Innovative safer composites made from recycled automotive parts materials

Abstract

This study investigates the mechanical properties of carbon fiber–reinforced epoxy composites reinforced with recycled textiles from the automotive industry—specifically airbag fabric (nylon 6.6) and seat-belt webbing (polyester). The objective was to assess the potential of these recycled reinforcements to enhance impact resistance, which is critical for automotive safety applications. Six composite configurations with different lay-ups were manufactured: a reference carbon composite (C), a carbon composite reinforced with a conventional aramid fabric (A), and four hybrid composites incorporating nylon 6.6 (AB), seat-belt webbing (P), an aramid/seat-belt combination (AP), and a nylon 6.6/seat-belt combination (PB). Specimens were produced by hand lay-up and subsequently subjected to mechanical testing. Tensile properties were evaluated in accordance with ČSN EN ISO 527-4, and Charpy impact performance was measured according to ISO 179-2 (unnotched). The results showed that incorporating recycled seat-belt reinforcement (P, AP, PB) led to a pronounced increase in impact strength (most notably for AP) compared with the reference carbon composite (C). However, this improvement in toughness was accompanied by a reduction in tensile strength. The aramid-only configuration (A) achieved the highest average tensile strength. The hybrid AP and PB laminates, combining carbon plies with aramid fabric/nylon 6.6 and seat-belt reinforcement, exhibited the best impact energy absorption capability. Overall, the findings suggest that recycled seat-belt textiles represent a promising alternative reinforcement for improving the toughness of carbon composites in applications requiring high impact resistance, provided that the trade-off between toughness and static strength is carefully considered.

Keywords

carbon fiber composite hybrid composite recycled materials seat belt airbag

1. Introduction

Advanced polymer composites represent a key materials solution for reducing vehicle mass and improving passive-safety efficiency, particularly in applications requiring controlled energy absorption under impact loading. In recent years, a clear trend has been the shift from “monolithic” materials toward hybrid and multilayer (multi-material) concepts that combine high specific stiffness and strength with increased toughness and a more progressive failure mode.¹

At the same time, the growing deployment of carbon-fibre-reinforced polymer (CFRP) structures has intensified pressure to improve sustainability and support circular-economy principles. Consequently, recycling of thermoset CFRP, as well as the reuse of secondary fibres and technical textiles, is being intensively investigated both in terms of processing technologies and with respect to the resulting properties and market feasibility.^2,3

From a structural-design perspective, a major limitation of CFRP is their relatively brittle behaviour under impact loading and their susceptibility to internal damage initiation (notably delamination and matrix cracking), which may be macroscopically barely visible yet can substantially reduce residual load-bearing capacity. Review studies indicate that the dominant damage mechanisms under impact and the subsequent property degradation depend on the reinforcement architecture, laminate thickness, fibre–matrix interfacial quality, and the applied loading regime.^4–6

Among the most effective strategies for improving impact resistance are reinforcement hybridisation (e.g., combining carbon and aramid fibres or technical textiles) and/or local toughening of the matrix or the fibre–matrix interface. Hybrid FRP systems can improve specific energy absorption by combining load-bearing fibres with fibres capable of higher deformation, enabling fibre bridging and controlled delamination.⁷ Impact response is also strongly influenced by laminate thickness and lay-up, as evidenced under ballistic and high-velocity regimes, where increasing thickness typically alters the balance between local matrix damage, delamination, and fibre fracture.⁸ Under low-velocity impact, carbon/aramid hybrid epoxy laminates have been experimentally shown to exhibit improved damage tolerance and altered failure mechanisms compared with purely carbon configurations.⁹

In addition to classical ply-level hybridisation, novel concepts based on structured reinforcements and architectures—such as auxetic structures—are also being explored to enhance energy dissipation and improve the response to localised loading characteristic of in-service damage in body panels.¹⁰ In this context, the metric of specific energy absorption (SEA) is essential for rational design, as it enables the comparison of different hybrid configurations under mass constraints.¹¹

Within advanced composite architectures, it has further been demonstrated that micro-hybridisation and targeted interfacial modifications can increase both the impact and post-impact strength of CFRP laminates.¹² For sandwich structures with aramid honeycomb cores and hybrid face sheets, enhanced resistance to localised damage has been reported, which is directly relevant to automotive body panels and protective structures.¹³ Composite–metal hybrid structures are also regarded as promising solutions, as they can improve damage tolerance and provide alternative load paths while maintaining low weight.¹⁴ Predictive approaches aimed at estimating composite impact behaviour from material parameters and reinforcement architecture have likewise been shown to support the design and optimisation of such systems.¹⁵

