymer network obtained by curing (referred to "co-polyester network"). The purified CM contains about 95 % non-polymerized 9(10),16-dihydroxyhexadecanoic acid and 5 % dicarboxylic fatty acids; curing them at 150 °C for 24 h resulted in a weakly crosslinked polyester network, with an estimated average number of 80 repeating units between crosslinking nodes [4]. As the tomato cutin monomers consist of saturated lipid chains which are resistant to oxidation, no degradation is expected for both materials stored in a desiccator at room temperature.

### 2.2. X-ray diffraction

Wide Angle X-ray diffraction (WAXD) measurements were performed on a Bruker D8 Advance X-ray diffractometer with a copper source ($\lambda_{Cu}$ = 0.154 nm) at room temperature within the angular range $2\theta$ = 5–40° and with a scanning step of 0.04°/s. The integrated intensities of the crystalline diffraction peaks and amorphous halo were

<figure><img src="image_2.png" /><figcaption>Fig. 1. Chemical structure of the 9(10),16-dihydroxyhexadecanoic acid.</figcaption></figure>

estimated using the DIFFRAC.EVA software. X-ray patterns have been recorded at room temperature, and after liquid $N_2$ quenching from the molten state.

### 2.3. Polarized-light optical microscopy

Polarized-light Optical Microscopy (POM) observations were performed using a universal Nikon EPI-illuminator with Nikon M plan lenses (x10; x20) connected to a Nikon DS-Fi2 camera, either in transmission of reflection mode. A sample was placed between two glass slides, then heated onto a Mettler Toledo HS82 Hot stage to 100 °C (above the melting temperature) and annealed for 3 min. A gentle pressure was then applied to the upper cover glass to squeeze the melt into a thin film before cooling down to room temperature at a controlled rate of 2 K min$^{-1}$. The hot stage was placed underneath the lenses for in-situ observations at different temperatures.

### 2.4. Stochastically modulated temperature DSC

Stochastically modulated temperature DSC (TOPEM) allows separating thermal events that are partially overlapping with other effects, such as crystallization and melting [13]. The total heat flow $\Phi_{tot}$ corresponds to the signal that one would measure with standard DSC, and is the algebraic sum of two components, the reversing heat flow $\Phi_{rev}$ and the non-reversing heat flow $\Phi_{non-rev}$. The first term is related to thermal events involving a heat capacity change, such as melting, whereas the second term is related to kinetic events, such as crystallization. From $\Phi_{rev}$ the quasi-static heat capacity $c_{p,0}$ is derived using an extrapolation to zero frequency. A correction based on sapphire is required as for other DSC techniques.

TOPEM measurements were performed on a DSC 3+ apparatus (Mettler Toledo) equipped with a FRS 6+ sensor and Huber intracooler TC100. Temperature, enthalpy and thermal lag calibrations were performed using indium and water standards. Three different heating rates were used: 2 K min$^{-1}$, 5 K min$^{-1}$ and 10 K min$^{-1}$, which allowed to compute the internal thermal lag of the device $\tau$.

Sample masses comprised between 3 and 10 mg were inserted in a 40 µL sealed aluminum pans. The experiments were performed under nitrogen (50 mL min$^{-1}$ flow rate), with a pulse height of 0.1 K and pulse widths between 15 and 30 s. TOPEM measurements consisted in heating segments from room temperature (20 °C) up to 100 °C with a heating rate of 2 K min$^{-1}$, followed by an isothermal holding step at 100 °C for 2 min, and then cooling from 100 °C down to –60 °C with a cooling rate of 2 K min$^{-1}$ followed by a subsequent heating from –60 °C up to 100 °C at the same rate. The upper limit of 100 °C was chosen to prevent further polymerization or crosslinking of the samples [4,14].

### 2.5. Fast scanning calorimetry

Fast Scanning Calorimetry (FSC) investigations were performed on a Flash DSC 2+ (Mettler Toledo) equipped with a Huber intracooler TC100 with conditioned and calibrated Multi STAR UFH1 sensors under a continuous argon flow of 60 mL min$^{-1}$. Each chip was conditioned according to the Mettler Toledo software-integrated conditioning procedure to remove internal stresses and solvent remains while recording the baseline. The samples were cut and placed at the center of the calibrated chip sensor under a binocular microscope, and their thermal contact with the chip membrane was ensured through several cycles of slow cooling and heating.

