Thermoforming Process Effect on Performances of Thermoplastic/Recycled Carbon Fiber Composites

Abstract

The reuse of recycled carbon fibers (rCF) is a response to growing environmental concerns associated with the composites industry. Recycling and reusing carbon fibers represents a more sustainable alternative by reducing waste at the end of the life cycle of composite materials and decreasing dependency on virgin raw materials. This study investigates the influence of process parameters on two different non-woven mats made by carding rCF and blending with thermoplastic filaments: Carbiso TM-PA6/60 and TM-MAPP/60. Two processing methods were examined—one-shot process (M1) and lamination (M2)—to fabricate multilayer coupons. The results indicate that the two-layer panels produced using M2 exhibited a lower porosity (9.9% for PA6/60 and 4.1 for MAPP/60) and superior mechanical performance. However, the differences in performance between the two methods diminished as the number of layers increased. Concerning matrix–fiber compatibility, MAPP/60 showed the best results due to the fiber’s roughness, matrix particles on the fibers, and the incorporation of maleic anhydride in polypropylene (PP), significantly enhancing adhesion.

1. Introduction

The development of carbon fiber-reinforced composites (CFRC) has attracted significant attention in recent decades, primarily because of their lightweight and superior mechanical properties like high specific strength and stiffness, corrosion and fatigue resistance [1,2]. Due to their excellent properties, they are widely used in many fields, such as in construction [3], the aeronautic sector [4], and the automotive sector [5], with both thermoset and thermoplastic matrices [6,7]. The growth in the production and use of CFs inevitably leads to high amounts of CF composite waste. Indeed, global CFRP waste is foreseen to reach up to 20 kilotons annually by 2025 [8]. The difficulty with the thermoset matrix is that it cannot be melted and reused, which represents a recycling problem. Most of this waste is either incinerated or landfilled, even though the European Waste Framework Directive states that landfills are the least preferred option for CF composite disposal [9]. As a solution, thermoplastic composites appear to be the most viable option because they can be melted and reshaped multiple times without significantly altering their chemical structure [10]. Several methods are used to reclaim CFs from thermoplastic composite waste, ranging from electromagnetic to thermal and thermochemical methods [11,12].

Recycled carbon fibers are typically in a discontinuous and random form due to the cutting process during composite manufacturing and the size reduction during recycling. From the recycling processes of carbon fibers, rCFs are obtained with mechanical properties that are not too far from those of virgin fibers.

Sun et al. [13] proposed an electrochemical method for recycling carbon fiber from carbon fiber-reinforced polymer (CFRP). The test on rCF shows that the maximum tensile strength of the reclaimed carbon fiber is 80% that of the virgin carbon fibers (VCF). Ye et al. [14] recovered carbon fiber from epoxy with the steam thermolysis method. The results show that reclaimed resin (or residue) free carbon fibers retained more or less over 90% of their original tensile strength. The tensile strength of rCF could reach approximately 80–95% of virgin carbon fibers (VCF) [15].

Among the manufacturing technologies applicable to the rCF, it is worth considering infusion molding, extrusion/compression, and 3D printing [16,17,18].

Developing recycled carbon fiber-reinforced composites (rCFRCs) and utilizing nonwoven fibrous mats infused with resin is possible. These mats can be produced using different techniques such as wet-laid, air-laid and carding processes. The air-laid process is a dry laying technique in which individual fibers are evenly dispersed in an airstream and then directed toward a permeable screen or conveyor, where the fibers are deposited randomly in the form of a web. The resulting nonwoven mat is essentially isotropic in nature with a three-dimensional structure where the areal weight ranges from 150 to 200 g/m² [19]. The process of producing wet-laid nonwovens is similar to papermaking. The fibers are mixed with water or a suitable solvent to form a fiber/water mixture. The mixture is then transferred to a head box, forming a continuous isotropic nonwoven web. The water is drained from the web and moved forward via a forming belt at this stage. Finally, the web is dried using the heaters and rolled up using a take-up system [20]. When fibers are randomly oriented, in-plane material isotropy can be expected. Carding is a mechanical process in which the discontinuous carbon fibers agglomerates are placed in a series of rotating cylinders covered with wire or pins that work to align and separate the fibers. During this phase, resin can be added to infuse the fibers. The resin-impregnated fibers are then laid in layers, forming mats with a specific fiber orientation [21]. Carding allows the production of anisotropic nonwoven composites, unlike the previous processes.

