# Poly(L-Lactide-co-ε-Caprolactone)/glass Fiber Composites: Toward Improved Mechanical Properties of Polylactide-Based Composites

Bernard Miranda Campos, Gaelle Fontaine, Serge Bourbigot, Gregory Stoclet, Fanny Bonnet

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Bernard Miranda Campos, Gaelle Fontaine, Serge Bourbigot, Gregory Stoclet, Fanny Bonnet. Poly(L-Lactide-co-ε-Caprolactone)/glass Fiber Composites: Toward Improved Mechanical Properties of Polylactide-Based Composites. Polymer Composites, 2025, Polymer Composites, 10.1002/pc.70652. hal-05376114

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RESEARCH ARTICLE

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# Poly(L-Lactide-Co-ε-Caprolactone)/glass Fiber Composites: Toward Improved Mechanical Properties of Polylactide-Based Composites

Bernard Miranda Campos<fnref n="1" /> ID | Gaëlle Fontaine<fnref n="1" /> ID | Serge Bourbigot<fnref n="1" />,² ID | Grégory Stoclet<fnref n="1" /> ID | Fanny Bonnet<fnref n="1" /> ID

**Correspondence:** Grégory Stoclet (gregory.stoclet@univ-lille.fr) | Fanny Bonnet (fanny.bonnet@univ-lille.fr)

**Received:** 25 July 2025 | **Revised:** 10 October 2025 | **Accepted:** 29 October 2025

**Funding:** This work was supported by Région Hauts-de-France. Centre National de la Recherche Scientifique. Université de Lille.

**Keywords:** glass fibers | mechanical testing | Polylactide | polymer matrix composites (PMCs) | thermoplastic resin transfer molding (TP-RTM) | ε-Caprolactone copolymers

## ABSTRACT

Inherent brittleness and limited toughness of polylactide (PLA) restrict its use in high-performance applications. To overcome these limitations, poly(L-lactide) (PLLA) is often blended with more ductile thermoplastics or chemically modified through copolymerization with other cyclic esters such as ε-caprolactone (ε-CL). This article reports the study of the mechanical properties of poly(L-lactide-co-ε-caprolactone)/glass fiber composites produced by thermoplastic resin transfer molding (TP-RTM) in one-step synthesis via in situ copolymerization of L-lactide (L-LA) and ε-CL into the mold with pre-disposed glass fibers. The influence of ε-CL content in the PLLA-based copolymer matrix on the resulting mechanical properties of these composites, compared to pure PLLA/glass fiber analogues, was evaluated. Mechanical testing, including tensile, three-point bending, and impact tests, was conducted. Composites achieved Young's modulus above 6.5 GPa and bending modulus over 10.3 GPa. An increase of 90% in the absorbed energy was observed for composites containing 30% of ε-CL in the matrix compared to pure PLLA-based composites. Finally, the morphology and fracture of the composites were observed by optical microscopy, revealing a good impregnation of fibers by the matrix.

## 1 | Introduction

The replacement of oil-based with bio-based and biodegradable plastics in composite materials is the first step to achieving sustainable composites production. Among bio-based and biodegradable polymers, polylactide (PLA) has been the most produced on an industrial scale and commercialized since the late 90s [1]. With a well-established production process, PLA is today worldwide

available [2, 3]. In 2024, the global production capacities of PLA represent 37% of biopolymer fabrication, and its production is expected to increase to almost 42% in 2029 [4]. The increase in global production of PLA is mainly due to the growing availability and productivity of the feedstocks (corn, sugar cane, beet, etc.), but also due to the legislation and policy incentives. The circular economy action instituted by the European Union calls for enhancing the production and use of bioplastics to reduce the environmental

**Abbreviations:** 10 PLCL/GF, Poly (L-lactide-co-ε-caprolactone) with 10%mol of ε-caprolactone reinforced with glass fibers; 20 PLCL/GF, Poly (L-lactide-co-ε-caprolactone) with 20%mol of ε-caprolactone reinforced with glass fibers; 30 PLCL/GF, Poly (L-lactide-co-ε-caprolactone) with 30%mol of ε-caprolactone reinforced with glass fibers; PCL, Poly (ε-caprolactone); PLCL, Poly (L-lactide-co-ε-caprolactone); PLLA, Poly (L-lactide); PLLA/GF, Poly (L-lactide) reinforced with glass fibers.

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

© 2025 The Author(s). *Polymer Composites* published by Wiley Periodicals LLC on behalf of Society of Plastics Engineers.

<footnote n="1">¹Univ. Lille, CNRS, INRAE, Centrale Lille, UMR 8207 - UMET - Unité Matériaux et Transformations, Lille, France | ²Institut Universitaire de France, Paris, France</footnote>
## Summary

* Mechanical tests of poly(L-lactide-co-ε-caprolactone)/ glass fiber produced by TP-RTM.

* ε-CL in PLLA matrix enhances impact strength of PLCL/glass fiber composites.

impact of plastics through recycling and promoting alternatives made from renewable sources [5].