The Sensor Support Temperature ($T_{ss}$) of the apparatus was set to –95 °C. The samples were first heated up to $T_m$ + 50 °C to erase their thermal history, then they were rapidly cooled to the selected crystallization temperature $T_c$ (10,000 K s$^{-1}$ for the purified CM and 5000 K s$^{-1}$ for the co-polyester network) and kept for a given time ($t_c$) in isothermal conditions ($T_c$) following the thermal protocol described in Fig. 2. More details about the choice of such crystallization protocol are given by
<figure><img src="image_3.png" /><figcaption>Fig. 2. Thermal protocols used for the isothermal crystallization of purified cutin monomer and its network. The sample is first heated up to $T_{max} = T_m + 50$ °C (twice) to erase its thermal history (rejuvenating segment), then rapidly cooled down to the selected crystallization temperature $T_c$ and held isothermally for a defined time $t_c$. After isothermal crystallization, the sample is again rapidly cooled down to $T_{min} = T_g - 50$ °C and finally heated up to $T_{max}$ for measurement. The same heating and cooling rates are applied ($\beta_h = |\beta_c|$).</figcaption></figure>

Fosse et al. [12]. Because the samples crystallize rapidly, the mass estimation relying on the ratio between the specific heat capacity step at the glass transition $\Delta c_p$ recorded using TOPEM (or standard DSC), and heat capacity step $\Delta C_p$ measured by FSC on the fully amorphous counterpart of the same material, could not be used here. Instead, the ratios of $C_p^{FSC}(T)$ and $C_p^{TOPEM}(T)$ measured at different temperatures in the glassy and in the liquid states were averaged to get the sample mass (Eq. (1)), which provided mass values comprised between 10 and 20 ng.

$$m = \frac{C_p^{FSC}(T) [J K^{-1}]}{c_p^{TOPEM}(T) [J g^{-1} K^{-1}]} \qquad \text{Eq 1}$$

Using fast scanning rates can lead to smearing effects and a significant thermal inertia, which in turn causes shifts and distortions of the observed thermal events, thus requiring specific procedures to correct the thermal lag, and therefore the recorded curves [15–18]. However, since the sample masses and thicknesses were found to be sufficiently small (few tenths of ng), no specific correction procedures were required for the measurements presented hereafter.

### 3. Results and discussion

#### 3.1. First thermal analyses of the purified monomers and its network

Preliminary thermal characterizations of the non-polymerized (purified CM) and polymerized (co-polyester network) samples were done using stochastically modulated temperature DSC (TOPEM). The reversing, non-reversing and total heat flows recorded for each step of the thermal protocol illustrated in Fig. 2 are represented in Fig. 3.

Different melting behaviors are observed for the non-polymerized (purified CM) and polymerized (co-polyester network) samples, during both the first and second heating ramps. The thermal parameters extracted from these heating and cooling curves are presented in Table 1. Fig. 3a and 3b shows that the purified CM is characterized by a complex endothermic peak with several shoulders from 40 °C to 92 °C in the reversing heat flow, whereas the co-polyester network manifests a main melting peak from 46 °C to 60 °C followed by another smaller

endothermic peak on the high-temperature side. The hatched areas show that recrystallization and melting occur simultaneously in both cases, proving the interest of TOPEM investigations for separating overlapping thermal events. Significant decreases in the melting enthalpy are observed on the total heat flow ($\Phi_{tot}$) between the first and second heating ramps, e.g. from 155 down to 87 J g$^{-1}$ for the purified CM, and from 74 down to 40 J g$^{-1}$ for the co-polyester network. This difference is due to a slow crystallization during the storage of the polyesters at room temperature, leading to a higher melting enthalpy during the first heating. The value of the total heat flow obtained during the second ramp is close to the value of 41.1 J g$^{-1}$ previously reported by Marc et al. [4] using standard DSC at 3 K min$^{-1}$ for the second heating ramp. Recording complex melting peaks made of several endotherms is not unusual for fatty acids due to the existence of different polymorphs ($\alpha$, $\beta'$ and $\beta$ forms) [19]. Each polymorphic form has a different crystalline structure with different thermodynamic stability, meaning that these polymorphs potentially have different lifetimes within a crystalline matrix as the crystalline structure (re)arranges. When striving for a greater stability, the $\alpha$ crystals would convert to $\beta'$ crystals, and then the $\beta'$ would further transform into $\beta$ crystals [20].

The observed low-temperature melting endotherm is generally associated with the least stable $\alpha$ phase, or with a sub-modification of the $\alpha$-form; the intermediate-temperature endotherms are associated with the $\beta'$-form or with other sub-modifications of the $\beta'$-form; the endotherm found at the highest temperature is generally attributed to the most stable polymorph $\beta^{20}$. The melting temperature associated with each of these polymorphs were found to depend on the fatty acid chain length; the longer is the chain length, the less defined is the separation between two adjacent melting peaks [21].