Subsequently, these mats can be thermoformed with traditional manufacturing techniques, such as compression molding, to produce composites with relatively high fiber volume fractions and good mechanical properties. Pimenta et al. [22] produced nonwoven mats based on rCFs with a paper-making process and then compression-molded them with an epoxy resin layer to manufacture rCFRCs. The samples’ nominal volume fraction was around 30%, resulting in the same stiffness and strength as the aerospace grade 2024-T4 aluminum alloy. Quan et al. [22] produced rCF/PPS composites from nonwoven mats consisting of commingled rCFs and polyphenylene-sulfide (PPS) fibers through compression molding. The experimental results demonstrated that interlaminar fracture energies and mechanical properties significantly increased. However, it also negatively affected the interlaminar fracture resistance. Therefore, studies of these mats’ processing, post-processing, and mechanical behavior are particularly interesting. Coupling Nonwoven mats made by rCF with polyamide (PA6) enables high abrasion resistance, heat resistance and excellent mechanical properties [23]. On the other hand, the use of MAPP, polypropylene modified with maleic anhydride (MA), allows the polymer to retain all the advantages of PP, including excellent chemical and mechanical properties, thermal stability, and low density [24]. Additionally, the presence of maleic anhydride enhances compatibility between the matrix phase and recycled carbon fibers (rCFs) [25].

The effect of MA in improving interfacial adhesion is primarily attributed to chemical interactions, such as the formation of covalent and hydrogen bonds, between the functional groups on the CF surface and the MA groups in the MAPP [25]. Fiber recovery is achieved through a modified pyrolysis process, where a furnace burns the resin. According to ELG Carbon Fibre, pyrolysis provides a recycled material that maintains 90% of the carbon fiber’s tensile strength. To convert the recycled carbon fiber, ELG Carbon Fibre use a carding line capable of providing up to 250 tonnes per year of carbon fiber or carbon fiber/thermoplastic nonwoven mats [26]. The recycled carbon fibers are initially mixed with thermoplastic fibers, polyamide PA6 and maleic anhydride-modified polypropylene (MAPP) to ensure homogeneous distribution. Subsequently, rCFs mixed with thermoplastic fibers are fed into the carding machine, which is equipped with toothed cylinders that align and distribute the fibers into a thin and uniform web. This fibrous web is then layered to achieve the desired basis weight. The resulting mats can be subjected to thermoforming processes [27].

This paper aims to investigate the optimal process condition of compression-molding for repurposing recycled carbon fibers nonwoven mats. Two types of pre-forms were considered: Carbiso TM-PA6/60 and Carbiso TM-MAPP/60, namely rCF/PA6-60 and rCF/MAPP, respectively. Two different approaches were explored: direct compression-molding in a one-shot fabrication (Method 1—M1) and lamination with stepwise consolidation (Method 2—M2). Differential scanning calorimetry (DSC) was used to evaluate how processing affected crystallinity. The influence of porosity on the elastic modulus, induced by the compression molding process, was assessed by dynamic mechanical analysis (DMA). The results show that with up to two layers, M2 yields better performance in terms of porosity and storage modulus; however, with more layers, performance differences between the two methods diminish. The viscoelastic properties also highlight differences in interphase behavior: PA6-based composites exhibit a reinforcement modulus of 60 GPa, while MAPP-based ones reach 90 GPa. The higher reinforcing efficiency in MAPP composites is attributed to a stronger interphase enabled by maleic anhydride, as evidenced by lower tanδ values, indicating reduced energy dissipation.