PLA which is biodegradable under industrial composting by anaerobic digestion, is also biodegradable in vivo and biocompatible. This polymer is currently used in the biomedical field and for short-lifetime applications such as packaging [2, 6]. In addition, PLA displays mechanical, optical, and thermal properties comparable to those of some commodity polymers, such as polypropylene (PP), polyethylene terephthalate (PET), and polystyrene (PS), which increases its applicability for technical use [2]. Nevertheless, PLA is a brittle polymer, with an elongation at break of only 2% to 6%. To address these drawbacks, notably the lack of toughness, and to allow the use of PLA in a wider range of extended lifespan applications, it is often used in blends with other polymers or as a matrix in composite materials [7]. Regarding composites, PLA has mainly been used as a matrix in composites reinforced with natural fibers that is, flax, kenaf, and cellulose, to produce fully compostable composites [8]. Overall, short natural fibers are incorporated into the PLA matrix via extrusion, and the composites are produced by compression molding or injection molding [9]. Due to the high viscosity of molten PLA, the production of PLA composites reinforced with continuous fibers is often conducted by commingled fibers or film stacking processes [9].

Some studies reporting the production of PLA composites reinforced with synthetic fibers were done, notably with glass and carbon fibers [10, 11]. Polylactic acid/glass fiber composites (PLA/GF) provide a combination of performance and sustainability when compared to other fiber-reinforced polymer systems, such as epoxy matrix composites, which are petro-based [10, 11]. In contrast to carbon fiber composites, which offer high stiffness and strength but have substantially higher cost and environmental impact, PLA/GF composites represent a more sustainable and cost-effective approach, especially for semi-structural applications [12, 13]. Compared to basalt fiber composites, which exhibit favorable thermal and mechanical resistance, PLA/GF composites remain advantageous due to the biodegradability of their matrix, simplifying their end-of-life management. Although thermosetting matrices with carbon or basalt fibers perform well in aggressive environments, their difficulty in being recycled is a drawback [12].

Advances in the production of PLA composites via liquid composite molding processes have been made, notably via thermoplastic resin transfer molding (TP-RTM) [13–18]. Unlike traditional composite manufacturing processes in which melted PLA is used as a raw material, TP-RTM involves the use of L-lactide (L-LA) monomer, which is directly polymerized in the mold containing the reinforcement. At the molten state, L-LA has low viscosity (<1 Pa·s), which allows its use in TP-RTM, leading to composites with high-quality impregnation [19]. One of the challenges to manufacturing

PLA composites using TP-RTM is linked to the inert character of the reinforcement regarding the ring-opening polymerization (ROP) catalyst to allow the formation of the matrix. In the case of synthetic fibers, glass and carbon fibers are known to be chemically inert, and they are a good choice for studying reactive thermoplastic systems for the production of composites.

Despite its properties, such as strength and stiffness comparable to engineering thermoplastics, PLA displays low elongation at break and small impact resistance [20, 21]. The mechanical properties of PLA can be improved by blending with other polymers, such as polypropylene (PP) [22, 23], acrylonitrile butadiene styrene (ABS) [23], polyamide (PA) [24], poly(ε-caprolactone) (PCL) [25, 26]. PCL is a polyester known for its very high elongation at break of 300% and is thus often used in blends with PLA. Despite the petro-based nature of PCL the latter is also biocompatible and compostable [27].

Another strategy to improve the elongation at break of PLA is to copolymerize L-lactide (L-LA) monomer with other cyclic esters such as ε-caprolactone (ε-CL) to form poly(L-lactide-co-ε-caprolactone) (PLCL) [28]. In that context, our research group recently published a study on the production of poly(L-lactide-co-ε-caprolactone)/glass fiber composites (PLCL/GF) in one-step synthesis from the co-monomers by TP-RTM. Glass fibers were chosen due to their inert character, ensuring that ring-opening copolymerization remains undisturbed. The resulting composites display matrices with high conversions of both comonomers (> 95%) along with $M_w$ as high as $9.07 \times 10^4$ g·mol⁻¹ and narrow dispersities. In opposition to previously reported PLLA/glass fiber by TP-RTM, which was observed to be brittle, these PLCL/GF composites display a rubbery character [13]. In the present work, we extend this approach by performing a comprehensive mechanical characterization (tensile, bending, and impact tests) of PLCL/GF composites with varying ε-CL content, supported by detailed failure analysis and thermal/morphological correlations. This makes it possible to describe the effects of ε-CL on both tensile and bending behavior, examine the shape-recovery response under bending, and note improvements in impact resistance on the bio-based composites.

## 2 | Materials and Methods

### 2.1 | Materials

A balanced 2/2 twill woven glass fabric (HexForce EGlass fabric, Hexcel, USA) with an areal weight of $390 \text{ g} \cdot \text{m}^{-2}$ was used as reinforcement. The fibers exhibit a density of $2.50 \text{ g} \cdot \text{cm}^{-3}$, a Young's modulus of 72 GPa, and a tensile strength of 5600 MPa. The fabric was dried at 100°C for 24 h before use to avoid moisture in the composites. Ten layers of glass fabric were placed into the TP-RTM mold to produce composite plates.

L-lactide (L-LA) monomer (assay >99.5% and water content <0.02%, $T_m=96^\circ\text{C}$) was provided by Corbion (Netherlands) and used as received. ε-caprolactone (ε-CL) monomer (assay >98.5%, $T_m=-1^\circ\text{C}$) and tin (II) 2-ethylhexanoate ($\text{Sn(Oct)}_2$) catalyst (assay >98.5%, $T_m=-1^\circ\text{C}$) were both delivered from Alfa Aesar (USA) and used as received. These monomers and the catalyst formed the reactive thermoplastic system used to form the matrix of the composites.
A commercial poly(ε-caprolactone) (CAPA 6500) was analyzed to compare the thermal properties with the synthesized matrices in this work. A mold sealer Zyvax Sealer GP (ChemTrend), and a mold release agent Zyvax Watershield (ChemTrend) were applied to the mold cavity to prevent imperfections at the surface of the composite plates and facilitate the demolding, respectively.

