# Influence of fiber/matrix interface on gas permeability properties of CF/TP composites

Gautier Allusse, Olivier de Almeida, Quentin Govignon, Monica Francesca

Pucci, Fabrice Schmidt

► To cite this version:

Gautier Allusse, Olivier de Almeida, Quentin Govignon, Monica Francesca Pucci, Fabrice Schmidt. Influence of fiber/matrix interface on gas permeability properties of CF/TP composites. Composites Part B: Engineering, 2025, 298, pp.112358. 10.1016/j.compositesb.2025.112358. hal-04993975v2

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# Influence of fiber/matrix interface on gas permeability properties of CF/TP composites

Gautier Allusse<sup>a,†</sup>, Olivier De Almeida<sup>a,∗</sup>, Quentin Govignon<sup>a,†</sup>, Monica Pucci<sup>b</sup>, Fabrice Schmidt<sup>a</sup>

## ARTICLE INFO

Handling editor: Ole Thomsen

**Keywords:**

Thermoplastic

Carbon fiber

Interphase

Hydrogen

Transcrystallinity

Permeability

## ABSTRACT

For hydrogen application, one of the most important material property required is low gas permeability. In composite materials, this property depends on the materials but also on the processing parameters. In particular the residual porosity, but also the quality of the fiber/matrix interface, play a crucial role. This is particularly the case in composites involving a thermoplastic matrix with carbon fibers as the lack of reactive groups on the fiber surface can limit the level of interfacial interactions between the reinforcement and the matrix. In this study, the role of the interface is analyzed through the investigation of the hydrogen permeability of carbon fiber reinforced thermoplastics (CF/PVDF and CF/PPS) using different polymers and carbon fibers. The hydrogen permeability of the composites was measured, and a correlation with the crystallization behavior of the matrix on the fiber surface was identified. Hydrogen permeability decreases when the fiber favors matrix nucleation. Nucleation is improved by increasing the surface roughness of the carbon fiber.

## 1. Introduction

The energy transition is the basis of development of numerous applications based on hydrogen, which has been identified as a necessity to reduce greenhouse gas emissions. Many technologies have been developed or are still being developed, whether for the production, storage, or use of hydrogen. Developments in Polymer Electrolyte Membrane Fuel Cells (PEMFC) are explored to create durable and lightweight solutions for transport applications. PEMFC consist of a stack of Bipolar Plate (BPP), and Membrane Electrode Assembly (MEA).

The MEA is the reaction site where the di-hydrogen is combined with di-oxygen to produce water and electrons. The BPP are employed for convey gases ($H_2$ and $O_2$) to the reaction site as well as circulating coolant fluid [1]. For efficient operation, the bipolar plates must be impermeable to gases, have sufficient electrical conductivity to carry the electrons generated during the electrochemical reactions, and have a good corrosion resistance to the PEMFC environment [2,3].

Thermoplastic composite materials are promising candidates for BPP. Thermoplastic matrices combine lightness, good corrosion and mechanical resistance and exhibit a low permeability to gas [4] which is crucial for hydrogen delivery to the MEA [1]. The addition of carbon

fibers improves the necessary thermal and electrical conductivity properties, as well as the mechanical properties of the composite BPP. However, the effect of fiber addition on the permeability of composite materials to gases [5] and in particular to hydrogen, which is a light gas of low steric hindrance [6], still needs to be investigated. Indeed, although the literature reports that gas permeability properties of thermoset matrix composites are in line with the properties of the matrix [7], very few data on thermoplastic matrix composites are available [8].

Thermoplastics (TP) have intrinsically good gas impermeability properties, of the order of $10^{-10} \text{ mol m}^{-1} \text{ s}^{-1} \text{ MPa}^{-1}$, close to thermosets [7,9]. Carbon and glass reinforcements also exhibit very low gas permeability, lower than $10^{-12} \text{ mol m}^{-1} \text{ s}^{-1} \text{ MPa}^{-1}$ [4,10]. But hydrogen permeation in thermoplastic composites is not necessarily a direct result of matrix and fiber properties. Indeed, interface issues are well known in composites, and reinforcements are generally sized to ease composite manufacturing and ensure good adhesion with the resin, improving mechanical strength [11,12]. For thermosets, the fiber/matrix interface can be well controlled by using chemical sizing. An appropriate sizing provides fiber protection for textile operations while promoting a better wettability and adhesion to the intended matrix. Other strategies have also been explored to modify the surface

<footnote><sup>a</sup> Institut Clément Ader, Université de Toulouse, CNRS, UMR 5312, IMT Mines Albi, UPS, INSA, ISAE-SUPAERO, Campus Jarlard, 81013, Albi, France</footnote>

<footnote><sup>b</sup> LMGC, Université de Montpellier, IMT Mines Ales, CNRS, Ales, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: olivier.dealmeida@mines-albi.fr (O. De Almeida).</footnote>
roughness and the chemical composition of carbon fiber surfaces (plasma treatment, heat treatment) [13,14], and in the end promote fiber/matrix interactions.