2. Materials and Methods

2.1. Sample Preparation

Carpet of Carbiso TM-PA6/60 and Carbiso TM-MAPP/60 (ELG Carbon Fibre Ltd., Coseley, UK, kindly supplied by AEROSOFT, Capua, CE, Italy) are made with 40% carbon fiber and a phase consisting of polycaprolactam (PA6) and maleic anhydride-modified polypropylene (MAPP) at 60%, respectively. In particular, MAPP is polypropylene modified with maleic anhydride (MA). For both materials, the nominal surface density is 500 GSM (grams per square meter). In this work, PA6/60 and MAPP/60 carpets were subjected to two different compression-molding processes to produce multilayer panels: one layer (1L), two layers (2L), and three layers (3L). The carpets were cut to obtain different layers measuring 50 × 50 mm. The thickness of the compression-molded layer was about 0.5 mm.

The first technique adopted was a one-shot process in which the layers were stacked and pressed simultaneously (Figure 1a). Conversely, the second method chosen was a lamination process. The layers were first pressed separately and then stacked and pressed again in a second phase to obtain the desired panel (Figure 1b). The two pressing processes will be called Method 1 and Method 2, respectively. Figure 2a shows the PA6/60 panel produced using the first method, while Figure 2b depicts the panel made with the second method. Similarly, in Figure 2c, the MAPP/60 panel produced using the one-shot technique is shown, whereas Figure 2d presents the panel obtained through lamination.

In the compression molding processes, the parameters adjusted for the two different materials were the process settings (temperature and time). The compression moulding was achieved using a Collin P200 E (Dr. Collin GmbH, Ebersberg, Germany) platen press. For the PA6/60 carpets, the one-shot process, shown in Figure 3a, was carried out with a preheating phase from room temperature to 210 °C, close to the melting temperature, by closing the press plates. Subsequently, pressing was performed first at 210 °C for 10 min and then at 220 °C for 20 min, applying a pressure of 100 bar in both cases.

Figure 3b shows the lamination process: the layers were preheated separately with the plates closed, from room temperature to 210 °C. Subsequently, the layers were pressed at 210 °C for 10 min and at 220 °C for another 10 min. After this phase, the press was opened, the layers were stacked, and a final pressing at 220 °C for 10 min was carried out. In all stages, the selected pressure was 100 bar.

For the MAPP/60 fibers, the one-shot process, shown in Figure 4a, was carried out as follows: the layers were preheated, with the plates closed, from room temperature to 170 °C, close to the melting point. Once the desired temperature was reached, the compression phase was performed for 20 min at 100 bar. On the other hand, for the lamination process, after the preheating process, the layers were first pressed separately at 170 °C. Then, the press was opened, and the layers were stacked and compressed for another 10 min at the same temperature. Compression-molding phases were carried out at 100 bar (Figure 4b).

2.2. Experimental Characterization

Thermogravimetric analysis (TGA) was performed using a TA Instruments Q5000 instrument (New Castle, DE, USA) to assess the thermal stability range of the polymer according to the ASTM E1131 standard [28]. The measurements were conducted under an inert atmosphere of nitrogen gas, with a temperature ramp of 10 °C/min from room temperature to 900 °C, and samples of 15 mg were considered. The weight loss was quantified at 600 °C. To calculate the actual fiber content (${\%w}_{f,reale}$), the residues obtained from thermogravimetric (${\%w}_c$) tests were reduced by dry residue of the only polymer ${(\%w}_{NR})$, in accord with Equation (1):

$${\%w}_{f,reale}={\displaystyle \frac{{\%w}_c-{\%w}_{NR}}{100-{\%w}_{NR}}}·100$$

(1)

All the residues were measured at 600 °C.

The thermal characteristics of the polymer were examined through differential scanning calorimetry (DSC) utilizing a TA Instruments DSC Q2000 instrument. Each test specimen of PA6/60 underwent two heating and cooling cycles from −50 to 280 °C at a rate of 10 °C/min in a nitrogen atmosphere. MAPP/60 specimens were subjected to two heating and cooling cycles from −50 to 220 °C at a rate of 10 °C/min in a nitrogen atmosphere. Before measurements, amounts of approximately 10 mg of the samples were enclosed in aluminum pans. The glass transition temperature (Tg) and the enthalpy of the reaction were derived from the DSC curves by ASTM D3418 standards [29].