From a sustainability perspective, technical textiles from the automotive industry (e.g., seat belts and airbag fabrics) constitute a significant yet still underutilised source of reinforcement materials. Their recycling is technologically and economically challenging and often results in downcycling; the current literature highlights the need to identify higher value-added applications while addressing issues of sorting, contamination, and consistent quality of recycled textiles.¹⁶ For practical implementation of textile reinforcements, adhesion and reliable interlaminar bonding are critical, and bonding is frequently a limiting factor in laminated textile composites.¹⁷ Published results on carbon/aramid hybrid epoxy composites indicate that an appropriate placement and fraction of the tougher layer can increase energy absorption and modify failure mechanisms while maintaining an acceptable level of static strength.¹⁸ For seat-belt webbings, energy absorption arises from a combination of high-elongation polyester fibres and the woven architecture, which is also reflected under dynamic loading in crash scenarios.¹⁹ In the broader crashworthiness context, composite tubes exhibiting progressive crushing are commonly cited as reference examples, where energy absorption is governed by a combination of local matrix damage, delamination, and fibre pull-out.^20,21 Airbag fabrics based on PA 6.6 are technical textiles designed to meet stringent strength and stability requirements; surface treatments and coating systems may influence their interaction with polymer matrices.²²

Against this background, the aim of the present work is to experimentally verify whether recycled automotive textiles (PA 6.6 airbag fabric and PES seat-belt webbing) can increase impact energy absorption in hybrid carbon laminates while preserving the manufacturing simplicity of hand lay-up, which is typical of bespoke or low-volume production of body panels. The research was motivated by the need to improve the safety performance of the composite body structure of the SC Titan prototype vehicle by enhancing the material’s capacity to absorb impact energy. The objective was to identify a structural and material solution that provides a higher level of occupant protection while maintaining low weight and practical manufacturability for one-off or small-series vehicles, which are often developed outside major automotive OEMs and rely on manual processing due to the lack of autoclaves or large-part vacuum-bagging capabilities.

2. Materials and methods

For the composite bodywork of the SC Titan prototype vehicle, the lay-up of hand lay-up moulded body panels was selected pragmatically based on three criteria: surface appearance (A-surface quality), resistance to in-service damage (impact, abrasion, stone chipping), and the required stiffness at the lowest possible mass.

2.1. Lay-up used for manufacturing SC titan body panels

1. Layer: CF245 (twill 2/2, 0/90, 245 g/m²) – “cosmetic” and forming ply

• Surface quality and appearance: Twill 2/2 provides a characteristic “carbon” aesthetic and conforms well to doubly curved geometries, thereby reducing the risk of wrinkling and fibre bridging.

• Reduced print-through risk: A finer surface fabric lowers the likelihood that the weave pattern of coarser backing plies will become visible through the lacquer/gelcoat over time (print-through).

• Improved drapability: A 600 g/m² twill is prone to handling issues on the outer surface, particularly in corners and small radii.

2. Layer: Aramid 173 g/m² (plain weave, 0/90) – toughening ply/mitigation of brittle failure

• Aramid (Kevlar/aramid fibres) is typically incorporated in body panels for the following reasons:

• Improved impact and perforation resistance (stone strikes, minor collisions, local indentation), owing to the high toughness of aramid fibres and the increased energy required for damage initiation and growth.

• Crack-arrest function: The aramid ply acts as an interlayer that constrains brittle damage propagation and limits delamination growth in carbon laminates.

• “Anti-splinter” effect: CFRP can fragment in a brittle manner; aramid helps retain fragments and reduces hazardous fraying.

• Practical placement away from the surface: Aramid fibres are difficult to sand and tend to fuzz; therefore, they are generally not used as the outer cosmetic ply.

3. Layer: CF600 (twill 2/2, 0/90, 600 g/m²) – rapid build-up of thickness and stiffness

• Primary load-carrying section: The 600 g/m² fabric increases laminate thickness efficiently, leading to a strong increase in flexural stiffness (approximately scaling with the third power of thickness).

• Hand lay-up efficiency: In manual lamination, achieving the target thickness with fewer plies is time-efficient, despite reduced drapability of heavier fabrics.

• Typical body-panel compromise: A lightweight cosmetic surface ply combined with robust backing plies is a common approach for non-primary load-bearing bodywork.

4. Layer: CF600 (twill 2/2, 0/90, 600 g/m²)

During body-panel manufacturing, the plies were always stacked with the same orientation. The total fibre areal weight was approximately 1618 g/m² (245 + 173 + 600 + 600). For manual lamination without vacuum consolidation, the resulting laminate thickness was 2.7 ± 0.05 mm.

For specimen manufacturing, commercially available composite materials used in SC Titan body-panel production were combined with automotive waste materials whose recycling is often technologically challenging and/or energy- and cost-intensive. As in the body-panel lay-up, the plies were always stacked with the same orientation.

2.2. Materials used

• Carbon fabric: Carbon fabrics (twill 2/2) with areal weights of 245 g/m² (designated CF245) and 600 g/m² (designated CF600).

• Aramid fabric: Aramid fabric (twill 2/2) with an areal weight of 173 g/m² (designated AF173).