However, in the case of thermoplastic matrices, the interfacial interactions at the surface of carbon fibers and the final properties of composites are more difficult to control [15], especially in the case of semi-crystalline thermoplastics where the presence of fibers can affect the crystallization process. Moreover, regarding hydrogen permeability, the role of the fiber/matrix interaction on gas barrier properties has rarely been explored and is still not fully understood.

This study therefore investigates the hydrogen permeability of carbon fiber reinforced composites materials involving semi-crystalline thermoplastic matrices of different nature in order to determine the factors that control the gas barrier properties. Poly(Vinylidene Fluoride) (PVDF) was first chosen for its good intrinsic gas permeability [4]. PVDF is a fluorinated semi-crystalline polymer known for its excellent chemical resistance, dielectric properties, excellent UV and moisture resistance, low coefficient of friction and very good thermal stability between 40 °C and 150 °C [16]. PVDF is particularly suitable for use in fluid storage and transport applications (piping), in membranes for the medical field or in piezoelectric applications [17]. However, because of the high electronegativity of the fluorine element, this polymer exhibits poor chemical interaction with its environment, which limits bond formation with carbon fibers in composite applications. For comparison purpose, Poly(Phenylene Sulfide) (PPS) was selected as this thermos stable thermoplastic has a hydrogen permeability of the same order of magnitude as PVDF [9]. This semi-crystalline thermoplastic polymer is widely used in composites for its high mechanical performance, high temperature stability, fire resistance and high chemical resistance [18], it has also a better compatibility with carbon fibers.

PVDF and PPS polymers were combined with two different CF fibers so as to prepare four thermoplastic composites with very different properties. After checking the impregnation quality and assessing the lack of observable porosity, the hydrogen permeability of the various materials is analyzed. The influence of the fiber/matrix interface being highlighted, thanks to the analysis of the different fiber/matrix pairs, an in-depth analysis of the fibers, the matrices and the interfacial interactions is then exposed. The main factors that control hydrogen permeability in carbon fiber reinforced thermoplastic composites are then discussed based on morphological features and crystallization behavior.

## 2. Experimental

### 2.1. Materials

To compare the influence of the fiber/matrix interface on the gas permeability properties, two grades of PVDF supplied by Arkema and one reference of PPS supplied by Toray were characterized and used to manufacture composites. The PVDF used are in the form of 100 µm thick films, while the PPS is presented as a 60 µm film. The first PVDF grade is a homopolymer (PVDF-H), while the second grade is a maleic anhydride (MAH) grafted PVDF (PVDF-G) for which the grafted function improves the chemical affinity of PVDF with carbon fibers. Studies on similar grafted-PVDF show a decrease in the contact angle of PVDF on carbon fibers, leading to an improvement in wettability as well as in composite interlaminar shear strength, indicating a stronger CF/PVDF interface [12,19]. The characteristics of the matrices are summarized in Table 1.

<table><caption>**Table 1**

Matrix characteristics.</caption><thead><tr><td>Matrix</td><td>Nature</td><td>Film thickness [µm]</td><td>Melting temperature</td></tr></thead><tbody><tr><td>PVDF-H</td><td>Homopolymer</td><td>100</td><td>180 °C</td></tr><tr><td>PVDG-G</td><td>MAH-grafted</td><td>100</td><td>180 °C</td></tr><tr><td>PPS</td><td>Homopolymer</td><td>60</td><td>285 °C</td></tr></tbody></table>

Composite plates were manufactured using the different thermoplastic polymers, in combination with two different carbon reinforcements in the form of non-woven mats with an areal weight of 465 g/m². For confidentiality purposes, references of carbon fiber materials cannot be disclosed. Nevertheless, the two non-woven mats were provided by the same supplier. The first reinforcement (CF-A) is an unsized reinforcement while the second one (CF-B), is a heat-treated version of the first one to modify the carbon fiber surface. This surface activation treatment aims to modify the chemistry and the roughness of the carbon fiber surface [13]. The characteristics of the reinforcements are summarized in Table 2.

### 2.2. Composite manufacturing

The composite plates were manufactured by thermocompression molding through a film stacking method. This method consists of heating and compressing a stack of alternating layers of polymer film and dry reinforcement (Fig. 1).

The conventional thermo-compression processing cycle consisted first in a low-load heat-up phase. The reinforcement/matrix system was then compacted at 10 MPa, once the matrix temperature was above its melting point to promote polymer flow within the fiber reinforcement (220 °C and 305 °C respectively for the PVDF grades and the PPS). The composite was finally cooled under pressure to set the structure of the impregnated composite and demolded when the temperature decreased below 50 °C. The composite plates manufactured are square-shaped, with sides measuring 135 mm and with an equivalent fiber volume ratio of 30 %, leading to a thickness of around 1.2 mm.