The exothermic crystallization peak recorded upon cooling is very sharp for both materials, and is located at a slightly higher temperature for the co-polyester network. A secondary crystallization peak is noticeable at about 0 °C for the purified CM. A low value of $T_g$ is obtained for both materials, resulting from the high molecular mobility related to the presence of long methylene sequences. A gap of about 35 °C is observed between the glass transition temperatures measured upon cooling for the semi-crystalline purified CM and the co-polyester network, due to the crosslinking nodes forming the network, as evidenced by Marc et al. [4]. In addition to the change observed in the $T_g$ value, the broadness of the glass transition differs as well when the purified CM and the co-polyester network are compared. This broadening is evidenced in the $\Delta T_g$ value (defined as $T_{g,max} - T_{g,min}$) shown in the insets to Fig. 3c and 3d. For the purified CM $\Delta T_g$ is only 3 °C, whereas for the co-polyester network $\Delta T_g$ reaches 25 °C, which is a clear sign of increased heterogeneity in the relaxation times associated with the molecular movements in the amorphous phase recorded as the temperature decreases, i.e. upon cooling. Since the purified CM seems to be more crystallized than the co-polyester network, this increase in the heterogeneity of the relaxation times cannot be imputed on a confinement effect due to a higher crystalline fraction, but is rather due to structural effects. Trubert et al. [14] drew a similar conclusion when assessing by dielectric relaxation spectroscopy the impact of the macromolecular network on the molecular mobility in this same co-polyester network, concluding that the mobility hindrance caused by crosslinking is stronger than the hindrance due to the crystalline phase. The values of melting enthalpies differ as well between the purified CM and the co-polyester network, since a larger value of $\Delta h_m$ is obtained for the purified CM both upon the first and the second heating ramps; the effects of crosslinking on the crystallization ability of the macromolecular segments are clearly evidenced by the value of melting enthalpy, that is double for the purified CM compared to the co-polyester network.

Striving to get a deeper understanding of the nucleation and growth process, the microstructures of both the purified CM and co-polyester network were observed under Polarized-light Optical Microscopy (POM). Only the observations made for the purified CM are shown here.
<figure><img src="image_4.png" /><figcaption>**Fig. 3.** Normalized total heat flow (black), reversing heat flow (red) and non-reversing heat flow (blue) curves obtained upon first heating for a) purified CM and b) co-polyester network with a heating rate of 2 K min⁻¹. Hatched areas indicate the melting (red) and recrystallization (blue) processes. Normalized total heat flow obtained for c) purified CM and d) co-polyester network upon cooling at 2 K min⁻¹. The insets show a magnification of the glass transition, with ΔT<sub>g</sub> indicating the broadness of the glass transition, and Δc<sub>p</sub> indicating the step height at the glass transition temperature. Normalized total heat flow (black), reversing heat flow (red) and non-reversing heat flow (blue) curves obtained upon second heating for e) purified CM and f) co-polyester network with a heating rate of 2 K min⁻¹. Hatched areas indicate the melting (red) and recrystallization (blue) processes.</figcaption></figure>

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

Glass transition temperature (T<sub>g</sub>), specific heat capacity step measured at T<sub>g</sub> (Δc<sub>p</sub>), temperature range corresponding to the melting process (ΔT<sub>m</sub>), melting enthalpy (Δh<sub>m</sub>) measured on the total (Φ<sub>tot</sub>) heat flow and on the reversing (Φ<sub>rev</sub>) heat flow, crystallization temperature (T<sub>c</sub>) and crystallization enthalpy (Δh<sub>c</sub>) measured upon cooling on the non-reversing heat flow (Φ<sub>non-rev</sub>).

<table><thead><tr><th>Sample</th><th></th><th>T<sub>g</sub> [°C]</th><th>Δc<sub>p</sub> [J g<sup>-1</sup> K<sup>-1</sup>]</th><th>ΔT<sub>m</sub> [°C]</th><th>Δh<sub>m</sub> [J g<sup>-1</sup>] Φ<sub>tot</sub></th><th>Δh<sub>m</sub> [J g<sup>-1</sup>] Φ<sub>rev</sub></th><th>T<sub>c</sub> [°C]</th><th>Δh<sub>c</sub> [J g<sup>-1</sup>] Φ<sub>non-rev</sub></th></tr></thead><tbody><tr><td rowspan="3">Purified CM</td><td>1st heating</td><td>/</td><td>/</td><td>40–92</td><td>155 ± 3</td><td>218 ± 5</td><td>/</td><td>/</td></tr><tr><td>cooling</td><td>-53 ± 1</td><td>0.19 ± 0.01</td><td>/</td><td>/</td><td>/</td><td>29 ± 1</td><td>57 ± 2</td></tr><tr><td>2nd heating</td><td>/</td><td>/</td><td>26–80</td><td>87 ± 5</td><td>156 ± 5</td><td>/</td><td>/</td></tr><tr><td rowspan="3">Co-polyester network</td><td>1st heating</td><td>/</td><td>/</td><td>46–74</td><td>74 ± 2</td><td>92 ± 2</td><td>/</td><td>/</td></tr><tr><td>cooling</td><td>-18 ± 1</td><td>0.15 ± 0.01</td><td>/</td><td>/</td><td>/</td><td>35 ± 1</td><td>37 ± 2</td></tr><tr><td>2nd heating</td><td>-22 ± 1</td><td>0.14 ± 0.01</td><td>33–76</td><td>40 ± 5</td><td>94 ± 5</td><td>/</td><td>/</td></tr></tbody></table>
Because of the crosslinking, the co-polyester sample kept its shape even at high temperature and could not be totally squeezed to a thin layer between the glass slides, although small spherulites could be observed. Nevertheless, the purified CM extract could be observed (Figs. 4 and 5).