## 2.2 | Experimental Set-Up and Composites Preparation

A single-component TP-RTM machine (CIJECT III) (Diatex (France)) was used to manufacture the composites (Figure S1 in the Supporting Information). The glass fabric was cut with a roll-cutter and placed into the cavity of a 3-part mold with dimensions of 120 × 120 × 5mm (Figure S2 in the Supporting Information). Ten layers of glass fabric were placed into the mold in the same orientation, and all composites displayed a fiber volume ratio of 30%. A pressure of 20 kN was applied to keep the mold closed. A quantity of 200g of co-monomers (L-LA and ε-CL) was used with ratios of ε-CL of 0, 10, 20, and 30 mol% [14]. The co-monomers and Sn(Oct)₂ catalyst, with a molar ratio of [L-LA+ε-CL]/[Sn]=2000, were placed into the injection tank under vacuum (−0.5 bar), kept for 30 min at 150°C until the mixture was completely melted, and then stirred for 1 min. The pressure in the tank was set at 1.5 bar using nitrogen, and the mixture was then injected into the mold containing pre-disposed glass fabric. After the injection, the mold was kept under pressure at 185°C for 2h for neat PLLA plates and 5h for the poly(L-lactide-co-ε-caprolactone) (PLCL) plates and then cooled at room temperature. The TP-RTM parameters, including the monomer to catalyst ratio, polymerization temperature, and polymerization time, were established according to previous work [13, 14]. In order to carry out the characterizations, three samples of each prototype were produced.

In a previous study, we reported the production of novel poly(L-lactide-co-ε-caprolactone)/glass fiber composites by TP-RTM with various ε-CL ratios in the copolymer matrix ranging from 10 to 30 mol% [14]. The resulting matrices displayed co-monomer conversions superior to 95% in all cases (Table S1 in the Supporting Information), indicating that the ring-opening copolymerization (ROP) took place successfully. The samples exhibited $M_w$ values ranging from 4.07 to 9.09×10⁴ g·mol⁻¹ with dispersities between 1.74 and 2.96 (Table S1 in the Supporting Information). Globally, the molecular weight is correlated with the ε-CL content, decreasing as the ε-CL concentration increases. Other research groups also observed this effect [29, 30]. The microstructure of the PLCL matrices was also studied by ¹H NMR analysis, revealing a statistical character of all PLCL matrices [14].

## 2.3 | Characterization of Composites

### 2.3.1 | Size Exclusion Chromatography (SEC)

Size exclusion chromatography (SEC) was used to determine the number average molecular mass ($M_n$), the weight average molecular weight ($M_w$) and the dispersity (D) of the poly(l-lactide) (PLLA) and poly(l-lactide-co-ε-caprolactone) (PLCL) matrices. The samples were prepared with a concentration of 1 mg·mL⁻¹,

using chloroform (CHCl₃) as a solvent for PLLA and tetrahydrofuran (THF) for PLCL and filtered through a 0.45 µm pore-size membrane. Given the widely acknowledged tendency of polystyrene-calibrated SEC measurements to overestimate the real molecular weight values of aliphatic polyesters, attributed to the relatively smaller hydrodynamic volume of polystyrene, the molecular weights of PLLA and PLCL were corrected using their respective correction factors.

Measurements of molecular weights of PLLA matrices were performed on an Agilent system (1260 Infinity II LC System) equipped with three columns (Shodex K-802.5, 803 and 804) placed in series and coupled with a differential refractometer (Wyatt Optilab). The PLLA samples were analyzed at 25°C in CHCl₃ and a flow rate of 1 mL·min⁻¹. It was established that both $M_n$ and $M_w$ would be corrected by a factor of 0.68 [31, 32].

The molecular weights of PLCL matrices were measured on an Agilent system (1260 Infinity II LC System) equipped with three columns (Waters HR1, HR3 and HR4) placed in series and coupled with a differential refractometer (Wyatt T-REX). The PLCL samples were analyzed at 35°C using THF as solvent and a flow rate of 1 mL·min⁻¹. The molecular weights ($M_n$ and $M_w$) of PLCL were corrected by combining the correction factors of PLLA and PCL in THF, according to the ratio of co-monomers in the PLCL matrix. The correction factor used for the PLLA is 0.58 and 0.56 for PCL [33, 34]. The correction factor applied to PLCL matrices was then: [0.58 × wt % LA units + 0.56 × wt % CL units].

### 2.3.2 | Thermogravimetric Analysis—TGA

The thermal degradation of the PLLA and PLCL matrices was studied by thermogravimetric analysis (TGA) using a TG 209 F1 Libra device. Approximately 10 mg of samples were placed in open alumina pans and submitted to an isotherm at 50°C for 10 min, followed by a heating ramp of 10°C/min until 800°C under $N_2$, with purge flow rates set at 15 and 100 mL·min⁻¹. Data collected from the TGA of PLLA and PLCL matrices were the onset degradation temperature ($T_{5\%}$), the endset temperature ($T_{95\%}$), the temperature at the maximum degradation rate, and the residual weight at 800°C. A commercial PCL (CAPA 6500) was also analyzed to compare the thermal properties with the matrices.