As the densities of PVDF (1.78 × 10⁶ g m³ [20]) and CF (1.68 × 10⁶ g m³, measured by pycnometer) materials are very similar, porosity analysis by micro-tomography was dismissed. In addition, PVDF solubilization requires the use of chemicals that are both dangerous to health and incompatible with the equipment available in the laboratory. It was therefore decided to characterize porosity by microscopic observation. The impregnation quality of composite panels was characterized by the presence or absence of porosity within the manufactured material. Samples taken from the plates in identical locations were polished using silicon carbide abrasive paper and then observed under a Zeiss Axio Imager M2M optical microscope.

In this study, the plates manufactured by thermocompression using the film stacking method exhibited a very low rate of residual porosity (<0.3 %). Microscopic observations (examples of which are presented in Fig. 2) confirmed the good impregnation quality of all the composite panels with no visible porosity.

### 2.3. Characterization techniques

#### 2.3.1. Hydrogen permeability

The hydrogen permeability properties of the plates were measured by the pressure differential method using a custom-designed bench (Fig. 3). The test bench is made up of 2 chambers. The measurement is based on Fick's first law. A hydrogen overpressure $p_{H_2}$ of 1 bar was applied on one side of the composite panel. The amount of hydrogen Q in ppm that had passed through the panel of area A of 54.8 cm² and thickness e of the tested material during time t was measured on the other side. Both chambers are equipped with pressure sensors to verify

<table><caption>**Table 2**

Reinforcement characteristics.</caption><thead><tr><td>Reinforcement</td><td>Type</td><td>Areal weight [g/m²]</td><td>Surface traitement</td></tr></thead><tbody><tr><td>CF-A</td><td>PAN based carbon felt</td><td>465</td><td>Not activated</td></tr><tr><td>CF-B</td><td>PAN based carbon felt</td><td>465</td><td>Thermally activated</td></tr></tbody></table>
<figure><img src="image_2.png" /><figcaption>Fig. 1. Thermocompression cycle using the film stacking method.</figcaption></figure>

<figure><img src="image_3.png" /><figcaption>Fig. 2. Example of microscopic observation of a) CF-A/PVDF-H and b) CF-B/PPS composites - black areas correspond to porosity and fiber pull-out due to polishing and represent 0.3 % of the image surface.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>Fig. 3. Schematic drawing of hydrogen permeability measurement and typical H₂ permeation curve.</figcaption></figure>

that there is no leak during the measurement. The tests were performed at ambient temperature and humidity. The permeability Pe coefficient [mol.m⁻¹.s⁻¹.MPa⁻¹] was then calculated with equation (1).

$$P_e = \frac{Qe}{tA p_{H_2}} \qquad (1)$$

Fig. 3 plots the quantity of gas Q permeating through the sample versus time. The permeability coefficient Pe is computed using the slope of the linear part of the curve as described in equation (1). For each fiber/matrix combination, 3 plates were characterized, with one measurement per plate. The uncertainty of the hydrogen measurement is of the order of 20 ppm, for measurements up to 2000 ppm, i.e. 1 % of
measurement uncertainty.

### 2.3.2. Differential scanning calorimetry

DSC measurements were carried out on polymers and composites using a PerkinElmer 8500 DSC equipment to determine the melting temperatures and enthalpies of fusion of matrices. Sealed aluminum sample pans were prepared with 10–20 mg of polymer and composite. The samples were then heated from 50 °C to 240 °C for PVDF and from 50 °C to 320 °C for PPS, at 10 °C/min. The melting temperature was recorded at the maximum of the heat flow peak, and the crystallinity degrees of the polymers and composites were calculated using equation (2):

$$X_c = \left( \frac{\Delta H_m}{\chi_m \times \Delta H_m^0} \right) \times 100 \quad (2)$$

In equation (2), $\chi_m$ and $\Delta H_m$ represent the mass fractions (equal to 100 % in the case of pure polymer) and the enthalpy of fusion of the sample, respectively. $\Delta H_m^0$ is the enthalpy of fusion of a pure polymer crystal, which is 104.6 J g⁻¹ for PVDF [21] and 76.5 J g⁻¹ for PPS [22]. For each matrix and fiber/matrix pair, 3 DSC measurements were performed on different samples.

### 2.3.3. Polarized Optical Microscopy (POM)

Isothermal crystallization from the molten state was studied using Polarized Optical Microscopy (POM) by means of a Linkam THMS-600P hot stage mounted under a Zeiss Axio Imager M2M optical microscope. Thin samples containing single carbon filaments were first prepared by thermocompression molding with two different polymer films (PVDF and PPS) and carbon fibers taken from the reinforcements. The polymers were heated to 200 °C and 295 °C for PVDF and PPS respectively over 3 min to avoid degradation and pressed to obtain samples with a thickness between 40 and 50 µm.