Dynamic mechanical analysis (DMA) was performed with a Dynamic Mechanical Analyzer Q850 from TA Instruments in the Single Cantilever mode (SC). Samples of a rectangular shape 35 mm in length, 6.0 ± 0.10 mm in width, and about 0.5, 1.0 and 1.5 mm in thickness were tested. The behavior of the PA6/60 and MAPP/60 samples was investigated with temperatures between −70 and 180 °C and −70 and 120 °C, respectively. The tests are performed considering a heating rate of 3 °C/min, a strain amplitude varying between 25 and 50 μm, depending on the sample thickness, and a frequency of 1 Hz. Data were elaborated according to the ASTM D790 standard [30] for the flexural behavior of unreinforced and reinforced plastics.

Surface roughness measurements were carried out using a digital optical microscope, specifically the VHX-X1 series from Keyence Corp. (Osaka, Japan). Samples were positioned vertically and carefully levelled on the stage of the 4K ultra-high-accuracy microscope. A separate ring light source and transmitted illumination were used, following ISO standards [31].

3. Results and Discussion

3.1. Thermal Characterization

The thermogravimetric analysis (TGA) of the PA6/60 and MAPP/60 samples, shown in Figure 5a,b, highlights the two materials’ thermal behavior as a temperature function. Both materials exhibit significant weight loss between 350 °C and 500 °C, indicating the thermal decomposition of the main polymer. However, the degradation profile differs between the two samples. PA6/60 shows the main decomposition phase with a peak in the derivative thermogravimetry curve around 460 °C, characteristic of polyamide 6 (PA6). On the other hand, MAPP/60 displays a primary decomposition occurring at slightly lower temperatures, with a DTG peak around 430 °C. This is likely due to the presence of maleic anhydride-grafted polypropylene (MAPP), which degrades at lower temperatures than PA6. Additionally, the solid residue at high temperatures (600 °C) is around 40% for both materials, a value consistent with the material datasheet [32].

Table 1. Results of TGA analyses.

Description	Weight Loss @ 250 °C [%]	T@ Maximum Decomposition [°C]
PA6/60	1.7 ± 0.5	440.9 ± 0.5
MAPP/60	0.5 ± 0.2	425.4 ± 0.5

reports the weight loss of the examined materials from room temperature up to 250 °C and the maximum degradation temperature observed. PA6/60 shows a higher weight loss at 250 °C, which could indicate the presence of components that volatilize or degrade at this temperature.

The differential scanning calorimetry (DSC) analysis of the PA6/60 and MAPP/60 samples, shown in Figure 6a,b, allows for the identification of key thermal transitions, including the glass transition temperature (Tg) and melting point (Tm). For PA6/60, the first heating cycle shows a glass transition temperature around 50–60 °C, indicated by a slight shift in the baseline of the heat flow. The primary melting peak is observed at approximately 220–225 °C, with a well-defined area characteristic of the crystalline phase of polyamide 6 (PA6). Additionally, in the second heating cycle, a similar glass transition is observed, confirming the stability of Tg after the first thermal treatment. The crystallization temperature during cooling (Tc) is identified around 180–190 °C, indicating the material’s ability to recrystallize after melting.

Instead, MAPP exhibits two transitions in the −50 °C to 100 °C range. The first transition at around 0 °C is β-relaxation, and it is associated with the glass–rubber transition of the fully amorphous phase. The temperature transition between 40 and 60 °C is α-relaxation [33]. The Tg associated with α-relaxation was selected for inclusion in

Table 2. Results of DSC, density and porosity analysis.

