Advancing Sustainable Materials Engineering with Natural-Fiber Biocomposites

Abstract

Natural-fiber biocomposites are increasingly viewed as promising materials for sustainable engineering. However, their broader adoption remains constrained by coupled challenges related to interfacial compatibility, moisture sensitivity, environmental durability, processing limitations, and end-of-life trade-offs. Rather than treating fiber selection, matrix chemistry, processing routes, durability, and sustainability as independent considerations, this review emphasizes their interdependence through the fiber–matrix interface, which governs stress transfer, moisture transport, and long-term property evolution. It provides a comprehensive and integrative analysis of natural-fiber–reinforced polymer composites, encompassing plant-, animal-, and emerging bio-derived reinforcements combined with bio-based, biodegradable, and selected synthetic matrices. Comparative analysis across the literature demonstrates that interfacial engineering consistently dominates mechanical performance, moisture resistance, and property retention, while mediating trade-offs among stiffness, toughness, recyclability, and biodegradability. Moisture transport and environmental ageing are examined using thermodynamic and diffusion-controlled frameworks that link fiber chemistry, interfacial energetics, swelling, and debonding to performance degradation. Fire behavior and flame-retardant strategies are reviewed with attention to heat-release control and their implications for durability and circularity. Processing routes, including extrusion, injection molding, compression molding, resin transfer molding, and additive manufacturing, are assessed with respect to fiber dispersion, thermal stability, scalability, and compatibility with bio-based systems. By integrating structure–property relationships, processing science, durability mechanisms, and sustainability considerations, this review clarifies how natural-fiber biocomposites can be designed to achieve balanced performance, environmental stability, and circular life-cycle behavior, thereby providing guidance for the development of systems suitable for near-term engineering applications.

1. Introduction

Biocomposites, comprising polymer matrices reinforced with naturally derived fibrous or bio-derived phases, have gained prominence as sustainable, high-performance alternatives to conventional synthetic composites [1]. By coupling renewable fibers with polymer matrices, these materials enhance stiffness, load-bearing capacity, and overall mechanical performance through improved interfacial interactions and efficient stress transfer from the matrix to the fibers [2,3,4]. Here, stress transfer denotes the transmission of mechanical load across the fiber–matrix interface via interfacial shear stresses, enabling effective load sharing and reinforcement. As a result, the fiber–matrix interface plays a central role in governing deformation mechanisms and long-term durability under environmental exposure. This review therefore focuses on how interface-controlled mechanisms link performance, durability, and sustainability across material classes.

Biocomposites may be classified as fully bio-based when both the matrix and reinforcement originate from renewable resources, or partially bio-based when natural fibers are incorporated into synthetic polymer matrices [5,6]. Although synthetic composites such as glass fiber–reinforced plastics have long dominated structural and transportation sectors due to their high stiffness and durability [7,8,9], increasing concerns regarding non-degradability, microplastic pollution, and energy-intensive production have driven growing interest in more sustainable material alternatives [10,11,12]. Importantly, however, the sustainability of biocomposites is not guaranteed by material selection alone; it is critically influenced by fiber–matrix interfacial compatibility, processing energy demands, durability under service conditions, and end-of-life management strategies.

Modern biocomposites offer the potential to reduce carbon footprint, improve end-of-life outcomes, and deliver mechanical performance suitable for transportation, construction, packaging, biomedical, and consumer applications [13,14,15,16]. Their design relies on informed selection of fiber and matrix systems, pairing plant- or animal-based reinforcements with thermoset, thermoplastic, or biodegradable polymers to achieve functional properties comparable to synthetic composites while enabling circular-material strategies. A comparison of key mechanical, environmental, and processing characteristics distinguishing natural and synthetic fiber reinforcements is provided in

Table 1. Comparison of natural and synthetic fiber reinforcements across mechanical, environmental, and processing/safety attributes. The “Primary Advantage” column indicates which fiber class provides a relative benefit for each criterion. Overall, natural fibers offer advantages in sustainability and specific weight, whereas synthetic fibers provide superior absolute mechanical performance and long-term durability.

Domain	Criterion	Natural Fibers	Synthetic Fibers	Primary Advantage
Mechanical	Density/specific weight	Low density; favorable specific properties	Higher density	Natural
	Strength/stiffness (typical)	Moderate	High	Synthetic
	Long-term durability (fatigue, ageing)	Moderate; moisture- and interface-sensitive	High; environmentally stable	Synthetic
Environmental	Feedstock origin	Renewable	Petrochemical	Natural
	Embodied energy/carbon footprint	Lower	Higher	Natural
	End-of-life pathways	Biodegradable and/or recyclable (system-dependent)	Persistent; recycling varies by system	Natural
Processing and safety	Manufacturing control and consistency	Simpler processing; higher feedstock variability	Tighter control; consistent properties	Synthetic
	Moisture sensitivity (processing and service)	Higher	Lower	Synthetic
	Health, cost, and availability (typical)	Lower hazard; lower cost; widely available	Higher cost; potential handling concerns; region-dependent	Natural

.

1.1. Natural Fiber Reinforcements

Natural fibers used in biocomposites are primarily classified as plant-based or animal-based, with plant fibers dominating due to their abundance, high cellulose content, low density, and favorable cost-to-performance ratio. Bast fibers such as flax, hemp, jute, kenaf, and ramie offer high tensile strength and aspect ratios that promote efficient stress transfer within composite matrices [17,18]. Extracted from the phloem of dicot stems through retting, scutching, and hackling, these fibers typically exhibit cellulose contents of 60–80% and are widely used in automotive, marine, and construction applications [19,20,21,22,23,24,25,26,27,28]. Other plant-fiber categories include seed fibers (cotton, kapok), leaf fibers (sisal, abacá), fruit fibers (coir, oil palm), grass and reed fibers (bamboo, bagasse, rice husk), and wood fibers (softwood and hardwood pulps), each possessing distinct anatomical origins and mechanical characteristics. The botanical classification of plant-based natural fibers and their representative applications is summarized in

Table 2. Botanical classification of plant-based natural fibers and their typical applications.

Category	Plant Part (Origin)	Examples of Fibers	Source Plants (Examples)	Typical Uses
Seed Fibers	Hair-like fibers from seed or fruit surface	Cotton, coir (from coconut husk), kapok	Gossypium spp., Cocos nucifera, Ceiba pentandra	Textiles, padding, ropes
Bast Fibers (Phloem Fibers)	Inner bark (phloem) of dicot stems	Flax, jute, hemp, kenaf, ramie, roselle	Linum usitatissimum, Corchorus capsularis, Cannabis sativa, Hibiscus cannabinus, Boehmeria nivea, Hibiscus sabdariffa	Ropes, sacks, composites, linen
Leaf Fibers	Fibers from vascular bundles of leaves	Sisal, abacá, henequen, pineapple (PALF)	Agave sisalana, Musa textilis, Agave fourcroydes, Ananas comosus	Cordage, mats, composites
Fruit Fibers	From fruit mesocarp or husk	Coir (coconut), oil palm fiber	Cocos nucifera, Elaeis guineensis	Brushes, mats, insulation
Grass and Reed Fibers	Stems of monocots, including grasses	Bamboo, bagasse, rice husk, wheat straw	Bambusa spp., Saccharum officinarum, Oryza sativa, Triticum aestivum	Pulp, boards, paper, composites
Wood Fibers	Xylem tissue of trees	Softwood and hardwood pulp	Pinus spp., Eucalyptus spp.	Paper, rayon, viscose
Other Specialty Fibers	Various plant parts	Banana pseudostem, nettle, kapok, mycelium (emerging)	Musa paradisiaca, Urtica dioica, Ceiba pentandra, fungal species	Textiles, biocomposites

.

Beyond plant-based reinforcements, a distinct class of natural fibers is derived from animal and marine biological systems, offering complementary mechanical and functional attributes. Animal fibers, including silk, wool, and horsehair, provide elasticity, toughness, and biocompatibility, making them valuable for biomedical, protective, and structural applications [29,30,31,32,33]. Keratin-based fibers such as hair, wool, and feathers are proteinaceous and sulfur-rich, exhibiting resilience, elasticity, and moisture absorption arising from extensive disulfide cross-linking within the keratin matrix. Silk fibers, secreted by silkworms and spiders, are renowned for their exceptional tensile strength and toughness, which derive from the hierarchical arrangement of $\beta$-sheet nanocrystals within fibroin, imparting a unique balance of stiffness and extensibility. Marine- and insect-derived fibers, including byssal threads and non-silkworm silks, are increasingly investigated for their adhesive, elastic, and biodegradable characteristics, offering opportunities for advanced biomaterial and composite applications. Collagen- and chitosan-based fibers, although reconstituted rather than naturally fibrous, are valued for their biocompatibility, tunable degradation, and applicability in wound healing, tissue engineering, and regenerative scaffolds. The principal categories, biological origins, and applications of animal-derived fibers are summarized in

Table 3. Classification of animal-based natural fibers by biological origin and structural composition.

Category	Biological Source	Examples	Principal Protein	Typical Applications
Hair and Wool Fibers	Mammalian body hair and fleece	Wool (sheep), cashmere (goat undercoat), mohair (Angora goat), alpaca, camel, llama, vicuña	$\alpha$-Keratin	Textiles, insulation, high-performance fabrics
Silk Fibers	Secretions from silkworm or spider glands forming continuous filaments	Mulberry silk (Bombyx mori), tussar silk, eri silk, muga silk, spider silk	Fibroin (core) and sericin (gum coating)	Luxury textiles, sutures, biomedical scaffolds
Avian Fibers (Feather, Down)	Epidermal appendages of birds	Chicken feather fibers, duck and goose down	$\beta$-Keratin	Thermal insulation, lightweight composites
Insect Fibers (Non-Silk)	Secreted structural or defensive fibers from insects	Honeybee cocoon silk, wasp silk (non-commercial)	Fibroin-like proteins	Bio-inspired materials, experimental fibers
Marine Animal Fibers (Byssus, Sponge, Shell Fibers)	Secretions or structural fibers from marine organisms	Byssus threads (Pinna nobilis), sponge spicules, chitinous crab or prawn fibers	Byssal proteins, spongin, chitin	Adhesives, composites, biomedical materials
Other Animal Secretions and Derived Fibers	Structural or processed animal biopolymers	Collagen, elastin, chitosan fibers	Collagen, elastin, chitosan (from chitin)	Wound dressings, tissue engineering, biodegradable composites

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1.2. Polymer Matrices for Biocomposites

Table 4. Classification of polymer matrices used in biocomposites based on origin and biodegradability.

Category	Origin/Description	Representative Polymers	Key Features and Applications
Bio-Based and Biodegradable	Derived from renewable biological feedstocks and degradable under microbial or enzymatic action	Polylactic acid (PLA), polyhydroxyalkanoates (PHB, PHBV, PHBHx), poly(butylene succinate) (PBS), thermoplastic starch (TPS)	Renewable and compostable; widely used in packaging, biomedical devices, and short-lifetime structural composites
Bio-Based and Non-Biodegradable	Derived partly or wholly from renewable resources but chemically durable	Bio-polyethylene (Bio-PE), bio-polypropylene (Bio-PP), bio-poly(ethylene terephthalate) (Bio-PET), epoxidized soybean oil-based resins	Renewable carbon feedstocks combined with long-term durability; used in automotive and structural biocomposites
Petrochemical-Based but Biodegradable	Synthetic origin but designed to degrade under environmental or biological conditions	Polycaprolactone (PCL), poly(butylene adipate-co-terephthalate) (PBAT), poly(vinyl alcohol) (PVA)	Flexible and biodegradable; common in biomedical, packaging, and agricultural composites
Petrochemical-Based and Non-Biodegradable	Conventional synthetic matrices combined with natural fibers for partial sustainability	Polypropylene (PP), polyethylene (PE), epoxy, unsaturated polyester, polyurethane (PU)	High mechanical and thermal performance; widely used in automotive, marine, and construction biocomposites
Natural Polymer Matrices	Polymers obtained directly from biological sources, with or without chemical modification	Cellulose, chitosan, alginate, gelatin, soy protein isolate	Fully renewable and biocompatible; applied in biomedical scaffolds, films, and environmentally friendly composites

categorizes polymer matrices used in biocomposites by feedstock origin (bio-based vs. petrochemical) and biodegradability, providing a concise framework for material selection under sustainability and end-of-life constraints. The choice of matrix—thermoplastic or thermoset, biodegradable or durable—fundamentally defines the attainable property envelope of natural-fiber composites, governing stiffness and strength, moisture sensitivity, thermal stability, and processing windows.

Figure 1 presents the repeat-unit structures of representative biodegradable and bio-based thermoplastic matrices, including polyhydroxyalkanoates (PHB, PHV, PHBV, PHBHx), other aliphatic polyesters (PLA, PBS, PCL), and the aromatic–aliphatic copolyester PBAT, along with acrylated epoxidized soybean oil (AESO) as a bio-derived thermoset precursor. Variations in side-chain length, aromaticity, and triglyceride functionality directly control crystallinity, polarity, and melt or cure behavior, thereby dictating moisture uptake, interfacial adhesion, and the mechanical response of the resulting composites.

1.3. Fiber–Matrix Interactions and Composite Performance

The performance of biocomposites arises from the coupled interaction between the reinforcement and the polymer matrix. Natural fibers provide strength, stiffness, and toughness, while the matrix governs stress transfer, cohesion, and environmental stability. Plant fibers, composed primarily of cellulose, hemicellulose, and lignin, dominate biocomposite applications due to their abundance, low density, and high specific strength. Protein-based animal fibers such as keratin and fibroin are used in specialized applications requiring elasticity or biocompatibility.

A persistent limitation of natural-fiber composites resides at the fiber–matrix interface. The intrinsic hydrophilicity of lignocellulosic fibers, together with surface impurities and thermal sensitivity, often leads to weak interfacial bonding with relatively hydrophobic polymer matrices. This mismatch manifests as inefficient stress transfer, moisture-induced swelling, fiber pull-out, and premature degradation. Consequently, interfacial engineering, through surface treatments, compatibilization, and interphase design, has emerged as a dominant strategy for achieving competitive mechanical performance and long-term environmental durability in biocomposites.

1.4. Interfacial Challenges and Modification Strategies

Challenges, including fiber hydrophilicity, property variability, thermal degradation, and limited compatibility with hydrophobic polymers, continue to constrain composite performance. To mitigate these limitations, chemical, physical, and enzymatic surface-modification strategies are widely employed to enhance interfacial bonding and dimensional stability, thereby improving composite effectiveness [8,34,35]. At the same time, improvements in mechanical performance are frequently accompanied by trade-offs in biodegradation rate, recyclability, or processing window, underscoring the need for an integrated evaluation framework.

1.5. Scope of This Review

Previous reviews have provided detailed coverage of fiber types, polymer matrices, life-cycle assessment frameworks, and processing routes for natural-fiber composites [1,8,13,14,36,37,38]. While these contributions have established a strong foundation, key aspects are often treated in isolation rather than as interdependent design variables that jointly govern performance, durability, and sustainability. Accordingly, this review adopts an integrative, interface-centered design framework.

The review is organized around four tightly coupled dimensions: (i) fiber–matrix interfacial chemistry, (ii) moisture transport and environmental durability, (iii) processing and thermal constraints, and (iv) sustainability and end-of-life trade-offs. Rather than providing an exhaustive catalog of materials, emphasis is placed on comparative trends and matrix-dependent behavior, supported by targeted tables that synthesize representative experimental results. Applications in automotive, construction, marine, aerospace, packaging, and biomedical sectors are discussed with attention to technological readiness, regulatory constraints, and realistic performance limitations.

Prior reviews have examined the economic performance and life-cycle impacts of natural-fiber composites, particularly in hybrid systems, consistently showing that meaningful sustainability and cost assessments require clearly defined material, processing, and application boundaries [39,40,41,42]. Even when formal life-cycle assessment (LCA), life-cycle costing (LCC), or techno-economic frameworks are applied, reported metrics depend strongly on baseline materials, fiber content, and manufacturing routes, as well as on interface-driven requirements such as fiber drying, compatibilizer addition, surface treatments, and thermal processing windows, which simultaneously influence both performance and cost [40,41,42,43]. Consequently, direct normalization of property gains, cost, and environmental impact across the broader literature remains impractical. Emphasis is therefore placed on comparative trends and matrix-dependent trade-offs rather than absolute metrics, enabling synthesis across a heterogeneous and rapidly evolving field. Figure 2 provides a structural guide to this perspective.

2. Materials for Bio-Based Composites

This section establishes the material descriptors governing subsequent interface design, moisture transport, and durability. Figure 3 summarizes the materials landscape addressed here, linking reinforcement and matrix families to their intrinsic physicochemical characteristics that govern compatibility and processing constraints.

2.1. Definitions

Because terminology related to material origin, degradability, and end-of-life pathways underpins the interpretation of durability, degradation, and disposal behavior throughout this review, three widely used descriptors are defined here. Bio-based refers to polymers or reinforcements derived partially or entirely from renewable biological feedstocks, irrespective of their degradability or disposal pathway. Biodegradable materials are those capable of being mineralized by microorganisms into carbon dioxide (or methane), water, and biomass under specified environmental conditions. Compostable materials constitute a regulated subset of biodegradable systems that must satisfy standardized criteria for disintegration rate, mineralization, and absence of ecotoxic residues under controlled composting conditions (e.g., ISO 17088 [44], ASTM D6400 [45]).

Crucially, these attributes are orthogonal rather than hierarchical: a material may be bio-based yet non-biodegradable (e.g., bio-PE), biodegradable yet fossil-derived (e.g., PBAT), both, or neither. Recognizing this decoupling is essential for rational matrix selection and for interpreting sustainability metrics, since feedstock origin, service durability, and end-of-life fate impose distinct and sometimes competing constraints on composite design.

2.2. Plant-Based Fibers

Plant fibers are widely used as reinforcements in biocomposites owing to their high specific stiffness, renewability, and low cost. Flax and hemp are therefore employed here as benchmark bast fibers (see Section 2.5 for justification), while other plant fibers are treated more concisely where the available data are comparatively limited.

2.2.1. Hierarchical Structure of Bast Fibers in Plant Stems

In bast-fiber plants such as flax and hemp, load-bearing fibers are located in the outer stem tissues rather than in the woody core. The stem is organized into concentric anatomical regions consisting of an outer epidermis, a cortex, a fibrous phloem (bast) zone, an inner xylem (woody core), and a central pith [46,47,48] (Figure 4). The epidermis is a thin, protective layer enriched in cutin and waxes. Beneath it lies the cortex, a relatively soft tissue composed primarily of parenchyma cells that provides metabolic, storage, and mechanical separation functions between the epidermis and the fibrous tissues [47]. Inside the cortex is the phloem, which forms the bast region and contains the fibers harvested for composite reinforcement.

The term bast fiber therefore refers to the anatomical origin of the fibers, namely their extraction from the phloem of dicotyledonous stems such as flax, hemp, jute, and kenaf, in contrast to leaf fibers (e.g., sisal) or seed fibers (e.g., cotton) [46,49]. Within the phloem, fibers occur as multicellular bundles. In flax, these are primary bast fiber bundles that form early during stem elongation, whereas hemp contains both primary and secondary bast fibers arranged in concentric rings as the stem thickens [46]. These fiber bundles are embedded in a pectin-rich matrix and separated from the xylem by a cambial zone that controls radial growth.

Each technical fiber used in composites is therefore a bundle of elementary fibers, which are individual elongated plant cells typically 10–30 µm in diameter and several millimetres to centimetres in length [47,49]. These cells are bonded together by the middle lamella, an intercellular layer rich in pectins, hemicelluloses, and minor amounts of lignin, which governs bundle cohesion, retting behavior, and separability during processing [46].

2.2.2. Cell-Wall Architecture of Elementary Bast Fibers

From a composite-engineering perspective, bast fibers are hierarchical natural composites: stiff, crystalline cellulose microfibrils embedded in a softer polysaccharide matrix at the nanoscale, assembled into multilayered cell walls, combined into elementary fibers, and further bundled into technical fibers. Each elementary fiber consists of a central lumen surrounded by a multilayered cell wall. From outside to inside, the wall comprises the middle lamella, a thin primary wall, and a secondary wall subdivided into S1, S2, and S3 layers [47,50]. The middle lamella contains pectins and hemicelluloses and acts as the adhesive between neighboring fibers. The primary wall is thin and relatively disordered, contributing little to axial stiffness but playing a role in flexibility and fiber separation.

The secondary wall, particularly the S2 layer, is the mechanically dominant component. In bast fibers, this layer, often termed the gelatinous or G-layer, constitutes most of the wall thickness and contains densely packed, highly oriented cellulose microfibrils aligned nearly parallel to the fiber axis [47]. The S1 and S3 layers are thin and contain microfibrils at high angles, whereas the S2 layer has a low microfibrillar angle, which is the primary reason for the exceptional stiffness and tensile strength of flax and hemp fibers [50,51]. Fiber extraction and retting processes act primarily by degrading the middle lamella and surrounding cortical tissues, thereby liberating these mechanically optimized elementary fibers for use as reinforcement in polymer composites [46,49].

2.2.3. Fiber Extraction and Retting Processes

Manian et al. [52] present a comprehensive review of flax and hemp fiber extraction, detailing plant anatomy, retting and delignification mechanisms, and emerging enzymatic and chemical processes that enhance fiber quality and sustainability for biocomposite applications. Retting remains the most widely adopted industrial process for extracting fibers from bast sources such as flax, hemp, and jute. In the native stem, elementary fibers are bound into multicellular bundles by a pectin-rich intercellular matrix that also contains hemicellulose and smaller amounts of lignin, particularly within the middle lamella between adjacent cells. Retting processes are therefore designed to selectively degrade or solubilize this non-cellulosic matrix, thereby releasing fiber bundles from surrounding tissues and reducing bundle cohesion for subsequent mechanical separation [53].

Water and dew retting represent the most common low-cost approaches, relying on microbial and enzymatic activity to hydrolyze pectins and partially remove hemicelluloses. Water retting, conducted under anaerobic conditions, typically yields cleaner, more uniformly retted fibers, whereas dew retting proceeds more slowly and heterogeneously under field conditions [54]. In both cases, excessive retting time can lead to progressive deterioration of tensile properties due to partial cellulose hydrolysis and microbial erosion of the fiber cell wall [55].