Fig. 4 depicts the crystallization of the purified CM as it was observed at specific temperatures during a cooling ramp at 2 K min⁻¹ from 100 °C down to 25 °C. At 53 °C sporadic nucleation occurs, and as the temperature decreases, those early nuclei grow in a radial fashion to form micron-sized particles. At about 33 °C, what could be a spontaneous and very rapid space-filling secondary nucleation occurs, closely followed up by growth until the temperature reaches 25 °C, corresponding to the main exothermic event observed in Fig. 3c. This rapid period of space-filling crystallization ends up masking the larger microstructures formed during the earlier stages of crystallization [20]. The resulting microstructure consists of a large number of small crystalline structures, likely quite heterogeneous in size and crystalline perfection. Fig. 5 shows the micrograph recorded once the cooling ramp is completed and after microstructure stabilization. Star-shaped structures are visible with a large size distribution, supposedly connected in a vast crystal network stabilized by van der Waals forces [22].

Fig. 6 shows the WAXD patterns recorded for the purified CM and the co-polyester network after a long storage (four years) at room temperature (red patterns) and after rapid cooling from the molten state (blue patterns). After a long storage at room temperature, the co-polyester network presents an orthorhombic β' polymorphic form, with two peaks at 21.1° (d = 4.2 Å) and 22.3° (d = 4.0 Å) (Fig. 6b). The presence of a Lorentzian peak tail in the X-ray spectrum could account for the presence of a paracrystalline phase (partially disordered crystals with short or medium-range order of the lattice) [23]. More peaks are visible in the WAXD pattern of the purified CM (Fig. 6a). The same

<figure><img src="image_5.png" /><figcaption>Fig. 5. Micrograph (magnification x20) recorded at 25 °C for the purified CM after cooling from 100 °C at 2 K min⁻¹.</figcaption></figure>

orthorhombic β' polymorphic form is observed through the peaks located at 21.6° (d = 4.2 Å) and 23.4° (d = 3.8 Å). An additional peak at 9° (d = 10.3 Å) is visible, which could indicate a longitudinal organization of the chains, as indicated in Marc et al. [4]. Two more peaks at 19.5° (d = 4.5 Å) and 25.2° (d = 3.5 Å) are observed, revealing the presence of a triclinic packing characteristic of the β form [24]. Interestingly, thermal analysis and X-ray diffraction both suggest the coexistence of β and β'-forms in the purified CM, but only the orthorhombic form β' is observed for the co-polyester network, indicating that the

<figure><img src="image_6.png" /><figcaption>Fig. 4. Micrographs recorded while cooling the purified CM from 100 °C down to 25 °C at 2 K min⁻¹.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Fig. 6.** X-ray patterns recorded at room temperature for a) purified CM and b) co-polyester network after 4-year storage at room temperature (red pattern) and after liquid N₂ quenching from the molten state (blue pattern). The low-angle deviation is due to the extinction of the incident beam.</figcaption></figure>

relative stability of the β and β'-form depends on topology. For both samples, fast cooling from the melt results in a hexagonal packing of the acyl chains (α polymorphic form) characterized by a single peak at 21° (d = 4.2 Å). Marc et al. [4] previously found that, after rapid cooling of the polyester network, as the sample is stored for a certain time at room temperature, a polymorphic transformation occurs, from the unstable α-phase to the more stable orthorhombic β' sub-cell within seven days of storage.

## 3.2. Isothermal crystallization kinetics of HFAs-derived samples

The procedure proposed by Cebe et al. [11] and Fosse et al. [12] relies on a linear regression of the normalized melting enthalpy against the specific heat capacity for different crystallization fractions assuming a two-phase model with reduced phase coupling. This requires, for as much as possible, the absence of a rigid amorphous fraction (RAF), whose formation depends on the temperature at which the crystallization is performed [25]. When the temperature chosen for isothermal crystallization is too low, the decreased chain mobility leads to the formation of imperfect crystals with a strong coupling with the neighboring amorphous phase, increasing the RAF content [26]. On the opposite, a high crystallization temperature results in a phase decoupling, although pushing it too high would lead to a competition between crystallization and melting. Although the RAF has scarcely been studied in HFAs-derived polyesters, it has been characterized in similar macromolecular systems, such as PHAs. Righetti et al. [27] investigated the evolution of tensile mechanical properties of different poly(3-hydroxybutyrate-co-3-hydroxyvalerate) copolymers as a function of the storage time at room temperature. In parallel, they quantified the crystalline, mobile amorphous (MAF) and rigid amorphous (RAF) fractions. They found that $X_c$ increases, $X_{MAF}$ decreases and $X_{RAF}$ increases during storage at room temperature, which they attributed to secondary crystallization, that progressively constrains adjacent amorphous chain segments. Another work by Di Lorenzo et al. [28] showed how the presence of a RAF influences the kinetics of cold crystallization in polyhydroxybutyrate (PHB). They found that RAF delays the cold crystallization onset because more energy is needed to mobilize the constrained chains before they can reorganize into crystals. Moreover, as heating proceeds, some of the RAF devitrifies, thus contributing to further crystallization. Esposito et al. [26] found that the glass transition is broadened and shifted to higher temperatures when poly(hydroxybutyrate-co-hydroxyvalerate) (PHBV) is isothermally crystallized from the glassy state (lower $T_{cryst}$) rather than from the melt (higher $T_{cryst}$), ascribing this phenomenon to a confinement effect and a stronger coupling between crystalline and amorphous regions.