Description	Tg, DSC [°C]	Tm, DSC [°C]	Density [g/cm3]	Porosity [%]
PA6/60 1L	50.2 ± 0.3	216.0 ± 0.5	1.2 ± 0.08	6.5 ± 0.02
PA6/60 2LM1	54.4 ± 0.3	220.4 ± 0.6	1.1 ± 0.15	13.2 ± 0.07
PA6/60 2LM2	54.3 ± 0.4	220.4 ± 0.5	1.2 ± 0.02	9.9 ± 0.02
PA6/60 3LM1	50.5 ± 0.5	221.5 ± 0.6	1.1 ± 0.11	14.0 ± 0.05
PA6/60 3LM2	48.8 ± 0.4	220.5 ± 0.7	1.1 ± 0.01	16.7 ± 0.01
MAPP/60 1L	46.4 ± 0.3	163.9 ± 0.4	0.9 ± 0.03	7.3 ± 0.01
MAPP/60 2LM1	51.4 ± 0.6	165.8 ± 0.3	0.9 ± 0.2	8.0 ± 0.03
MAPP/602LM2	53.1 ± 0.4	163.9 ± 0.5	1.0 ± 0.02	4.1 ± 0.02
MAPP/60 3LM1	46.5 ± 0.6	164.4 ± 0.7	0.9 ± 0.02	4.8 ± 0.02
MAPP/60 3LM2	45.3 ± 0.5	164.9 ± 0.4	1.0 ± 0.01	5.0 ± 0.01

. The melting peak (Tm) appears between 160 and 170 °C, characteristic of polypropylene. The second heating cycle confirms these values, indicating good reproducibility of the material’s thermal structure. The crystallization temperature (Tc) is around 120 °C, reflecting the material’s recrystallization behavior. In the first heating cycle, all the analyzed samples show two melting peaks (Figure 6b). The peak on the left is associated with the melting of the β-phase, while the one on the right corresponds to the α-phase [34]. After further heating, the rCF/MAPP60 at 10 °C/min exhibits a single peak corresponding to the crystallization of the β-phase at 160 °C.

The geometric density was calculated using Equation (2):

ρ = m/V

(2)

where ρ is the density (g/cm³), m is the sample’s weight (g), and V is the measured volume (cm³). The obtained density values are reported in Table 2. The surface density of samples was measured by averaging the weight of coupons to their surface. Densities of 530 ± 10 gsm and 560 ± 30 gsm were measured for rCF/MAPP and rCF/PA6 carpets, respectively. By the calculated density, it is possible to obtain the average porosity of each sample. The average porosity, P, was obtained by applying Equation (3):

P = 1 − ρ/ρ_th

(3)

where ρ_th is the theoretical density obtained by considering the rule of mixture [35]. The values of ρ_th for PA6/60 and MAPP/60, considering fiber volumes (v_fs) of 29.4% and 25.1%, equal 1.33 g/cm³ and 1.13 g/cm³, respectively.

Table 2 summarizes the results obtained from differential scanning calorimetry (DSC), density and porosity analyses.

3.2. Viscoelastic Characterization

The viscoelastic behavior of the sample with a temperature from −70 to 180 °C is reported in Figure 7. Figure 7d shows a scanning electron microscopy (SEM) image of the fractured surface of the PA6/60 one-layer panel. The microstructure reveals a heterogeneous distribution of recycled carbon fibers embedded in the PA6 matrix. The image displays clear interfacial bonding between the matrix and the fibers, although some fiber pull-outs and voids are visible, indicating localized debonding or porosity.

DMA experiments were performed in the linear viscoelastic deformation range to investigate how the molding process influences the storage modulus, loss modulus and tanδ of the manufactured panels, the number of overlapping layers and the porosity.

Table 3. Results of DMA analysis at 50 °C.

Description	E′ @ 50 °C [GPa]	E″ @ 50 °C [GPa]	Tanδ @ 50 °C [-]
PA6/60 1L	16.2 ± 1.10	0.40 ± 0.03	0.025 ± 0.0029
PA6/60 2LM1	10.0 ± 1.06	0.36 ± 0.04	0.036 ± 0.0029
PA6/60 2LM2	15.2 ± 1.06	0.34 ± 0.05	0.023 ± 0.0014
PA6/60 3LM1	11.4 ± 1.22	0.32 ± 0.05	0.028 ± 0.0020
PA6/60 3LM2	11.6 ± 0.31	0.32 ± 0.01	0.028 ± 0.0005

, instead, shows the results of DMA at 50 °C for the PA6/60 panels. The selected temperature was chosen close to the glass transition temperature of the hosting matrices to compare both materials under the same mobility conditions.