• Nylon 6.6 (airbag): Fabric recovered from recycled automotive airbags (designated PA 6.6) with an areal weight of 150 g/m². Prior to use, the fabric was degreased and mechanically roughened to improve adhesion.

• Seat-belt webbing: Fabric recovered from recycled automotive seat belts, assumed to be polyester (designated PES), with an areal weight of 1250 g/m². The surface was mechanically abraded (brushed) to enhance adhesion.

• Matrix: Two-component epoxy resin LH 288 with hardener H 146. The resin-to-hardener mixing ratio was 100:27 by mass.

• Release agents: Release wax Oskar’s M 700/C-WAX and a poly(vinyl alcohol) (PVA) release agent Polivaks EKO PVA.

2.3. Pre-treatment and pre-selection of recycled textiles

The recycled PA 6.6 airbag fabric and PES seat-belt webbing used in this study were recovered from end-of-life Škoda Octavia II passenger vehicles obtained from a scrapyard, specifically from vehicles without prior crash involvement. Thus, the reinforcements did not represent virgin materials or industrial production scrap, but automotive components previously exposed to real service conditions. Prior to lamination, both recycled textile reinforcements were visually pre-inspected, and regions exhibiting visible fraying, local damage, contamination, severe crease marks, or obvious inhomogeneity were excluded from specimen preparation. Both recycled textile reinforcements were degreased before use; subsequently, the PA 6.6 fabric was mechanically roughened and the PES webbing was mechanically abraded (brushed) in order to improve interfacial adhesion to the epoxy matrix. Nevertheless, no separate intrinsic mechanical characterization of the recovered textile reinforcements was performed before composite fabrication; therefore, the possible effects of prior service history, ageing-related degradation, and feedstock variability could not be quantified and are acknowledged as a limitation of the present study.

2.4. Specimen configuration and designation

Six different composite configurations were prepared, along with a reference set of neat epoxy resin specimens designated EPOX. The ply sequences and specimen designations are summarised in Table 1. All laminates used CF245 as the first ply and CF600 as the third and fourth plies (or as the second and third plies in configuration C). The second ply consisted of AF173 or was replaced by PA 6.6. The outermost ply was PES, selected as the final layer because it enables local placement in critical regions of the body panel, with the aim of increasing impact resistance and improving energy absorption under loading.

Table 1.

Composition and designation of tested composite laminates.

Designation	Layer 1	Layer 2	Layer 3	Layer 4	Layer 5
C	CF245	CF600	CF600	-	-
AB	CF245	PA 6.6	CF600	CF600	-
P	CF245	CF600	CF600	PES	-
A	CF245	AF173	CF600	CF600	-
AP	CF245	AF173	CF600	CF600	PES
PB	CF245	PA 6.6	CF600	CF600	PES

The material parameters of the composite laminates for each configuration are reported in Table 2. The bulk density was calculated from the measured mass and geometric dimensions of the specimens. The fibre volume fraction (FVF) was also determined by calculation for the respective laminate configurations. These values represent bulk calculation-based parameters and do not fully account for local inhomogeneities related to manual hand lay-up processing, such as local thickness variation, resin-rich areas, or minor fabrication-induced imperfections.

Table 2.

Material parameters of the composite laminates for the individual configurations.

Designation	Thickness (mm)	Bulk density (kg/m³)	Fiber volume fraction (%)
EPOX	2.0 ± 0.2	1082	0
C	2.0 ± 0.2	1283	40
AB	2.1 ± 0.2	1276	45
P	3.8 ± 0.2	1108	43
A	2.7 ± 0.2	961	36
AP	4.5 ± 0.2	1131	37
PB	5.2 ± 0.2	1089	38

The surfaces of PA 6.6 (Figure 1) and PES (Figure 2) were roughened using a brushing process to enhance binder adhesion to the substrate. This treatment was implemented based on the results of previous adhesion tests evaluating the bonding performance of these materials with the epoxy matrix.

Figure 1.

Mechanical surface roughening of PA 6.6 airbag fabric by brushing. Bottom-left micrograph: prior to roughening (16× magnification). Bottom-right micrograph: after roughening (16× magnification).

Figure 2.

Mechanical surface roughening of PES seat-belt webbing by brushing. Bottom-left micrograph: prior to roughening (16× magnification). Bottom-right micrograph: after roughening (16× magnification).

Figure 3 illustrates the five-ply lay-up of specimen AP (left) and specimen PB (right), in which the second aramid ply (AF173) is replaced by the airbag fabric PA 6.6.

Figure 3.

Reinforcement layup of the specimens: set AP (left) and set PB (right).

2.5. Specimen preparation

Test specimens for mechanical characterisation were machined from the manufactured composite plates. The laminates were manufactured by hand lay-up under controlled ambient conditions of (23 ± 2) °C and (50 ± 10) % relative humidity.