Mechanical decortication provides rapid physical separation of bast fibers from the woody core but does not remove the pectin-rich middle lamella, resulting in fibers that retain higher levels of non-cellulosic components and exhibit strong bundle cohesion and limited fibrillation [54]. Chemical and physicochemical degumming processes, including alkaline, oxidative, or osmotic treatments, accelerate the extraction of pectins and hemicelluloses and can partially remove lignin, producing fibers with higher cellulose content, greater individualization, and improved surface accessibility for polymer impregnation. However, these treatments must be carefully controlled to avoid excessive degradation or surface modification that could adversely affect fiber strength or interfacial compatibility.

Enzymatic retting, based primarily on pectinolytic and hemicellulolytic enzymes, offers a more selective and environmentally benign route to fiber separation, enabling controlled removal of the middle-lamella matrix and improved consistency of fiber quality, although large-scale implementation is still limited by cost and process complexity. Hybrid strategies, such as short-duration microbial or enzymatic retting followed by mechanical decortication, represent promising compromises between fiber quality, throughput, and sustainability, and are increasingly explored for the production of high-performance bast-fiber reinforcements for polymer composites.

2.2.4. Chemical Composition of Bast Fibers

The chemical composition of bast fibers is not a fixed material property but varies systematically with both botanical origin and extraction history, which in turn governs moisture uptake, interfacial chemistry, and durability in biocomposites. For hemp, comprehensive datasets are available from Zimniewska and Romanowska [56] and Liu et al. [57]. Zimniewska and Romanowska provide cultivar- and process-resolved measurements showing that retting and degumming methods strongly modulate fiber chemistry: water-retted hemp typically contains about 71–73 wt% cellulose with low lignin (2–3 wt%) and pectin (<1.5 wt%), whereas dew-retted, decorticated, or osmotically degummed fibers retain higher fractions of hemicellulose, pectin, and lignin, yielding cellulose contents closer to 66–68 wt% and lignin levels approaching 4–6 wt%. These non-cellulosic components are concentrated in the middle lamella and cell-wall matrix and are primarily responsible for moisture sensitivity and interfacial polarity. Liu et al. [57] further demonstrate that hemp fiber composition varies widely across cultivars and retting conditions, with reported cellulose contents ranging from ≈60 to 85 wt%, hemicellulose from 2 to 18 wt%, and lignin from 1 to 21 wt%, highlighting the strong coupling between agronomy, processing, and materials performance. These component fractions are typically determined using detergent-fiber (Van Soest) analysis, which separates cellulose, hemicellulose, and lignin by sequential neutral- and acid-detergent extractions [58].

Flax fibers show similar processing sensitivity but differ in their internal distribution of components. Zimniewska and Romanowska [56] report cellulose contents of ≈65–77 wt% for technical flax fibers, accompanied by 14–30 wt% hemicellulose, 4–9 wt% lignin, and 2–5 wt% pectin depending on cultivar and degumming route. At the cell-wall level, Goudenhooft et al. [47] reveal that these bulk values arise from a highly heterogeneous architecture dominated by a gelatinous (G) layer that occupies roughly 90% of the cell-wall area and is extremely cellulose-rich (typically 75–90 wt%), with 15–20 wt% hemicelluloses and 5–10 wt% pectins but negligible lignin. This cellulose-rich, low-microfibril-angle G-layer is the principal load-bearing phase of flax fibers, whereas the smaller, more lignified S1 layer contributes disproportionately to moisture resistance and interfacial chemistry.

Jute bast fibers are chemically distinct, retaining substantially higher lignin and hemicellulose contents than flax or hemp. The literature compiled by Islam et al. [59] indicates that jute fibers typically contain 58–72 wt% cellulose, 12–24 wt% hemicellulose, and 12–20 wt% lignin, along with minor amounts of pectin and surface waxes. The elevated lignin fraction imparts relatively high stiffness and moisture resistance, but reduces surface polarity and chemical reactivity, limiting interfacial adhesion and ductility in polymer composites. These datasets demonstrate that the mechanical properties, moisture sensitivity, and interfacial behavior in bast-fiber composites are governed not simply by cellulose content, but by the balance and spatial distribution of hemicellulose, pectin, and lignin, which are strongly shaped by cultivar selection and, more importantly, by retting and degumming history. Tables of chemical compositions are also available in [60,61], but these do not specify the cultivar or the fiber extraction method.

The microfibrillar angle, the inclination of cellulose microfibrils in the S₂ layer of the secondary cell wall relative to the fiber axis, is a central structural parameter governing stiffness, tensile behavior, and moisture response. Low MFAs (5–10°) correspond to highly aligned microfibrils and high axial modulus, whereas larger MFAs (20–40°) are associated with increased extensibility and reduced stiffness. Substantial variation in MFA is evident across fiber types: high-MFA fibers such as coir and sisal tend to be less stiff, while low-MFA fibers including flax, hemp, kenaf, and jute exhibit superior mechanical performance.

2.2.5. Microfibrillar Angle and Structural Anisotropy

A widely used method for determining MFA in plant fibers is X-ray diffraction (XRD). When an X-ray beam interacts with the oriented cellulose crystallites, the azimuthal spread of diffracted intensity around the (002) or (040) reflection directly reflects the underlying microfibril orientation [50]. Because cellulose microfibrils exhibit well-defined crystallographic symmetry, an explicit relationship can be derived between their orientation and the angular position of the diffracted beam [62].

Experimentally, the azimuthal intensity distribution $I\left(\varphi \right)$ of the (002) reflection is recorded, and the widely adopted Cave method estimates the MFA from the peak half-width at half-maximum (HWHM), $\beta$ [63]:

$$\mathit{MFA}=\arctan (\frac{1}{2}tan\beta )$$

(1)

which assumes a symmetric distribution of microfibril orientations around the fiber axis. A more general approach computes the MFA as the intensity-weighted centroid of the azimuthal profile,

$$\mathit{MFA}={\displaystyle \frac{{\displaystyle {\int }_{{\varphi }_1}^{{\varphi }_2}\varphi I(\varphi )\mathrm{d}\varphi }}{{\displaystyle {\int }_{{\varphi }_1}^{{\varphi }_2}I(\varphi )\mathrm{d}\varphi }}}$$

(2)

providing improved accuracy for broad or asymmetric orientation distributions.

Conventional analyses rely primarily on the (002) reflection, whose azimuthal intensity is dominated by contributions from the S₂ layer. Consequently, MFA estimates based solely on (002) lose precision when fibers exhibit nonuniform cross-sections, heterogeneous MFA across the wall, or overlapping signals from the S₁, S₂, and S₃ layers. Cave [62] demonstrated that the (040) reflection provides complementary geometric sensitivity. Joint analysis of the (002) and (040) peaks, each with a distinct angular dependence, enables resolution of MFA gradients across the full cell-wall thickness, separation of overlapping layer-specific orientation distributions, and substantially enhanced accuracy for fibers with complex or irregular morphologies.

2.2.6. Influence of Composition and Microfibrillar Angle on Fiber Mechanics

Table 5 provides a unified quantitative view of how cellulose content, crystallinity, and MFA jointly govern the mechanical performance of plant fibers across botanical classes. Although flax, hemp, jute, and several leaf fibers all contain 60–80 wt% cellulose, their Young’s moduli span nearly two orders of magnitude, from less than 4 GPa for coir to more than 70 GPa for flax. This wide dispersion demonstrates that cellulose content alone does not determine stiffness; instead, the orientation and crystalline quality of the cellulose microfibrils are the controlling factors.

The elementary-fiber data of Bourmaud et al. [50] show that flax cultivars, which exhibit MFAs of 8.5–9.5°, achieve Young’s moduli of 41–75 GPa and tensile strengths up to 1.45 GPa, whereas hemp fibers with a higher mean MFA of 11.2° exhibit a much lower modulus (19.1 ± 11.3 GPa) and strength (685 ± 590 MPa) despite comparable cellulose contents (Table 5). This disparity is consistent with differences in cell-wall architecture and microstructure rather than bulk cellulose content alone.

Micromechanical models developed by Gassan et al. [64] and Xu and Liu [65] rationalize the strong sensitivity of axial stiffness to microfibril orientation by treating the fiber wall as a multilayer composite of helically wound cellulose microfibrils embedded in a softer matrix. For a microfibril with intrinsic axial modulus $E_{\mathrm{l}}$ (typically 130–150 GPa for cellulose I [66]), geometric projection of both stress and strain leads to the limiting relation

$$E_{\mathrm{f}}\left(\theta \right)=E_{\mathrm{l}}{cos}^4\alpha$$

(3)

where $\alpha$ is the microfibrillar angle [65]. Importantly, for the low MFAs typical of bast fibers (single-digit to low-double-digit degrees), ${cos}^4\alpha$ varies only weakly; for example, increasing MFA from ≈9° to ≈11° changes ${cos}^4\alpha$ by only a few percent. Therefore, the large modulus and strength differences observed between flax and hemp cannot be attributed to MFA alone, but must also reflect differences in the fraction and continuity of the load-bearing cellulose-rich layer (S2/G), cellulose crystallinity, and effective reinforcing volume fraction, and microstructural defects introduced by growth conditions and processing.

The broader fiber classes in Table 5 follow the same trend. Ramie and pineapple leaf fibers, with MFAs near 8–14° and high crystallinity, exhibit Young’s moduli of ≈25–30 GPa, whereas sisal fibers with MFAs in the range of 10–25° show much lower stiffness (about 8–10 GPa) despite comparable cellulose contents. Highly lignified fibers such as coir combine very high MFAs (30–49°) with low crystallinity (27–33%), yielding extremely low stiffness (3–5 GPa) and tensile strength (<150 MPa) [64,67].

Table 5. Structure–property relationships for selected natural fibers.

Fiber	Young’s Modulus (GPa)	Tensile Strength (MPa)	MFA (°)	Cellulose (wt%)	Crystallinity (%)	References
Flax (elementary)	41.0–75.0	663–1454	8.5–9.5	64.6–77.4	50–90	[50,51,56]
Hemp (elementary)	$19.1\pm 11.3$	$685\pm 590$	11.2	66.0–72.5	50–90	[50,51,56,57]
Jute	28.4	331–414	≈8	58.0–71.5	50–80	[59,64,68]
Ramie (bundle)	$28.4\pm 3.6$	$849\pm 108$	≈8	≈76	≈64	[64,67]
Kenaf (bundle)	$25.1\pm 2.0$	$473\pm 46$	≈8–10	55–65	≈60	[64,67]
Sisal (bundle)	$9.1\pm 0.8$	$375\pm 38$	10–25	66–78	50–70	[64,67]
Banana (bundle)	–	–	≈11	44–64	45–55	[64]
Pineapple leaf (bundle)	$27.0\pm 2.3$	$654\pm 46$	12–14	70–82	44–60	[67]
Coir (bundle)	$3.7\pm 0.6$	$137\pm 11$	30–49	32–43	27–33	[64,67]

Notes. Mechanical properties of flax and hemp elementary fibers (Young’s modulus and tensile strength) are taken from Bourmaud et al. [50]. For flax, the reported ranges correspond to the minimum and maximum values among eight cultivars; for hemp, values are reported as mean ± standard deviation. Bundle-scale mechanical properties for ramie, kenaf, sisal, pineapple leaf, and coir are from Munawar et al. [67]. Cellulose contents for flax and hemp are from Zimniewska et al. [56], with broader cultivar-dependent ranges for hemp summarized by Liu et al. [57]. Jute cellulose content is from Islam et al. [59]. Microfibrillar angles for flax and hemp are from Bourmaud [50], while MFAs, cellulose crystallinity indices, and composition ranges for the other fibers are compiled from Gassan et al. [64]. Mechanical properties of jute fibers are from Biswas et al. [68].

Table 5. Structure–property relationships for selected natural fibers.

Fiber	Young’s Modulus (GPa)	Tensile Strength (MPa)	MFA (°)	Cellulose (wt%)	Crystallinity (%)	References
Flax (elementary)	41.0–75.0	663–1454	8.5–9.5	64.6–77.4	50–90	[50,51,56]
Hemp (elementary)	$19.1\pm 11.3$	$685\pm 590$	11.2	66.0–72.5	50–90	[50,51,56,57]
Jute	28.4	331–414	≈8	58.0–71.5	50–80	[59,64,68]
Ramie (bundle)	$28.4\pm 3.6$	$849\pm 108$	≈8	≈76	≈64	[64,67]
Kenaf (bundle)	$25.1\pm 2.0$	$473\pm 46$	≈8–10	55–65	≈60	[64,67]
Sisal (bundle)	$9.1\pm 0.8$	$375\pm 38$	10–25	66–78	50–70	[64,67]
Banana (bundle)	–	–	≈11	44–64	45–55	[64]
Pineapple leaf (bundle)	$27.0\pm 2.3$	$654\pm 46$	12–14	70–82	44–60	[67]
Coir (bundle)	$3.7\pm 0.6$	$137\pm 11$	30–49	32–43	27–33	[64,67]

Notes. Mechanical properties of flax and hemp elementary fibers (Young’s modulus and tensile strength) are taken from Bourmaud et al. [50]. For flax, the reported ranges correspond to the minimum and maximum values among eight cultivars; for hemp, values are reported as mean ± standard deviation. Bundle-scale mechanical properties for ramie, kenaf, sisal, pineapple leaf, and coir are from Munawar et al. [67]. Cellulose contents for flax and hemp are from Zimniewska et al. [56], with broader cultivar-dependent ranges for hemp summarized by Liu et al. [57]. Jute cellulose content is from Islam et al. [59]. Microfibrillar angles for flax and hemp are from Bourmaud [50], while MFAs, cellulose crystallinity indices, and composition ranges for the other fibers are compiled from Gassan et al. [64]. Mechanical properties of jute fibers are from Biswas et al. [68].

The data in Table 5 demonstrate that a high cellulose fraction is necessary but not sufficient for high stiffness. Only when high crystallinity is coupled with very low microfibrillar angle, as in flax, can the intrinsic stiffness of cellulose microfibrils be efficiently projected into macroscopic fiber performance.

2.3. Animal-Based Fibers

Animal fibers represent a complementary class of protein-based reinforcements characterized by tunable elasticity, intrinsic damping, and biocompatibility. Their performance arises from hierarchical protein architectures, primarily based on $\alpha$-keratin in hair and wool and fibroin in silks, which form multiscale networks stabilized by disulfide and hydrogen bonding. These molecular interactions enable large reversible deformation, energy dissipation, and thermal stability, distinguishing animal fibers from cellulose-based plant reinforcements.

2.3.1. Protein Fiber Architecture

Keratin-based animal fibers exhibit a hierarchical morphology consisting of a scaly cuticle surrounding a cortex of macro- and microfibrils embedded in a sulfur-rich amorphous matrix [69,70,71]. The cortex contains intermediate filaments of $\alpha$-keratin organized as coiled-coil structures, which are stabilized by cystine-derived disulfide cross-links and confer elasticity, resistance to permanent deformation, and chemical durability. In some animal fibers, a medulla may be present as a discontinuous or continuous internal channel, contributing to low density and thermal insulation [71].

Silk fibers, particularly from Bombyx mori, are structurally distinct from keratin fibers and consist of a fibroin core coated with sericin. Repetitive $\beta$-sheet nanocrystalline domains within fibroin act as reinforcing motifs that provide high tensile strength-to-weight ratios and controlled extensibility [31,33]. Alpaca fibers belong to the class of fine animal fibers derived from South American camelids and are chemically similar to wool, being composed primarily of keratin. Compared with sheep wool, alpaca fibers exhibit smoother cuticle scale morphology, lower grease content, and often non-circular cross-sections with internal channels, contributing to reduced density and distinct tactile properties [30].

2.3.2. Hygroscopicity, Damping, and Functional Behavior

Protein-based fibers are moderately hygroscopic due to polar amino acid side chains and can absorb substantial moisture while maintaining thermal insulation and structural integrity [71]. In contrast, lignocellulosic fibers such as flax and hemp exhibit greater moisture-induced swelling and stiffness loss, largely associated with their hemicellulose content and hydrogen-bonded cellulose microfibrils [71]. Bast fibers possess high axial stiffness arising from elevated cellulose crystallinity and low microfibril angles, whereas keratin fibers display greater extensibility and recoverable deformation due to their helical protein architecture and disulfide-bonded network [70].

Rodopoulos et al. [72] demonstrated that plant-based fiber structures (e.g., manila and coconut) exhibit substantially higher stiffness and static strength than animal-fiber structures (goat hair), while animal fibers show markedly higher ductility and hysteretic energy dissipation under cyclic loading. Specifically, plant fibers achieved peak stresses exceeding 70 MPa, whereas animal fibers exhibited lower strength but larger strain capacity and greater damping efficiency. These findings reinforce a functional distinction between natural fiber classes: lignocellulosic fibers are better suited for stiffness- and load-bearing applications, while keratin fibers are more appropriate for energy-absorbing, damping, and comfort-driven functions [33,72].

2.3.3. Composite Performance of Animal Fibers

Comparative performance studies highlight the reinforcing potential of animal fibers. Composites of unidirectionally aligned Bombyx mori silk/polypropylene (PP) composites fabricated by compression molding, containing 20 wt% fiber, exhibit tensile strength, flexural strength, and flexural modulus values of 55.6 MPa, 57.1 MPa, and 3.32 GPa, respectively, which are substantially higher than those reported for unidirectionally aligned jute/PP composites processed under identical conditions (41.3 MPa, 44.2 MPa, and 2.20 GPa) [73]. After 12 weeks of soil burial, silk/PP composites retained ≈90% of their initial tensile strength, compared with 68% retention for jute/PP, underscoring the superior resistance of silk fibers to hydrolytic and microbial degradation within a controlled matrix, fiber content, and processing framework [73].

Keratin-rich fibers also show considerable promise. Waste sheep wool/epoxy composites produced by vacuum-assisted resin transfer molding (VARTM) have been evaluated across woven, needle-punched, and felted reinforcement architectures [74]. The needle-punched configuration delivers the highest tensile strength due to efficient stress transfer within its randomly oriented fiber network, whereas the woven and felted variants exhibit low thermal conductivities (≈0.22 W m⁻¹ K⁻¹), underscoring their suitability for lightweight thermal-insulating components in construction and transportation. Chicken-feather keratin likewise improves dispersion and impact strength in polyolefins when compatibilized with benzoyl peroxide or maleic anhydride (MA), while cow and goat hair fibers enhance polyester and linear low-density polyethylene (LLDPE) composites through their mixed hydrophilic–hydrophobic surface chemistry and reactive amine, carboxyl, and disulfide functionalities [32].

2.3.4. High-Performance and Interfacial Engineering

At the high-performance frontier, spider dragline silk exhibits tensile strengths up to ≈1.1 GPa and moduli near 10 GPa at a density of about 1.3 g cm⁻³, corresponding to among the highest specific strengths reported for natural fibers [29]. Untreated animal fibers generally exhibit weaker interfacial adhesion than alkali-treated plant fibers due to their smoother surfaces. However, Talabi et al. [29] showed that fiber–matrix bonding can be significantly enhanced through surface-modification strategies including enzymatic treatment, oxidative processes (e.g., peroxide), plasma activation, and coupling-agent grafting such as maleic-anhydride systems.

In addition, keratin- and silk-based fibers absorb less moisture than lignocellulosic fibers because of their lower hydroxyl-group content, leading to improved dimensional stability and retention of mechanical properties in humid environments [32]. Together, their exceptional specific strength, tunable interfacial chemistry, and reduced moisture sensitivity establish keratin and silk fibers as promising reinforcements for hybrid and functional biocomposites.

2.4. Fungal Fibers and Bio-Based Fillers

2.4.1. Mycelium-Based Structural Composites

Emerging biological reinforcements broaden the concept of natural fibers to include fungal mycelium networks and bio-sourced particulate fillers. Fungal mycelium differs fundamentally from plant fibers in both chemical composition and load-bearing architecture. The vegetative hyphae that form fungal fibers possess a multilayered cell wall rich in chitin, $\beta$-glucans, and associated structural proteins, rather than cellulose [75]. In mycelium-based composites, cellulose and lignin are confined primarily to the lignocellulosic growth substrate, which is colonized and mechanically integrated by the fungal network but does not determine the intrinsic chemistry of the mycelial filaments themselves [76]. The resulting chitin–glucan framework imparts mechanical integrity, elasticity, and environmental responsiveness, distinguishing mycelium-derived reinforcements from both cellulose-based plant fibers and protein-based animal fibers.

Mycelium-based composites (MBCs) are produced by cultivating fungal hyphae through particulate lignocellulosic substrates, yielding a self-assembled biological matrix that binds substrate particles without the need for synthetic resins. Following densification by pressing or molding, these materials exhibit low density, biodegradability, and sufficient mechanical performance for applications such as packaging and thermal or acoustic insulation [76,77]. Scalability, however, remains constrained by inherently slow biological growth rates.

2.4.2. Bio-Derived Particulate and Nanoscale Fillers

Bio-based fillers, such as cellulose nanocrystals, lignin, or starch nanoparticles, powdered protein fibers, and chitosan microparticles, enhance stiffness, barrier performance, and antimicrobial activity [78,79,80,81]. Incorporating hyperbranched, bio-derived polymers synthesized from castor oil or citric acid further improves matrix toughness and interfacial adhesion [82]. These bio-derived particulate reinforcements reflect a shift toward multifunctional composites from renewable resources, in which micro- and nanoscale bio-additives complement traditional fibers to define both structural and functional performance.

2.5. Relative Depth of the Available Literature

The depth of discussion devoted to different reinforcement classes in this review reflects the maturity and breadth of the available literature rather than preferential emphasis. Plant fibers such as flax, hemp, jute, and sisal are supported by an extensive body of experimental and industrial research, driven by their availability, established processing routes, and demonstrated performance in automotive and construction applications. By contrast, animal- and fungal-derived fibers remain emerging reinforcement systems, with a more limited but rapidly growing literature focused primarily on fundamental properties, bio-functionality, and proof-of-concept applications. Their comparatively shorter treatment in this review, therefore, reflects the current state of research activity rather than a limitation of their long-term potential.

2.6. Bio-Based Polymer Matrices

The polymer matrix governs load transfer, mechanical stability, and degradation behavior in biocomposites [83]. Bio-based matrices derived from agricultural, microbial, or vegetable-oil feedstocks provide renewable alternatives to petrochemical resins. Widely used systems, including poly(lactic acid), polyhydroxybutyrate (PHB), polycaprolactone, starch-based polymers, and vegetable-oil-derived thermosets, differ substantially in crystallinity, molecular architecture, and thermal stability. These differences lead to distinct trade-offs between stiffness, ductility, processability, and biodegradability. Achieving balanced performance, therefore, requires chemical compatibility with the reinforcement and careful control of processing conditions to minimize chain scission, thermal degradation, and poor interfacial adhesion.