Based on the thermographs recorded with TOPEM (Fig. 3), a temperature of 20 °C is chosen for isothermal crystallization, i.e. a

temperature below the first melting event, but still high enough for the crystallization to occur with as little phase coupling as possible. There is a second reason for choosing this crystallization temperature, i.e. previous investigations showed that the mechanical behavior of the co-polyester network at room temperature is determined by its semicrystalline structure [4], which is something deserving further investigations to properly estimate the crystalline fraction in the tested samples. Indeed, during storage, a number of post crystallization processes occur that can affect properties; hardness, for instance, often noticeably increases because of the so-called sintering, a phenomenon related to the formation of solid bridges between crystals, eventually forming a percolating network [29]. In any case, it should be pointed that applying the procedure proposed by Cebe et al. [11] and Fosse et al. [12] to slightly crosslinked polymer networks may be challenging, as it is quite hard to predict the influence of topological constraints other than entanglements. The method may seem more legit in thermoplastic polymers compared to crosslinked networks, for the molecular mobility is only temporarily constrained by entanglements in the former case, whereas crosslinks are irreversible chemical bonds in the latter. However, having a fully crystalline microstructure in a macromolecular material is equally unlikely in both thermoplastic and slightly cross-linked samples; in this regard, it appears reasonable to apply the same extrapolation method.

Fig. 7 shows the normalized heat flow curves measured by FSC for the purified CM and the co-polyester network following different times of isothermal crystallization at 20 °C. The main melting peak increases from 40 to 60 °C for the purified CM, and to 60–70 °C for the co-polyester network. More complex phenomena are observed for the purified CM, such as cold crystallization and additional endothermic peaks resulting from the formation of nuclei upon cooling. Further investigations at increasing cooling rates could probably help to understand what really happens during cooling, as well as to predict what would happen during processing [30]; the melting enthalpy resulting from cold crystallization is anyway neglectable in comparison with the melting of the crystals formed at 20 °C. For each isothermal crystallization, the specific heat capacity $c_p$ and the melting enthalpy $\Delta h_m$ are estimated from the curves recorded during the subsequent heating.

An increase of $T_g$ from –16 °C for the purified CM, to 0 °C for the co-polyester network, is observed in the absence of crystallization. Compared to the values reported in Table 1, these values of $T_g$ are shifted towards higher temperatures because of the very high scanning rates involved with FSC [31,32]. The glass transition temperature $T_g$ remains largely unaffected by crystallization, which is a sign of reduced coupling between the crystalline and the amorphous fractions. However, as the crystallized fraction increases, $\Delta c_p$ decreases, and for long crystallization times the glass transition is not visible anymore for purified CM. The
<figure><img src="image_8.png" /><figcaption>**Fig. 7.** Heat flow curves measured upon heating at 10,000 K s⁻¹ for **a)** purified CM after crystallization at 20 °C from 0 to 100 s. Heat flow curves measured upon heating at 5000 K s⁻¹ for **b)** co-polyester network after crystallization at 20 °C from 0 to 1000 s.</figcaption></figure>

shape of the melting curve is also affected, and the broadness of the melting peak increases due to crosslinks. A shift of the melting endotherm towards higher temperatures is progressively observed for crystallization times beyond 1s; a partially semi-crystalline structure is obtained for short crystallization times, and the final semi-crystalline structure is obtained for long crystallization times.

The decrease in $\Delta c_p$ with the crystallization time is represented in Fig. 8 for both samples. The purified CM shows a larger and sharper drop of $\Delta c_p$ at 0.5 s. The crystallization process is hindered in the co-polyesters network because of crosslinking, with a decrease in $\Delta c_p$ starting at 0.5 s until reaching a constant value after a crystallization time of 10 s. For the purified CM, the amorphous phase gets so constrained that no heat capacity step is observed during the glass transition. This result corroborates the POM observations (Figs. 4 and 5) showing a high density of tightly packed spherulites.