• Tensile test: Specimens were prepared in accordance with ISO 527-4 (overall length 200 mm, width in the reduced section 10 mm, width in the gripping sections 20 mm, transition radius R = 60 mm).

• Charpy impact test (unnotched): Rectangular specimens with dimensions 80 × 20 mm (length × width) were prepared according to ISO 179-2. Specimen thickness corresponded to the thickness of the produced hybrid composites (2.0–5.2 ± 0.2 mm). Tests were performed on unnotched specimens (type 1eU).

2.6. Mechanical testing

Mechanical testing was performed at (23 ± 2) °C and (50 ± 10) % relative humidity (standard atmosphere 23/50, ISO 291).

• Tensile test: Tensile testing was carried out on a universal testing machine SHIMADZU AG-X plus (50 kN) at a constant crosshead speed of 10 mm/min, in accordance with ČSN EN ISO 527-4. The force–extension response was recorded. For each material set, seven specimens were tested. The mean tensile strength (σ_M) and the mean strain at break (ε_M) were determined.

• Charpy impact test (unnotched): Impact testing was performed using a LabTest CHK 30 J Charpy pendulum in accordance with ISO 179-2. The energy required to fracture an unnotched specimen (W) was measured. For each material set, seven specimens were tested. The unnotched impact strength (a_cU) was calculated as a_cU = 100·W/(b·h), where b is the specimen width (mm) and h is the specimen thickness (mm); the factor 100 converts the result from J/mm² to J/cm². Each specimen was evaluated individually, and the resulting values are reported in the tables and figures.

3. Results

3.1. Tensile properties

The mean values of tensile strength and strain at break for each composite configuration are summarised in Table 3. A graphical comparison of tensile strength, including standard deviations, is presented in Graph 1.

Table 3.

Average results of tensile tests.

Designation	Average tensile strength σ_m (MPa)	Standard deviation σ_m (MPa)	Average elongation ε_m (%)	Standard deviation εM (%)	Tensile modulus E (GPa)
EPOX	19	1.41	0.7	0.05	2.74 ± 0.12
C	423	51	1.5	0.15	33.93 ± 3.52
AB	403	18	1.5	0.07	34.06 ± 1.08
P	275	13	1.5	0.16	18.27 ± 0.65
A	450	34	2.8	0.13	28.98 ± 0.44
AP	290	21	2.3	0.14	14.60 ± 1.14
PB	175	21	2.0	0.08	13.05 ± 1.66

Graph 1.

Comparison of results of all tensile tests.

The highest mean tensile strength (450 MPa) was achieved by specimen A, reinforced with aramid fabric. This configuration also exhibited the highest mean strain at break (2.8%). Specimens C (all-carbon) and AB (carbon/nylon 6.6) showed comparable, slightly lower strengths (approximately 400–420 MPa) and low elongation (around 1.5%). In contrast, all configurations incorporating recycled seat-belt reinforcement (P, AP, PB) exhibited a pronounced reduction in tensile strength. The lowest strength (175 MPa) was recorded for specimen PB, which combined nylon 6.6 with seat-belt webbing. Specimens P and AP reached tensile strengths of approximately 275–290 MPa. The strain at break of the seat-belt-containing specimens was slightly higher than that of C and AB, but lower than that of A. Specimen P exhibited the lowest standard deviation in tensile strength, indicating the best repeatability. Representative fracture appearances after tensile testing are shown in Figure 5.

Figure 5.

Representative failure modes after tensile testing.

3.2. Impact strength (charpy)

The mean values of absorbed impact energy and unnotched Charpy impact strength for each composite configuration are summarised in Table 4. A graphical comparison of the mean impact strength values, including standard deviations, is presented in Graph 2.

Table 4.

Average results of unnotched Charpy impact tests (ISO 179-2, type 1eU).

Designation	Average impact work W (J)	Average notched impact strength a^cᵁ (J/cm²)	Standard deviation a^c (J/cm²)
EPOX	1.1	2.6	0.1
C	3.6	8.9	0.7
AB	5.5	12.7	2.2
P	8.5	11.4	2.1
A	5.6	10.4	1.4
AP	14.1	15.7	1.7
PB	13.9	13.4	1.1

Graph 2.

Average unnotched Charpy impact strength for individual laminate configurations.

Among the CFRP laminates, the lowest unnotched Charpy impact strength was observed for the all-carbon composite (C), reaching 8.9 J/cm². The incorporation of an aramid ply (A) resulted in a modest increase to 10.4 J/cm². A more pronounced improvement was obtained for the seat-belt-reinforced configuration (P, 11.4 J/cm²) and for the laminate reinforced with PA 6.6 (AB, 12.7 J/cm²). The highest impact strength values were achieved by PB (13.4 J/cm²) and AP (15.7 J/cm²), which combined carbon plies with either aramid fabric or PA 6.6 and additionally included recycled seat-belt reinforcement. These two configurations exhibited an increase in impact toughness of more than 60% compared with the reference carbon composite C. Standard deviations were generally higher for the tougher configurations (P, AB, AP), indicating increased variability in their energy-absorption capability. Representative fracture appearances after the impact test are shown in Figure 7.