2.6.1. PLA as a Benchmark Bio-Based and Biodegradable Thermoplastic

PLA is produced via ring-opening polymerization of lactide, which is derived from lactic acid obtained by fermentation of renewable resources such as corn starch, sugarcane, or cassava [84]. It is the most widely used commercially available bio-based, industrially compostable thermoplastic matrix for natural-fiber composites due to its availability, high stiffness, and melt-processability. It exhibits a glass-transition temperature $T_{\mathrm{g}}$ of ≈55–65 °C and a melting temperature $T_{\mathrm{m}}$ of 150–175 °C [85]. Flax- and hemp-reinforced PLA composites with randomly oriented short fibers at ≈40 wt% loading, processed by injection molding, have been reported to reach tensile moduli exceeding 5 GPa under quasi-static testing at ambient conditions [86], compared with 3.5–4.2 GPa for neat PLA [85]. Despite this favorable stiffness, hydrolytic degradation during composting accelerates chain scission and embrittlement, while alkali-treated fibers, although effective in increasing stiffness, can reduce biodegradation rates by increasing matrix crystallinity [87], making interfacial engineering essential for balancing mechanical performance with controlled environmental degradation. Consequently, although PLA serves as a benchmark biodegradable thermoplastic, its broader industrial deployment in natural-fiber composites remains constrained by intrinsic brittleness and moisture sensitivity [8], a relatively narrow processing window and high cost [1], and unresolved challenges in closed-loop recyclability [13].

2.6.2. Polyhydroxyalkanoates (PHB, PHBV, and PHBHx)

PHB is a semicrystalline bacterial polyester produced through microbial fermentation of renewable feedstocks such as glucose, sucrose, or plant oils. As a member of the polyhydroxyalkanoate family, PHB is fully bio-based and biodegradable, offering a sustainable alternative to polypropylene. Its high crystallinity (60–80%) and melting temperature (170–180 °C) enable tensile strengths of 30–40 MPa and moduli of 2–3 GPa, but also result in brittleness (typically <5% elongation) and slow crystallization kinetics that limit toughness and processing flexibility [88]. Coats et al. [89] demonstrated that unpurified PHB-rich bacterial biomass from Azotobacter vinelandii can be directly blended with wood flour to form thermoplastic composites, eliminating the need for polymer purification. These PHB–biomass composites achieved bending strength and stiffness comparable to those of purified PHB, as residual cellular material acted as a natural compatibilizer and processing aid. Improved ductility is commonly achieved through copolymerization, most often by incorporating 3-hydroxyvalerate to form poly(hydroxybutyrate-co-hydroxyvalerate), or by blending with PLA or PCL, which reduces crystallinity and enhances toughness while preserving biodegradability.

PHBV, an important PHB copolymer, combines reduced crystallinity and melting temperature with increased ductility and impact resistance, providing a wider processing window and more balanced mechanical behavior while retaining full biodegradability and bio-based origin. Both PHB and PHBV can be produced from lignocellulosic biomass, agricultural residues, and food-processing wastes rather than refined sugars, which dominate current commercial routes, thereby significantly reducing feedstock-driven production costs [90]. The hydroxyvalerate content in PHBV can also be tuned during fermentation—through metabolic control or propionic acid feeding—enabling systematic adjustment of crystallinity and mechanical properties [90]. In natural-fiber biocomposites, PHBV is particularly attractive because fibers and fillers simultaneously reinforce stiffness and strength while accelerating biodegradation; PHBV and PHB/PHBV composites derived from agricultural and fruit-waste feedstocks retain full biodegradability in soil and compost while achieving mechanical performance suitable for packaging, agricultural products, and other short-lifetime structural applications [90]. Consistent with this behavior, PHBV reinforced with high-stiffness natural fibers such as Spartium junceum L. exhibits reduced crystallinity together with improved thermal stability and mechanical performance relative to neat PHBV, offering a promising route to overcome the intrinsic brittleness of fully biodegradable matrices [91].

Poly(3-hydroxybutyrate-co-3-hydroxyhexanoate) (PHBHx) represents a major evolution within the PHA family, in which incorporation of medium-chain-length 3-hydroxyhexanoate (3HHx) units disrupts the highly crystalline PHB lattice, yielding markedly higher ductility, lower melting temperatures, and a broader melt-processing window than PHB and most PHBV grades [92,93]. This effect reflects the intrinsically flexible nature of the 3HHx comonomer: while the PHB homopolymer exhibits a $T_{\mathrm{g}}$ of ≈3 °C, poly(3-hydroxyhexanoate) (PHHx) has a much lower $T_{\mathrm{g}}$ near −28 °C [92]. In PHBHx copolymers, the 3HHx molar fraction therefore acts as a molecular “plasticizer” that systematically reduces melting temperature, modulus, and crystallinity while increasing elongation at break and enzymatic degradability [93]. These tunable structure–property relationships allow PHBHx to span behaviors ranging from polypropylene-like thermoplastics at low 3HHx contents to elastomeric materials at higher 3HHx fractions, making it especially attractive for melt-processed biodegradable products and fiber-reinforced composites.

Early work demonstrated effective reinforcement of PHBHx with flax fibers, especially when long fiber mats or surface-modified fibers were used, leading to improved mechanical performance and transcrystalline interphases at the fiber–matrix interface [94]. More recently, PHBHx reinforced with Agave americana fibers exhibited significant increases in stiffness and storage modulus with increasing fiber content while retaining good thermal stability and processability [95]. PHBHx has also been combined with nanocellulose and other bio-based fillers to produce fully biodegradable bionanocomposites with enhanced mechanical and barrier properties [96]. In parallel with these materials developments, metabolic and enzymatic engineering strategies now allow precise control of 3HHx incorporation during fermentation, providing a scalable route to tailor PHBHx composition, processing window, and end-of-life behavior for targeted composite applications [93].

2.6.3. Vegetable-Oil-Derived Thermosets

Vegetable-oil-based thermosetting resins, including acrylated epoxidized soybean oil (AESO), epoxidized linseed oil (ELO), and maleated castor oil (MACO), have emerged as versatile bio-based matrix systems for natural-fiber composites owing to their high renewable content, tunable functionality, and ability to form densely cross-linked networks [97]. AESO, produced by acrylating epoxidized soybean oil, readily undergoes photopolymerization, enabling rapid curing. Ultraviolet curing enables rapid, room-temperature processing of natural-fiber-reinforced AESO composites, achieving full cure through sections up to 10 mm thick within ≈10 min (5 min per side) [98]. For unidirectional flax-reinforced AESO laminates (fiber mats compressed to ≈1 mm thickness), UV curing yields tensile strengths of $51\pm 6$ MPa, approximately twice that of thermally cured counterparts under comparable fiber architecture, attributable to improved fiber wet-out and reduced interfacial gap formation during short-duration curing.

Linseed- and castor-oil-derived thermoset matrices offer complementary toughness and processing flexibility. In fully plant-derived MACO/linseed-oil networks reinforced with randomly oriented, NaOH-treated short banana fibers at an optimal loading of 50 wt%, compression-molded at 130 °C and 500 psi, tensile strengths of $16.9\pm 4.7$ MPa were reported together with high compressive resistance, showing no fracture up to a nominal compressive stress of about 45 MPa at a strain of ≈25%, and low thermal conductivity (≈0.19 W m⁻¹ K⁻¹) [99]. These properties support potential use in structural-insulating panels and packaging applications rather than primary load-bearing components.

While AESO offers high cross-link density and energy-efficient curing, MACO/ELO formulations provide greater ductility and fully petrochemical-free matrices; however, the tightly cross-linked architectures typical of these systems can limit microbial accessibility and slow biodegradation despite their renewable origin. Although derived from renewable feedstocks, plant-based thermosetting matrices are generally not recyclable in a closed-loop sense due to their permanently cross-linked network structures, and their sustainability benefits therefore arise primarily from bio-based content rather than end-of-life reprocessability.

2.6.4. Other Biodegradable Thermoplastics (TPS, PCL, PBS, PBAT)

Starch-based polymers offer highly degradable, hydrophilic matrices suited for short-lifetime applications [100,101,102,103,104]. Thermoplastic starch, produced by plasticizing native starch under heat and shear, forms a melt-processable matrix that remains biodegradable due to its starch content. TPS can be blended with other thermoplastics [105] to tailor mechanical performance, although high water uptake and limited stiffness generally constrain its use in load-bearing biocomposites.

Polycaprolactone provides high flexibility, toughness, and low processing temperatures (≈60 °C), making it particularly attractive for composites containing moisture- or temperature-sensitive natural fibers. Flax/PCL and hemp/PCL composites achieve moduli up to 1.5 GPa and show improved impact resistance when formulated with peroxides or coupling agents [106,107]. Jute/PCL composites retain structural integrity during prolonged marine exposure and undergo gradual biodegradation, supporting their use in controlled-degradation applications [108]. However, the low glass-transition temperature of PCL limits its use in elevated-temperature service environments.

Poly(butylene succinate) occupies an intermediate position between rigid bioplastics such as PLA and highly flexible systems such as PCL and PBAT. It is a semicrystalline, fully biodegradable aliphatic polyester with a low $T_{\mathrm{g}}$ in the range of $-45$ to −10 °C and a melting temperature $T_{\mathrm{m}}$ in the range of 90–120 °C [85,109], making it readily processable by conventional extrusion, molding, and compounding methods. In natural-fiber composites, PBS provides greater toughness and strain accommodation than PLA, reducing interfacial stress concentrations and suppressing fiber debonding, while remaining significantly less hydrophilic than starch- or protein-based matrices. As a result, PBS is widely used in biodegradable packaging, agricultural products, and semi-structural biocomposites where a balance of stiffness, ductility, and moisture tolerance is required.

PBAT occupies a complementary niche as a flexible, fully biodegradable aromatic–aliphatic copolyester widely used to toughen brittle bio-based matrices such as PLA and PHBV. Although petrochemical in origin, PBAT is compostable and exhibits low modulus, high elongation, and excellent melt processability, making it particularly effective in natural-fiber composites and PLA-based blends where impact resistance and ductility are required. In fiber-reinforced systems, PBAT enhances stress transfer and crack resistance while maintaining compostability, supporting its use in packaging films, agricultural mulches, and short-lifetime molded products. For example, melt-mixed PBAT composites reinforced with Amazonian natural fibers such as Croton lanjouwensis, Malvastrum tomentosum, and Trema micrantha showed modulus increases of 48–72% relative to neat PBAT, demonstrating that lignocellulosic fibers can significantly enhance stiffness without compromising biodegradability [110]. More broadly, PBAT–natural fiber systems consistently exhibit improved mechanical performance and interfacial stress transfer while retaining compostability, confirming PBAT’s role as a practical biodegradable matrix for fiber-based sustainable composites [111].

2.6.5. Matrix Selection as a Design Variable

Processing conditions strongly influence final composite performance. PLA and PHB require tightly controlled extrusion and molding temperatures to avoid chain scission and excessive brittleness [86,87,88,89]. Vegetable-oil thermosets rely on in situ polymerization and optimized curing protocols [97,99], whereas starch-based matrices are typically processed by casting or extrusion [112]. PCL’s low melt viscosity enables simple melt blending and film stacking [106,107,108]. Coordinated optimization of matrix chemistry, fiber polarity, and processing temperature is therefore essential for developing durable, high-performance, and environmentally responsible biocomposites.

Different polymer matrices occupy distinct roles within the natural-fiber composite literature, reflecting their industrial relevance and the specific challenges they address. PLA functions primarily as a benchmark biodegradable thermoplastic and is therefore discussed across processing, moisture sensitivity, biodegradation, recyclability, and brittleness contexts, often in multiple modified or blended forms. Polypropylene serves as the dominant industrial reference for durability, compatibilization strategies, and performance benchmarking against bio-based systems. PHB and PHBV are emphasized in discussions of biodegradation kinetics, moisture-accelerated degradation, processing constraints, and cost–performance trade-offs among fully bio-based matrices. Epoxy and polyester systems are primarily treated as high-performance and structural baselines, particularly in durability and weathering comparisons. Other biodegradable and bio-based polymers, including PBS, PCL, PBAT, thermoplastic starch, PVA, and bio-based polyolefins, are discussed more selectively, where their distinctive processing windows, degradation behavior, or application niches provide complementary insight. In this context, the prominence of specific matrices reflects their functional role as reference systems rather than preferential material selection.

Within the framework of this review, fiber chemistry and polymer matrix selection are treated not as independent material choices, but as coupled design variables whose interactions at the fiber–matrix interface govern mechanical performance, moisture sensitivity, durability, and end-of-life behavior.

3. Fiber–Matrix Interfacial Engineering

Figure 5 summarizes the causal chain that governs composite performance in this section, showing how surface chemistry controls wetting, interfacial shear strength, stress transfer, and ultimately durability in natural-fiber composites. Building on the material classes introduced in Section 2, this section focuses on interfacial engineering as the key link between fiber chemistry, matrix selection, processing conditions, and long-term composite behavior. Because the fiber–matrix interface mediates load transfer and moisture ingress, it plays a dominant role in mechanical reliability under service conditions. Natural fibers, which are rich in hydroxyl groups and often coated with lignocellulosic impurities, can exhibit limited compatibility with many polymer matrices, leading to weak adhesion, fiber pull-out, and premature debonding during mechanical loading.

3.1. Why Interfaces Control Performance and Durability

The sensitivity of natural-fiber composites to interfacial quality is clearly illustrated by benchmark systems in which the polymer matrix and processing conditions are held constant while fiber type and morphology vary. Injection-molded plant-fiber/PLA composites, for example, exhibit lower tensile and impact strengths than composites reinforced with synthetic Cordenka fiber, a high-modulus regenerated cellulose benchmark [113]. SEM observations of cryo-fractured surfaces indicate more favorable fiber distribution and fiber–matrix interaction in Cordenka-reinforced PLA compared with flax- and hemp-reinforced systems, whereas natural-fiber composites exhibit greater heterogeneity in fiber distribution. These observations suggest that the apparent underperformance of bio-based reinforcements reflects not only intrinsic differences in fiber properties but also processing-related limitations associated with fiber dispersion and interfacial efficiency. Together, these findings motivate strategies that concurrently enhance fiber dispersion and fiber–matrix interactions to more fully realize the reinforcing potential of natural fibers.

3.2. Interfacial Modification Strategies in Natural-Fiber Composites

Because interfacial quality governs both load transfer and environmental stability, deliberate modification of the fiber surface has become a central strategy for enhancing the performance and durability of natural-fiber composites. To address the dispersion, wetting, and adhesion limitations discussed above, a range of surface-modification approaches, including alkali treatment, acetylation, silane coupling, and maleic anhydride (MA) grafting, are widely employed to tailor fiber surface chemistry and introduce reactive functional groups, as comprehensively reviewed by Arunachalam et al. [114].

The foundational importance of interfacial engineering in natural-fiber composites was recognized in early studies, which highlighted the role of fiber surface chemistry in governing adhesion and mechanical performance [115]. Subsequent work has provided converging evidence that interfacial modification is a dominant lever controlling composite behavior. At the level of mechanical response and fracture morphology, alkali cleaning, cyanoethylation, and polymer grafting have been shown to substantially improve interfacial bonding and stress-transfer efficiency in lignocellulosic fibers, with tensile strength increases often exceeding 40% and microscopy revealing reduced fiber pull-out and more cohesive fracture surfaces [115]. Beyond initial strength, several surface treatments also influence durability by modulating biodegradation and moisture sensitivity, thereby enabling control over environmental stability in service.

At the material-system level, focused reviews of hemp and other bast-fiber composites consistently report that alkali treatment, silane coupling, and maleic-anhydride-based compatibilization improve fiber–matrix adhesion and, in many cases, mechanical performance in both thermoplastic and thermoset matrices [116]. However, these same studies emphasize that treatment effectiveness is highly sensitive to fiber species, treatment severity, and matrix chemistry, with over-treatment or chemical mismatch frequently offsetting gains in strength or toughness. Broader surveys of natural-fiber–reinforced composites reinforce this conclusion, identifying interfacial engineering as a primary lever for mitigating moisture sensitivity and stress-transfer limitations, while also highlighting the strong coupling between interfacial chemistry and processing-induced microstructural heterogeneity [117].

Across these reviews and related syntheses, a consistent mechanistic picture emerges: chemical modification strategies act principally by reducing fiber hydrophilicity and promoting more effective interfacial stress transfer, rather than by altering the intrinsic stiffness or strength of the fibers themselves. At the same time, reported property gains span a wide range even for nominally similar fiber–matrix systems, reflecting differences in treatment severity, compatibilizer chemistry, processing history, and baseline normalization. Detailed experimental studies demonstrate that such factors can dominate measured reinforcement efficiency, limiting the validity of direct cross-study comparisons based on aggregated percentage improvements alone. Consequently, no single surface modification strategy can be regarded as universally effective, and sequential or hybrid treatments—often combined with matrix-side functionalization—are increasingly employed to address multiple interfacial constraints simultaneously and achieve balanced improvements in strength, moisture resistance, and long-term durability.

Table 6. Representative chemical and physical modification reactions enhancing fiber–matrix interfacial bonding in natural fiber composites. (a) Alkali treatment (NaOH): ionization of –OH and removal of surface impurities; (b) acetylation: esterification of –OH with acetic anhydride, reducing hydrophilicity; (c) silane coupling (e.g., KH570): hydrolysis and condensation to Si–O–cellulose, with organofunctional copolymerization to the matrix; R = alkyl, R^* = matrix-reactive group; (d) maleic anhydride–grafted polymers: covalent linkage between fiber –OH and anhydride groups; (e) peroxide oxidation (keratin): formation of polar sulfoxide/sulfonic groups, increasing roughness and reactivity; (f) silane treatment of wool/silk: Ker–O–Si interfacial bonds and improved compatibility with polymer matrices; (g) silk degumming: sericin removal and surface fibrillation; (h) plasma or corona discharge: radical formation and oxygenation for subsequent grafting; (i) thermal deacetylation of acetylated hemicelluloses (e.g., O-acetylxylan, glucomannan), releasing acetic acid and exposing hydroxyl groups that increase surface reactivity; (j) enzymatic retting: selective pectin removal, preserving cellulose microfibrils; (k) stearoylation: esterification of –OH with long-chain fatty acids to create hydrophobic surfaces; and (l) citric acid cross-linking: formation of multifunctional ester bridges between fiber –OH groups.

Label	Representative Reaction Equation
(a)	$\mathrm{Fiber}–\mathrm{OH}+\mathrm{NaOH}\to {\mathrm{Fiber}–\mathrm{O}}^{-}{\mathrm{Na}}^{+}+{\mathrm{H}}_2\mathrm{O}$
(b)	$\mathrm{Fiber}–\mathrm{OH}+{(\mathrm{CH}}_3{\mathrm{CO})}_2\mathrm{O}\to {\mathrm{Fiber}–\mathrm{OCOCH}}_3+{\mathrm{CH}}_3\mathrm{COOH}$
(c)	$\mathrm{Fiber}–\mathrm{OH}+{\left(\mathrm{RO}\right)}_3{\mathrm{Si}–\mathrm{R}}^{*}+{\mathrm{H}}_2\mathrm{O}\to {\mathrm{Fiber}–\mathrm{O}–\mathrm{Si}\left(\mathrm{OH}\right)}_2{–\mathrm{R}}^{*}+\mathrm{ROH}$
(d)	$\mathrm{Fiber}–\mathrm{OH}+\mathrm{Polymer}–\mathrm{MA}\to {\mathrm{Fiber}–\mathrm{O}–\mathrm{C}(=\mathrm{O})–\mathrm{CH}}_2–\mathrm{CH}\left(\mathrm{COOH}\right)–\mathrm{Polymer}$
(e)	$\mathrm{Keratin}–\mathrm{S}–\mathrm{S}–\mathrm{Keratin}+5{\mathrm{H}}_2{\mathrm{O}}_2\to 2{\mathrm{Keratin}–\mathrm{SO}}_3\mathrm{H}+4{\mathrm{H}}_2\mathrm{O}$
(f)	$\mathrm{Keratin}–\mathrm{OH}+{\left({\mathrm{H}}_3\mathrm{CO}\right)}_3\mathrm{Si}{–\mathrm{R}}^{*}+2{\mathrm{H}}_2\mathrm{O}\to {\mathrm{Keratin}–\mathrm{O}–\mathrm{Si}\left(\mathrm{OH}\right)}_2{–\mathrm{R}}^{*}+3{\mathrm{CH}}_3\mathrm{OH}$
(g)	$\mathrm{Raw}\mathrm{silk}(\mathrm{Fibroin}+\mathrm{Sericin})+{\mathrm{OH}}^{-}$
	$\to \mathrm{Degummed}\mathrm{silk}\left(\mathrm{Fibroin}\right)+\mathrm{Soluble}\mathrm{sericin}\mathrm{fragments}$
	$\mathrm{Sericin}–\mathrm{CO}–\mathrm{NH}–\mathrm{Sericin}+{\mathrm{OH}}^{-}\to {\mathrm{Sericin}–\mathrm{COO}}^{-}+{\mathrm{H}}_2\mathrm{N}–\mathrm{Sericin}$
(h)	$\mathrm{Fiber}–\mathrm{H}\underset{{\mathrm{O}}_2,{\mathrm{H}}_2\mathrm{O}}{\overset{\mathrm{plasma}/\mathrm{corona}}{\to }}\mathrm{Fiber}–\mathrm{OH},\mathrm{Fiber}–\mathrm{C}=\mathrm{O},\mathrm{or}\mathrm{Fiber}–\mathrm{COOH}$
(i)	${\mathrm{Hemicellulose}–\mathrm{OCOCH}}_3+{\mathrm{H}}_2\mathrm{O}\stackrel{\Delta }{\to }\mathrm{Hemicellulose}–\mathrm{OH}+{\mathrm{CH}}_3\mathrm{COOH}$
(j)	$\mathrm{Pectin}–\mathrm{COOR}+{\mathrm{H}}_2\mathrm{O}\stackrel{\mathrm{pectinase}}{\to }\mathrm{Pectin}–\mathrm{COOH}+\mathrm{ROH}$
(k)	$\mathrm{Fiber}–\mathrm{OH}+{\mathrm{C}}_{17}{\mathrm{H}}_{35}\mathrm{COCl}\to {\mathrm{Fiber}–\mathrm{O}–\mathrm{C}(=\mathrm{O})–\mathrm{C}}_{17}{\mathrm{H}}_{35}+\mathrm{HCl}$
(l)	${\mathrm{HOOC}–\mathrm{CH}}_2{–\mathrm{C}\left(\mathrm{OH}\right)\left(\mathrm{COOH}\right)–\mathrm{CH}}_2–\mathrm{COOH}+2\mathrm{Fiber}–\mathrm{OH}$
	$\to {\mathrm{Fiber}–\mathrm{O}–\mathrm{C}(=\mathrm{O})–\mathrm{CH}}_2{–\mathrm{C}\left(\mathrm{OH}\right)(\mathrm{COO}–\mathrm{Fiber})–\mathrm{CH}}_2–\mathrm{COOH}+2{\mathrm{H}}_2\mathrm{O}$

summarizes the principal chemical mechanisms by which these treatments strengthen fiber–matrix interfaces. Such modifications alter fiber polarity and surface reactivity while often removing weak boundary layers and increasing surface roughness. As a result, fiber–matrix wettability and interfacial shear strength are improved, stress transfer becomes more efficient, and composite performance is enhanced, with concurrent reductions in moisture sensitivity and improvements in long-term durability.