### 2.3.3 | Differential Scanning Calorimetry (DSC)

The thermal properties of PLLA and PLCL matrices were determined on a differential scanning calorimetry (TA Instrument Discovery DSC). At the first scan, samples of 10 mg were cooled at—80°C and heated to 200°C at 10°C min⁻¹. The melting enthalpy ($ΔH_m$) and the melting temperature ($T_m$) were obtained from the first scan. The second scan was made from −80°C to 200°C at 10°C min⁻¹ to determine glass transition temperature ($T_g$). The degree of crystallinity ($\chi_c$) was calculated according to Equation (1), assuming that only lactide units can crystallize.

$$ \chi_c = \frac{100 (\Delta H_m)}{\Delta H_m^0 \cdot w} \qquad (1) $$
where $\Delta H_m^0 = 93 J \cdot g^{-1}$ is the standard melting enthalpy of a 100% crystalline PLLA and $w$ is the lactide mass fraction in the copolymer [35].

### 2.3.4 | Tensile Test

Tensile tests were performed to evaluate the mechanical properties of the composites, specifically the Young's modulus (E), tensile strength ($\sigma_t$), and elongation at break ($\epsilon$).

The experiments were carried out using a universal testing machine (Instron 4466) equipped with a 10 kN load cell. The tests were conducted at a constant crosshead speed of 10 mm·min$^{-1}$. A total of five specimens exhibiting valid failure modes were tested. The dimensions of the specimens were adapted from the ISO 527-4 standard (100 mm × 5 mm × 5 mm).

### 2.3.5 | Three-Point Bending Test

Three-point bending test was used to determine the flexural strength ($\sigma_f$) and the bending modulus (E$_B$). A universal testing machine (Instron 5867) was used, and the specimen was placed on two support anvils and bent through an applied force at 10 mm·min$^{-1}$. Five test specimens with valid failures were tested. The dimensions of the specimens were adapted from ISO 14125 standard with a length of 60 mm, a width of 15 mm, and a thickness of 5 mm.

A cyclic three-point bending test was performed on 30 PLCL/CF composites. A load (5 mm·min$^{-1}$) was applied to the specimens until 5 mm of deformation, followed by an unload at 0.5 mm·min$^{-1}$. This cycle was repeated 5 times, and the shape recovery (Rr) of composites was calculated according to the equation below:

$$R_r = \frac{\Delta U_d}{\Delta L_d}$$

where $\Delta U_d$ is the displacement (mm) measured at unload. $\Delta L_d$ is the displacement (mm) measured at load.

### 2.3.6 | Impact Test

Charpy impact test was carried out according to the ISO 179-1 standard to determine impact strength ($\alpha_{cU}$). Tests were performed with unnotched specimens with dimensions 80 × 15 × 5 mm$^3$ using a 50 J hammer on a pendulum impact testing machine (CEAST 9050—Instron). The specimens were positioned edgewise to the blow, and five specimens were tested. The test specimens for all mechanical tests were prepared with a water jet cutter from the produced plates.

### 2.3.7 | Digital Microscopy

Digital microscopy (Keyence VHX-7000 series) was used to evaluate the quality of the impregnation, the presence of voids, and the morphology of the fractures in the specimens tested in mechanical tests. The morphology of the fractures was evaluated

directly in the specimens after mechanical tests, while the quality of the impregnation and the presence of voids were determined in mounted samples. Cross sections between 25 × 5 × 5 mm$^3$ and 10 × 5 × 5 mm$^3$ were cut from 120 × 120 × 5 mm$^3$ composite plates using a bandsaw.

The cross sections were cold mounted into a 30 mm diameter mold, ground with abrasive papers, flushed with water (P80, P180, P600, P1200), and polished with silicon carbide papers (800/P2400 and 1200/P400).

## 3 | Results and Discussion

### 3.1 | Thermal Behavior of the Matrices

The thermal behavior of the composite matrices was evaluated by TGA. The degradation of the PLLA matrix begins at 237°C, with the maximum degradation rate occurring at 294°C. Regarding the PLCL matrices, the degradation temperature increases with the rise in $\epsilon$-CL content (Figure 1a). The $T_{5\%}$ increases by 2°C, 7°C, and 25°C for 10 PLCL, 20 PLCL, and 30 PLCL matrices, respectively, compared to the pure PLLA matrix.

Compared to the maximum DTG of the PLLA matrix, the DTG of 10 PLCL, 20 PLCL, and 30 PLCL was increased by 4°C, 10°C and 63°C, respectively. The pure commercial poly($\epsilon$-caprolactone) (PCL) displays $T_{5\%}$ at 377°C. Similarly, $T_{95\%}$ increases from 314°C for PLLA to 406°C for the 30 PLCL matrix. Overall, as the $\epsilon$-CL content increases, PLCL matrices show a slower degradation rate and a higher maximum DTG. Specifically, the 10 PLCL matrix degrades at a rate of 1.84% weight·°C$^{-1}$·min$^{-1}$, with a DTG max of 298°C, whereas the 30 PLCL matrix degrades more slowly at 0.78% weight·°C$^{-1}$·min$^{-1}$, with a DTG max of 358°C (Figure 1b).

These results corroborate earlier studies, indicating that PCL exhibits greater resistance to thermal degradation compared to PLLA and that their copolymers exhibit intermediate thermal stability characteristics. The $T_{95\%}$ values for 10 PLCL, 20 PLCL, and 30 PLCL were reported as 336°C, 347°C, and 406°C, respectively. Additionally, these works also reported a decrease in the degradation rate with an increasing $\epsilon$-CL content in the copolymer [29].