The viscoelastic behavior of the sample with a temperature from −70 to 120 °C is reported in Figure 8. Figure 8d presents an SEM image of the fracture surface of an MAPP/60 one-layer panel. Several fibers appear to be well embedded, while others exhibit pull-out features, suggesting partial debonding or fiber–matrix interfacial failure. The presence of voids and resin-rich areas can also be observed, indicating heterogeneities likely introduced during the compression-molding process. These morphological characteristics are consistent with the mechanical behavior observed in the DMA results, particularly in terms of storage modulus and damping performance.

DMA experiments were performed in the linear viscoelastic deformation range to investigate how the molding process influences the storage modulus, loss modulus and tanδ of the manufactured panels, the number of overlapping layers and the porosity.

Table 4. Results of DMA analysis at 20 °C.

Description	E′ @ 20 °C [GPa]	E″ @ 20 °C [GPa]	Tanδ @ 20 °C [-]
MAPP/60 1L	11.2 ± 1.41	0.2 ± 0.03	0.021 ± 0.0003
MAPP/60 2LM1	10.8 ± 1.36	0.3 ± 0.05	0.027 ± 0.0014
MAPP/602LM2	16.7 ± 0.57	0.4 ± 0.02	0.021 ± 0.0007
MAPP/60 3LM1	15.4 ± 0.40	0.3 ± 0.01	0.022 ± 0.0009
MAPP/60 3LM2	14.7 ± 1.31	0.3 ± 0.03	0.023 ± 0.0011

, instead, shows the results of DMA at 20 °C for the MAPP/60 panels. This temperature was chosen because it is near the glass transition temperature of the maleic anhydride-grafted polypropylene (MAPP).

3.3. Effect of Molding Process on Crystallinity

Since the properties of composites based on semicrystalline polymers are partly determined by the crystalline structure and the degree of crystallinity of the thermoplastic matrix, it was necessary to study the effect of the recovery process, in particular the molding, on the degree of crystallinity. Measuring the melting enthalpy of a polymer allows for estimating its degree of crystallinity. The amorphous or semi-crystalline nature affects the melting enthalpy value and, consequently, the degree of crystallinity (Xc). In particular, the Xc can be estimated using the following equation [36]:

$$X_C={\displaystyle \frac{{\mathsf{\Delta }H}_m-{\mathsf{\Delta }H}_c}{{\mathsf{\Delta }H}_m^0}}·100 \left[\%\right]$$

(4)

where ${\mathsf{\Delta }H}_m$ is the enthalpy of neat resin obtained by normalizing the weight of the tested sample on the percentage of fiber content; ${\mathsf{\Delta }H}_c$ is the enthalpy associated with the cold crystallization phenomenon, and $\mathsf{\Delta }{H}_m^0$ is the melting enthalpy of a perfect crystal.

The melting enthalpy of 100% crystallized PA6 is 230 J/g [37,38].

Table 5. Results of DSC and crystallinity analysis of PA6/60 panels.

Description	1st Reaction Enthalpy [J/g]	2nd Reaction Enthalpy [J/g]	1st DOC [%]	2nd DOC [%]	Δ DOC [%]
PA6/60 1L	74.5 ± 1.0	63.4 ± 1.2	32.4	27.6	4.8
PA6/60 2LM1	73.1 ± 1.4	58.2 ± 1.6	31.8	25.3	6.5
PA6/60 2LM2	71.0 ± 1.3	64.2 ± 1.8	30.9	27.9	3.0
PA6/60 3LM1	72.1 ± 1.2	65.1 ± 1.2	31.3	28.3	3.0
PA6/60 3LM2	70.0 ± 1.1	66.4 ± 1.0	30.4	28.9	1.6

reports the enthalpy values of the two heating cycles of DSC and the respective crystallinities calculated using Equation (4) for the manufactured PA6/60 panels. Additionally, the difference between the two crystallinity values was determined by subtracting the crystallinity of the second cycle from that of the first cycle. The data analysis shows that for Carbiso TM-PA6/60, Method 1 leads to a greater crystallinity reduction than Method 2, resulting in more amorphous material.

The melting enthalpy of 100% crystallized PP is 209 J/g [39,40]. In

Table 6. Result of DSC and crystallinity analysis of MAPP/60 panels.