3.3. Interfacial Engineering of Lignocellulosic Fibers

3.3.1. Alkali Treatment (Mercerization)

Alkali treatment is one of the most widely adopted strategies for enhancing the interfacial activity of lignocellulosic fibers, owing to its simplicity, scalability, and effectiveness in activating cellulose surfaces [118]. When used as a pretreatment prior to grafting or esterification, NaOH removes weak boundary layers such as waxes, hemicellulose, and lignin, while increasing surface roughness, fibrillation, and the accessibility of reactive –OH groups. Sodium hydroxide converts accessible hydroxyls to alkoxide species (–O⁻Na⁺) [Reaction (a), Table 6], thereby disrupting intermolecular hydrogen bonding, increasing surface energy, and generating sites for subsequent chemical modification. While these effects generally improve wettability and interfacial bonding, overly strong or prolonged alkali exposure can reduce cellulose crystallinity and compromise fiber integrity.

Recent fiber-level studies confirm the existence of a narrow alkali-treatment window for bast fibers. Flax fibers treated with moderate NaOH concentrations exhibit increased tensile strength, reduced moisture uptake, and improved thermal stability due to hemicellulose removal and fibrillation, whereas higher NaOH concentrations lead to fiber damage and strength loss [119]. These results highlight that enhanced surface activation does not monotonically translate into improved reinforcement performance when fiber integrity is degraded.

Despite these limitations, alkali pretreatment remains highly effective as a foundation for advanced surface-modification strategies. For example, Kovuru et al. [120] demonstrated that a 4 wt% NaOH pretreatment removes surface impurities, reduces interfacial voids under comparable processing conditions, and improves fiber wetting and impregnation quality, resulting in stronger interfacial bonding and lower water absorption relative to untreated fibers. Consistent trends in fiber surface chemistry are widely reported across fiber types: NaOH refinement reduces hemp fiber contact angles from 95–101° (untreated) to 54–60° after treatment with 10 M NaOH [121], and alkalized coir, sisal, bamboo, and banana fibers have been shown to yield polypropylene and epoxy composites with improved tensile, flexural, and impact properties relative to untreated systems [122,123,124,125,126]. Alkaline–peroxide variants have been reported to further increase cellulose crystallinity and interfacial shear strength (IFSS) [127]. However, excessive alkali exposure can lead to cellulose degradation and deterioration of fiber mechanical integrity [128].

Building on the activated fiber surfaces generated by alkali pretreatment, fungal (mycelial) treatment has been demonstrated as an additional route to interfacial reinforcement in hybrid natural-fiber composites. In epoxy-matrix composites reinforced with woven hemp fabric (15 wt%) and randomly oriented chopped sisal fibers, controlled mycelial colonization for 12–15 days forms an interfacial hyphal network that enhances mechanical interlocking and crack bridging [120]. Relative to untreated laminates, the alkali- and fungal-treated hybrid composite containing 25 wt% sisal (total fiber volume fraction ≈ 41%) exhibited increases in tensile strength from 32 to 55 MPa, tensile modulus from 2.27 to 3.23 GPa, flexural strength from 36 to 55 MPa, flexural modulus from 2.15 to 3.26 GPa, and impact strength from 5.0 to 8.6 kJ m⁻², accompanied by reduced water absorption and a 15–20% increase in strain to failure under quasi-static loading [120]. To interpret these stiffness gains in terms of reinforcement efficiency and fiber orientation, simple rule-of-mixtures (ROM) bounds provide a useful physical framework. Using an orientation-corrected ROM, Kovuru et al. [120] predicted a composite modulus of ≈3.22 GPa for randomly oriented hemp–sisal/epoxy laminates with 25% fiber volume fraction and ${\eta }_0=0.33$, in close agreement with the experimentally measured modulus of 3.23 GPa.

Modeling Note: Rule-of-Mixtures Bounds

ROM models provide a physically transparent framework for estimating the elastic properties of composite materials and for defining meaningful upper and lower bounds based on the assumed mode of load sharing between phases [129]. For a two-phase system consisting of a reinforcing filler (fiber) and a polymer matrix, the Reuss (iso-stress) model defines a lower bound on the composite modulus,

$$E_{\mathrm{c}}={\left({\displaystyle \frac{{\varphi }_{\mathrm{f}}}{E_{\mathrm{f}}}}+{\displaystyle \frac{{\varphi }_{\mathrm{m}}}{E_{\mathrm{m}}}}\right)}^{-1}$$

(4)

whereas the Voigt (iso-strain) model provides an upper bound corresponding to perfectly aligned, continuous fibers,

$$E_{\mathrm{c}}={\varphi }_{\mathrm{f}}{E}_{\mathrm{f}}+{\varphi }_{\mathrm{m}}{E}_{\mathrm{m}}$$

(5)

with ${\varphi }_{\mathrm{f}}$ and ${\varphi }_{\mathrm{m}}$ denoting the fiber and matrix volume fractions and $E_{\mathrm{f}}$ and $E_{\mathrm{m}}$ their respective elastic moduli.

For realistic composites containing discontinuous or non-aligned fibers, the Voigt expression is commonly modified using the Krenchel formulation by introducing an orientation efficiency factor ${\eta }_0$ to account for fiber misalignment [130],

$$E_{\mathrm{c}}={\eta }_0{\varphi }_{\mathrm{f}}{E}_{\mathrm{f}}+{\varphi }_{\mathrm{m}}{E}_{\mathrm{m}}$$

(6)

This expression reflects the reduction in reinforcement efficiency arising from imperfect fiber orientation. Additional microstructural effects, particularly finite fiber length and porosity, can be incorporated through further efficiency factors, as proposed by Madsen et al. [131],

$$E_{\mathrm{c}}=\left({\eta }_0{\eta }_{\mathrm{l}}{\varphi }_{\mathrm{f}}{E}_{\mathrm{f}}+{\varphi }_{\mathrm{m}}{E}_{\mathrm{m}}\right){\left(1-{\varphi }_{\mathrm{p}}\right)}^n$$

(7)

where ${\eta }_{\mathrm{l}}$ accounts for finite fiber length, ${\varphi }_{\mathrm{p}}$ is the porosity volume fraction, and n is a porosity sensitivity exponent.

In real short-fiber polymer composites, an additional efficiency factor is required to capture the quality of stress transfer across the fiber–matrix interface. This interfacial efficiency factor, ${\eta }_i$, reflects the interfacial shear strength and the ability of the interface to transmit load from the matrix to the reinforcing fibers. Incorporating this effect yields

$$E_{\mathrm{c}}={\eta }_0{\eta }_{\mathrm{l}}{\eta }_i{\varphi }_{\mathrm{f}}{E}_{\mathrm{f}}+{\varphi }_{\mathrm{m}}{E}_{\mathrm{m}}$$

(8)

where ${\eta }_i$ depends on fiber surface chemistry, surface roughness, and the presence of chemical coupling or interphase layers introduced by surface treatments.

Matrix-Dependent Interfacial Response

Recent studies further demonstrate that the effectiveness of alkalization is strongly matrix-dependent [132], highlighting why results from nonpolar matrices such as PP cannot be directly transferred to polar biopolyesters such as PLA.

Matrix-dependent effects are particularly evident in PHBV-based biocomposites, where the response to alkali treatment depends strongly on both fiber type and treatment severity. Moderate NaOH treatment enhances interfacial adhesion and mechanical performance in PHBV composites reinforced with hemp fibers, whereas excessive treatment of flax fibers results in fiber degradation and reduced composite strength [133]. Beyond mechanical properties, alkali treatment also influences PHBV crystallinity, melt-processing behavior, and water absorption, underscoring the narrow optimization window required for PHBV-based biocomposites.

NaOH-treated lignocellulosic fibers exhibit substantially improved interfacial adhesion and thermal stability in polypropylene matrices, whereas similar treatments lead to limited adhesion enhancement in PLA matrices, where chemical compatibility is already high and interfacial failure is instead governed by interphase morphology and stress-transfer efficiency rather than surface polarity. These findings indicate that alkalization alone may be insufficient for polar biopolyesters and that additional compatibilization strategies are often required.

Table 7. Matrix-dependent effectiveness of alkaline (NaOH) fiber treatment in natural-fiber composites.

Matrix Class	Fibers Studied	Typical Response to Alkalization	Mechanical Trend and Key Limitations	Ref.
Polyolefin (PP)	Kenaf, hemp, bagasse, softwood	Mild NaOH treatment (5% for 1 h) removes surface impurities and increases fiber surface roughness; SEM of fracture surfaces shows no statistically significant morphological differences between treated and untreated fiber–PP composites under the applied conditions, despite improved fiber cleanliness and contact. All PP composites were compatibilized with 3 wt% PP-g-MA.	At 25 wt% fiber loading, tensile strength increased from 23.4 MPa (neat PP) to 40.1 MPa (kenaf), 30.1 MPa (softwood), 28.7 MPa (bagasse), and 28.2 MPa (hemp); however, no statistically significant differences in tensile strength or Young’s modulus were observed between treated and untreated fiber–PP composites at this fiber content.	[132]
Biopolyester (PLA)	Natural fiber mats (same as PP comparison)	Alkali-treated fiber mats were incorporated without compatibilizers; although surface chemistry was modified, SEM revealed only limited improvement in interfacial adhesion, with matrix stiffness and crystallinity dominating composite response and constraining gains achievable through NaOH treatment alone.	Neat PLA tensile strength and modulus were 71.2 MPa and 3630 MPa; incorporation of treated fibers produced only slight modulus increases relative to untreated-fiber composites, indicating a narrow optimization window.	[132]
Biopolyester (PHBV)	Hemp and flax fibers	Mercerization (2–10 wt% NaOH, 1 h) strongly altered fiber geometry and surface chemistry in a fiber-dependent manner; SEM showed progressive surface etching and fibrillation, with excessive treatment causing severe fiber thinning and degradation of interfacial integrity, particularly for flax fibers.	For hemp/PHBV (15 wt% fiber), tensile strength increased by ≈1.5% and modulus by ≈5% at 10 wt% NaOH, with impact tensile strength increasing by ≈12% and water absorption decreasing by ≈16%. For flax/PHBV, 10 wt% NaOH reduced fiber diameter by ≈51%, decreased impact tensile strength by up to ≈62%, and increased water absorption by up to ≈133%.	[133]
Thermoset (epoxy)	Kenaf fibers	Alkali treatment (6 wt% NaOH, room temperature) for 48 h produced cleaner fiber surfaces and altered surface morphology as observed by SEM; prolonged treatment (144 h) caused surface cracking and severe fiber damage. No chemical composition analysis was performed; effects are inferred from morphology rather than directly quantified chemical removal.	Relative to composites reinforced with untreated kenaf fibers, 48 h alkali treatment increased tensile strength by 36% (randomly oriented fiber mat laminates) and 11% (UD laminates), with modulus increases of 12% and 3.5%, respectively. Fibers treated for 144 h showed degraded single-fiber properties, but composites using these fibers were not fabricated.	[128]

provides a qualitative summary of reported matrix-dependent responses to alkalization. The trends presented are intended to highlight commonly observed behaviors; specific outcomes remain highly system-dependent and are influenced by fiber type, alkali concentration, treatment duration, compatibilization, and processing conditions.

3.3.2. Acetylation

Acetylation reduces hydrophilicity and improves compatibility with hydrophobic polymer matrices by converting fiber hydroxyl groups into ester-linked acetyl groups [–O–C(=O)–CH₃] via reaction with acetic anhydride [Reaction (b), Table 6]. This decreases polarity and interfacial tension, promoting improved wetting and interfacial adhesion, in contrast to alkali treatment, which primarily enhances roughness and hydroxyl accessibility. Representative experimental studies illustrating acetylation-based surface modification of natural fibers, together with their processing conditions and reported interfacial and mechanical outcomes, are summarized in

Table 8. Representative studies employing esterification-based surface modification of natural fibers, including acetylation and benzoylation.

Mechanism	Representative Conditions	Reported Effects	Ref.
Esterification of accessible hydroxyl groups increases fiber hydrophobicity and alters surface chemistry.	Solvent-free acetic or propionic anhydride treatment (120 °C, 30–180 min) applied to flax, hemp, and wood fibers.	ATR–FTIR and XPS confirmed ester formation; SEM showed removal of non-crystalline surface components; no composite mechanical testing.	[134]
Partial acetylation reduces fiber hydrophilicity; excessive acetylation degrades cellulose integrity and induces fiber cracking.	Flax fibers acetylated using acetic anhydride/toluene with perchloric-acid catalyst; polypropylene composites contained 30 wt% fiber and MAPP (5 wt% of fiber) as coupling agent.	Tensile and flexural strengths increased with degree of acetylation to an optimum at ≈18% acetyl content, then decreased at higher levels due to fiber degradation. Charpy impact strength decreased monotonically with acetylation. MAPP increased tensile and flexural strengths by ≈20–35% relative to uncoupled composites.	[135]
Esterification limits moisture-induced swelling and enhances interfacial stress transfer under wet conditions.	Wood fibers acetylated in boiling acetic anhydride; washed, dried, and compounded with PLA (20 wt% fiber).	Tensile strength of water-soaked composites increased by >30%; X-ray microtomography showed reduced fiber pull-out.	[136]
Substitution of fiber –OH groups with acetyl groups reduces moisture affinity and enhances fiber–epoxy interfacial adhesion.	Flax fiber mats were soaked for 60 min at 20 °C in aqueous acetic anhydride solutions (1–4 wt%), dried at 105 °C for 12 h, and vacuum-bagged into unidirectional epoxy laminates.	At optimal treatment, tensile strength increased by 55%, modulus by 58%, and bond shear strength by 7%; moisture resistance improved by ≈65%.	[137]
Chemical modification alters fiber constituents and surface morphology; moderate treatment improves tensile behavior, while excessive treatment degrades properties.	Fibers treated at room temperature for 3 h in dilute acetic anhydride media (2, 4, and 6 vol%; solvent not specified in the original study).	Optimal concentration reported as 4 wt% for most fibers; Combretum racemosum reached tensile strength 155 MPa and strain 0.046 at 4 wt%; tensile strength typically decreased beyond 4 wt%.	[138]
Acetylation induces fiber swelling, surface roughening, and void formation, enhancing kenaf–starch interfacial adhesion primarily through mechanical interlocking.	Kenaf powder (<1 mm) acetylated in acetic anhydride/toluene with perchloric acid catalyst at 60 °C for 3 h; composites compression molded at 130 °C and 3000 psi.	Interfacial adhesion and mechanical performance improved; PVA and PEG plasticizers did not enhance properties. Biodegradability was highest under moist soil (soil + water) conditions, reaching ≈80%.	[139]
Partial esterification of cellulose and CNF hydroxyl groups reduces hydrophilicity and improves compatibility with polyethylene matrices.	Aqueous alkaline acetylation using NaOH and acetic anhydride (70–80 °C, ≈120 min); treated kraft fibers and CNFs compounded with HDPE.	Composite tensile strength increased from ≈19 to 30–40 MPa; water absorption reduced to ≈0.6 wt% after 10 days; improved ductility.	[140]
Fiber acetylation improves fiber–epoxy adhesion; acetylated cellulose nanofibers provide nanoscale reinforcement and further enhance interfacial binding.	Hand lay-up followed by hot pressing; acetylated sisal fabric composites with acetylated CNF (0.5–1 wt%).	With 0.5 wt% acetylated CNF, tensile/flexural/impact strengths increased by 331%/118%/265% vs. neat epoxy.	[141]
Bulky aromatic ester groups reduce fiber hydroxyl density and enhance compatibility with PVC.	Alkali pretreatment followed by benzoyl chloride treatment; 2 h at 140 °C; 10 wt% fiber loading.	Tensile strength increased from 11.3 to 12.8 MPa; tensile modulus increased from 112 to 305 MPa; water uptake reduced to ≈0.65%.	[142]
Alkali pretreatment enhances surface roughness and cellulose exposure; subsequent acetylation reduces fiber polarity and improves compatibility with polypropylene.	Unidirectionally oriented nonwoven BBF/PP composites. BBF mercerized in 5% NaOH (RT, 2 h; solution–BBF = 15:1 w/w), then acetylated using acetic anhydride/toluene with perchloric acid catalyst (60%) at 70 °C for 3 h.	For acetylated (alkali-pretreated) BBF/PP, tensile strength and modulus increased up to 40 wt% fiber, then decreased at higher loading. Alkali-treated BBF/PP showed a similar trend, reaching 38.3 MPa tensile strength at 40 wt% fiber (+33.2% vs. neat PP).	[126]

.

Across plant fiber systems, acetylation has been reported to improve stiffness, strength, and moisture resistance by modifying fiber surface polarity and fiber–matrix compatibility. In acetylated banana bunch fiber (BBF)/PP composites, significant increases in flexural strength and modulus were observed, together with an 83% increase in IFSS, reaching 7.5 MPa, and a 62.7% reduction in moisture absorption [126]. Moderate acetylation levels (≈4%) were found to provide a balance between improved interfacial performance and preservation of fiber structure in the system studied by Oladele et al. [138]. Kenaf–starch composites also exhibited improved stiffness and moisture resistance after acetylation, which was attributed to increased nucleation and improved fiber–matrix compatibility [139].

For unidirectional stitched sisal fabric–reinforced epoxy laminates fabricated by hand lay-up followed by hot pressing (33 wt% fabric content), Sukmawan et al. [141] evaluated tensile, flexural, impact, and interlaminar shear strength (ILSS) and reported that both sisal-fabric acetylation and the addition of acetylated cellulose nanofibers (ACNF) improved interfacial performance. The highest ILSS was obtained for hybrid laminates containing acetylated sisal fabrics together with 0.5 wt% ACNF, rather than for acetylated fabrics alone, indicating that interfacial strengthening arose from the combined effects of fiber acetylation and nanofiber reinforcement [141].

In the same study, incorporation of acetylated sisal fabrics together with 0.5 wt% ACNF increased the tensile, flexural, and impact strength of the epoxy matrix by 331%, 118%, and 265%, respectively, relative to neat epoxy (corresponding to ≈4.3-fold, 2.2-fold, and 3.6-fold increases) [141]. These enhancements were attributed to improved interfacial adhesion and more effective load transfer arising from the combined presence of acetylated fibers and dispersed nanofibers [141]. These results indicate that acetylation can be used to tune fiber polarity and improve durability while retaining the biodegradable character of lignocellulosic reinforcements.

Beyond acetylation, benzoylation has been explored as an esterification-based surface modification to improve fiber–matrix compatibility in PVC composites. Benzoyl chloride–treated date palm fiber/PVC composites exhibited improved interfacial adhesion, higher tensile strength and modulus (12.8 MPa and 304.5 MPa, respectively, at 10 wt% fiber), and reduced water uptake compared with untreated systems, as confirmed by FTIR analysis and fracture morphology [142]. Although moisture-induced degradation was not fully eliminated, benzoylation significantly improved stiffness and interfacial stability under hydrothermal exposure [142].

3.3.3. Coupling Agents

Coupling agents create durable chemical bridges between hydrophilic fibers and hydrophobic polymer matrices, thereby improving stress transfer, reducing moisture sensitivity, and stabilizing long-term performance [143]. Unlike alkali or acetylation treatments that primarily modify surface energy, coupling agents form covalent bonds across the interface, yielding much stronger molecular-scale adhesion.

Silane-Based Interfacial Coupling

Silanes remain the most extensively studied class of coupling agents for natural-fiber composites owing to their versatile chemistry and effectiveness across a wide range of matrices. Their mechanism involves hydrolysis of alkoxy groups to form silanols, condensation with fiber surface hydroxyls to generate covalent Si–O–cellulose linkages [Reaction (c), Table 6], and subsequent interaction or copolymerization with the polymer matrix [144]. Through this dual reactivity, silanes reduce interfacial polarity mismatch, suppress fiber pull-out, and enhance stress transfer. Representative experimental studies employing silane coupling agents, together with their treatment chemistries, processing routes, and resulting interfacial and mechanical outcomes, are summarized in

Table 9. Representative studies employing silane coupling agents to modify natural-fiber surfaces and enhance fiber–matrix interfacial performance.

Interface Modification	Representative Conditions	Reported Effects	Ref.
Methacryloxy silane (waru bark fiber/polyester): Sequential alkali–silane treatment improves fiber–polyester interfacial bonding, as indicated by SEM-observed cleaner, denser fiber surfaces; tensile behavior is governed primarily by continuous-fiber orientation.	Waru (Hibiscus tiliaceus) bark fibers treated with 6 % NaOH (120 min), then immersed in ≈0.75 wt% methacryloxypropyltrimethoxysilane solution (pH 3.4–4.5) for 4 h and oven-dried at 70 °C; polyester composites fabricated by vacuum-assisted resin infusion.	Interfacial bonding improved relative to alkali-only fibers. Tensile strength was strongly orientation-dependent, reaching 401.4 MPa for unidirectional composites and decreasing to 65.2 MPa for 45° /45° laminates.	[144]
KH570 silane (jute/PLA): Methacryloxy silane coupling enhances fiber–PLA adhesion and reduces fiber hydrophilicity.	KH570 at 1–9 wt% in ethanol–water (2:3); 1 h immersion; oven-dried at 80 °C.	Tensile strength increased by 10–15%; flexural strength by 40–60%; moisture uptake reduced.	[145]
Glymo silane (flax/epoxy): Epoxy-functional silane enables covalent coupling with the epoxy matrix and suppresses moisture sensitivity.	Dynasylan® Glymo at 20–60 g/L; pad–dry–cure process (120 °C drying, 150 °C curing).	Tensile strength increased by 15%; flexural strength by 117%; impact strength by 20%; moisture uptake reduced.	[146]

.