A progressive decrease in the glass transition temperature as the $\epsilon$-CL content increases in PLCL is observed (Figure 2). PLLA exhibits the highest $T_g$ around 52°C, which drops down to ~40°C for 10 PLCL, ~30°C for 20 PLCL, and ~20°C for 30 PLCL (Tables S2 and S3 in the Supporting Information). This trend confirms that $\epsilon$-CL units increase the flexibility of the polymer chains, reducing intermolecular interactions and lowering $T_g$.

A similar effect is observed in the melting temperature ($T_m$) and crystallinity. Pure PLLA matrix displays a $T_m$ of 169°C and a crystallinity of around 50%. With the increasing $\epsilon$-CL content in the copolymer, $T_m$ decreases down to 143°C for 10 PLCL; a small melting peak was detected for some 20 PLCLC matrices, and no melting peak for 30 PLCL matrices, indicating a loss of crystallinity. This suggests that $\epsilon$-CL units disrupt the crystalline
<figure><img src="image_2.png" /><figcaption>**FIGURE 1** | (a) TGA curves and (b) normalized DTG curves of PLLA, 10 PLCL, 20 PLCL, 30 PLCL, and PCL matrices.</figcaption></figure>

<figure><img src="image_3.png" /><figcaption>**FIGURE 2** | DSC curves of PLLA, 10 PLCL, 20 PLCL, and 30 PLCL matrices at 1st and 2nd scans.</figcaption></figure>

structure of PLLA, leading to fully amorphous copolymers at higher ε-CL contents.

In the first scan, PLLA and 10 PLCL showed a double peak, which can be attributed to the mixture of less perfect (α') and more perfect (α) lamellae (Figure 2). During recrystallization, PLLA recrystallizes more uniformly under cooling, resulting in a single dominant melting endotherm peak from the melt/reorganization. In contrast, 10 PLCL retains a double peak because ε-CL units act as defects, leading to a more heterogeneous and partly defective lamellar structure and chain sequence distribution. This behavior is commonly reported for PLCL and copolymers [29, 30, 36, 37].

## 3.2 | Voids: Formation and Characteristics

The increase of ε-CL content in the PLLA-based copolymer did not impact the impregnation quality of the glass fibers by the matrix, since no micro-voids were observed between the

fiber and the matrix in all composites on optical micrographs (Figure 3). In addition, the absence of debonding in the fiber-matrix interface proves the optimal impregnation of the fiber by the matrix, indicating that additional surface treatments of the glass fiber are unnecessary to improve the impregnation quality in this case. Generally, high resin velocity enables the formation of micro-voids as the flow in the channel between the bundles is faster than in between the fibers in a bundle, due to the capillarity effect [38]. In this way, the resin flow with 1.5 bar of N₂ during injection is appropriate to avoid micro-voids in these composites.

The formation of meso and micro-voids is controlled by the competition between the resin viscous flow and capillarity flow. The hydrodynamic forces between the bundles control the resin viscous flow, while the surface tension of the fibers drives the capillarity flow [39]. The dominant factor driving the resin flow between the bundles is the resin viscous flow, whereas the capillarity flow drives the resin flow between the fibers.
<figure><img src="image_4.png" /><figcaption>**FIGURE 3** | Optical micrographs of the composite containing 20% of ε-CL showing the absence of micro-voids between the fibers.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**FIGURE 4** | Optical micrograph of composite plate containing 20% of ε-CL showing meso-voids between fiber bundles.</figcaption></figure>

A compromise between these forces must be established to obtain composites with a minimum voids content. Meso-voids can also appear due to air entrapment arising from inhomogeneous glass fabric architecture.

Meso-voids present between the fiber bundles were observed in all composites (Figure 4). These types of voids can be formed due to gaps between fiber tows, fiber waviness, local resin-rich regions, or fiber misalignment [39, 40]. The presence of these types of voids is undesirable since they can create local stress concentrations and contribute to delamination or interlaminar failure, affecting the composite stiffness. The presence of meso-voids in the PLLA/GF and PLCL/GF composites could be attributed to fiber waviness and misalignment. The stacking of ten layers of glass fabric could provoke fiber misalignment, which contributes to creating empty zones between the bundles. These imperfections can cause local deviations in fiber orientation, forming gaps or empty zones between adjacent fiber bundles. These gaps tend to be filled with excess resin, creating resin-rich areas. Furthermore, meso-voids can appear due to the intrinsic nature of woven glass fabrics, where the crimp in the weave pattern introduces undulations, and so promotes the formation of voids at the mesoscale. The resin could not fully impregnate

areas where fiber undulations overlap, leaving pockets of air entrapped during the manufacturing process.

## 3.3 | Tensile Test

The tensile properties of the PLLA/GF and PLCL/GF composites exhibit a marked tendency governed by the ε-CL content in the matrix. The Young modulus shows a small variation between PLLA/GF and 10 PLCL/GF and 20 PLCL/GF, suggesting that these quantities do not significantly affect stiffness (Figure 5 and Table S2 in the Supporting Information). However, with 30% of ε-CL in the matrix, the modulus drops sharply to nearly half that of the other compositions, indicating a substantial softening effect due to the increased flexibility of the matrix.