Description	1st Reaction Enthalpy [J/g]	2nd Reaction Enthalpy [J/g]	1st DOC [%]	2nd DOC [%]	Δ DOC [%]
MAPP/60 1L	108.0 ± 1.6	92.2 ± 1.1	51.7	44.1	7.6
MAPP/60 2LM1	109.6 ± 1.0	97.6 ± 1.2	52.4	46.7	5.7
MAPP/602LM2	104.1 ± 1.2	93.2 ± 1.5	49.8	44.6	5.2
MAPP/60 3LM1	105.8 ± 1.3	94.2 ± 1.7	50.6	45.1	5.6
MAPP/60 3LM2	103.1 ± 1.1	92.9 ± 1.3	49.3	44.5	4.9

, the results of the DSC analysis of MAPP/60 panels and the respective crystallinity are reported. The results highlight that the Carbiso TM-MAPP/60 lamination process leads to lower crystallinity with respect to the one-shot method. This result is opposite to that of PA6/60.

3.4. Effect of Porosity on the Storage Modulus

The intricate relationship between porosity and the mechanical properties of composite materials is a critical factor in assessing their performance. The presence of pores within the polymer matrix can drastically impact both stiffness and viscoelastic behavior, making it essential to understand this dynamic.

Generally, increased porosity leads to a reduction in the storage modulus and an elevated damping factor. This underlines the importance of selecting optimal hot-pressing methodology and processing parameters to effectively control porosity, influencing the material’s properties and potential for advanced structural applications. Figure 9 illustrates the correlation between the elastic modulus and porosity in PA6/60 panels. As porosity rises, the storage modulus (E′) value declines. Notably, for samples obtained by two layers and compression-molded with Method 2 emerged as the better choice, overcoming Method 1. The 2LM1 sample achieved an average E′ value of 10 GPa, while the 2LM2 sample excelled with a remarkable 15.2 GPa.

This striking contrast in modulus values is directly linked to the higher porosity of the one-shot process (13.2%) compared to the lamination process (9.9%), highlighting the significance of compression molding process parameters. However, as the number of layers increases, the differences in compaction methodologies decrease. In the case of three-layer panels, the storage modulus for the 3LM1 sample reaches 14.4 GPa, while the 3LM2 sample holds at 11.6 GPa. Additionally, porosity values begin converging at approximately 14% for the first sample and 16.7% for the second, emphasizing the complex interplay between processing conditions and material performance.

Figure 10 shows the microscopic analysis on a PA6/60 2LM1 sample in proximity of a porosity, while Figure 11 is the same analysis for a PA6/60 2LM2 sample.

The defect in the sample made with Method 1 has an average thickness of 44.7 µm, while the defect in the sample produced with Method 2 has an average thickness of 39.2 µm. Micrographs confirm that the samples prepared according to Method 1 (M1) reproduce samples with a higher degree of porosity and weak materials.

Figure 12 illustrates the variation of Elastic modulus as a function of porosity for the MAPP/60 panels. Consistent with previous findings, Method 2 yielded better results for the two-layer panels. Specifically, the elastic modulus for the 2LM2 sample is 16.7 GPa, compared to 10.8 GPa for the 2LM1 sample. This difference is linked to the variation in porosity between the two methods, which is 4.1% for the lamination process and 8.0% for the one-shot method. Also, in the case of MAPP/60, increasing the number of layers minimizes the quality difference between the two compression-molding methods. For instance, the 3LM1 sample recorded an elastic modulus (E′) of 15.5 GPa with a porosity of 4.8%, while the 3LM2 sample exhibited an E′ of 14.7 GPa with a porosity of 5.0%.