In jute/PLA composites, $\gamma$-methacryloxypropyltrimethoxysilane (KH570) treatment significantly improved fiber–matrix interfacial bonding, as evidenced by fracture SEM showing a marked reduction in fiber pull-out voids and the formation of a resin-philic silane layer on the jute surface. Tensile and flexural properties increased with KH570 concentration and reached optimal performance at ≈5 wt% silane loading [145].

In flax/epoxy laminates, silane treatment produced particularly pronounced improvements, increasing flexural strength from about 7.0 to 15.3 MPa while substantially reducing moisture uptake. Moisture regain decreased to ≈7.1% at the reinforcement level and to about 0.42% in the resulting composites, indicating effective suppression of fiber hydrophilicity and enhanced interfacial sealing [146].

Beyond chemistry alone, fiber architecture strongly modulates the benefits of silane treatment. In continuous natural-fiber composites, fiber alignment and surface modification together govern tensile performance: unidirectional alkali–silane–treated waru (Hibiscus tiliaceus) fiber/polyester laminates achieved tensile strengths exceeding 400 MPa, whereas cross-ply (0° /90°) and ±45° configurations exhibited substantially lower strengths due to off-axis load transfer inefficiencies [144].

Maleic Anhydride Coupling

Maleic anhydride coupling is a widely used strategy for improving the compatibility of lignocellulosic fibers with nonpolar and moderately polar polymer matrices. Through grafting of anhydride-functional units onto the polymer backbone [Reaction (d), Table 6], MA facilitates stronger interfacial interactions with fiber hydroxyl groups and promotes more efficient stress transfer across the fiber–matrix interface. Owing to this interfacial reactivity, MA-based compatibilizers are extensively employed in PLA-, PP-, and PCL-based composite systems. Representative primary studies employing MA functionalization or maleated compatibilizers, together with their processing conditions and reported interfacial and mechanical outcomes, are summarized in

Table 10. Representative studies employing maleic anhydride functionalization or maleated compatibilizers (e.g., PP-g-MA, PLA-g-MA) to enhance interfacial bonding in natural-fiber–reinforced polymer composites, with performance changes reported relative to uncompatibilized systems at identical fiber loadings.

Interface Modification	Representative Conditions	Reported Effects	Ref.
MA and glycerol triacetate (GTA) compatibilizers (PALF/PLA): MA forms polar interactions with fiber hydroxyl groups, whereas GTA primarily plasticizes the PLA matrix.	Compression-molded PALF/PLA composites with MA or GTA additives.	MA increased tensile strength by ≈3.9% and impact strength by ≈3.3%, while GTA increased impact strength by ≈6.5% but reduced tensile strength by ≈1.4%; SEM confirmed improved adhesion only with MA.	[147]
PLA-g-MA: Covalent grafting introduces anhydride groups that chemically couple PLA to cellulose fibers.	PLA-g-MA (MA grafting ≤ 3.5%) compounded with 10 wt% cellulose fibers.	Cellulose increased tensile strength by ≈15%, while addition of 4 wt% PLA-g-MA increased tensile strength by up to ≈24% relative to neat PLA, with reduced fiber pull-out.	[148]
MA coupling-agent treatment of jute and kenaf fibers (hybrid fiber/PLA biocomposites): MA-modified fibers used to improve fiber–matrix interaction.	Chopped jute and kenaf fibers (≈2 mm) treated in 5% MA solution with 0.5% NaOH catalyst (1 h; MA:fiber = 1:20, w/w), dried at 60 °C; hot-compressed PLA composites with 50 wt% PLA + 25 wt% jute + 25 wt% kenaf.	Tensile strength +35%, tensile modulus +15%, and impact strength +20% relative to untreated composites; FTIR showed increased C=O intensity after MA treatment and SEM indicated improved fiber dispersion/adhesion; water absorption was reduced (qualitative).	[149]
MA-grafted polyolefin compatibilizer in agave-fiber/LLDPE composites: Anhydride groups improve fiber wetting and interfacial stress transfer in nonpolar matrices.	LLDPE composites with ≈70 wt% agave fiber and MA-grafted compatibilizer.	Tensile strength increased to ≈18.5 MPa and impact strength to ≈51.6 kJ m−2, accompanied by reduced interfacial voids observed by SEM.	[150]
MAPP in hemp-fiber/PP composites: MAPP improves fiber dispersion and chemical coupling via anhydride–hydroxyl interactions.	Extruded and compression-molded PP/hemp composites with 3–5 wt% MAPP.	Optimal MAPP (3–5 wt%) increased impact resistance and hardness, suppressed PP $\beta$-phase formation, and improved dispersion; excessive MAPP (10 wt%) reduced hardness.	[151]
MAPP grade selection in highly filled kenaf/PP boards: Compatibilizer efficiency depends on anhydride content and molecular weight.	Hot-pressed PP boards with 85 wt% kenaf fiber and 5 wt% MAPP.	Optimized MAPP grade yielded flexural strength ≈ 24 MPa and flexural modulus ≈ 3.5 GPa, demonstrating effective stress transfer at ultra-high fiber loading.	[152]
MAPP in bamboo-fiber/PP composites under natural ageing: Improved interfacial bonding limits damage accumulation during environmental exposure.	PP/bamboo-fiber composites (30 wt% fiber) with 3 wt% MAPP, aged outdoors for 6–12 months.	After ageing, MAPP-containing composites showed reduced losses in Young’s modulus (≈23%), tensile strength (≈18%), and impact strength (≈6%) relative to uncompatibilized systems.	[153]
PP-g-MA in recycled-PP/rice-husk composites (3D printing): Compatibilization improves particle dispersion and interfacial load transfer while mitigating shrinkage-driven warping via reduced crystallinity.	Fused-filament fabrication using rPP with 10 wt% rice-husk particles (sieved to 150 and 250 $\mathsf{\mu }$m) and 5 wt% PP-g-MA; particle sizes below 250 $\mathsf{\mu }$m were identified as suitable for filament production.	For 150 $\mathsf{\mu }$m rice-husk composites, PP-g-MA increased tensile strength from 11.67 to 20.36 MPa and improved warping by 62% relative to rPP; the 250 $\mathsf{\mu }$m rice-husk composite without compatibilizer could not be printed into specimens due to filament nonuniformity and nozzle blockage.	[154]
MA-grafted PCL with particulate cellulose (MCC): One-pot MA grafting introduces anhydride functionality on PCL, improving polymer–particle compatibility and encapsulation of microcrystalline cellulose.	One-pot MA grafting of PCL followed by incorporation of MCC; solvent-cast composite films.	Tensile strength increased by ≈77.8%; surface roughness reduced by ≈66.7%; water contact angle increased to ≈87.5°; DSC and TGA show melting temperature (≈59 °C) and thermal stability unchanged relative to neat PCL.	[155]

.

In PLA-based composites, MA compatibilization has been shown to significantly enhance tensile performance. Oliver-Ortega et al. [148] reported that incorporation of a PLA-g-MA coupling agent increased the ultimate tensile strength of PLA/cellulose composites by up to 24% at a 4 wt% compatibilizer loading relative to neat PLA, with fracture morphology indicating reduced fiber pull-out and more efficient stress transfer. Similarly, in hybrid jute/kenaf–reinforced PLA composites, MA treatment increased tensile strength by 35% and tensile modulus by 15%, while also improving impact resistance [149].

For polyolefin systems, the effectiveness of MA coupling depends strongly on compatibilizer chemistry and dosage. In hemp/PP composites, Da Silveira et al. [151] observed improved impact strength at intermediate MAPP loadings of 3–5 wt%, whereas higher loadings (10 wt%) led to a decline in performance, attributed to over-plasticization of the matrix. In highly filled kenaf/PP boards containing 85 wt% fiber, Sanadi and Stelte [152] found that optimum mechanical performance occurred at ≈5 wt% MAPP, with low-molecular-weight, high-anhydride-content grades producing the highest flexural strength (≈24 MPa) and flexural modulus (≈3.5 GPa) under hot-pressing conditions.

Beyond improvements in as-fabricated composites, MA-based compatibilization can also enhance durability under environmental exposure. In PP/bamboo-fiber composites subjected to natural ageing, Fajardo Cabrera de Lima et al. [153] reported that MA-modified formulations retained stiffness and strength more effectively than uncoupled composites. For example, Young’s modulus of the coupled composite decreased from 1.63 to 1.45 GPa after six months of exposure, compared with a decrease from 1.40 to 1.16 GPa for the uncoupled system. Morales et al. [154] further demonstrated that PP-g-MA compatibilization enables stable fused-filament fabrication of recycled polypropylene filled with rice-husk particles by improving interfacial adhesion and reducing crystallization-driven warping, thereby recovering tensile strength and printability in particle-filled filaments.

MA grafting has also been applied in biodegradable polyester systems beyond PLA. In polycaprolactone–microcrystalline cellulose (PCL–MCC) composites, Li et al. [155] showed that a one-pot MA grafting approach reduced surface roughness by 66.7%, increased hydrophobicity (water contact angle up to 87.5°), and improved tensile strength by 77.8% without altering the thermal stability or melting behavior of the PCL matrix. FTIR analysis confirmed successful MA grafting without backbone degradation, indicating that the mechanical enhancements arise primarily from improved interfacial adhesion and heterogeneous nucleation rather than polymer chain scission.

3.3.4. Hybrid and Nanostructured Interphases in Composite Systems

Hybrid interphases that combine chemical modification with micro- or nanoscale structural features can substantially enhance fiber–matrix bonding by improving wetting, increasing effective contact area, and suppressing fiber pull-out. The representative experimental studies discussed in this section are summarized in

Table 11. Advanced interface modification approaches, including nanostructured coatings, bio-based coupling agents, enzymatic, plasma, and thermal treatments, and hybrid interphases. Symbols ↑ and ↓ indicate increases and decreases, respectively, in the reported parameters.

Interface Modification	Primary Mechanism	Representative Conditions	Reported Property Changes	Ref.
Acetylated sisal fabric + acetylated cellulose nanofiber (ACNF) hybrid interphase	Acetylation reduces fiber hydrophilicity and improves chemical compatibility with epoxy; dispersed ACNF forms a nanostructured interphase that enhances stress transfer and shifts failure from fiber pull-out toward fiber breakage.	Sisal fabrics acetylated with acetic anhydride; ACNF added at 0–1 wt% (optimum 0.5 wt%); laminates fabricated by hand lay-up and hot pressing (90 °C, 1 MPa) followed by post-curing at 105 °C; fiber content ≈ 33 wt%.	Tensile strength increased from 112.0 to 126.6 MPa after acetylation and further ↑ by ≈22% at 0.5 wt% ACNF; flexural strength and ILSS increase relative to untreated laminates (ILSS ≈ 21.4 MPa); impact strength ↑ by ≈15%; fracture mode shifts from fiber pull-out to fiber breakage at optimal ACNF loading.	[141]
Silane–TiO2 nanocoating (dip-coated flax fibers)	Silane coupling generates Si–O–C (cellulose), Si–O–Ti (TiO2), and Si–O–Si linkages, enabling nanoparticle grafting and enhanced interfacial mechanical interlocking with epoxy.	Dip coating in ethanol–silane = 10:1 (vol.) with dispersed TiO2; TiO2 = 0.2–2 wt% (optimum 0.2–0.8 wt%); 30 min immersion, oven-dried at 110 °C for 8 h; laminates compression-molded (5 MPa, 100 °C), ${\varphi }_{\mathrm{f}}\approx 28\%$.	Tensile strength +22% (max. at 0.4 wt% TiO2); flexural strength +24% (73→91 MPa) and ILSS +16% (14.29→16.64 MPa) at 0.6 wt% TiO2; water diffusion coefficient reduced by ≈42% (0.6 wt% TiO2).	[156]
Graphene oxide coating (after alkali treatment)	Oxygen-rich GO nanosheets introduce hydrogen bonding and mechanical interlocking, increasing surface energy and crack-bridging capability.	5 wt% NaOH (100 °C, 1 h) followed by 1 wt% GO dip coating (30 min).	Tensile strength +113%; flexural strength +93%; increased impact resistance; water absorption reduced to ≈2%.	[157]
Citric-acid esterification (wood fibers)	Citric acid carboxyl groups esterify fiber –OH groups (FTIR C=O at 1720 cm−1; –OH ↓), reducing hydrophilicity and strengthening PCL–fiber interfacial adhesion via improved wetting and stress transfer.	Wood fibers soaked in citric-acid aqueous solution (0.05–1.00 M, 1 h, RT), dried (40 °C, 24 h, vacuum), thermally treated for esterification (30 min; 80 °C identified as optimum), washed to neutral pH, and re-dried; composites melt compounded (80 °C, 9 min, 50 rpm) and compression molded (80 °C, 6 min, 2 MPa).	Yield strength ↑ up to ≈30% at 0 °C and ↑ ≈21% at 25 °C relative to PCL composites with identical 10 wt% untreated wood-fiber content; Vicat softening temperature ↑ ≈6 °C vs. neat PCL; water uptake ↓ (90-day immersion: 7.44%→3.22%) and water contact angle ↑ (51.2°→61.2°).	[158]
Fluorocarbon treatment	Application of a fluorocarbon finish to reduce fiber hydrophilicity and modify the fiber–matrix interphase.	Commercial fluorocarbon textile finish (Ruco Guard® AFR6, Rudolf Group) applied by pad–dry–cure at concentrations of 80–120 g L−1.	Reinforcement moisture regain reduced to ∼3%; reinforcement tensile strength increased by ≈35%. In epoxy composites, moisture regain decreased to ≈0.4% with tensile, flexural, and impact strength increases of ≈23%, ≈149%, and ≈31%, respectively.	[146]
Laccase-mediated enzymatic treatment	Laccase-induced oxidation modifies lignin-rich surface regions, increasing interfacial reactivity and compatibility with polymer matrices.	Aqueous laccase treatment under controlled temperature and pH conditions.	Enhanced fiber–matrix adhesion and improved interfacial performance, supported by spectroscopic evidence and reduced interfacial defects observed in microscopy.	[159]
Multienzyme retting (xylanase, laccase, lipase)	Selective enzymatic removal of pectin and hemicellulose improves surface cleanliness and fiber–matrix contact while preserving cellulose microfibrils.	Mixed enzyme system applied under mild aqueous conditions (≈17 h).	Flexural strength increased by ≈10–15%, accompanied by reduced water absorption and improved surface resistance to biological degradation relative to untreated fibers.	[160]
Ionic-liquid-modified lignin (bio-coupling agent)	[Emim][OAc] modifies lignin chemistry, increasing interfacial reactivity and dispersion in recycled PP matrices.	Lignin modified with [Emim][OAc] (20:1–30:1); composites with 60 wt% bagasse fiber and 1–5 wt% IL-lignin.	Flexural strength up to 41 MPa; flexural modulus up to 2.75 GPa; tensile strength ≈ 25 MPa; water absorption reduced by 8–34%.	[161]
Heat treatment	Thermally induced stress relaxation, reduced void content, and partial reduction in accessible hydroxyl groups.	80 °C for 8 h with controlled cooling.	Tensile and flexural strength increases of ≈20–30%; improved thermal and hydrothermal stability.	[162]
Low-temperature plasma treatment and biodegradation response	Plasma-induced surface oxidation enhances fiber–matrix adhesion and increases biological accessibility.	Low-temperature plasma applied to natural fibers; biodegradation assessed by 30-day soil burial; SEM analysis.	Plasma treatment amplified tensile strength gains from fiber reinforcement and accelerated biodegradation onset relative to untreated composites.	[163]
In-line plasma strategies during compounding	Plasma activation improves interfacial bonding, with effectiveness dependent on direct fiber treatment.	Plasma treatment of flax fibers, PA6 matrix, or both during melt compounding.	Direct plasma treatment of flax fibers or in-line treatment during compounding significantly increased tensile and flexural strength, whereas plasma treatment of PA6 alone had limited effect.	[164]

.

In continuous natural-fiber laminates, Sukmawan et al. [141] showed that acetylation of unidirectional stitched sisal fabrics, followed by incorporation of a small fraction (0.5 wt%) of acetylated cellulose nanofibers into the epoxy matrix, generated a nanostructured hybrid interphase that markedly improved stress transfer. Compared with untreated sisal/epoxy laminates, tensile strength increased from 112.0 to 126.6 MPa after acetylation, with further gains upon nanofiber addition, and fracture morphologies shifted from fiber pull-out to fiber breakage, indicating more effective interfacial load transfer.

Nanostructured surface coatings provide a complementary strategy for interfacial enhancement by introducing both chemical coupling and nanoscale roughness. Prasad et al. [156] demonstrated that dip coating flax fibers with silane-functionalized TiO₂ nanoparticles produced a grafted hybrid interphase characterized by Si–O–C, Si–O–Ti, and Si–O–Si linkages. At optimal nanoparticle loadings (0.2–0.8 wt%), tensile, flexural, and ILSS increased by up to 22%, 24%, and 16%, respectively, while the water diffusion coefficient decreased by ≈42%, reflecting improved load transfer and reduced moisture ingress.

Similarly, Islam et al. [157] reported that graphene oxide (GO) nanosheets deposited on alkali-treated natural fibers introduced oxygen-rich functional groups that enhanced hydrogen bonding and mechanical interlocking with the epoxy matrix. The resulting composites exhibited large gains in tensile and flexural strength (up to 113% and 93%, respectively), along with increased impact resistance and a pronounced reduction in water absorption to ≈2%, confirming the effectiveness of two-dimensional nanocoatings in suppressing fiber pull-out and improving environmental resistance.

A related hybrid-interphase concept was reported by Simonini et al. [158], who employed citric-acid-functionalized wood fibers in polycaprolactone (PCL) matrices. Ester bridges formed during citric acid treatment enhanced interfacial cohesion and load transfer while simultaneously reducing water uptake, demonstrating that chemically induced hybrid interphases can provide concurrent improvements in mechanical performance and environmental resistance without the need for inorganic nanofillers.

Although fluorocarbon treatments do not introduce an explicit nanostructured or hybrid interphase, they may be regarded as a simplified form of interphase engineering in which surface chemistry and surface energy are tailored to influence fiber–matrix interactions. Rauf et al. [146] reported that fluorocarbon treatment of flax fibers deposits low–surface-energy fluorinated chains on the cellulose surface, substantially reducing fiber hydrophilicity and improving apparent interfacial performance. At the highest treatment level, the moisture regain of the treated flax reinforcement decreased to ≈3% relative to untreated fibers, while reinforcement tensile strength increased by about 35%. When incorporated into epoxy composites, fluorocarbon-treated fibers produced composites with very low moisture regain (≈0.4%) and pronounced mechanical enhancements, including increases of about 23% in tensile strength, 149% in flexural strength, and 31% in impact strength compared with untreated flax–epoxy composites. However, because fluorination lowers surface polarity and reduces the availability of reactive functional groups, it may decrease true fiber–matrix interfacial shear strength, even when bulk mechanical properties improve through moisture suppression and reduced defect formation.

3.3.5. Bio-Based and Biomolecular Coupling Strategies

Bio-based and biomolecular surface treatments (see Table 11) provide selective, low-energy routes to interfacial modification while preserving fiber integrity and cellulose microstructure. Martin et al. [159] demonstrated that enzymatic pretreatment using laccase significantly enhanced interfacial bonding in lignosulfonate/wood-fiber composites when both the matrix and fibers were treated simultaneously. Under optimized conditions (17 h laccase treatment), flexural strength increased from 6.55 to 17.53 MPa and flexural modulus from 1.34 to 3.51 GPa, underscoring the importance of coordinated interphase formation rather than isolated fiber modification.

Selective enzymatic removal of weak boundary layers has also proven effective in thermoplastic systems. Meng et al. [160] reported that xylanase-treated bamboo fibers in polyethylene composites exhibited a 15.45% increase in flexural strength and a 13.31% increase in flexural modulus relative to untreated fibers, while simultaneously reducing water uptake and improving resistance to fungal growth. These improvements were attributed to targeted modification of hemicellulose-rich surface regions, which enhanced surface cleanliness and fiber–matrix contact without compromising cellulose microfibrils.

Bio-derived coupling agents offer an additional biomolecular strategy for interfacial reinforcement. Younesi-Kordkheili and Pizzi [161] demonstrated that ionic-liquid-modified lignin functions as an effective bio-coupling agent in recycled polypropylene/bagasse-fiber composites. Increasing the content of modified lignin from 1 to 5 wt% increased flexural strength from 22 to 41 MPa and flexural modulus from ≈1.8 to 2.7 GPa, achieving mechanical performance comparable to or exceeding that obtained using conventional maleated polypropylene.

3.3.6. Thermal Surface Restructuring and Interfacial Stabilization

Thermal treatments modify fiber–matrix interfaces by promoting matrix consolidation, reducing void content, and stabilizing interfacial morphology. Shifa et al. [162] reported that post-cure heat treatment at 80 °C for 8 h significantly enhanced interfacial performance in cotton–glass/epoxy hybrid laminates. For G–C–G stacking sequences, tensile strength increased from 176 to 213 MPa and flexural strength from 203.5 to 259.2 MPa following heat treatment, with scanning electron microscopy revealing reduced porosity and improved fiber wetting.

Although subsequent water immersion degraded mechanical properties through fiber swelling and partial interfacial weakening, heat-treated laminates consistently retained higher stiffness and strength than untreated counterparts after 30 days of exposure. These results demonstrate that thermal processing can stabilize interfacial morphology and partially mitigate moisture-induced degradation, particularly in hybrid reinforcement architectures where matrix consolidation plays a dominant role.