Figure 7.

Representative fracture appearances after unnotched Charpy impact testing.

3.3. Analysis of fracture surfaces

Fracture surfaces of the carbon-fibre composite laminates after the Charpy impact strength test were documented using an Olympus SZX7 stereomicroscope at 16× magnification. The image series captures the damage morphology at the crack initiation region and along the fracture-propagation path, enabling qualitative identification of the dominant failure mechanisms of the laminates—particularly delamination, matrix cracking, and local ply splitting—which are subsequently discussed in relation to the measured impact response.

3.4. Specimen C

The fractographic image of the carbon laminate fracture surface (Figure 8) after the Charpy impact test shows a pronounced layered (“leaf-like”) failure morphology, characterised by multiple delaminations along interply interfaces accompanied by longitudinal ply splitting into thin lamellae. The fracture surface appears predominantly brittle, with no pronounced long fibre pull-out, while the dominant mechanisms include matrix cracking/matrix shear initiating damage on the tensile side, followed by interlaminar propagation (mixed-mode I/II). The observed peeling-type morphology is consistent with crack initiation on the tensile face and progressive delamination opening through the laminate thickness. Local pitting features may indicate matrix heterogeneity or porosity, which is typical of hand lay-up processing at a fibre volume fraction of approximately 40%.

Figure 8.

Fracture-surface detail of specimen C (16× magnification).

The crack-initiation detail (Figure 9) after the Charpy impact test for the hybrid laminate with a PA 6.6 interlayer reveals failure dominated by interlaminar delamination, with a pronounced opening of the interply gap and a clearly lifted laminate “flap” above the delaminated region. Along the interface, a continuous, lighter, wavy band is visible, indicating preferential crack propagation along a materially dissimilar interface (most likely within the CF245/PA 6.6 or PA 6.6/CF600 region). Local fibre bridging and signs of fibre pull-out span the opened gap, whereas the carbon plies exhibit a stepped, shear-dominated fracture morphology; extensive areal rupture of carbon bundles appears secondary in this region relative to interlaminar separation. Local surface irregularities and occasional void-like features may be associated with non-uniform consolidation/impregnation inherent to hand lay-up processing.

Figure 9.

Fracture-surface detail of specimen AB (16× magnification).

The crack-initiation detail (Figure 10) after the Charpy impact test shows failure dominated by interlaminar splitting and delamination, with a pronounced step-like fracture morphology in the carbon plies (CF245/CF600/CF600). In the upper region, a zone corresponding to the PES layer (seat-belt webbing) is evident, exhibiting abundant fibre bridging and fibre pull-out/fibrillation during crack opening, whereas extensive areal rupture of carbon fibre bundles appears secondary in this view relative to interlaminar separation. Locally observed voids/porosity within the carbon region may be associated with limited consolidation inherent to hand lay-up processing and with variable impregnation/adhesion at the interfaces. The observed combination of delamination and PES bridging is consistent with increased energy dissipation under impact loading, while the relatively higher scatter suggests that the response is sensitive to local heterogeneities in the damage processes.

Figure 10.

Fracture-surface detail of specimen P (16× magnification).

3.5. Specimen A

The crack-initiation detail (Figure 11) after the Charpy impact test for the hybrid laminate incorporating an AF173 aramid interlayer shows failure dominated by interlaminar delamination and ply splitting. Crack propagation is preferentially guided along the interface adjacent to the aramid layer, which appears in the section as a light band exhibiting pronounced fibrillation, fibre pull-out, and fibre bridging. The carbon plies (CF245/CF600/CF600) display a predominantly stepped, shear-driven fracture morphology with local peeling of individual plies, whereas extensive “clean” rupture of carbon fibre bundles is less dominant in this region. The presence of intense aramid pull-out/bridging is consistent with increased energy dissipation under impact loading, while the lower fibre volume fraction (FVF = 36%) may promote a higher matrix contribution and a greater extent of interlaminar damage.

Figure 11.

Fracture-surface detail of specimen A (16× magnification).

The crack-initiation detail (Figure 12) after the Charpy impact test for the hybrid laminate CF245/AF173/CF600/CF600/PES shows failure dominated by interlaminar delamination and ply splitting, with locally stepped, shear-driven fracture features in the carbon plies. In the toughening layers, pronounced activation of fibre bridging and fibre pull-out is evident, most notably in the lower region corresponding to AF173, where light-coloured, “fuzzy” extracted filaments indicate extensive aramid fibrillation and progressive tearing along the interface. In contrast, extensive areal “clean” rupture of carbon fibre bundles appears less dominant in this view. The fracture surface also exhibits local resin-rich regions and voids/porosity, which may be associated with non-uniform consolidation and impregnation inherent to hand lay-up processing. Overall, the observed combination of delamination and intensive bridging is consistent with increased energy dissipation under impact loading.