The crystallization study under hot stage was achieved from the molten state by first heating the polymers at 50 °C/min to temperatures of 220 °C and 320 °C for PVDF and PPS respectively. The samples were then held at this temperature for 3 min to remove the thermal history. The samples were finally cooled at a rate of 50 °C/min to the set crystallization temperature (160 °C for PVDF and 260 °C for PPS) and images of the crystallization process were taken at a frequency of 0,03 Hz. The direct observation of spherulitic growth during polymer crystallization under polarized optical microscope allowed identification of the bulk nucleation density and fiber-induced nucleation density. For each fiber/matrix combination, at least 6 measurements have been performed.

### 2.3.4. Scanning electron microscopy analysis of fiber surface

CF surface morphology was observed using scanning electron microscopy (FEG NovaNanoSEM 450) and apparent fiber surface roughness was measured using laser scanning microscopy (3D LEXT OLS5100, Olympus). The apparent surface roughness of the fibers was obtained using the surface roughness module of the Gwyddion software (v2.62). At least 20 fibers from each reinforcement were analyzed. In this case, the apparent surface roughness values are used to perform a comparative study of the impact of heat treatment on carbon fiber topography. To determine the real surface roughness value, a method developed by Jeantet et al. [23] can be used.

### 2.3.5. Contact angle measurements on matrix

The surface tension $\sigma_s$ of PVDF-H, PVDF-G and PPS films at room temperature was evaluated using a goniometer (DSA 30, Germany). By using measured contact angles and two testing liquids with known surface energies, the dispersive $\sigma_s^D$ and polar $\sigma_s^P$ components of polymers can easily be calculated according to Owens–Wendt method [24]. Deionized water ($\sigma = 72.8$ N/m, $\sigma_d = 21.8$ mN/m, $\sigma_p = 51.0$ mN/m) and diiodomethane ($\sigma = 50.8$ mN/m, $\sigma_d = 50.8$ mN/m, $\sigma_p = 0$ mN/m) were

chosen as testing liquids. At least 6 measurements were achieved for each sample, and an average value was used for the surface energy analysis.

### 2.3.6. Atomic force microscopy analysis of fiber/matrix interface

AFM measurements were carried out on zones with an interface between the fiber and the matrix, with the fiber oriented perpendicular to the sample surface. This avoids any influence from the presence of the fiber below the indentation.

The setup used in this study was an MFP3D Infinity from Asylum Research (Oxford Instruments). Images of 256x256 pixels were generated by means of AFM scans with a size varying from 20 × 20 µm² (as shown in Fig. 5) up to 3 × 3 µm² (as in Fig. 6). A silicon probe (AC55TS-R3 from Olympus) with a tetrahedral tip at the free end of the cantilever, a resonance frequency of ~1600 kHz and a spring constant of ~85 N/m was used for these tests. A Fast Force Mapping (FFM) mode was set. It consists in a contact mode technique in which the probe is oscillated while in contact with the sample surface. For each pixel, a force curve is recorded and analyzed to extract information like surface topography and mechanical properties. In this study, the surface topography was obtained as well as the local contact modulus using Hertz theory [25]. It is important to observe that, in order to obtain quantitative mechanical results, the tip radius parameter should be adjusted by calibrating the probe with a suitable substrate. For this preliminary study, no calibration was made, and qualitative results are shown.

## 3. Results

### 3.1. Gas permeability

Fig. 4 presents the hydrogen permeability measurements of the pure polymers (PVDF-H, PVDF-G and PPS) and their respective composites with the two carbon fiber references (CF-A and CF-B). The pure polymers have a low hydrogen permeability in accordance with data from literature [9]. The results also show that PVDF has a slightly lower permeability than PPS which confirms the renewed interest in PVDF for hydrogen applications.

After coupling the polymers with the CF-A reinforcement, the composites show a strong increase in hydrogen permeability of several orders of magnitude. As expected, the CF-A/PVDF-G composite has a slightly lower hydrogen permeability than the CF-A/PVDF-H composite because the addition of the grafted functions improved the quality of the fiber/matrix interface. However, the decrease of hydrogen permeability

<figure><img src="image_5.png" /><figcaption>Fig. 4. Hydrogen permeability of polymers (PPS, PVDF-H and PVDF-G) and their corresponding composites with CF-A and CF-B reinforcements.</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>Fig. 5. AFM cartography of CF-A/PVDF-H obtained in FFM mode: a) Height mapping and b) Local modulus of elasticity.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>Fig. 6. Local AFM cartography (FFM mode) of a fiber/matrix area in CF-A/PVDF-H: a) Height mapping, b) Distribution of local elastic modulus, c) Elastic modulus along profiles plotted in b).</figcaption></figure>

induced by the addition of MAH grafting is less significant than the effect of PPS matrix.

These observations are confirmed by the hydrogen permeability measurements performed on the CF-B/PVDF-H, CF-B/PVDF-G and CF-B/PPS composites (Fig. 4). These composites, involving the thermally activated reinforcement CF-B, exhibit a lower hydrogen permeability than in the case of CF-A reinforcement. However, the coupled effect of the grafted functions in PVDF-G with the thermally activated reinforcement CF-B decreases the gas permeability properties of the CF/PVDF composites to a higher extend than grafting or fiber activation individually. Finally, the combination of PPS with CF-B fiber recovers and even slightly lowers hydrogen permeability properties compared to neat PPS, indicating a synergy between the two constituents.