3.3.7. Plasma-Based Interfacial Activation and In-Line Processing

Plasma and corona discharge treatments modify fiber surfaces through mild oxidation and physical etching, increasing surface energy and polar functionality without altering bulk fiber structure. As summarized in Table 11, studies show that plasma activation can markedly enhance interfacial bonding and mechanical performance when processing conditions are carefully controlled.

Nyssanbek et al. [163] examined low-temperature plasma treatment of natural fibers in cellulose-reinforced polypropylene composites and demonstrated that plasma activation amplifies the mechanical benefits of fiber reinforcement. Relative to the neat polymer (18 MPa), tensile strength increased to 21 MPa with untreated fibers, to 25 MPa when plasma-treated fibers were used, and to 29 MPa when plasma-treated fibers were incorporated into the composite, corresponding to increases of 16.7%, 33.3%, and 50%, respectively. Plasma-treated specimens also exhibited an earlier onset of biodegradation (day 12 versus day 19 for untreated composites), consistent with increased biological accessibility associated with surface oxidation.

Pitto et al. [164] investigated in-line plasma strategies during compounding of flax-fiber–reinforced PA6 composites. Plasma treatment of flax fibers prior to compounding increased tensile strength by 28.1% and flexural strength by 31.7%, while simultaneous plasma treatment of both fibers and matrix during melt blending yielded flexural strength increases of up to 33.5%. In contrast, plasma treatment of the PA6 matrix alone produced negligible changes in neat polymer properties, underscoring that plasma effectiveness depends primarily on direct interfacial activation of the reinforcement rather than bulk matrix modification.

These findings are consistent with the broader assessment by Ramachandran et al. [117], who emphasize that plasma and corona treatments are highly effective for improving wettability and interfacial shear strength but require careful integration with processing to mitigate surface aging effects. They are most effective when applied directly to fibers or implemented in-line during composite fabrication.

3.4. Interfacial Strategies for Protein-Based (Animal) Fibers

Unlike lignocellulosic plant fibers, which require chemical treatment to mitigate hydrophilicity from abundant surface hydroxyl groups, protein-based animal fibers are inherently more reactive due to amino (–NH₂), carboxyl (–COOH), and thiol (–SH) functionalities along their keratin or fibroin backbones [165]. These functional groups enable potential interactions with diverse polymer matrices through hydrogen bonding, electrostatic attraction, or covalent coupling. However, their cuticular morphology, characterized by overlapping surface scales with relatively smooth interscale regions, provides limited surface roughness, restricting mechanical interlocking and hindering efficient stress transfer across the fiber–matrix interface [166]. As a result, untreated animal fibers often display poor interfacial adhesion and reduced load-bearing efficiency, compromising composite strength and durability [167].

Targeted surface modification strategies have been employed to enhance both chemical reactivity and physical adhesion. Silane treatment of wool fibers has been shown to improve compatibility with poly(methyl methacrylate) (PMMA), increasing impact strength and stabilizing interfacial bonding [166]. Peroxide oxidation increases wool fiber roughness and introduces additional polar groups, improving adhesion with PLA matrices without damaging the keratin backbone [168]. For silk-based reinforcements, alkali treatment of filature silk waste hydrolyzes amide linkages within the outer sericin layer, converting it into soluble peptide fragments and exposing the fibroin core [169]. The resulting increase in surface roughness and polarity enhances interfacial wetting, mechanical interlocking, and tensile modulus in epoxy composites. Alkaline modification of chicken feather and cow hair fibers similarly increases surface wettability and interfacial affinity with unsaturated polyester resin, yielding composites with higher strength and toughness [167].

3.5. Matrix Modification and Molecular Design Strategies

Matrix modification strategies complement fiber surface treatments by enhancing cohesion, flexibility, and stress dissipation within the polymer phase. Plasticizers such as glycerol, sorbitol, or vegetable oils improve the ductility and water resistance of hydrophilic biopolymers like starch, PLA, and PHAs [170,171,172]. Functional additives, including flame retardants, UV stabilizers, and impact modifiers (typically rubbery or elastomeric phases that absorb impact energy and suppress brittle fracture) have also been incorporated into wood–plastic and other biocomposites to tailor durability and environmental resistance [173,174,175]. Controlled phase separation through block copolymerization has further emerged as a molecular design approach to toughen bio-based matrices. Bio-derived modifiers such as acrylated epoxidized soybean oil (AESO), epoxidized linseed oil (ELO), lignin derivatives, and poly($\alpha$-methylene-$\gamma$-butyrolactone) (PMBL) introduce nanoscale domains that inhibit crack initiation and propagation, albeit often at the expense of stiffness and glass-transition temperature [82]. These copolymers can be tailored through monomer selection and chain architecture, offering a versatile route to engineer mechanical, thermal, and interfacial properties across diverse bio-based matrix systems.

3.6. Mesostructural Strategies for Interfacial Enhancement

Beyond chemical surface modification, interfacial performance can be enhanced through mesostructural control of the reinforcement phase. Needle-punched nonwoven architectures improve stress transfer by converting loose fiber webs into three-dimensional entangled networks, increasing fiber–fiber friction and micromechanical interlocking in addition to conventional fiber–matrix adhesion.

Kenned et al. [176] demonstrated this approach using untreated banana pseudostem fibers (Musa sapientum) formed into needle-punched mats and impregnated with unsaturated polyester resin. Mechanical properties peaked at ≈40 wt% fiber content, where tensile and flexural strengths, hardness, and indentation resistance approached those of glass-fiber-reinforced polyester, while thermal stability was retained up to ≈260 °C. Fractography showed reduced fiber pull-out and improved matrix continuity, confirming that mechanical interlocking within the needle-punched network enables efficient load transfer without chemical surface treatment.

3.7. Reinforcement Effects of Nanofillers in Natural-Fiber Composites

Nanofillers can improve the performance of natural-fiber composites, but the magnitude and even the direction of change depends on filler chemistry, loading level, and how well the filler is dispersed and coupled to the fiber/matrix interface [177]. For sisal/unsaturated polyester laminates containing 3 wt% nanoclay, Venkatram et al. [178] reported compressive strength values of 165–173 MPa versus 151–158 MPa without nanoclay, and impact energy increased from 1.5–2.0 J (control) to 2.0–3.1 J (with nanoclay).

Bagasse/HDPE composites containing 2–5 wt% nano-SiO₂ exhibit substantial tensile-strength gains of roughly 40–70% [179], while adding 2 wt% nano-TiO₂ to bagasse/EVA matrices increases tensile strength by 10% (11.4 MPa to 12.5 MPa) [180]. Ramie/epoxy composites modified with 0–0.6 wt% CNTs show 34% and 37% increases in flexural strength and modulus, along with higher interlaminar shear strength and fracture toughness arising from CNT-mediated crack bridging and enhanced interfacial adhesion [181]. Nanographene reinforcement likewise improves performance: in bagasse/PP composites, 0.1 wt% nanographene increases tensile, flexural, and notched-impact properties before agglomeration at higher loadings limits effectiveness [182], while in jute/epoxy laminates, 0.3–3 wt% graphene reduces drilling-induced delamination and increases tensile strength by 32–49%, reflecting greater matrix stiffness, improved fiber–matrix adhesion, and enhanced crack resistance [183].

In polypropylene/wood-flour composites, 3 wt% montmorillonite raises tensile strength by about 20% and flexural strength by 13%, while lowering 24 h water absorption through reduced interfacial voids and improved fiber wetting [184]. Sisal/epoxy composites show similarly pronounced gains: adding 2–5 wt% nanoclay increases tensile strength and modulus by 27% and 47%, and a 5 wt% loading reduces water absorption by roughly one-third via platelet-induced barrier effects and microvoid filling [185]. However, these enhancements are not universal. In hemp/polyester bio-based hybrids, 1.5 wt% nanoclay improves toughness but decreases tensile strength by about 20%, indicating that resin embrittlement and weakened fiber–matrix adhesion can outweigh the benefits of reduced moisture uptake [186].

Hierarchical nanoscale surface engineering provides an additional, highly effective pathway for interfacial reinforcement. Depositing graphene oxide nanoparticles onto jute fibers produces a dense, multiscale interfacial architecture that dramatically improves load transfer [187]. The interfacial shear strength, a measure of the maximum shear stress transferable across the fiber–matrix interface during debonding, rises by nearly 245% relative to untreated fibers due to combined chemical bonding, mechanical interlocking, and increased surface roughness.

3.8. Summary of Interfacial and Nanoscale Reinforcement Strategies

The mechanical performance and environmental stability of natural-fiber composites depend fundamentally on the chemistry and morphology of the fiber–matrix interface. Surface-modification strategies that strengthen bonding and reduce interfacial voids greatly improve stress transfer, moisture resistance, and long-term durability. Nanofillers provide a complementary route to these goals: well-dispersed clays, metal oxides, and carbon nanomaterials enhance stiffness, strength, fracture resistance, and barrier properties through microvoid filling, crack bridging, and multiscale interlocking. Together, interfacial engineering and nanoscale reinforcement offer the most effective pathway for overcoming the intrinsic hydrophilicity and structural heterogeneity of natural fibers, enabling composites with greater mechanical reliability and environmental robustness.

4. Moisture Resistance and Transport Behavior in Natural Fiber Composites

Figure 6 summarizes the coupled mechanisms governing moisture transport and environmental durability in natural-fiber composites. The framework links moisture exposure and transport pathways to interfacial degradation and macroscopic property loss, highlighting how fiber chemistry, interface quality, and composite microstructure jointly control long-term performance.

Within this interface-centric framework, moisture transport and environmental degradation arise primarily from coupled effects of fiber surface chemistry, interfacial energetics, and processing-induced interphase structure, rather than from bulk polymer properties alone. Hygrothermal cycling, repeated exposure to alternating temperature and humidity, subjects the fiber–matrix interface to repeated swelling and shrinkage. For lignocellulosic reinforcements, whose cellulose and hemicellulose phases are highly hygroscopic, these dimensional changes generate interfacial shear stresses, microcracking, and progressive debonding, making hygrothermal cycling a particularly sensitive probe of long-term interfacial stability.

Moisture ingress into natural-fiber composites occurs through a combination of molecular diffusion through the polymer matrix and capillary transport along interfacial defects, voids, and microcracks generated by swelling and plasticization. These coupled mechanisms degrade tensile, flexural, and hardness properties through fiber swelling, matrix softening, and interfacial debonding [188]. Depending on fiber volume fraction, porosity, temperature, and interfacial quality, the resulting uptake kinetics may appear approximately Fickian under mild conditions but become increasingly non-Fickian as transport becomes coupled to damage and stress-relaxation processes [189,190]. Surface modification mitigates these effects by improving interfacial bonding, reducing surface polarity, and limiting void formation, thereby restricting moisture pathways and enhancing dimensional stability [78,191].

In PLA-based biocomposites, moisture uptake is governed largely by the hydrophilicity and morphology of the lignocellulosic reinforcement. Alothman et al. [192] reported that PLA composites containing 10 wt% natural fibers absorb 0.4–1.9% water over 3–30 days, accompanied by 2.8–3.4% thickness swelling. SEM analysis revealed that moisture transport is dominated by interfacial voids and localized debonding, with palm-leaf-sheath composites exhibiting the highest void content and the greatest strength degradation after conditioning.

Hygrothermal durability studies of flax-fiber-reinforced epoxies further demonstrate that moisture uptake and mechanical degradation follow distinct kinetics. Although moisture sorption may be Fickian, mechanical performance degrades through irreversible microstructural damage [193,194]. Scida et al. [194] showed that quasi-unidirectional flax–epoxy laminates exhibit up to ≈55% reductions in Young’s modulus under hygrothermal ageing, despite comparatively smaller losses in tensile strength. Acoustic emission and SEM revealed microfibril reorientation, fiber–matrix debonding, matrix microcracking, and fiber pull-out, indicating that long-term degradation is controlled by progressive damage accumulation rather than reversible plasticization alone.

Fully recyclable flax-fiber composites based on bio-derived epoxies and acrylic thermoplastics show similar sensitivity to liquid-water exposure [195]. Under immersion, flax–bio-epoxy laminates absorb ≈ 11 wt% water and suffer severe mechanical losses, including 57% reductions in tensile modulus and 64–70% losses in flexural properties. By contrast, ageing at 75% RH and 45 °C produces only ≈2.6% uptake with negligible mechanical degradation, demonstrating that direct liquid-water ingress, not humid air, dominates durability loss. Hygroscopic swelling under immersion induces interfacial stresses and microcracking that accelerate capillary diffusion, a trend widely observed in flax- and hemp-fiber composites [189,193,194].

At the fiber level, the origins of this moisture sensitivity lie in the hierarchical, multi-wall architecture of lignocellulosic reinforcements. Moisture uptake in plant fibers is governed by three coupled mechanisms—adsorption on hydroxyl sites, capillary filling of lumens and pores, and clustering of water molecules at high humidity—with most of the absorbed water residing in amorphous cellulose, hemicellulose, pectins, and lignin [196]. Hygrothermal exposure, therefore, induces transverse fiber swelling, plasticization of amorphous phases, and progressive decohesion of the cell-wall layers, generating internal stresses that promote microcracking, lumen collapse, and interfacial debonding. Although such ageing can increase the apparent crystallinity of fibers by leaching hemicellulose and other amorphous constituents [196], cyclic hygrothermal loading leads to severe mechanical degradation. For example, Cadu et al. [197] observed a ≈30% reduction in the tensile strength of flax fiber bundles and an ≈40% decrease in Young’s modulus during cyclic exposure at 55 °C and high humidity. These results demonstrate that moisture-induced chemical and microstructural evolution of the fibers directly drives interfacial instability and long-term composite property loss, as depicted in Figure 6.

4.1. Diffusion-Based Descriptions of Moisture Transport

4.1.1. Fickian Analysis of Moisture Uptake

Before introducing the governing diffusion equations, it is useful to clarify their practical interpretation in natural-fiber composites. The apparent diffusion coefficient, $D_{\mathrm{app}}$, characterizes the rate of moisture ingress, while the equilibrium uptake, $M_{\infty }$, reflects the accessible hydrophilic fraction of the composite. Both parameters are meaningful only when transport is diffusion-controlled and the microstructure remains intact. As discussed below, departures from Fickian behavior signal the onset of swelling-induced damage, interfacial debonding, and matrix plasticization, which ultimately govern irreversible hygrothermal ageing and mechanical degradation.

Calabrese et al. [198] quantified moisture uptake in flax fibers, epoxy resin, and flax/epoxy composites using dynamic vapor sorption (DVS) coupled with Fickian diffusion analysis. The moisture content at each humidity step was defined as

$$w={\displaystyle \frac{m^{*}-m_0}{m_0}}$$

(9)

where $m^{*}$ is the equilibrium mass and $m_0$ is the dry mass. Transient uptake was modeled by Fick’s second law,

$${\displaystyle \frac{\mathsf{\partial }C}{\mathsf{\partial }t}}=D{\displaystyle \frac{{\mathsf{\partial }}^{2}C}{\mathsf{\partial }{x}^2}}$$

(10)

with the classical slab solution [199]:

$${\displaystyle \frac{M_t}{M_{\infty }}}=1-{\displaystyle \frac{8}{{\pi }^2}}\sum _{n=0}^{\infty }{\displaystyle \frac{1}{{(2n+1)}^2}} \exp [-{\displaystyle \frac{{(2n+1)}^2{\pi }^{2}Dt}{H^2}}]$$

(11)

For short times ($M_t/M_{\infty }<0.6$),

$${\displaystyle \frac{M_t}{M_{\infty }}}={\displaystyle \frac{4}{H}}\sqrt{{\displaystyle \frac{Dt}{\pi }}}$$

(12)

giving the apparent (1D slab) diffusion coefficient

$$D_{\mathrm{app}}={\displaystyle \frac{{\alpha }^2\pi H^2}{16}}$$

(13)

A geometric correction then yields the specimen-corrected diffusion coefficient

$$D_{\mathrm{real}}={\displaystyle \frac{D_{\mathrm{app}}}{{\left(1+\frac{H}{L}+\frac{H}{W}\right)}^2}}$$

(14)

where L and W denote the in-plane dimensions (length and width) of the specimen, respectively, and H is the through-thickness dimension.

Although such diffusion models accurately describe short-term uptake, long-term hygrothermal ageing is governed by irreversible mechanisms, namely, interfacial debonding, fiber swelling–shrinkage cycling, matrix plasticization, and microcracking, that control mechanical degradation but are not captured by Fickian fits alone [193,194,200,201,202]. Consequently, reliable durability prediction requires combining diffusion modeling with damage-based ageing assessments.

4.1.2. Diffusion Coefficients and the Role of Fiber Content and Orientation

Moisture transport is commonly parameterized by an apparent diffusion coefficient, $D_{\mathrm{app}}$, extracted from Fickian fits, with reported values for plant-fiber composites typically spanning ${10}^{-9}$–${10}^{-12}{\mathrm{m}}^2{\mathrm{s}}^{-1}$ depending on matrix chemistry, fiber type, surface treatment, and architecture. Increasing fiber volume fraction raises both $D_{\mathrm{app}}$ and $M_{\infty }$ by increasing hydrophilic surface area and interfacial connectivity. Muñoz and García-Manrique [203] showed that flax–bio-epoxy laminates with 55 wt% flax exhibit substantially higher diffusivity and equilibrium uptake than 40 wt% laminates: $M_{\infty }$ increased from ≈6.2–6.6% at 40 wt% flax to 8.7–9.8% at 55 wt%, while the apparent diffusion coefficient rose from ≈1.47–1.63 $\times {10}^{-12}$ to ≈1.85–2.32 $\times {10}^{-12}$ m² s⁻¹. These increases reflect swelling-induced microcracking and enhanced capillary transport along the fiber–matrix interface at elevated fiber contents.

Although interfacial regions often serve as preferential moisture pathways in natural-fiber composites, their role depends on whether transport occurs by liquid water or by water vapor. In DVS, vapor uptake is governed by molecular dissolution and diffusion through polymer free volume rather than by capillary flow. Under these conditions, the epoxy matrix acts as an effective permeability barrier that partially shields hydrophilic fiber surfaces, reducing the composite diffusion coefficient relative to neat flax fibers (from $\approx 4.6\times {10}^{-12}$ to $1.8\times {10}^{-12}$ m² s⁻¹ at $P/P_0=0.7$) [198]. Humidity-induced plasticization and water clustering further modulate vapor transport kinetics.

By contrast, in liquid-water immersion experiments such as those of Assarar et al. [193] and Muñoz et al. [203], fiber swelling, interfacial debonding, and microcracking create capillary pathways that markedly accelerate transport, causing both the apparent diffusion coefficient and equilibrium uptake to increase with fiber content. Thus, epoxy “shielding” can suppress vapor diffusion but cannot prevent rapid liquid-water ingress once interfacial damage and capillary connectivity develop.

Fiber orientation introduces true anisotropy in moisture transport: diffusion parallel to the fibers ($D_{\Vert }$) generally exceeds transverse diffusion ($D_{\perp }$) because fiber surfaces and interphases act as preferential transport pathways, whereas fibers themselves act as obstacles in the transverse direction. This behavior is formally described by the three-dimensional Fickian diffusion model for orthotropic media,

$${\displaystyle \frac{\mathsf{\partial }C}{\mathsf{\partial }t}}=D_{11}{\displaystyle \frac{{\mathsf{\partial }}^{2}C}{\mathsf{\partial }{x}^2}}+D_{22}{\displaystyle \frac{{\mathsf{\partial }}^{2}C}{\mathsf{\partial }{y}^2}}+D_{33}{\displaystyle \frac{{\mathsf{\partial }}^{2}C}{\mathsf{\partial }{z}^2}}$$

where $D_{11}$, $D_{22}$, and $D_{33}$ represent the principal diffusion coefficients along the fiber, transverse, and through-thickness directions, respectively.

Zhang et al. [204] quantified this anisotropy using a fiber-scale finite-element homogenization framework that yields an effective orthotropic diffusion tensor from a heterogeneous interfacial-phase model. They constructed a representative volume element (RVE) containing fibers, bulk epoxy, and a distinct interfacial phase whose diffusivity varies from the fiber surface to the bulk matrix. Directional moisture-uptake experiments were performed with diffusion constrained either parallel or perpendicular to the fiber axis, and the interfacial diffusion parameters in the RVE were adjusted until the simulated uptake curves matched the experimental data. From the resulting moisture flux and concentration gradients, the homogenized tensor components $D_{11}$ and $D_{22}$ were extracted, yielding $D_{\Vert }=D_{11}$ and $D_{\perp }=D_{22}$. These quantities represent intrinsic microstructural transport properties and are distinct from the apparent slab diffusivity obtained from one-dimensional gravimetric tests.

Their analysis showed that $D_{\Vert }>D_{\mathrm{bulk}}>D_{\perp }$ in glass/epoxy composites when interfacial microcracking is taken into account, reflecting the competing effects of fast interfacial transport channels and transverse fiber obstruction. Here, $D_{\mathrm{bulk}}$ denotes the moisture diffusion coefficient of the neat polymer matrix, measured independently from resin-only specimens and used as the reference diffusivity in the fiber-scale homogenization model. At the laminate scale, these directional effects propagate upward: through-thickness diffusivities $D_{33}$ vary by orders of magnitude across flax-, glass-, and carbon-fiber epoxies and between chopped-strand and roving layers [205], mirroring the underlying fiber-scale anisotropy.

The diffusion coefficients summarized in

Table 12. Representative moisture diffusion coefficients illustrating the effects of fiber content, matrix shielding, and anisotropy in fiber-reinforced composites.