Figure 12.

Fracture-surface detail of specimen AP (16× magnification).

The crack-initiation detail (Figure 13) after the Charpy impact test for the hybrid laminate CF245/PA 6.6/CF600/CF600/PES shows failure dominated by interlaminar delamination with pronounced opening of interply gaps, together with ply splitting in the carbon layers and locally stepped, shear-driven fracture features. In the upper region, tensile filaments originating from the PES layer are visible, forming fibre bridging and local pull-out/fibrillation, whereas the carbon region exhibits marked separation of individual plies and longitudinal splitting of fibre bundles. The PA 6.6 interlayer likely acts as a preferential pathway (and/or initiator) for delamination growth along the adjacent interfaces. The combined occurrence of delamination and activation of ductile bridging mechanisms is consistent with increased energy dissipation under impact loading, while local voids and surface irregularities may be attributed to non-uniform impregnation and consolidation inherent to hand lay-up processing.

Figure 13.

Fracture-surface detail of specimen PB (16× magnification).

3.6. SEM examination of representative fracture surfaces

Following the stereomicroscopic analysis of the fracture surfaces, representative specimens were selected for more detailed SEM investigation. Laminates C, AP, and PB were chosen as characteristic examples of the reference and hybrid configurations. SEM imaging was performed at 120× magnification (500 µm field of view) in a sequence extending from the top surface towards the PES layer, allowing the fracture morphology to be examined across the laminate thickness. The corresponding figures and descriptions below present the dominant failure features observed in these selected laminates.

3.7. Specimen C

BSE-SEM observations of the reference carbon laminate C (CF245/CF600/CF600) after unnotched Charpy impact show a fracture morphology dominated by interlaminar, interface-parallel separation, manifested as large, relatively smooth lamellar fracture surfaces separated by a continuous open gap. Only limited evidence of progressive fibre-related mechanisms (bridging/pull-out) is apparent at this magnification; instead, the damage is consistent with predominantly matrix-assisted cracking and local ply splitting accompanying extensive delamination. This brittle, interface-controlled failure scenario aligns with the lower impact performance of C (acU = 8.9 ± 0.7 J/cm²), where energy absorption is mainly associated with delamination surface formation rather than frictional dissipation from tough textile interlayers. Bright BSE speckles cannot be assigned to a specific phase without EDS verification. From a larger SEM dataset, representative fields of view were selected to document the dominant fracture morphology of the reference laminate on Figure 14.

Figure 14.

BSE-SEM fracture surface of specimen C showing dominant interlaminar delamination; 120×, 500 µm.

3.8. Specimen AP

BSE-SEM images show that the AP laminate (CF245/AF173/CF600/CF600/PES) (divided step by step into the Figure 15) failed predominantly by interlaminar damage, characterized by extensive delamination, repeated crack deflection and branching, and local ply splitting within the carbon plies. In the region of the aramid interlayer (AF173), fibrillation and fiber bridging associated with progressive pull-out are evident, whereas towards the PES layer, exposed bundles/yarns with pronounced pull-out and interfacial debonding predominate, indicating substantial frictional energy dissipation during the progressive opening of delaminations. This combination of damage mechanisms is consistent with the high impact response of the AP laminate (acU = 15.7 ± 1.7 J/cm²), where energy absorption occurs primarily through the formation of delamination surfaces, crack arrest at hybrid interfaces, and fiber bridging/pull-out. Bright BSE spots/areas cannot be unambiguously assigned to a specific phase without EDS analysis; they may correspond either to inclusions or to topographic/charging effects, and their interpretation therefore remains uncertain at this stage. From the broader SEM image series, representative regions capturing the dominant failure mechanisms across the laminate thickness were selected.

Figure 15.

BSE-SEM fracture surface of specimen AP showing dominant interlaminar delamination; 120×, 500 µm.

3.9. Specimen PB

BSE-SEM images of the PB laminate (120×; 500 µm; sequence from the top surface towards the PES layer) (divided step by step into the Figure 16) reveal predominantly interlaminar failure of the hybrid CF245/PA6.6/CF600/CF600/PES laminate, characterized by extensive interlaminar delamination, repeated crack deflection and branching, and local ply splitting within the carbon plies. Towards the bottom side, the failure mode transitions to a more textile-controlled character: in the PES layer region, exposed bundles/yarns with pronounced pull-out and interfacial debonding predominate, promoting frictional energy dissipation during the progressive opening of delaminations. This combination of mechanisms (delamination + crack deflection/branching + PES pull-out/debonding) is consistent with the high impact response of the PB laminate (acU = 13.4 ± 1.1 J/cm²), where energy absorption occurs primarily through the formation of delamination surfaces and friction associated with bundle pull-out. Local bright BSE bands/areas cannot be unambiguously assigned to a specific phase without EDS analysis; they may correspond either to inclusions or to topographic/charging effects, and their interpretation therefore remains uncertain at this stage. From the larger BSE-SEM image series, representative regions across the laminate thickness were selected for inclusion in the main text.