As shown previously, these results cannot be attributed to an effect of the void content since all composite panels showed high impregnation quality with very low porosities. And since fibers and polymer matrices individually have lower permeability coefficients, differences in gas permeability behavior are necessarily linked to the fiber/matrix

interface: either this interface provides a favorable path for gas to pass through, or bulk material properties are modified by the interfaces. To understand the differences in the hydrogen permeability behavior of composites, the matrix, fiber, as well as the interface and interphase were investigated.

## 3.2. Matrix crystallization

The degree of crystallinity of the semi-crystalline matrices is the first property that has been investigated as it can affect the permeability of TP composites. Indeed, Monson et al. [26] showed that a higher proportion of crystalline phase leads to an improvement in gas impermeability. As in the case of fillers, the presence of crystalline phases increases the effective path length of gases through the composites, i.e. the tortuosity of diffusion path. The semi-crystalline microstructure of the bulk polymer and of the composite matrices were therefore analyzed, since the presence of fibers may have led to different degrees of crystallinity.
After checking that the cooling rate (from 1 °C.min⁻¹ to 100 °C. min⁻¹) had a low influence on matrix crystallization, the degree of crystallinity of the different materials (neat polymers, CF-A and CF-B reinforced composites) were measured by DSC during the heating phase, subsequent to a cooling step at 10 °C/min. Mean values of 3 experiments per material and related standard deviations are detailed in Table 3. Results show that the crystallinity of the semi-crystalline polymers is similar and comparable, and that the addition of CF-A and CF-B fibers does not significantly affect the crystallinity of the polymer matrices in the composites.

With degrees of crystallinity ranging from 47.7 % to 54.1 %, differences of hydrogen permeability behavior cannot be directly correlated to the amount of crystals within the polymer matrices in composite materials. The fact that the CF-B/PPS sample has the lowest crystallinity in the matrix, but also the lowest hydrogen permeability, effectively confirms that the degree of crystallinity is not a first-order hydrogen permeability factor for these composite materials.

### 3.3. Interphase investigation

Since the degree of crystallinity of the matrices is not a liable reason of permeability differences between the composite materials, the local properties of the interface was investigated. To deeply investigate the local variations of matrix density close to the fiber, nanomechanical characterizations using Atomic Force Microscopy (AFM) were performed. In this preliminary exploration, only samples taken from CF-A/ PVDF-H composites were studied.

Fig. 5 shows the topography and the local elastic modulus of the composite on a cross-sectional area of a plate. Three fibers can be clearly identified in the height mapping (a) as well as in the cartography of elastic modulus (b) thanks to their greater stiffness relative to the matrix. This large difference in elastic behavior makes it easy to identify the presence of fiber on the AFM cartography.

Thanks to this nanoscopic analysis technique, the local properties in the vicinity of a fiber could be investigated so as to explore the transition zone between the matrix and the fiber in the CF-A/PVDF-H (Fig. 6). As shown in Fig. 6b, the elastic modulus mapping reveals two main zones, namely the polymer with low modulus and the carbon fiber with higher modulus properties. These distinct moduli are also well defined in Fig. 6c along three horizontal profiles defined in Fig. 6b. Both phases exhibit well-defined moduli and the change in elastic modulus in the transition zone between fiber and matrix is characterized by a stiffness gradient over a length of around 0.26 µm. This gradient length is consistent with measurements performed on an acrylic matrix composite material reinforced with glass fibers, although it is to be noticed that the acrylic matrix (Elium® resin) is an amorphous polymer [27].

This gradient indicates that the transition zone does not exhibit a lower density than that of the two constituents, which could suggest the presence of voids at the fiber/matrix interface. Therefore, although the CF/PVDF-H composite exhibits the highest hydrogen permeability, its permeability behavior does not appear to result from a lack of contact or bonding between matrix and fibers.

<caption>**Table 3**</caption>

Degree of crystallinity of polymer and composite materials resulting from a cooling phase at 10 °C.min⁻¹ and determined upon heating by DSC at 10 °C. min⁻¹.

<table><thead><tr><th>Crystallinity degree<br>[%]</th><th>Pure polymer</th><th>Composite CF-A/<br>TP</th><th>Composite CF-B/<br>TP</th></tr></thead><tbody><tr><td>PVDF-H</td><td>53.49 ± 0.88</td><td>49.36 ± 0.68</td><td>52.34 ± 0.48</td></tr><tr><td>PVDF-G</td><td>54.17 ± 0.44</td><td>52.81 ± 2.86</td><td>50.35 ± 3.31</td></tr><tr><td>PPS</td><td>51.90 ± 0.68</td><td>53.21 ± 2.26</td><td>47.74 ± 1.24</td></tr></tbody></table>

### 3.4. Crystalline phase organization

In the absence of a low-properties zone close to the fibers, the existence of a interphase around the fibers with distinct characteristics was examined and in particular the organization of the crystalline phase. At the microscopic scale, the crystallization of semi-crystalline thermoplastics in composites is indeed influenced by the presence of the reinforcement. The fiber nature, its surface roughness, its thermal conductivity or the CF surface treatment can actually induce the formation of a transcrystalline microstructure around the fibers [28], and accelerate crystal nucleation kinetics [29].