Tensile strength shows a similar tendency, reaching its peak at 10% of ε-CL (275 MPa) but decreasing progressively with higher ε-CL content, with a significant decrease at 30%. Elongation at break remains relatively stable, indicating that ε-CL incorporation does not significantly enhance ductility. Thus, the presence of glass fiber limits the elongation at break. The decrease observed in glass transition temperature and crystallinity with increasing ε-CL content highlights that the presence of the co-monomer plasticizes PLLA, leading to a softer material. Another factor that could justify the decrease in stiffness is the decrease in the weight-average molecular weight (M~w~) of the matrices when increasing the ε-CL content from about 85,000 for 10 PLCL down to 40,000 g/mol for 30 PLCL (Table S1 in the Supporting Information). This decrease of M~w~ can be attributed to secondary reactions of intramolecular reactions leading to cyclic polymer structures with different hydrodynamic radii [29, 41].

Following the tensile test, the tested specimens were analyzed by digital microscopy to determine the type of failure and the morphology of the fracture. PLLA/GF, 10 PLCL/GF, and 20 PLCL/GF specimens broke into two parts, whereas the 30 PLCL/GF specimen remained in one piece after testing (Figure S3 in the Supporting Information).

All specimens showed multiple fractures after the tensile test. The specimens PLLA/GF and 10 PLCL/GF exhibit interlaminar delamination, intralaminar matrix cracks, and fiber pull-out. As
<figure><img src="image_6.png" /><figcaption>**FIGURE 5** | Tensile properties of PLLA/GF and PLCL/GF composites.</figcaption></figure>

expected, the PLLA/GF specimen shows a brittle character, due to the low elongation at break of the PLLA matrix and the reinforcement, which favors the delamination as well as the fiber pull-out (Figure 6a). The addition of 10% of ε-CL was not enough to reduce the brittleness of the matrix since the 10 PLCL/GF specimen displays the same types of failures and similar values of Young's modulus and tensile strength as the PLLA/GF (Figure 6b).

The 20 PLCL/GF did not show delamination fracture (Figure 6c). However, a translaminar crack, an intralaminar matrix crack, a translaminar crack, and fiber pull-out are observed. The plasticization of the matrix caused by the presence of 20% ε-CL seems to avoid a premature fracture by delamination.

The 30 PLCL/GF composite exhibited the most favorable performance, with less extended failure compared to other composites.

The primary failure observed was an intralaminar matrix crack (Figure 6d). A minor fiber pull-out is observable at the specimen base, and translaminar cracks are also present. However, the last two specimens exhibited a negligible impact on the global failure, as evidenced by the 30 PLCL/GF specimen's inability to separate into two distinct parts following the tensile test, attributable to the cohesion of translaminar fibers. The intralaminar matrix exhibits signs of cracking throughout the specimen's longitudinal axis (Figure 7). Conversely, the translaminar crack and fiber pull-out are observed at specific points. The specimen's failure was attributed to a cumulative intralaminar matrix crack.

Despite the plasticizing effect of ε-CL, the tensile behavior is still primarily governed by the fibers. Failure occurs via intralaminar matrix cracking, interlaminar delamination, and fiber pull-out. An increase in ε-CL reduces early delamination and promotes a transition toward intralaminar matrix cracking, resulting in enhanced post-fracture cohesion at 30% ε-CL.
<figure><img src="image_7.png" /><figcaption>**FIGURE 6** | Optical micrographs after tensile test of PLLA (a), PLCL 10 (b), PLCL 20 (c), and PLCL 30 (d) composite plates cut longitudinally.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>**FIGURE 7** | Detailed micrographs of 30 PLCL/GF specimens showing (a) different types of fracture and (b) intralaminar matrix crack.</figcaption></figure>

## 3.4 | Three-Point Bending Test

As observed for the tensile properties, the bending properties of the composites show a significant dependence on the ε-CL content (Figure 9 and Table S3 in the Supporting Information). The bending modulus decreases drastically as ε-CL content increases, from 10.34 GPa in PLLA/GF to 0.30 GPa in 30 PLCL/GF (Figure 8). This confirms that higher ε-CL fractions make the material significantly more flexible, reducing its stiffness in bending. Similarly, bending strength follows the same trend, peaking at 273 MPa for PLLA/GF and then dropping to 21 MPa for 30 PLCL/GF, indicating a severe loss of mechanical resistance as ε-CL content increases (Figure 8).

In contrast, bending strain exhibits the opposite behavior, remaining relatively low for PLLA/GF and 10 PLCL/GF but increasing abruptly for 20 PLCL/GF and 30 PLCL/GF, both around

22%. This suggests that as the material becomes more flexible, it can undergo much larger deformations before failure. The decrease in glass transition temperature and crystallinity with increasing ε-CL content further supports this trend, as lower $T_g$ and crystallinity correlate with reduced stiffness and strength. In addition, the decrease in molecular weight helps to explain the decrease in stiffness and strength.