Matrix/Fiber	Fiber vol% (wt%)	Measurement	Diffusion Coefficient (m2 s−1)	Ref.
None/flax bundle	100 (100)	DVS, $P/P_0=0.7$, 30 °C	$D\approx 4.6\times {10}^{-12}$	[198]
Epoxy/flax	n.r.	DVS, $P/P_0=0.7$, 30 °C	$D_{33}\approx 1.8\times {10}^{-12}$	[198]
Bio-epoxy/flax	≈35 (40)	Through-thickness, water immersion, 23 °C	$D_{33}=(1.55\pm 0.08)\times {10}^{-12}$	[203]
Bio-epoxy/flax	≈50 (55)	Through-thickness, water immersion, 23 °C	$D_{33}=(2.09\pm 0.24)\times {10}^{-12}$	[203]
Epoxy/glass (UD rod)	80–85 (≈76–82)	Fiber-parallel diffusion ($D_{\Vert }=D_{11}$)	$D_{11}=1.463\times {10}^{-12}$	[204]
Epoxy/glass (UD rod)	80–85 (≈76–82)	Fiber-transverse diffusion ($D_{\perp }=D_{22}$)	$D_{22}=0.649\times {10}^{-12}$	[204]
Epoxy/none (neat resin)	0 (0)	Bulk resin reference ($D_{\mathrm{bulk}}$)	$D_{\mathrm{bulk}}=0.827\times {10}^{-12}$	[204]
Epoxy/flax (UD laminate)	51 (≈56.5)	Through-thickness diffusion ($D_{33}$)	$D_{33}=1.05\times {10}^{-12}$	[193]
Epoxy/glass (UD laminate)	33 (≈51)	Through-thickness diffusion ($D_{33}$)	$D_{33}=1.38\times {10}^{-13}$	[207]
Epoxy/carbon (UD laminate)	69 (≈77)	Through-thickness diffusion ($D_{33}$)	$D_{33}=7.7\times {10}^{-15}$	[206]
Epoxy/none (neat resin)	0 (0)	Through-thickness diffusion (neat resin)	$D_{33}=1.41\times {10}^{-13}$	[207]
Polyester (EL-400)/glass CSM layer	n.r.	Through-thickness layer ($D_{33}$)	$D_{33}=9.52\times {10}^{-13}$	[205]
Polyester (EL-400)/glass roving (UD layer)	n.r.	Through-thickness layer ($D_{33}$)	$D_{33}=2.72\times {10}^{-14}$	[205]

For [204], the pultruded GFRP core rod contained 80–85 vol% glass fiber and 15–20 vol% epoxy resin; the corresponding fiber content (76–82 wt%) was estimated using the reported constituent densities. For the unidirectional (UD) laminates in [193,206,207], fiber weight fraction values were calculated from the reported volume fraction, $v_{\mathrm{f}}$ using $w_{\mathrm{f}}={\displaystyle \frac{v_{\mathrm{f}}{\rho }_{\mathrm{f}}}{v_{\mathrm{f}}{\rho }_{\mathrm{f}}+(1-v_{\mathrm{f}}){\rho }_{\mathrm{m}}}}$ with ${\rho }_{\mathrm{m}}=1.20$ g cm⁻³ (epoxy) and ${\rho }_{\mathrm{f}}=1.50$ g cm⁻³ (flax), 2.55 g cm⁻³ (glass), and 1.80 g cm⁻³ (carbon). Ref. [205] did not report fiber volume or weight fractions for the CSM and roving layers; only layer densities and diffusion coefficients were measured. n.r. = not reported.

quantitatively support the framework in Figure 6, illustrating how fiber content, matrix shielding, and interfacial architecture govern moisture transport pathways and thereby control the progression from ingress to microstructural damage and property degradation. Although not all systems are natural-fiber composites, the included data provide useful benchmarks for comparing trends across reinforcement classes. Notably, water diffusion coefficients in natural-fiber–reinforced composites are at least two orders of magnitude higher than those in carbon-fiber–reinforced systems [206], reflecting the hydrophilic nature and high sorption capacity of plant fibers.

4.2. Multi-Year Hygrothermal Ageing and Lifetime Prediction

Long-duration hygrothermal durability data for plant-fiber composites extending beyond one year remain comparatively scarce, largely because of the time and cost associated with controlled environmental exposures. Nevertheless, a limited number of studies do extend into the multi-year regime under moisture–thermal environments relevant to service conditions and provide important insight into both long-term degradation behavior and approaches for lifetime prediction.

Le Duigou et al. [208] reported a two-year natural seawater immersion study of injection-molded flax/PLA biocomposites (20 wt% flax; ≈16 vol%), conducted at 5 m depth in Kernevel harbor (Lorient, France), where water temperature typically varies between 8 and 19 °C. Specimens were periodically retrieved for gravimetric water-uptake measurements and mechanical testing. Cyclic load–unload tensile tests were employed to define a stiffness-based damage parameter, $d=1-E/E_0$, where E is the Young’s modulus measured for a given exposure state, and $E_0$ is the initial modulus of the unaged composite. Wet versus dried characterization enabled the separation of reversible matrix plasticization from irreversible mechanical damage. Fracture surfaces were examined by SEM. Although moisture uptake reached near-saturation within the first months of immersion, stiffness-based damage continued to accumulate over extended exposure times. The combined use of cyclic mechanical testing and wet/dry comparisons demonstrated that a portion of stiffness loss was reversible, while irreversible damage increased progressively with ageing. SEM fractography revealed increased fiber pull-out and interfacial decohesion with immersion time, consistent with progressive moisture-induced weakening of the fiber–matrix interface.

In contrast to the sustained saturation regime examined by Le Duigou et al. [208], Cadu et al. [197] showed that cyclic hygrothermal exposure produces pronounced and largely irreversible mechanical degradation at the flax fiber-bundle level. Repeated moisture–temperature cycling led to substantial reductions in tensile strength and Young’s modulus over relatively short exposure times. Although they investigated fiber bundles rather than composite materials, their results provide a mechanistic reference for intrinsic, cell-wall-scale microstructural degradation induced by cyclic swelling–shrinkage, associated with hemicellulose loss and reduced load transfer between cellulose microfibrils. By comparison, Le Duigou and coworkers demonstrated how moisture-driven effects manifest at the composite scale under continuous saturation, where reversible matrix plasticization and progressive fiber–matrix interfacial degradation govern the long-term mechanical response.

Complementary long-horizon accelerated ageing programs have been developed within the civil and infrastructure durability literature. Chlela et al. [209] presented an experimental ageing framework for unidirectional flax/bio-epoxy flax fiber–reinforced polymer (FFRP) laminates and FFRP-strengthened concrete blocks, spanning climatic-chamber exposures (20–60 °C, 50–75% RH) and fully wet immersion conditions (100% RH). Periodic tensile, short-beam (interlaminar shear), and pull-off tests were conducted at 3, 6, and 12 months, with an additional 12-month outdoor exposure condition [209]. The results showed that immersion in water was the dominant degradation driver, leading to pronounced reductions in tensile stiffness and adhesive bond strength due to matrix plasticization and interfacial weakening, whereas non-immersed elevated-temperature conditions primarily induced post-curing and did not cause mechanical deterioration. Bond strength at the composite–concrete interface was identified as the most critical durability-limiting property, motivating the development of a coupled temperature–humidity performance evolution model for lifetime prediction.

Building on this experimental framework, Yan et al. [210] reported a two-year dataset for FFRP laminates, based on destructive tensile tests conducted at $t_0$, 3, 6, 12, and 24 months under six combined temperature–humidity $(T,\mathrm{RH})$ environments. To interpret these sparse but long-duration data, the authors developed a probabilistic lifetime-prediction framework specifically adapted to destructive testing, with the objective of quantifying hygrothermal durability and extrapolating service performance.

A key experimental observation was the presence of non-monotonic behavior at early exposure times, attributed to post-curing and transient property recovery, followed by a well-defined monotonic ageing regime. Consequently, only the latter stage was considered in the degradation modeling. The measured tensile capacity was expressed as the difference between an effective initial performance and an accumulated degradation variable,

$$y_k\left(t\right)={\xi }_k-x(t;T_k,{\mathrm{RH}}_k)$$

(15)

where ${\xi }_k$ denotes the effective performance at the onset of monotonic ageing under the k-th environmental condition. Here, tensile capacity refers to the maximum tensile force sustained at failure, normalized by specimen width (and number of layers), as commonly adopted in FRP structural design to minimize uncertainties associated with laminate thickness.

The mean evolution of degradation was described using a shifted power-law dependence,

$$\mathbb{E}\left[x\left(t\right)\right]=\mu (T,\mathrm{RH}){(t-b_k)}^q,t\ge b_k$$

(16)

where the environment-dependent time shift $b_k$ accounts for post-curing effects and q controls the nonlinearity of the ageing kinetics. This representation captures the experimentally observed degradation trends while providing a flexible empirical description of long-term performance loss.

Environmental acceleration of degradation was introduced through a generalized Eyring model,

$$\mu (T,\mathrm{RH})=a\exp (-{\displaystyle \frac{c_1}{T}}+c_2\mathrm{RH}+{\displaystyle \frac{c_3\mathrm{RH}}{T}})$$

(17)

where T is the absolute temperature and $\mathrm{RH}$ is the relative humidity. The formulation explicitly incorporates thermally activated degradation, direct moisture effects, and a pronounced temperature–humidity coupling, which was shown to be essential for reproducing the observed acceleration of degradation under combined hygrothermal exposure.

Beyond mean trends, Yan et al. demonstrated that the scatter in tensile capacity increases with ageing time. This behavior was incorporated by allowing the variability of degradation to increase with accumulated damage,

$$\mathrm{Var}\left[x\left(t\right)\right]\propto {(t-b_k)}^q$$

(18)

reflecting the growing uncertainty associated with aged natural-fiber composites.

Material lifetime was defined as the earliest time at which the accumulated degradation reaches a prescribed threshold $\omega$ (i.e., the first time $x\left(t\right)\ge \omega$), with $\omega$ selected based on reduction factors from existing FRP design codes. This probabilistic definition allows lifetime to be expressed in terms of the likelihood of failure over time, rather than as a single fixed failure time.

Yan et al. [210] showed that hygrothermal ageing of flax fiber composites is governed by strongly coupled temperature–humidity effects, non-monotonic early behavior associated with post-curing, and increasing variability with ageing time. Their results highlight the limitations of deterministic or separable Arrhenius-type models and demonstrate that existing fiber-reinforced polymer durability models and design-code provisions, largely calibrated for carbon- and glass-fiber systems, are not directly applicable to natural-fiber composites without careful reassessment.

4.3. Fiber–Water and Fiber–Polymer Interfacial Energy Analysis and Correlation to Adhesion

Thermodynamic incompatibility between the fiber and the polymer matrix can promote the formation of interfacial voids, which act as pathways and reservoirs for water ingress in natural fiber–polymer composites (Figure 7). The relative tendency of water versus polymer to wet the fiber surface can be assessed by comparing the corresponding interfacial energies. Within the Owens–Wendt framework, the surface free energy of each phase is expressed as the sum of its polar and dispersive components,

$${\gamma }_i={\gamma }_i^{\mathrm{p}}+{\gamma }_i^{\mathrm{d}}$$

(19)

where ${\gamma }_i^{\mathrm{p}}$ and ${\gamma }_i^{\mathrm{d}}$ denote the polar and dispersive components, respectively, for phase i (fiber “f” or matrix “m”).

The fiber–matrix interfacial energy is estimated using the geometric–mean rule,

$${\gamma }_{\mathrm{f}/\mathrm{m}}={\gamma }_{\mathrm{f}}+{\gamma }_{\mathrm{m}}-2(\sqrt{{\gamma }_{\mathrm{f}}^{\mathrm{p}}{\gamma }_{\mathrm{m}}^{\mathrm{p}}}+\sqrt{{\gamma }_{\mathrm{f}}^{\mathrm{d}}{\gamma }_{\mathrm{m}}^{\mathrm{d}}})$$

(20)

where lower ${\gamma }_{\mathrm{f}}/\mathrm{m}$ generally corresponds to improved wetting and stronger interfacial adhesion. The spreading parameter, which characterizes the thermodynamic driving force for matrix spreading on the surface, is given by

$$S={\gamma }_{\mathrm{f}}-\left({\gamma }_{\mathrm{m}}+{\gamma }_{\mathrm{f}}/\mathrm{m}\right)$$

(21)

where $S>0$ indicates spontaneous spreading and $S<0$ indicates partial wetting. Less negative values of S correspond to more favorable wetting and stronger potential adhesion.

Table 13. Surface-energy components (mJ m⁻²) of jute, jute–SiO₂, PP, and water (${\gamma }_i={\gamma }_i^{\mathrm{p}}+{\gamma }_i^{\mathrm{d}}$).

Phase	${\mathit{\gamma }}_{\mathit{i}}^{\mathbf{p}}$	${\mathit{\gamma }}_{\mathit{i}}^{\mathbf{d}}$	${\mathit{\gamma }}_{\mathit{i}}$
Jute	30.0	4.4	34.4
Jute–SiO2	43.4	6.6	50.0
PP	1.0	29.9	30.9
Water	51.0	21.8	72.8

lists the polar and dispersive components of the surface energies for jute, SiO₂-functionalized jute, and PP as reported by Liu et al. [211], together with standard reference values for water. These values were used to compute the interfacial energies and spreading parameters from Equations (20) and (21), summarized in

Table 14. Calculated interfacial energies ${\gamma }_{\mathrm{f}/\mathrm{m}}$ and spreading parameters S for jute and jute–SiO₂ in contact with PP or water (units: mJ m⁻²).

Interface	${\mathit{\gamma }}_{\mathbf{f}/\mathbf{m}}$	S
Jute/PP	31.4	$-27.9$
Jute/Water	9.4	$-47.8$
Jute–SiO2/PP	39.6	$-20.5$
Jute–SiO2/Water	4.7	$-27.5$

.

The results indicate that unmodified jute has a strong thermodynamic affinity for water, as reflected by the low jute/water interfacial energy (${\gamma }_{\mathrm{f}/\mathrm{m}}\approx 9.4$ mJ m⁻²; Table 14). However, the highly negative spreading parameter for water on jute ($S\approx -47.8$ mJ m⁻²) shows that water does not spontaneously form a continuous film on the bare fiber surface in air, but instead exhibits partial wetting with a finite contact angle. Thus, although water is strongly attracted to jute, complete thermodynamic spreading is not favored. Moisture ingress in jute/PP composites is therefore expected to proceed primarily through diffusion and capillary penetration into interfacial defects and microvoids, rather than by spontaneous interfacial wetting alone.

SiO₂ coating substantially increases the fiber surface energy (from 34.4 to 50.0 mJ m⁻²), mainly through an increase in the polar component (Table 13). Consequently, the jute–SiO₂/water interfacial energy decreases (${\gamma }_{\mathrm{f}/\mathrm{m}}=4.7$ mJ m⁻²), indicating a stronger affinity for water. In contrast, at the fiber/PP interface the spreading parameter becomes markedly less negative after modification ($S=-20.5$ versus $-27.9$ mJ m⁻²), implying a reduced penalty for PP spreading on the modified fiber surface and thus improved polymer wetting.

Despite the higher calculated jute–SiO₂/PP interfacial energy relative to jute/PP (Table 14), the concurrent increase in fiber surface energy ${\gamma }_{\mathrm{f}}$ reduces the driving force opposing PP spreading. From the standpoint of spreading energetics, SiO₂ modification therefore promotes improved PP wetting even though the nominal interfacial energy is higher.

Liu and coworkers employed Equation (21) for their thermodynamic analysis. However, moisture propagation along a buried fiber/PP interface is more appropriately assessed using the three-phase (displacement) spreading parameter

$$S_{\mathrm{w}\mathrm{at}(\mathrm{f}/\mathrm{PP})}={\gamma }_{\mathrm{f}/\mathrm{PP}}-\left({\gamma }_{\mathrm{f}/\mathrm{w}}+{\gamma }_{\mathrm{w}/\mathrm{PP}}\right)$$

(22)

which quantifies the tendency of water to displace PP from the fiber surface. Using the interfacial energies in Table 14 and the water/PP interfacial energy from Table 13 (${\gamma }_{\mathrm{w}/\mathrm{PP}}\approx 38.4$ mJ m⁻²), one obtains $S_{\mathrm{w}\mathrm{at}(\mathrm{jute}/\mathrm{PP})}\approx -16.3$ mJ m⁻² and $S_{\mathrm{w}\mathrm{at}(\mathrm{jute}–{\mathrm{SiO}}_2/\mathrm{PP})}\approx -3.4$ mJ m⁻². Both values remain negative, indicating that complete interfacial spreading of water is not favored. Notably, the substantially less negative value after SiO₂ modification shows that PP displacement by water becomes less unfavorable at the modified interface, in contrast to the two-phase analysis, which predicts improved PP wetting.

These apparently opposing trends arise because the two-phase spreading parameter [Equation (21)] evaluates PP wetting on the fiber surface in air, whereas $S_{\mathrm{w}\mathrm{at}(\mathrm{f}/\mathrm{PP})}$ describes the tendency of water to displace PP at a buried interface. SiO₂ functionalization therefore improves initial PP wetting while simultaneously lowering the penalty for water displacement if moisture reaches the interface. The experimentally observed improvement in water resistance thus arises primarily from microstructural effects rather than surface thermodynamics. SEM observations by Liu et al. [211] show that SiO₂ nanoparticles fill surface grooves (Figure 7) and eliminate interfacial microvoids, increasing real contact area. This microstructural sealing restricts moisture access and suppresses capillary transport pathways, outweighing the mixed energetic trends.

5. Processing Techniques for Natural-Fiber–Reinforced Biocomposites

5.1. Process–Microstructure–Performance Framework

Figure 8 illustrates how different processing routes impose distinct thermal, mechanical, and moisture histories on natural fiber composites, which in turn govern fiber morphology, dispersion, porosity, and interphase development. These microstructural features control interfacial quality and ultimately determine mechanical performance, anisotropy, and environmental durability.

Lignocellulosic fibers are both thermally and chemically sensitive: exposure to elevated temperatures and long residence times induces hemicellulose degradation, cellulose depolymerization, and progressive loss of fiber integrity. Hemicellulose and cellulose exhibit peak thermal-decomposition rates at ≈280 °C and 344 °C, respectively, with the onset of significant mass loss occurring at temperatures as low as ≈230 °C [212]. Consequently, composite processing temperatures are generally minimized and, where possible, maintained below about 230 °C [213,214].

In practical manufacturing, melt-based processes such as extrusion, injection molding, and fused-filament additive manufacturing impose the most severe thermal and residence-time histories (typically 170–230 °C, from seconds in injection molding to minutes in extrusion and melt compounding), accelerating chemical degradation of the fiber cell wall. In contrast, liquid composite molding routes such as resin transfer molding (RTM) and vacuum-assisted RTM (VARTM) operate at substantially lower temperatures and better preserve fiber morphology and interfacial chemistry.

Processing selection, therefore, exerts a first-order influence on stiffness, strength, dimensional stability, and moisture resistance in natural-fiber biocomposites. Process-dependent shear and flow histories further control fiber breakage, deagglomeration, orientation, and matrix wetting, directly governing composite stiffness, toughness, and moisture-dependent performance [37,215].

5.2. Classification of Processing Routes

Melt-based routes such as extrusion, injection molding, compression molding, and thermoforming dominate thermoplastic natural-fiber composites because they are compatible with existing industrial infrastructure and provide high throughput and geometric flexibility [216,217]. By contrast, liquid composite molding routes, including RTM and continuous pultrusion, are mainly applied to thermosetting and reactive matrices reinforced with continuous or woven natural fibers, enabling higher fiber volume fractions, improved impregnation control, and structural laminate architectures [215,218,219,220]. Solution casting remains a low-shear, solvent-based method for films and coatings, whereas fused-filament additive manufacturing enables complex geometries but introduces layer-wise anisotropy, porosity, and moisture-sensitive interlayer bonding [221].

5.3. Solvent-Based Processing

Solution casting involves dissolving the polymer matrix in a suitable solvent, dispersing natural-fiber fillers, casting the suspension, and removing the solvent by evaporation or drying to form a film [222]. Because processing occurs in solution rather than under melt-compounding conditions, shear-induced fiber damage is minimized and fiber length and morphology are largely preserved. The method is therefore widely used for thin films and coatings and for laboratory-scale studies examining the effects of filler loading, particle size, and surface treatment [9].

Table 15. Solution casting—operating window, structure–property trends, and applications.

Processing Strategy	Matrix/Fiber	Processing Conditions	Structure–Processing–Property Response	Typical Applications	Ref.
Solution casting of PLA composite films reinforced with milled flax or hemp fillers.	PLA/flax or hemp (particle sizes < 75 $\mathsf{\mu }$m and 149–210 $\mathsf{\mu }$m; loadings 2.5–30 wt%).	Ambient drying after casting.	Low filler loadings (2.5–5 wt%) gave the most favorable balance of mechanical and barrier properties; higher loadings increased moisture sensitivity and brittleness.	Sustainable packaging films and valorization of agricultural by-products.	[223]
Solvent casting of antimicrobial PHBV films incorporating ethyl lauroyl arginate (LAE) to control spoilage bacteria after package opening.	PHBV/LAE (5 wt% LAE, based on polymer).	Solvent casting from chloroform or formic acid; film thickness ≈ 35–42 $\mathsf{\mu }$m.	LAE-containing films showed strong inhibition of Listeria monocytogenes and Pseudomonas sp. in broth and in an almond beverage system; antimicrobial performance depended on release time and food–polymer interactions.	Active biodegradable packaging for extension of secondary shelf life of plant-based beverages.	[224]

presents representative examples.

Pokharel et al. [223] prepared PLA/flax and PLA/hemp composite films by solution casting. PLA was first dissolved in chloroform at a polymer–solvent ratio of 1:12 and mixed using a magnetic stirrer; milled biomass fillers were then added at 2.5, 5, 10, 20, and 30 wt% (based on PLA), and the suspension was stirred for 48 h. The biomass was knife-milled and sieved into two particle-size fractions (<75 $\mathsf{\mu }$m and 149–210 $\mathsf{\mu }$m) and surface-modified by alkali and acetylation treatments prior to blending. The resulting film-forming solution was cast onto glass plates and dried under ambient conditions; neat PLA films were prepared following the same protocol.

Mechanical performance depended strongly on filler loading, surface treatment, and particle size. Alkali-treated PLA/flax composites exhibited significant increases in tensile strength, elongation at break, and Young’s modulus at low filler contents (2.5–5 wt%), whereas increasing the filler loading to 10–30 wt% led to reduced elongation at break and increased brittleness without proportional gains in tensile strength. Moisture absorption increased with filler content, and water-vapor permeability likewise rose at higher loadings, with the highest values reported at 30 wt% fiber content. Overall, PLA-based composite films containing 5 wt% flax or hemp provided the most favorable balance between mechanical performance and barrier properties [223].