To highlight the CF effect on crystallization, the crystallization process of representative samples with an isolated fiber was investigated using the POM technique, and numerous pictures were taken during the spherulite growth process. Nucleation kinetics are usually determined by direct counting of nuclei per unit volume for bulk crystallization and per unit fiber area for fiber-induced nucleation. Unfortunately, this method can only be used for relatively low matrix nucleation densities and large spherulite sizes. Here, the spherulites associated with PPS nucleation are too small to be counted (Fig. 7).

Therefore, the choice was made in this study to switch to an analysis of the image surface covered with spherulites in distinguishing the area around the fiber from the rest of the images. The CF influence was assessed by measuring the percentage of an area close to the fiber covered by spherulites. To define the size of this zone, it was assumed that the carbon fiber could only influence crystallization within a distance smaller than the size of the spherulites. The distance around the fiber was arbitrary defined as the size of a spherulite at 10 % of image coverage in bulk zone. At this stage of crystallization, the size of spherulite can be easily measured and is representative of half the size of spherulites at the end of crystallization. In this case, a distance of 25 µm around the fiber was used for the PVDF matrix, while a distance of 11 µm was used for PPS. The surface coverage in the rest of the image, i.e. in an area far from the fiber, was also calculated, representing the bulk crystallization process. The image coverage in bulk and fiber zones as a function of acquisition time are plotted in Fig. 8b along with the ratio of fiber to bulk coverage. As it can be observed in Fig. 8b, the ratio of fiber to bulk coverage stabilizes when the curves of image fraction covered by spherulites in the two zones intercept. This particular point corresponds to a transition between a nucleation-controlled crystallization and a spherulitic growth regime. The mean ratio of fiber to bulk coverage beyond this intercept was thus used as an indicator of fiber influence and

<figure><img src="image_8.png" /><figcaption>**Fig. 7.** Observation of growing spherulites using the POM technique (CF-B/ PPS, t = 13 min).</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Fig. 8.** Influence of the fiber zone on the crystallized surface of CF-B/PVDF-G sample: a) Observation of growing spherulites by POM; b) Ratio of fiber to bulk coverage and image surface fraction covered as a function of acquisition time.</figcaption></figure>

values collected for all material combinations are reported in Fig. 7. It is to be noted that the distance from the fiber that defined the fiber zone had a low influence on the results displayed in Fig. 9.

A value close to unity demonstrates an absence of fiber influence on the crystallization process. A value greater than one indicates that the fiber has a nucleating effect on the matrix, while a value lower than one indicates that the fiber has a negative effect on polymer nucleation at the interface.

As shown by ratios lower than unity in Fig. 9, the CF-A reinforcement has a low influence on the PVDF crystallization and even a negative effect. On the other hand, the presence of this fiber favors PPS crystallization, as demonstrated by the greater nucleation effect on the fiber surface. By changing to CF-B with an activated surface, the crystallized surface ratio increases for all matrices, indicating a higher surface-induced nucleation effect of the CF-B surface. In the case of PPS matrix, both fibers favor crystal nucleation, but with a more significant effect of CF-B. On the contrary, for PVDF, the fiber influence on crystal nucleation can be positive or negative, which suggest that this matrix is highly influenced by the fiber surface properties.

Although fiber surface nucleation does not significantly modify the degree of crystallinity (Table 3), a direct correlation can be identified between surface-induced nucleation density (Fig. 9) and permeability properties displayed in Fig. 4. The use of CF-B fiber improves the gas barrier property of composites whatever the matrix involved, just as it increases surface nucleation. The existence of a fiber-induced crystalline interphase, whether it generates a transcrystalline or a shish-calabash structure (Fig. 7), thus seems to coincide with the level of hydrogen permeability. This correlation also seems to be confirmed in the case of

the PPS matrix. Despite its lower nucleating ability, CF-A still generates surface nucleation of PPS matrix (ratio of fiber to bulk coverage above unity). This nucleating ability of CF-A fiber on PPS could therefore be the reason for the lower permeability increase when CF-A fibers are added to PPS (Fig. 4).

This relationship between fiber nucleating ability and barrier properties suggests that the organization of the crystalline lamellae around the fibers may exhibit a specific behavior responsible for permeability properties. Indeed, it has already been shown that gas permeation in polymeric materials depends on crystal tortuosity or on the behavior of the amorphous phase in the vicinity of the crystalline lamellae (Rigid Amorphous Fraction) [30]. Hence, the oriented lamellar organization resulting from nucleation on the surface of the fibers is likely to influence the gas path in the vicinity of the fibers. However, the results obtained do not demonstrate that it is this crystalline microstructure around the fibers that exhibits particular barrier properties. The presence of a specific crystalline organization may simply be indicative of a particular behavior of the interface at which the gas could permeate.