While the general trend (decreasing stiffness and strength with increasing ε-CL content) is consistent between bending and tensile tests, the magnitude of change differs, with bending properties showing a more drastic loss in stiffness and a greater increase in strain at high ε-CL content. This suggests that ε-CL incorporation affects flexural performances more than tensile performances, probably due to differences in stress distribution and failure mechanisms in bending versus tensile loading.
<figure><img src="image_9.png" /><figcaption>**FIGURE 8** | Bending properties of PLLA/GF and PLCL/GF composites.</figcaption></figure>

After the bending test, the specimens were analyzed by digital microscopy to determine the type of failure and the morphology of the fracture. The specimens of PLLA/GF and 10 PLCL/GF kept a V shape after the test, while the 20 PLCL/GF and 30 PLCL/GF returned to their initial shape. The PLLA/GF specimens showed different types of failure: compressive fracture (Figure 9a), tensile fracture of fiber (Figure 9b), and tensile fracture (Figure 9c). All 10 PLCL/GF specimens showed tensile fracture and, in some cases, tensile fracture of the fiber. These specimens showed the same V shape as the PLLA/GF specimen after the bending test (Supporting Information).

No failure was observed in the 20 PLCL/GF and 30 PLCL/GF specimens. In the 20 PLCL/GF specimens, only compressive damage centered in the first four outer layers is observed, visible to the naked eye. This damage was caused by the central loading member. The micrograph of the cross-section of the specimen

shows the damage in longitudinal fibers, where these became whitish (Figure 10).

Three-point bending creates a tension-compression gradient but also interlaminar shear, making the response sensitive to both the matrix and interface. As $\epsilon$-CL increases, the external plies undergo early compressive yielding and shear, producing a pronounced drop in bending modulus and strength. Microscopy shows compressive damage confined to external layers at 20% $\epsilon$-CL and no catastrophic failure at 30% $\epsilon$-CL with shape recovery, explaining the more substantial effect of $\epsilon$-CL content on bending than on tensile properties.

After the bending test, the 30 PLCL/GF composite nearly returned to its original shape before deformation. This behavior was anticipated since poly(L-lactide-co-$\epsilon$-caprolactone) is recognized for its shape memory property [42–46]. However, the
<figure><img src="image_10.png" /><figcaption>**FIGURE 9** | Digital micrographs showing the different types of failure observed in the PLLA/GF specimens.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>**FIGURE 10** | Digital micrographs of the (a) 20 PLCL/GF and (b) 30 PLCL/GF specimens.</figcaption></figure>

significant deformation of these composites could lead to damage to the reinforcement, which was not observed in the 30 PLCL/GF composites (Figure 10b). To assess the shape recovery properties of this composite, cyclic three-point bending tests were performed.

In the first cycle, a recovery of 72.9% was observed. This value decreases throughout the cycles, reaching 61.5% of recovery at the last cycle (Table 1 and Figure S4 in the Supporting Information).

The ability of PLCL copolymers to regain their shape is mostly attributed to their semi-crystalline character, but also due to the $T_g$ being close to room temperature. The crystalline regions, created by L-LA, act as anchors that hold the material together. At the same time, the amorphous part made of ε-CL allows the material to stretch and return to its original shape [42]. Whereas the 30 PLCL matrix is completely amorphous, due to its statistical copolymer character, the shape recovery effect was still observed. This is due to the fact that shape recovery effects are not
<table><caption>**TABLE 1** | Shape recovery value for 30 PLCL/GF composite submitted to cyclic three-point bending test.</caption><thead><tr><th>Sample</th><th>Cycle</th><th>ΔL<sub>d</sub><br>(mm)</th><th>ΔU<sub>d</sub><br>(mm)</th><th>Recovery<br>(%)</th></tr></thead><tbody><tr><td rowspan="5">30 PLCL/<br>GF</td><td>1</td><td>5.32</td><td>3.88</td><td>72.9</td></tr><tr><td>2</td><td>5.32</td><td>3.64</td><td>68.4</td></tr><tr><td>3</td><td>5.32</td><td>3.48</td><td>65.4</td></tr><tr><td>4</td><td>5.32</td><td>3.38</td><td>63.5</td></tr><tr><td>5</td><td>5.32</td><td>3.27</td><td>61.5</td></tr></tbody></table>

<figure><img src="image_12.png" /><figcaption>**FIGURE 11** | Impact strength values of PLLA/GF, 10 PLCL/GF, 20 PLCL/GF, and 30 PLCL/GF composite.</figcaption></figure>

exclusively dependent on crystallinity. In fully amorphous polymers, shape memory can also arise from the entropic elasticity of the polymer network and the frozen-in stresses within the amorphous matrix [47–49]. When the PLCL is deformed during the bending test, polymer chains are temporarily stretched and aligned. After the stress relief, the chains relax into a random coil structure, leading to the return of the initial form. The amorphous PLCL, particularly rich in flexible ε-CL segments, can store elastic energy during deformation and release it during recovery, leading to shape memory behavior even in the absence of crystalline domains [42].

### 3.5 | Impact Test

PLLA is known for its low-impact strength (47 kJ/m²) due to its brittle character, and manufacturing processes can drastically decrease the impact strength of PLLA. According to previous studies, it has been proven that its impact strength decreases down to 12 kJ/m² after compression molding [50]. A substantial increase in impact strength was observed for all composites compared to the neat PLLA and PLLA/GF composite (Figure 11). This increase is caused by the presence of glass fiber and by the presence of ε-CL in the PLCL matrix. Due to the high stiffness of

glass fibers, they can fracture and absorb a substantial amount of energy, dissipating the impact force. The impact strength is closely related to the $T_g$ and is therefore directly influenced by the ε-CL content. When the $T_g$ is close to room temperature, very high impact strengths are measured [50]. The increased chain mobility caused by the presence of ε-CL segments in the PLCL copolymer structure improves the flexibility of the PLCL composites, allowing for better energy dissipation during the impact.