Figure 16.

BSE-SEM fracture surface of specimen PB showing dominant interlaminar delamination; 120×, 500 µm.

4. Discussion

The tensile results indicate that the reference carbon laminate C achieves a tensile strength of 423 MPa, whereas the configuration with an aramid interlayer (A) exhibits the highest strength (450 MPa) together with a higher strain at break, consistent with the expected contribution of the tougher aramid component to improved damage tolerance. In contrast, hybrids incorporating recycled automotive textiles show a pronounced load-carrying trade-off: AP reaches 290 MPa and PB 175 MPa, and the final failure is accompanied by progressive ply-by-ply damage and delamination. The strength reduction can be attributed to a combination of locally poorer impregnation of woven/recycled structures and, most importantly, limited interfacial shear stress transfer across the textile–epoxy interface, which is typical for hybrid FRP systems when the interface is not deliberately optimised.^7,18

The most significant benefit of the recycled textiles is reflected in the impact response. Charpy testing demonstrated that AP and PB achieve the highest unnotched impact strength values of 15.7 and 13.4 J/cm², respectively, compared with only 8.9 J/cm² for the reference C (i.e., an increase of approximately 76% and 51%). In terms of the absolute absorbed energy W, AP (14.1 J) and PB (13.9 J) also differ markedly from C (3.6 J), confirming the enhanced energy-dissipation capacity based on direct comparison of the measured energy. This increase can be interpreted as a combined effect of (i) a higher extent of interply delamination, (ii) fibre bridging and pull-out within the tougher textile layers, and (iii) local plastic deformation and frictional interactions within the woven reinforcements—mechanisms commonly associated with increased SEA and more progressive failure in hybrid laminates.^9,20,21

SEM examination of representative fracture surfaces corroborates this interpretation of the Charpy response. The reference laminate C exhibits predominantly interface-parallel interlaminar separation, appearing as relatively smooth, lamellar fracture surfaces separated by a continuous open gap, with only limited evidence of progressive fibre-related dissipation at the investigated scale. In contrast, the hybrid laminates show a markedly more progressive, multi-path fracture morphology dominated by extensive delamination accompanied by repeated crack deflection/branching and local ply splitting within the carbon plies. For AP, the AF173 interlayer displays fibrillation with pronounced fibre bridging and progressive pull-out, while toward the PES belt side exposed bundles/yarns with pull-out and interfacial debonding are evident; PB shows an analogous transition toward a textile-controlled region near PES, where bundle pull-out/debonding accompanies delamination growth. These features are consistent with the higher absorbed energies measured for AP and PB, as additional energy is dissipated through delamination-surface formation and frictional interactions associated with bridging and pull-out at hybrid interfaces. Moreover, the comparatively larger scatter observed for textile-containing configurations is plausibly linked to local variability in impregnation and interfacial bonding inherent to hand lay-up, particularly for thick textile plies.

The present results are in qualitative agreement with previously published studies on hybrid carbon-based laminates. The increase in impact toughness observed for AP and PB, together with the reduction in tensile strength relative to the reference carbon laminate, is consistent with the behaviour reported for carbon/aramid hybrid epoxy systems, in which tougher interlayers promote more progressive failure, delamination, fibre bridging, and enhanced energy absorption, while the static load-carrying capability remains sensitive to interfacial quality and layer placement.^9,18 For example, Pincheira et al.¹⁸ reported increases of 37.9% in energy absorption, 12.7% in fracture resistance at crack initiation, and 43% in steady-state fracture resistance, together with an increase in ultimate strain of up to 19.5% and through-thickness compression strength of 8.3%, although accompanied by slight reductions in stiffness. Similarly, Ying et al.⁹ showed that interply carbon–aramid hybrid laminates with carbon fabric on the impacted surface exhibited improved impact damage resistance and damage tolerance under the same impact energy. Although direct quantitative comparison is limited by differences in laminate architecture, fibre volume fraction, and manufacturing route, the same overall trend is evident in the present study: tougher hybrid interlayers improve impact-related performance, whereas static tensile properties remain more sensitive to interface quality and structural arrangement. This is further supported by comparison with AB and P, where a higher fibre volume fraction (AB: 45%; P: 43%) did not by itself result in superior impact performance. By contrast, AP and PB, despite their lower FVF (37–38%) and greater thickness (4.5–5.2 mm), provided more interfaces capable of delamination and energy dissipation, albeit at the expense of tensile strength. Among the studied hybrids, AP showed the most favourable balance between strength and toughness, whereas the weaker tensile response of PB suggests a stronger dependence of the PA 6.6/PES-based laminate on impregnation quality, interlaminar bonding, and the surface condition of the recycled nylon-based reinforcement.^17,22 More specifically, the present findings for the PA 6.6- and PES-containing laminates are consistent with the broader behaviour reported for textile-based interlayers in carbon fibre composites, where improved impact-related response is often achieved at the expense of static tensile performance when the textile phase is more compliant or when interfacial bonding is not fully optimised.^17,22 In this sense, the present AP and PB results broaden the existing understanding derived mainly from aramid-based hybrid systems toward recycled nylon- and polyester-derived textile reinforcements used in carbon/epoxy laminates.