Since crystal nucleation on fiber surface is related to permeability properties, key factors influencing surface nucleation may also give new insights into the origin of gas barrier properties of semi-crystalline thermoplastic matrix composites. In this study, two main characteristics were analyzed, the fiber surface microstructure and the surface tension of the matrices [29].

### 3.5. Carbon fiber surface

Fig. 10 shows scanning electron microscopy images of the carbon fibers from the two carbon reinforcements. On a macroscopic scale (a and c), the surface of the two carbon fibers remains comparable: both fibers have a rough surface, characterized by a fibrillar structure that can be related to the structure of the PAN precursor and the process to obtain carbon fiber. However, at nanoscale and high magnification (b and d), numerous pores appear on the surface of activated CF-B fiber. These pores result from the thermal treatment applied to the fibers that modified the surface properties [31].

The surface roughness induced by the presence of these pores was measured by laser scanning microscopy. It increases from 0.681 ± 0.240 µm for CF-A fiber to 1.219 ± 0.461 µm for CF-B fiber. To attest this observation, specific surface areas were determined by the BET method using N₂ at 77 K with an ASAP 2460 Micromeritics apparatus. The specific surface area of the fiber is increased by the presence of nano-pores, rising from 1.2 m²/g for CF-A fiber to 1.5 m²/g for CF-B fiber.

As reported by Wang et al. [29], roughness induces an increase in overall fiber surface area, which may be responsible for greater fiber nucleation efficiency. Moreover, roughness can also induce thermal stress at the fiber surface which in turn can induce local orientation of matrix chain segments and thus a lower free-energy barrier for nucleation. The greater nucleation ability of CF-B relative to CF-A whatever

<figure><img src="image_10.png" /><figcaption>**Fig. 9.** Ratio of the fiber to bulk coverage for different combinations of reinforcements and matrices.</figcaption></figure>
<figure><img src="image_11.png" /><figcaption>Fig. 10. SEM observation of the surface morphology of a) CF-A (x 10 000); b) CF-A (x 50 000); c) CF-B (x 10 000); d) CF-B (x 50 000).</figcaption></figure>

the matrix can therefore be directly related to its greater roughness. Though, in addition to changing the CF morphology, the heat treatment applied to CF for activation also modifies the chemical functions present on the fiber [13,31,32]. This could have increased the fiber surface energy [33] and contributed to the nucleation ability of CF-B.

### 3.6. Matrix surface tension

The surface tension $\sigma_s$ of PVDF-H, PVDF-G and PPS films at room temperature was evaluated using a goniometer. This method was only applied to the polymeric matrices as it is not suitable for fiber surface energy determination [34]. Although standard deviations are large, in particular in the case of PVDF-H and most probably due to film inhomogeneity which led to drop instability, significant differences between the polymer films could be identified (Table 4).

The surface free energy measurements indicate a higher surface tension for PPS, along with the highest dispersive component. The barrier energy for the formation of PPS nuclei is therefore much lower than for PVDF grades. In the case of PVDF, both polymers have similar surface tension, but the addition of the MAH functional group increases the dispersive component of surface tension, to the detriment of the polar component. The correlation of the surface tension data with the carbon fiber surface nucleation ability then suggests that the dispersive contribution is detrimental for nucleation onto carbon fibers. The same correlation can be observed for hydrogen permeability, with the higher

<table><caption>**Table 4**

Surface free energy of polymer matrices and corresponding polar and dispersive components.</caption><thead><tr><th>Matrix</th><th>σ<sub>s</sub><sup>D</sup> [mN/m]</th><th>σ<sub>s</sub><sup>B</sup> [mN/m]</th><th>σ<sub>s</sub> [mN/m]</th></tr></thead><tbody><tr><td>PVDF-H</td><td>27.77 ± 2.76</td><td>8.64 ± 3.20</td><td>36.41 ± 5.95</td></tr><tr><td>PVDF-G</td><td>32.42 ± 0.57</td><td>4.20 ± 0.25</td><td>36.62 ± 0.82</td></tr><tr><td>PPS</td><td>40.21 ± 2.55</td><td>4.10 ± 0.17</td><td>44.31 ± 2.45</td></tr></tbody></table>

the surface tension, the greater the barrier effect, again with a potentially non-negligible contribution from the dispersive component.

### 4. Discussion

The topological analysis of carbon fibers and the surface tension measurements of polymer matrices are perfectly in line with the literature on the conditions for the development of a transcrystalline or shish-calabash phase [29]. PPS with its high surface tension, and to some extent PVDF-G functionalized with MAH groups, are matrices that favor nucleation on the surface of the fibers, and this seems to contribute well to improving the barrier performance of composites. On the other hand, the surface topology of the carbon fibers is also a crucial factor, not only for crystallization nucleation, but above all for the reduction in permeability as demonstrated by the benefit of using the rougher CF-B fiber.