3.5. Effect of Compatibility Between Polymer Matrix and rCF

In composite materials, the compatibility between the matrix and the fiber plays a crucial role in determining the mechanical properties of the final material. The strong adhesion between these two components ensures effective stress transfer, enhancing the composite’s strength, durability, and overall performance. Conversely, poor compatibility can lead to reduced mechanical resistance. Among the numerous models present in the literature, the Cox–Krenchel model is suitable for the evaluation of the elastic modulus for carbon fiber-reinforced composites (CFRC) and is found to yield good agreement with the experimental modulus for a range of carbon fiber lengths and volume fraction [41,42,43]. When stress is applied parallel to the fibers, the mechanical elastic modulus of the composite (E_c) is given by [44]:

$$E_c=v_f{E}_f^{\mathrm{*}}+v_m{E}_m$$

(5)

where $E_f^{\mathrm{*}}$ is the effective elastic modulus of the recycled fibers (rCF), $E_m$ is the elastic modulus of PA6 and MAPP matrices, $v_f$ is the volume of the fibers expressed as a percentage and $v_m$ is the volume of the matrix as a percentage. Making the necessary substitutions, the final equation becomes:

$$E_c=v_f\eta E_{rCF}+{(1-v}_f)E_m$$

(6)

where η is the reinforcement efficiency factor, which depends on several factors, including the fiber aspect ratio (length to diameter ratio), the packing and orientation of the fibers, the compatibility between fibers and matrix, and the shear strength [45].

For PA6/60 and MAPP/60, E_c values of 19.3 GPa and 23.8 GPa were obtained, respectively. The significant difference recorded among the two materials is mainly related to the compatibility between the matrix and fibers. In particular, the effective elastic modulus ($E_f^{*}$) is 60 GPa for PA6/60 and 90 GPa for MAPP/60.

The actual reinforcement of the rCF is mitigated by the parameter η. By comparing the two systems, it is possible to note that the fibers are the same, with the same average aspect ratio and similar random distribution. The main difference is the interaction between fibers and the hosting matrix due to the surface roughness and chemical compatibility. Rough surfaces increase frictional force and interfacial adhesion between the reinforcement and the polymer, allowing us to obtain good interfacial shear strength values [45] for both the systems.

Furthermore, in MAPP/60, the hydrophilic anhydride groups of the sizing reacted with the hydroxyl groups on the CF surface. This resulted in a stronger interfacial interaction and thus stronger adhesion with the matrix [46].

Evidence of the strength at the interphase between fibers and reinforcement is the damping capacity (proportional to tanδ): the rCF/PA6-60 has a tanδ of 0.3. At the same time, the rCF/MAPP60 shows a value of 0.2, which means a dissipation capacity of less than −30% compared to similar conditions (polymer chain mobility). This means that the PA6-60-based composite dissipates more energy, which implies weak bonding with the reinforcement. Conversely, the MAPP/60 composite acts as a stiffer system with good adhesion between the matrix and reinforcement, thanks to the grafting of fibers promoted by maleic anhydride.

4. Conclusions

This paper investigates the optimal process conditions for repurposing recycled carbon fiber-based materials. Two different nonwoven mats made from recycled carbon fiber (rCF) are selected, each using a different polymer MAPP and PA6. First, the process layout is discussed according to two different manufacturing protocols: one-shot fabrication (M1) and lamination consolidation (M2). The processing temperatures are chosen based on the melting points of the respective polymers for both systems. A reduction in crystallinity is observed in the PA6-60 produced by M1 compared to the case of the MAPP60 system produced by M2.

A critical parameter to manage during the process is the porosity of the final materials. Increasing the number of layers does not impact the processing procedure, resulting in an average porosity of 15% for PA6 and 5% for MAPP, with elastic moduli of 11 GPa and 15 GPa, respectively. The impact of porosity on the elastic modulus for both systems was studied, revealing different reinforcement efficiencies due to the distinct polymer compositions. The effective fiber modulus is measured at 90 GPa for MAPP/60 compared to 60 GPa for PA6/60. This difference is attributed to the presence of maleic anhydride in the matrix, which introduces polar groups along the polypropylene chain, enhancing adhesion with the rCFs. Additionally, the average tan δ values for the two materials are analyzed. The PA6/60 has a tan δ value of 0.3, while the MAPP/60 shows a value of approximately 0.2. This indicates that PA6/60 dissipates more energy, leading to greater viscoelastic deformation and reduced effectiveness in transferring forces from the matrix to the fiber. In contrast, the MAPP/60, with lower energy dissipation, demonstrates more stability and less viscoelastic deformation.