### 3.3. Theoretical melting enthalpy of an ideally 100% crystalline sample
$\Delta h_m^0$

Fig. 9 shows the melting enthalpy estimated from the thermograms in Fig. 7 as a function of the specific heat capacity step at the glass

<figure><img src="image_9.png" /><figcaption>**Fig. 8.** Change in the specific heat capacity step at the glass transition $\Delta c_p$ as a function of the crystallization time $t_c$ varying from $10^{-3}$ to $10^3$ s for purified CM (red) and co-polyester network (blue). Vertical bars give the measurement uncertainty. Dashed lines are a guide for the eye.</figcaption></figure>

<figure><img src="image_10.png" /><figcaption>**Fig. 9.** Melting enthalpy $\Delta h_m$ against the specific heat capacity step at the glass transition $\Delta c_p$ normalized by the sample mass and the heating rate for both the purified CM (red) and the co-polyester network (blue) crystallized for different crystallization times $t_c$ at a crystallization temperature $T_c = 20$ °C. A linear regression (solid line) extrapolated to $\Delta c_p = 0 \text{ J g}^{-1} \text{ K}^{-1}$ ($X_c = 0\%$) gives the theoretical melting enthalpy $\Delta h_m^0$ corresponding to an ideally fully crystallized material under the assumption of a two-phase model [11,12]. Hatched areas give the uncertainty of $\pm 10 \text{ J g}^{-1}$ for the estimation of $\Delta h_m^0$.</figcaption></figure>

transition. From Fig. 9, the melting enthalpy $\Delta h_m^0$ (calculated by a linear regression extrapolated to $\Delta c_p = 0 \text{ J g}^{-1} \text{ K}^{-1}$) reaches a value of $102 \pm 10 \text{ J g}^{-1}$ and $77 \pm 10 \text{ J g}^{-1}$ for the purified CM and the co-polyester network, respectively. The uncertainty is mainly due to the arbitrary baseline used to integrate the melting peaks, as well as to the mass estimation of the samples analyzed by FSC. Because the thermal protocol is similar to the one employed for the WAXD measurements after rapid cooling from the melt by liquid nitrogen, i.e. a magnitude order of $10^{-3} \text{ K s}^{-1}$ for the cooling rate, it seems reasonable to assume that the crystals formed correspond to the metastable $\alpha$-crystals with a hexagonal packing.

The X-ray diffraction patterns recorded at room temperature for the co-polyester network showed the formation of a hexagonal packing of the acyl chains ($\alpha$ polymorphic form) after rapid cooling from the melt [4]. However, after seven days of storage at room temperature, the sample was found to be in an orthorhombic $\beta'$ sub-cell. Moreover, the
thermogram recorded with TOPEM showed a broad endotherm with a melting enthalpy $\Delta h_m = 92$ J g⁻¹ for the co-polyester network, corresponding to the melting of the β'sub-cell, as evidenced by WAXD, which is a larger value compared to the theoretical value of $\Delta h_m^0$ estimated by FSC measurements for the α-form. To estimate the value of $\Delta h_m^0$ of the β'-form, a cross-comparison of the WAXD and the calorimetric results was performed. The crystalline fraction was estimated from WAXD patterns using Eq. (2):

$$X_c = \frac{I_{cryst}}{I_{cryst} + I_{am}} \qquad \text{Eq 2}$$

where $I_{cryst}$ is the integrated intensity of the crystalline diffraction peaks and $I_{am}$ is the integrated intensity of the amorphous halo. It is worth mentioning that this approach could underestimate the total crystalline volume fraction, due to the inability to separate overlapping peaks and correctly model the diffuse signal arising from the paracrystalline fraction [23]. The data obtained using calorimetry and XRD can thus differ, as previously reported in the literature [9]; it can still be useful, though, to get an estimation. Based on calorimetric data, $X_c$ was estimated from Eq. (3):

$$X_c = \frac{\Delta h_m - \sum \Delta h_c}{\Delta h_m^0} \qquad \text{Eq 3}$$

Where $\Delta h_m$ is the melting enthalpy, $\Delta h_c$ is the cold crystallization enthalpy, and $\Delta h_m^0$ corresponds to the melting enthalpy of a theoretical fully-crystallized sample. In the absence of cold crystallization, inserting Eq. (3) into Eq. (2) provides the following equation:

$$\Delta h_m^0 = \frac{I_{cryst} + I_{am}}{I_{cryst}} \Delta h_m \qquad \text{Eq 4}$$

From the WAXD spectrum it is found that the co-polyester network in the β'-form has a crystallinity fraction of 78 %, quite a high amount despite of the crosslinking nodes. Such a high crystalline content could have been a consequence of the purification process, which reduces the phenolic content and therefore reduces any further crosslinking process, as otherwise evidenced by Marc et al. [3] and Trubert et al. [14]. The theoretical melting enthalpy associated with the melting of the β' crystals of the co-polyester network was estimated at 92 J g⁻¹ (Table 1), and using Eq. (4) a value of 118 J g⁻¹ is found for the β'-form.

Concerning the β-phase, such procedure is not so straightforward, since the WAXD revealed a co-existence of two different crystalline phases (the β' and β forms) for the purified CM. Gradys et al. [33] used the method developed by Turner-Jones [34] to estimate the content of one crystalline form in relation to another for polypropylene, however this method requires well separated peaks. Flores et al. [35] relied on the molar contribution function of van Krevelen [36] to estimate the theoretical melting enthalpy $\Delta h_m^0$ of new aliphatic polyethers. When applying van Krevelen function to the saturated 9(10)-16-dihydroxyhexadecanoic acid) whose structure is presented in Fig. 1 (major component of the purified CM), $\Delta h_m^0$ would be:

$$\Delta h_m^0 = \frac{(\Delta h_{CH_2} n_{CH_2} + \Delta h_{OH} n_O + \Delta h_{CH(OH)} n_{CH(OH)} + \Delta h_{OH} n_{OH})}{M_i} \qquad \text{Eq 5}$$

Where $n_x$ is the amount of a given group x in the repeating unit and $M_i$ is the molar mass.