Since the addition of fiber to polymer matrices does not significantly alter their crystallinity, the higher hydrogen permeability of composites compared to matrices alone supports the hypothesis that the properties of the fiber-matrix interface are crucial for gas barrier properties. If hydrogen permeation occurs at the interface, then the presence of nanopores and the higher roughness of the CF-B surface are naturally favorable factors, since they increase the tortuosity of the gas path. This could explain the lower permeability of CF-B/PPS composite compared to PPS. However, this requires the matrix to wet the fiber topology perfectly, which is not necessarily the case for PVDF due to its low surface tension. For this matrix, the combination with fibers does not allow recovering the actual permeability of the PVDF and gas diffusion in the fiber/matrix interface is therefore facilitated. The benefit of carbon fiber roughness in increasing the overall tortuosity of the gas diffusion path at the interface then seems limited for this matrix.

However, the AFM results did not indicate that the interface is marked by the presence of voids or a zone of lower density than the
matrix. It is therefore possible that the level of chemical interaction at the interface, and in particular the level of adhesion, also controls the transport of gas along the interface. The low steric hindrance of the hydrogen molecule allows it to take advantage of the free volume available at the interface and accommodate small-scale cooperative molecular movements to permeate. Regardless of fiber surface roughness, it is thus possible that the molecule may have more difficulty permeating an interface when adhesion is the result of stronger chemical interactions. Direct measurements of the surface tension of the molten matrix [35,36] and on the carbon fiber [37] would enable to confirm these hypotheses and better understand interfacial interaction phenomena.

Nevertheless, although the fiber/matrix interface seems to have a strong influence on hydrogen permeability properties, the role of the crystalline interphase cannot be totally ruled out. POM observations clearly indicate that surface-induced nucleation correlates with permeability properties. Even if this could not be demonstrated, the lamellar microstructure and orientation may influence the gas path around the fibers in addition to the effect of interface properties. A short transition length between fiber and matrix could be observed in the case of the CF-A/PVDF-H sample and this measurement is consistent with the low nucleation ability of the fiber in this composite system (Fig. 9a). AFM technique therefore appears to be a relevant approach for studying density and stiffness changes in a transcrystalline phase. Measurement around the fibers in the CF-B/PPS composite could reveal changes in local properties in the vicinity of fibers and therefore support a potential role of the interphase on hydrogen permeation.

## 5. Conclusion and perspectives

Hydrogen permeability measurements on neat polymers and CF/TP composites of various compositions have highlighted the importance of the fiber/matrix interaction on this property. The ability of the fiber surface to induce matrix crystallization appears as a clear indicator of low gas permeability in composite materials. A matrix like PPS with a high surface tension and in particular a high dispersive component therefore seems beneficial for improving impermeability properties. As well, the activated carbon fiber with a rough and nanoporous surface has a beneficial effect by increasing the nucleation ability of the fiber surface.

Even if this study gives information about the correlation between some interfacial mechanism and hydrogen permeability, questions remain open about the origin of changes and key parameters affecting permeability properties. On the one hand, increasing matrix wettability and increasing fiber roughness obviously increase the tortuosity of the fiber/matrix interface, and are the most probable reason of the decrease in permeability. Nevertheless, the transcrystalline or shish-calabash microstructure around fiber can also increase the matrix density and therefore improve the impermeability of composites because of the interphase formed.

For a better understanding of the role of the interface on permeability, an in-depth study of the fiber surface is necessary. Indeed, the characterization of the chemical properties and surface tension of fiber surface would provide interesting information about the interfacial interaction. The analysis of the density of the semi-crystalline phase around the fibers by means of nanomechanical investigations should also be further investigated. Preliminary AFM measurements showed that it is a promising technique to investigate the presence of an interphase and to better understand its role on permeability. A more comprehensive study involving a larger number of fiber/matrix pairs would also enable to draw clearer conclusions about the mechanisms controlling hydrogen permeability in thermoplastic composites.

### CRediT authorship contribution statement

Gautier Allusse: Writing – original draft, Visualization, Validation,

Methodology, Investigation, Formal analysis, Conceptualization. **Olivier De Almeida:** Writing – review & editing, Supervision, Project administration, Methodology, Funding acquisition, Formal analysis, Conceptualization. **Quentin Govignon:** Writing – review & editing, Supervision, Methodology, Formal analysis, Conceptualization. **Monica Pucci:** Writing – review & editing, Investigation. **Fabrice Schmidt:** Writing – review & editing, Supervision.

### Declaration of competing interest

The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: Olivier De Almeida reports financial support was provided by HYCCO. If there are other authors, they declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgements

The authors acknowledge HYCCO for funding this study and Arkema for supplying the materials.

### Data availability

The data that has been used is confidential.

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