From a practical standpoint, the results indicate that recycled seat-belt and airbag layers can be used as targeted energy-absorbing interlayer reinforcements—particularly in regions exposed to localised impact—without the need to sacrifice the static load-bearing capacity of the entire component. Such textile reinforcements are not intended for full-area substitution across an entire body panel, but rather for local incorporation into specifically selected zones where increased energy absorption and damage tolerance are required. In practical vehicle structures, local safety reinforcements are commonly used only in designated regions, for example as reinforcing members inside doors or as energy-absorbing elements in frontal impact areas. From this perspective, the limited width of seat-belt webbing does not necessarily represent a disadvantage, but rather defines its most suitable use as a local reinforcing layer or insert. Although such reinforcement cannot provide uniform strengthening over the full panel area, its local incorporation into side or front body regions may increase local impact energy dissipation and thereby contribute to improved occupant protection. To improve repeatability while mitigating the tensile-strength penalty, further work should focus on (i) surface treatment of the recycled textiles (degreasing, activation/plasma treatment, compatible sizing, and/or chemical modification) and moisture control of PA 6.6 prior to lamination, (ii) optimisation of ply sequencing and local placement of recycled interlayers, and (iii) process measures to reduce variability inherent to hand lay-up (improved consolidation, vacuum bagging, and control of thickness and resin content).

5. Conclusions

This work demonstrates that recycled technical textiles from the automotive industry—specifically airbag fabric (PA 6.6) and seat-belt webbing (PES)—can be employed as functional reinforcing plies in hybrid carbon/epoxy laminates manufactured using a simple hand lay-up process, which is typical of bespoke or low-volume production of body panels.

From a passive-safety (crashworthiness) perspective, two configurations with the highest energy-absorption capability were identified: AP and PB. These laminates exhibited the highest unnotched Charpy impact strength (15.7 and 13.4 J/cm²) and the highest absorbed fracture energy (W = 14.1 and 13.9 J), outperforming the reference all-carbon laminate C (8.9 J/cm²; W = 3.6 J). The results indicate that an appropriately selected combination of carbon plies with tougher textile interlayers promotes energy-demanding damage mechanisms (delamination, fibre bridging, and pull-out), leading to a more progressive impact failure response.

At the same time, the expected trade-off was confirmed: the improved impact strength in configurations AP and PB was achieved at the expense of reduced tensile strength (290 MPa and 175 MPa, respectively, compared with 423 MPa for C). From a practical standpoint, it is therefore reasonable to consider local placement of recycled plies only in regions with a high risk of impact damage (e.g., exposed edges and impact-prone zones), while the global load-carrying capacity of the component can continue to be provided by the primary carbon load-bearing plies.

BSE-SEM fractography of selected representative specimens (C, AP and PB) further supports these trends by linking the measured impact toughness gains to a transition from predominantly interface-parallel delamination in the reference laminate to a more progressive, multi-path damage mode in the hybrids. In AP and PB, extensive delamination is accompanied by crack deflection/branching and by textile-mediated bridging and pull-out (AF173 fibrillation/bridging in AP and pronounced PES bundle pull-out/debonding in AP and PB), providing additional frictional energy dissipation during impact.

The innovativeness of the proposed approach lies in the upcycling of components that typically have no direct reuse after end-of-life (airbags and seat belts). Their conversion into functional reinforcing layers within composite laminates can offer a dual benefit for specific applications: reduced environmental burden and improved capacity of body panels to absorb energy under local impact loading, i.e., enhanced passive safety—particularly in bespoke or low-volume vehicle production.

Future work should focus on optimising the adhesion of recycled textiles to the epoxy matrix (surface treatments, compatible sizing) and on quantifying property scatter arising from variability in the recycled feedstock. It will also be beneficial to extend the test matrix to include post-impact residual strength assessment and to validate performance in realistic structural joints (e.g., bonded/bolted connections) relevant to composite body structures.

Footnotes

Acknowledgements

This article was supported by project CZ.02.01.01/00/22_008/0004631, “Materials and Technologies for Sustainable Development”, funded by the European Union and the state budget of the Czech Republic within the framework of the Jan Amos Komensky Operational Program.

ORCID iD

Marek Beseda

Funding

The authors received no financial support for the research, authorship, and/or publication of this article.

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

Data Availability Statement

The datasets generated and/or analysed during the current study are available.*

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