In the repeating unit there are 14 -CH₂- groups (4.0 kJ mol⁻¹ per group), 2 -OH- groups (1.0 kJ mol⁻¹ per group), one -C=O- group and one CH(OH)- group (1.3 kJ mol⁻¹ per group) [4]. The molar mass $M_i$ of the repeating unit is 288 g mol⁻¹. Assuming a deviation of 11 % associated with van Krevelen function [36], the estimation gives a theoretical melting enthalpy of 206 ± 23 J g⁻¹, which may correspond to the most stable polymorph β, and which is close enough to the melting enthalpy of 208.2 J g⁻¹ reported by Cedeno et al. [37] for pure palmitic

acid (also containing 16 carbon atoms in the chain).

Fig. 10 shows the evolution of the melting enthalpy $\Delta h_m$ with crystallization times. Based on calorimetric and WAXD results, it seems that both the purified CM and the co-polyester network initiate crystallization rapidly at room temperature under the α-form (hexagonal packing) before transforming into the β'-form (orthorhombic packing).

The co-polyester network crystallizes under the orthorhombic β' form when stored during sufficiently long times, whereas the purified CM is characterized by the coexistence of different crystalline forms, including the α and β' forms. Different metastable states are reached before transforming from the hexagonal into the orthorhombic phase. In the literature, secondary crystallization can be activated through different processes, with RAF participation [38] or through cross-nucleation, with the initial crystalline phase providing the nucleation sites for secondary crystallization. POM experiments seem to suggest that this second mechanism is to be privileged in the case of the purified CM. The α-crystals have a short lifetime, which was expected for a metastable crystalline form, and convert rapidly (in the span of minutes) into β'-crystals. As the initial crystallization process reaches its end, secondary crystallization occurs for the co-polyester network, involving a solid-solid transformation from the initial α-form (hexagonal packing) to the β'-form (orthorhombic packing). In the samples crystallized at 20 °C, a mesophase forms as an intermediate transient stage within the path of transformation of the melt into β'-crystals, following Ostwald's rule of stages. With this crystallization pathway, the isotropic melt first converts into a Liquid Crystalline phase (LC-phase) and then undergoes a monotropic transition into a more ordered crystalline phase. Such crystallization pathway is not uncommon in aromatic polyesters, where the stiff aryl rings act as mesogenic units [39], or in other polymers with high flexibility such as polyethylene, for which it was shown that during the initial process of growth, a crystal goes through thermodynamically stable and metastable states before a transformation from the hexagonal to the orthorhombic phase occurs [40]. In-situ optical microscopy and

<figure><img src="image_11.png" /><figcaption>Fig. 10. Dependence of the melting enthalpy $\Delta h_m$ on the crystallization time $t_c$ at room temperature for both the purified CM and the co-polyester network. The square symbols represent the FSC experimental data. The full and empty triangles represent the experimental data from the total heat flow $\phi_{tot}$ and the reversing heat flow $\phi_{rev}$ measured with TOPEM, respectively. The blue diamond symbol indicates the melting enthalpy measured for the co-polyester network by Marc et al. [4] after one-week storage using standard DSC. Dotted gray lines are added to indicate the storage time at room temperature before calorimetric measurements. The dashed lines are a guide for the eye. The experimental points at $t_c < 10^3$ s correspond to in-situ crystallization performed within FSC, whereas the experimental points at $t_c > 10^6$ s are issued from ex-situ crystallization.</figcaption></figure>
transmission electron microscopy showed that the newly formed crystals actually act as nucleation centers for secondary crystallization [40].

### 3.4. Application of the Avrami model for isothermal crystallization

To better understand the formation of the α-form during the initial growth stage, the Avrami model for isothermal crystallization was employed [41]:

$$ \alpha(t) = 1 - \exp(-k(t - t_0)^n) \qquad \text{Eq 6} $$

where $\alpha(t)$ is the relative crystalline fraction transformed as a function of time, n is the Avrami index, k is the crystallization rate constant, and $t_0$ is the induction time at which the crystallization process initiates. In the logarithmic form, Eq. (6) becomes:

$$ \log[-\ln(1-\alpha)] = \log(k) + n \log(t-t_0) \qquad \text{Eq 7} $$

The evolution of the relative fraction α with the isothermal crystallization time $t_c$ is plotted in Fig. 11a. The increase of $\alpha(t)$ with the crystallization time follows a typical sigmoidal shape. However, the blue curve related to the co-polyester network reveals two crystallization modes. The primary crystallization is associated with a rapid increase of the relative crystalline fraction (initial crystal growth) and is followed by a secondary crystallization, which is characterized by a linear increase of the relative crystalline fraction $\alpha(t)$ with respect to time in logarithm scale (crystal thickening). A similar behavior was observed in other polyesters, such as PLA [42] and PCL [43]. When the molar mass increases (co-polyester network vs. purified CM), the curves are shifted to longer times, with the branched chains and the crosslinks inhibiting the crystallization as the motions of the macromolecular segments is hindered. Fig. 11b shows the Avrami plots obtained for the purified CM and the co-polyester network. Besides a time-independent transition rate k and a possibly different Avrami exponent n, the evolutions of $\log(-\ln(1-\alpha))$ as a function of $\log(t-t_0)$ should appear linear. However, the results clearly show nonlinearity at the early stage of crystallization in the case of the purified CM. One possible reason of such behavior could be the time-dependent effects arising while approaching the steady-state regime at early stages of crystallization, as mentioned by Gradys et al. [44].

The Avrami exponents n informs about the nucleation mode for the subsequent growth of the crystalline domains. Here, the n values range from 2.7 for the purified CM down to 1.6 for the co-polyester network. The Avrami exponent depends on the nucleation mechanism as well as on the dimensionality of growth of the crystalline fraction. Using high cooling rates results in higher degrees of undercooling $\Delta T$ and thus stronger driving forces to crystallization, with consequently lower free energy barriers to overcome for the nucleation process [22]. When the

Avrami index n decreases, the nucleation process is less sporadic and the dimensionality of growth is lower [22]. A value of 1.6 suggests that 2D crystalline structures are formed, such as axialites, which is not uncommon for polymers in case of low undercooling [45,46].

### 4. Conclusion

Co-polyesters based on hydroxy fatty acids present complex microstructures depending on their processing conditions and their crosslinking state. Different ordered structures are formed depending on the cooling rate and the crystallization time. After rapid cooling, both purified cutin monomers and their co-polyester networks undergo a monotropic transition from a metastable crystal phase (α-phase) to a more ordered crystal phase (β'-phase) when maintained at room temperature for a short time. Purified cutin monomers can eventually form the β-phase given sufficiently long time. The formation of the stable β polymorph was not observed for the co-polyester network, possibly due to crosslinking nodes and chain branching hindering the formation of a triclinic phase. The use of extremely high cooling rates (several thousands of K s⁻¹) leads to the formation of metastable α-crystals characterized by a hexagonal packing with an interplanar distance d = 4.1 Å in the co-polyester networks as well. From a combination of calorimetry and WAXD experiments, the melting enthalpy of a theoretically fully crystallized sample only containing the α-phase is found equal to 77 J g⁻¹, whereas for the β'-phase and the β-phase the estimated values are 118 J g⁻¹ and 206 J g⁻¹, respectively. When maintained at room temperature, the purified cutin monomers form a complex microstructure with micron-sized spherulites, and crystallize to a full extent. On the other hand, the co-polyester network reaches a crystalline fraction of about 78 % after a four-year storage at room temperature. The present data suggest that the polyester macromolecular segments retain enough mobility despite the weakly crosslinked network. Previous works have demonstrated the major role of crystallinity in the mechanical properties of polyesters and polyester networks. This work delves deeper into their crystallization behavior, which could eventually allow a better control over the properties. Such approach represents a further step towards the design of original, fully bio-based, high-performance materials, in an effort to exploit the biomass to the best of its potentialities.

### CRediT authorship contribution statement

**Kylian Hallavant:** Writing – review & editing, Writing – original draft, Software, Methodology, Investigation, Formal analysis, Data curation. **Antonella Esposito:** Writing – review & editing, Writing – original draft, Validation, Supervision, Methodology, Funding acquisition, Formal analysis. **Bénédicte Bakan:** Writing – review & editing, Writing – original draft, Resources. **Denis Lourdin:** Writing – review &

<figure><img src="image_12.png" /><figcaption>Fig. 11. a) Evolution of the relative crystalline fraction issued from primary crystallization $\alpha(t)$ as a function of the isothermal crystallization time $t_c$ for purified CM (red) and co-polyester network (blue). The dashed lines are guide for the eye. b) Avrami plots for purified CM (red) and co-polyester network (blue). The lines correspond to the fitting of the experimental data following Eq. (6).</figcaption></figure>
editing, Writing – original draft, Validation, Resources, Methodology. **Allisson Saiter-Fourcin:** Writing – review & editing, Writing – original draft, Visualization, Validation, Supervision, Project administration, Methodology, Funding acquisition, 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.

Acknowledgments

Authors acknowledge Mettler Toledo for lending the DSC 3+ and LabEx EMC3 for financing the FSC apparatus DSC 2+. The authors are grateful to Nicolas Couvrat from the laboratory Sciences et Méthodes Séparatives (SMS EA 3233) of the University of Rouen Normandy for WAXD experiments.

Data availability

Data will be made available on request.

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