In contrast to the previous tensile and bending strength, the impact strength of PLLA and PLCL composites increases with the increase of ε-CL content. Indeed, an increase in impact strength of about 90% is observed between the PLLA/GF ($\alpha_{Cu}$ = 138.2 ± 8.3 kJ/m²) and the 30 PLCL/GF ($\alpha_{Cu}$ = 262.3 ± 16.4 kJ/m²). After the impact test, the specimens PLLA/GF, 10, and 20 PLCL/GF were divided into a hinge shape, while 30 PLCL/GF kept its initial bar shape (Figure 11).

The failure of damaged composite specimens was evaluated for the front side (impacted) and backside (non-impacted) by digital microscopy. The failure observed after the impact test in PLLA/GF and 10 PLCL/GF specimens was a hinge break with tensile fracture. In both samples, the impacted surfaces (Figure 12a-I,b-I) show damage caused by compression, and the non-impacted surfaces (Figure 12a-II,b-II) are fractured by tension. The 20 PLCL/GF specimens exhibit a partial break with tensile fracture, while the 30 PLCL/GF ones showed a non-break failure with buckling damage and some fiber pullout (Figure 12c-II,d-II).

### 4 | Conclusion

In this study, the mechanical properties of poly(L-lactide-co-ε-caprolactone) reinforced with glass fiber composites (PLCL/GF) produced by TP-RTM, with varying ε-caprolactone ratios in the statistical copolymer PLCL matrix, were evaluated. Tensile, bending, and impact tests were conducted, and the effect of the ε-CL ratio in the matrix on the resulting properties was determined. The most relevant effect on the properties of the introduction of ε-CL in the PLLA matrix on the overall composite was observed in impact properties.

The impact strength of PLCL/GF composites with 30% ε-CL was substantially increased by about 90% when compared to the PLLA/GF composites. These results confirm that incorporating ε-CL in the PLLA chain decreases the brittleness of PLLA. The microscopy analysis reveals the presence of meso-voids between the bundles, but no micro-voids between the glass fibers in a tow were detected. The absence of micro-voids and debonding in the fiber-matrix interface indicates a strong adhesion. This shows that the studied reactive thermoplastic system provides an optimal fiber-matrix bond without requiring prior surface treatment of the glass fibers, unlike systems using melted polymers as the matrix. Looking at crystallinity ratios along with the mechanical properties allows us to conclude that glass fibers play a more significant role in the mechanical properties than the crystallinity of the matrix. In addition, the elongation at break in the tensile test of composites is governed by glass fibers since the movement of the copolymer chain is blocked by the fibers. Glass transition temperature significantly affects
<figure><img src="image_13.png" /><figcaption>**FIGURE 12** | Digital micrograph showing the failure observed in the impacted surface of PLLA/GF (a-I), 10 PLCL/GF (b-I), 20 PLCL/GF (c-I), 30 PLCL/GF (d-I), and opposed impact surfaces of PLLA/GF (a-II), 10 PLCL/GF (b-II) 20 PLCL/GF (c-II), and 30 PLCL/GF (d-II).</figcaption></figure>

mechanical properties, especially in tensile and bending tests, more than crystallinity. The decrease in T<sub>g</sub> made the material less brittle and so less resistant to bending. The composites with

20% and 30% of ε-CL in the copolymer composition showed elastic deformation during the bending test, and they returned to their initial shape after the test.
The PLCL/GF composites presented in this study could be
used in applications where relatively high-impact absorption
is required, such as sports equipment (helmets, shin guards),
automotive components (dashboards, door trim panels), and
durable goods (suitcase shells). This study represents an ini-
tial strike on the greener production of PLA-based composites
by TP-RTM in one-step synthesis from the co-monomers in
the absence of solvent. The use of biobased and biodegrad-
able matrix is the first approach for the production of greener
composites.

Author Contributions

**Bernard Miranda Campos:** data curation, investigation, writing – original draft, methodology, formal analysis. **Gaëlle Fontaine:** investigation, writing – review and editing, supervision. **Serge Bourbigot:** investigation, writing – review and editing, supervision. **Grégory Stoclet:** investigation, supervision, data curation, writing – review and editing, validation. **Fanny Bonnet:** conceptualization, investigation, funding acquisition, writing – review and editing, supervision, project administration, validation.

Acknowledgments

The authors would like to thank Aurélie Malfait for the SEC analysis, Johan Sarazin, and Pierre Bachelet for their contribution to the TP-RTM process.

Conflicts of Interest

The authors declare no conflicts of interest.

Data Availability Statement

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

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### Supporting Information

Additional supporting information can be found online in the Supporting Information section. **Figure S1:** Schematic representation of TP-RTM machine. **Figure S2:** Aluminum mold used to produce composite plates. **Table S1:** Poly(L-lactide-co-ε-caprolactone)/glass fabric composites analyzed in this study. **Table S2:** Tensile properties, glass transition temperature, and crystallinity of glass fabric reinforced composites. **Table S3:** Bending properties, glass transition temperature and crystallinity of glass fabric reinforced composites. **Table S4:** Impact properties of glass fabric reinforced composites. **Figure S3:** Specimens after tensile test (a) PLLA/GF (b) 10 PLCL/GF (c) 20 PLCL/GF (d) 30 PLCL/GF. **Figure S4:** Strain x stress curve on cyclic three-point bending test of 30 PLCLC/GF.