Nicosia et al. [224] developed solvent-cast active films based on PHBV incorporating ethyl lauroyl arginate (LAE) as an antimicrobial agent to inhibit spoilage microorganisms after package opening. LAE was selected in preference to nisin because of its broader activity against mixed spoilage cultures and its stability during solvent-based processing. The minimum bactericidal concentration of LAE was determined against Listeria monocytogenes and Pseudomonas sp., and PHBV–LAE films exhibited pronounced antimicrobial activity in both broth and a real food system (almond beverage). Reductions of up to 7.6 log units for Pseudomonas sp. and complete inhibition after sufficient release time were reported, indicating that antimicrobial effectiveness is governed by additive release kinetics and food–polymer–additive interactions rather than by the polymer matrix alone. These results demonstrate the potential of solvent-cast PHBV films as biodegradable active-packaging materials for extending secondary shelf life and mitigating post-opening food spoilage.

Owing to long drying times, solvent recovery requirements, and limited throughput, solution casting is primarily applied to thin coatings, biodegradable packaging films, and controlled laboratory-scale comparisons rather than high-volume structural components [9,222].

5.4. Melt Processing of Thermoplastic Natural-Fiber Composites

Twin-screw extrusion is a widely used route for processing thermoplastic natural-fiber composites because it enables continuous melt processing, simultaneous mixing and shaping, and compatibility with pelletized feedstocks [216,217]. Representative examples are summarized in

Table 16. Extrusion and melt compounding of NFRCs: operating windows and structure–processing–property relationships.

Processing Strategy	Matrix/Fiber	Processing Conditions	Reported Property Trends	Ref.
Twin-screw compounding analyzed by thermomechanical extrusion modeling to define a feasible processing window for PP/flax composites.	PP/PP-g-MA + flax (20 wt%).	Barrel setpoint ≈ 180 °C; constrained to $T_{max}<200$ °C.	Optimized condition ($Q=3.6$ kg h−1, $N=95$ rpm) yields a predicted die-exit flax fiber length of 1.46 mm (from ≈4 mm initial) and diameter of ∼100 $\mathsf{\mu }$m.	[225]
Single- and twin-screw extrusion compared for PHBV biocomposites to quantify the effect of processing route on fiber dimensions and mechanical response.	PHBV/hemp or flax (15 wt%).	Twin-screw extrusion with barrel zones at 145–160 °C and feed zone at 50 °C; single-screw extrusion operated with a head temperature of 170 °C at 100 rpm.	Single-screw processing yielded higher stiffness (31–36%) and tensile strength (2.5–9%) than twin-screw processing.	[226]
Twin-screw extrusion used to produce recycled polypropylene/wood-dust filaments for evaluating the influence of surface treatments on filament properties.	Recycled PP/wood dust (3 wt%).	Extruder barrel temperatures of 160–200 °C and nozzle temperature of 180–190 °C; filament diameter ≈ 1.75 mm.	Silane-treated formulations exhibited the highest wire-pull strength (35.2% above untreated) and lower water uptake.	[227]
Two-step pellet extrusion followed by single-screw film extrusion used to fabricate PHBV-based blend films for transport-property evaluation.	PHBV/PHB (20 wt%)/BC-TOCN-AMD C-18 (5 wt%).	Four-zone extrusion profile: 170/185/200/210 °C.	Ternary PHBV/PHB/BC-TOCN-AMD C-18 films exhibited lower water-vapor and oxygen permeability (10.47% and 9.54% below neat PHBV).	[228]

. During extrusion, material fed through the hopper is melted, conveyed, and mixed along the screw channel before being forced through a die to form the desired profile. The intense shear and extensional flows generated within the screw flights and kneading elements impose repeated mechanical deformation on the fibers, so that fiber dimensions and morphology evolve progressively with processing history rather than remaining fixed.

Berzin et al. [225] used thermomechanical flow modeling to quantify how twin-screw compounding governs fiber attrition in polypropylene/flax composites. Their one-dimensional extrusion simulations (Ludovic^®), coupled with experimentally calibrated, strain-based fiber-breakage laws, demonstrate that fiber dimensions are controlled by the cumulative deformation history imposed by screw configuration and operating conditions. For an initial flax fiber length of ≈4 mm, the optimized processing window ($Q=3.6$ kg h⁻¹, $N=95$ rpm) yields a predicted die-exit length of 1.46 mm and a diameter of ∼100 $\mathsf{\mu }$m while maintaining $T_{max}<200$ °C to avoid thermal degradation. Screw speed and throughput regulate the residence-time distribution and the cumulated strain (scaling with $N/Q$), thereby defining a fundamental trade-off between distributive mixing and fiber preservation. This analysis establishes extrusion as a mechanically driven microstructural filter in which processing history, rather than initial fiber geometry alone, sets the attainable fiber length scale in thermoplastic natural-fiber composites.

Extrusion configuration and interfacial chemistry further modulate composite structure and mechanical performance. Janowski et al. [226] compared single- and twin-screw extrusion of PHBV biocomposites reinforced with hemp or flax fibers and showed that the single-screw route preserved larger fiber dimensions, resulting in substantial increases in elastic modulus (by ≈31–36%) and modest gains in tensile strength (by ≈2.5–9%), accompanied by reduced ductility. By contrast, twin-screw extrusion produced shorter fibers but yielded lower porosity and lower water absorption, which the authors attributed to enhanced degassing and melt homogenization. Nafis et al. [227] demonstrated that, in recycled polypropylene/wood-dust filaments compounded by twin-screw extrusion, silane surface treatment increased wire-pull strength by 35.2% relative to untreated formulations and reduced water uptake. SEM observations in that study revealed fewer interfacial voids and improved fiber–matrix adhesion after treatment, directly linking filament strength and moisture resistance to interfacial quality.

For polyhydroxyalkanoate-based systems, Gamboa-Suárez et al. [228] investigated film extrusion of PHBV blends containing low-molecular-weight PHB and hydrophobically modified bacterial cellulose nanofibers (BC-TOCN-AMD C-18). Films were produced by a two-step process involving pellet extrusion followed by single-screw film extrusion. PHB addition (20 wt%) slightly reduced water-vapor transmission relative to neat PHBV by altering crystallization behavior and reducing surface porosity. In contrast, incorporation of BC-TOCN-AMD C-18 alone increased water and oxygen permeability due to heterogeneous morphology and aggregate formation. When PHB and BC-TOCN-AMD C-18 were combined, dispersion of the nanocellulose improved and barrier performance was enhanced, with water-vapor and oxygen permeability reduced by 10.47% and 9.54%, respectively, relative to neat PHBV. These results indicate that transport behavior in extruded PHBV films is governed primarily by blend chemistry and microstructural homogeneity rather than by processing-induced fiber damage.

These studies show that extrusion primarily governs fiber attrition, dispersion, and interfacial quality during feedstock generation, thereby defining the initial microstructural state entering downstream shaping operations. Injection molding then acts on this preconditioned microstructure and can further modify it through additional shear, extensional flow, and mold-filling dynamics. Although most injection-molding studies focus on glass- or carbon-fiber thermoplastics, they provide transferable insights into processing–structure relationships that are directly relevant to natural-fiber composites. Specifically, injection molding introduces an additional fiber-attrition step beyond compounding and generates strong flow-induced fiber-orientation gradients through the part thickness, which lead to localized stress concentrations and govern where failure initiates under load [229,230]. Process variables including mold orientation, flow direction, and local wall thickness therefore act as effective levers for homogenizing fiber alignment and reducing peak stresses [230]. Because natural fibers are generally more susceptible to shear-induced damage and orientation effects than synthetic fibers, these principles are expected to be even more critical in natural-fiber composite moldings.

5.5. Continuous Fiber Liquid Molding

5.5.1. Pultrusion

Pultrusion is a continuous liquid-composite molding process for manufacturing fiber-reinforced profiles with constant cross section and high fiber volume fraction. In a typical pultrusion process, continuous reinforcements (e.g., rovings, stitched fabrics, or mats) are drawn from a creel through a resin-impregnation stage and subsequently through a heated die, where curing under controlled temperature and pulling speed consolidates the composite into a rigid profile that is continuously extracted [218,219]. Impregnation may be achieved using an open resin bath or a closed injection chamber, followed by guiding and preforming elements that promote wet-out, remove excess resin, and stabilize fiber alignment prior to die entry [218].

Pultrusion integrates impregnation, shaping, and curing into a single continuous operation, enabling high fiber volume fractions and excellent dimensional repeatability. Reviews report that fiber contents exceeding 50–70 vol% are routinely achievable, which underpins the high stiffness and strength of pultruded profiles [218,219]. Process quality is governed by coupled parameters including resin viscosity, fiber fraction, die temperature and curing kinetics, and pulling speed, which together determine impregnation quality, void content, and degree of cure [218]. Continuous fiber tension and steady pulling promote uniform resin distribution and reproducible laminate architecture over long production lengths [219].

Relative to melt extrusion of short-fiber thermoplastics, pultrusion preserves fiber continuity and produces strongly anisotropic, load-bearing architectures. Typical profiles such as plates, beams, rods, channels, and gratings employ unidirectional rovings to carry axial loads, while surface mats or veils enhance impregnation permeability, surface finish, and transverse integrity [218,219]. The resulting combination of high fiber volume fraction and continuous reinforcement distinguishes pultrusion as a primary route for structural composite components [219,220].

From a sustainability perspective, pultrusion offers low scrap rates and efficient material utilization, and it is compatible with bio-based fibers and resins provided that viscosity, wetting, and curing requirements are satisfied [219]. These attributes make pultrusion a promising manufacturing route for long, repeatable structural elements in applications demanding high stiffness, low weight, and controlled architecture.

Recent studies demonstrate that pultrusion is increasingly viable for natural-fiber composites incorporating bio-derived resin systems (

Table 17. Pultrusion of NFRCs: processing constraints and structure–processing–property responses.

Processing Strategy	Matrix/Fiber	Processing Conditions	Reported Property Trends	Ref.
Moisture diffusion and durability of pultruded multilayer composite plates, linking laminate architecture to transport behavior.	Unsaturated polyester reinforced with E-glass fibers; inner unidirectional roving core with chopped-strand-mat surface layers.	Water immersion with diffusion analysis using combined experiments and finite-element modeling.	Outer chopped-strand layers exhibited higher effective diffusivity than the roving core; overall moisture uptake governed by layer sequence and thickness rather than bulk-average properties.	[205]
Laboratory-scale pultrusion of solid natural-fiber round profiles to assess stiffness and elastic bendability as functions of fiber form and processing conditions.	Hemp bast fibers (kemafiled strips, 3800–4500 tex) in a bio-based thermosetting resin system (with mineral filler/flame retardant); fiber volume fraction ≈ 32%.	Pultrusion; solid circular profile, 6 mm diameter (continuous length).	Density ≈ 1.2 g cm−3; flexural strength 270 MPa; bending modulus 21.5 GPa; elastic bending to radii ≈ 0.5 m demonstrated.	[231]
Industrial-scale evaluation of pultruded natural-fiber biocomposite profiles for structural applications, including mechanical performance and moisture-durability screening.	Natural-fiber-reinforced bio-based resin system; round hollow profile geometry optimized for axial and bending loads.	Pultrusion; round hollow profile, 25 mm outer diameter, 4 mm wall thickness (continuous length).	Compression strength 31.2 kN; compression modulus 118 MPa; flexural strength 300 MPa; bending modulus 30 GPa; minimum bending radius ≈ 2.4 m. Water immersion (up to 4 months): ∼2% mass gain for clean-cut surfaces vs. up to ∼60% for scratched specimens, indicating strong sensitivity to surface integrity.	[232]

). Pultruded natural-fiber profiles combine high stiffness with substantial elastic curvature capacity, enabling both load-bearing and bending-active structural applications. Laboratory-scale demonstrations show that continuous hemp-fiber reinforcements can be consolidated into mechanically efficient solid profiles with pronounced elastic bendability, while industrial-scale hollow profiles retain comparable flexural performance together with axial load-bearing capability [231,232]. These results indicate that pultrusion can integrate stiffness-dominated load transfer with geometric flexibility in continuous biocomposite members.

Durability is governed primarily by laminate architecture and surface condition. Multilayer pultruded plates with unidirectional roving cores and chopped-strand surface layers exhibit strong through-thickness gradients in moisture transport, with diffusion behavior controlled by layer sequence and thickness rather than bulk-average properties [205]. Complementary immersion tests on natural-fiber pultruded profiles further show that surface integrity is critical: moisture uptake remains limited for intact cut edges but increases dramatically when surfaces are damaged, accompanied by visible fiber exposure and degradation [232]. Together, these findings identify laminate design and cut-edge protection as key durability controls for pultruded natural-fiber composites intended for structural and architectural deployment.

5.5.2. Resin Transfer Molding (RTM)

RTM is a closed-mold liquid composite molding process in which a low-viscosity resin is injected under pressure into a stationary porous natural-fiber preform placed within a matched mold. Impregnation quality is governed by tightly coupled flow and cure phenomena, controlled by injection pressure, mold temperature, resin viscosity, preform permeability and architecture, and gate–vent design, as established in prior analyses of RTM process physics [215]. As shown in

Table 18. Resin transfer molding and RTM-derived processing routes for NFRCs: impregnation mechanisms and structure–processing–property relationships.

Processing Strategy	Matrix/Fiber	Processing Conditions	Reported Property Trends	Typical Applications	Ref.
Direct comparison of RTM and VARI highlighting effects of pressure and temperature on impregnation uniformity.	Partially bio-based epoxy/flax textiles.	Tool temperature 60 or 100 °C; cure time 30 or 180 min; RTM at 6 bar; VARI under vacuum.	RTM increases tensile and flexural properties; VARI improves impregnation homogeneity and ILSS.	Structural laminates requiring reproducibility.	[233]
TP-RTM via in situ ring-opening polymerization enabling thermoplastic matrices from low-viscosity monomers.	PLLA (in situ)/flax fabrics (UD, twill).	Polymerization at 185 °C for 2 h.	High conversion and stiffness; meso-scale voids persist due to flax heterogeneity.	Fully compostable thermoplastic composites.	[234]
Light RTM (L-RTM): vacuum-assisted closed-mold process for smooth dual-surface laminates.	Unsaturated polyester/flax (glass-fiber comparison).	Room-temperature processing; vacuum ≈ 0.8 bar; cure ≥ 24 h.	Void-free laminates with moderate strength; moisture sensitivity limits durability.	Outdoor panels; water-slide structures.	[235]

, experimental studies demonstrate how these coupled variables translate into system-specific processing windows and structure–property responses for natural-fiber composites. RTM, therefore, does not have a single set of processing conditions that applies to all material systems. Instead, the resin must remain sufficiently fluid long enough to fully impregnate the fiber preform before curing begins, which requires system-specific selection of temperature and injection conditions.

Natural-fiber preforms introduce defect mechanisms that are particularly critical in RTM. Clearance between the preform and mold edges promotes preferential edge flow (race-tracking), which disrupts flow uniformity and leads to dry regions and incomplete wetting; this effect becomes more pronounced as preform permeability decreases [215]. Excessive injection pressure can deform the mold or displace fibers, while elevated mold temperatures accelerate cure and increase the risk of premature gelation and short-shot filling. In addition, microscale permeability heterogeneity promotes void formation, which directly degrades mechanical performance. These sensitivities explain the frequent need for vacuum assistance and the strong dependence of RTM outcomes on preform quality, placement, and moisture control.

Recent experimental studies quantify how these sensitivities translate into laminate quality and mechanical response. For flax/epoxy laminates, Kirschnick et al. [233] systematically compared resin transfer molding (RTM) and vacuum-assisted resin infusion (VARI) using identical partially bio-based epoxy systems and flax textile preforms. Two curing windows were examined: 60 °C for 180 min and 100 °C for 30 min, with RTM operated at an injection pressure of 6 bar and VARI under vacuum. Increasing tool temperature lowered resin viscosity and shortened filling times but led to greater density variability in RTM, indicating less uniform impregnation. In contrast, VARI produced more homogeneous laminates with consistently lower density scatter (Table 18). Mechanically, RTM yielded higher tensile and flexural properties, whereas VARI exhibited superior interlaminar shear strength, reflecting stronger fiber–matrix adhesion and more uniform through-thickness impregnation. These results demonstrate a clear trade-off between in-plane stiffness and interlaminar integrity that depends on the impregnation mode and pressure regime.

RTM concepts have also been extended to thermoplastic systems via thermoplastic RTM (TP-RTM). Campos et al. [234] employed in-mold ring-opening polymerization of L-lactide to generate PLLA directly within flax-fiber preforms, achieving monomer conversions exceeding 96 % and mass-average molar masses above $1.4\times {10}^5$ g mol⁻¹. Bending moduli of ≈9.3 GPa and strengths of 177 MPa were obtained, comparable to compression-molded PLLA/flax laminates (Table 18). However, digital microscopy revealed persistent meso-scale voids (≈500–2000 $\mathsf{\mu }$m) between yarns, attributed to intrinsic flax heterogeneity and non-uniform permeability, indicating that low initial resin viscosity alone does not eliminate preform-driven flow defects.

A related closed-mold variant, Light-RTM (L-RTM), incorporates vacuum assistance to improve surface quality on both laminate faces. For flax/polyester systems, Türkhan et al. [235] demonstrated that L-RTM yields essentially void-free laminates with fiber volume fractions of ≈30–35%, tensile strengths near 150 MPa, flexural strengths near 120 MPa, and Charpy impact energies around 30 kJ m⁻². While surface finish and impregnation quality were suitable for outdoor panels and water-slide structures, moisture sensitivity and reduced mechanical performance relative to glass-fiber systems were identified as key limitations for long-term service.

RTM and its derivatives are best suited to medium- and large-area natural-fiber composite components where surface finish, dimensional control, and repeatable impregnation are critical. Their effective deployment remains highly sensitive to preform permeability heterogeneity, edge-flow control, and the pressure–temperature window that couples resin flow to cure progression, as highlighted by both review analyses and recent experimental studies [215,233].

5.6. Mat and Fabric Consolidation

Compression Molding

Compression molding is a closed-mold hot-pressing route well suited to woven, stitched, and mat-based natural-fiber reinforcements because it preserves fiber length, minimizes shear-induced damage, and enables consolidation at relatively high fiber contents. Process quality is governed primarily by the coupled control of temperature, pressure, and moisture content [236]. Natural fibers must be dried from typical ambient levels (6–12%) to below ≈3% prior to molding to suppress vapor-driven void formation. Mold temperatures are commonly maintained in the range of 20–80 °C, while composite processing temperatures are constrained to about 150–220 °C by the thermal stability of lignocellulosic fibers. Consolidation pressures on the order of 6 MPa are typically required to achieve adequate densification and surface quality, with cycle times of ≈30–60 s reported for non-isothermal thermoplastic forming [236].

Laminate quality is therefore determined by the coupled effects of processing temperature, applied pressure, charge configuration, and moisture content. Excessive temperature or pressure accelerates fiber degradation and lumen collapse, reducing effective load transfer, whereas insufficient heat or consolidation pressure leads to incomplete impregnation and residual defects. When these variables are properly balanced, direct platen consolidation yields tight thickness tolerances and good surface finish, as reflected in representative studies summarized in

Table 19. Compression molding of NFRCs: processing constraints and structure–processing–property responses.

Processing Strategy	Matrix/Fiber	Processing Conditions	Reported Property Trends	Ref.
Thermoplastic laminate compression molding from discrete pellets and woven fabrics; laminate quality governed by charge configuration and melt infiltration without fiber damage.	PLA/jute woven fabric laminates	Pellets and fabrics dried 24 h at 80 °C; molding at 190 °C and 1000 psi for 3.5 min; cooling under pressure 15 min; optimal sandwich charge (PLA/fabric stack/PLA).	Charge configuration governs laminate quality: interleaved pellets → incomplete melting and fabric damage; pellets beneath the stack → buckling; sandwich layout (PLA/fabric/PLA) → uniform fiber distribution and void-free laminates. Stiffness ↑ (up to +145% vs. neat PLA), ductility ↓, and impact strength ↓ at low fabric content but ↑ (≤9%) at higher fabric layer counts.	[237]
Process optimization of plant-fiber felt laminates; preheating and pressure- hold stages control consolidation, strength, and moisture resistance; integrated with life-cycle assessment (LCA).	PP/bamboo fiber felt (45:55 bamboo–PP; 2 mm panels)	Orthogonal design: preheat 200–240 °C; preheat time 180–300 s; hold 20–60 s at 20 MPa. Optimum: 220 °C, 210–240 s preheat, 40–50 s hold.	At optimum: tensile strength 35 MPa; flexural strength 45 MPa; ↓ water absorption by ≈15%. LCA identifies energy use, compression molding, and material composition as dominant environmental impact contributors.	[238]
Hybrid laminate compression molding combining natural and synthetic fibers; mechanical response dictated by hybrid architecture and fiber fraction.	PLA/bidirectional hemp–glass fiber mats	PLA dissolved in dichloromethane and impregnated into mats; consolidation at 180 °C and 10 MPa for 20 min; cooling under pressure for 40 min.	↑ Glass-fiber fraction → strong gains in mechanical performance. Hybrid laminate (25% hemp/75% glass): tensile strength 49 MPa; impact strength 7.21 kJ m−2; moisture absorption 0.36%.	[239]
Manufacturing-oriented optimization of wet-formed nonwoven fiber/PP laminates; response-surface modeling quantifies effects of fiber content and molding variables.	PP fibers/wet-formed paper fiber mats (market pulp)	Central composite design: fiber content 45–70 wt%; temperature 185–195 °C; pressure 0.69–2.07 MPa; molding time 2–8 min.	Wide property envelope across design space: water uptake (24 h) 4–117%; flexural modulus 320–3800 MPa; yield strength 1.8–30.5 MPa; Izod impact strength 12.2–41.4 kJ m−2. Fiber content dominates response, followed by molding time, pressure, and temperature.	[240]

.

In PLA/jute fabric systems, Cabrera et al. [237] show that increasing fabric content progressively increases stiffness, reaching up to ≈145% above neat PLA, while ductility decreases by as much as ≈60%, consistent with a transition toward fiber-dominated load bearing. Tensile strength exhibits a non-monotonic dependence on reinforcement content, reflecting competing effects of improved stress transfer and impregnation-related defects. These observations highlight the strong sensitivity of compression-molded laminates to charge architecture and melt-flow control.

Beyond single-fiber thermoplastic laminates, recent studies substantially broaden both the processing latitude and material architectures accessible by compression molding. Li et al. [238] show that mechanical performance and moisture resistance in bamboo fiber/PP felt composites are governed by the coupled effects of preheating temperature, preheating time, and pressure-hold duration, with optimized conditions simultaneously increasing tensile and flexural strength while suppressing water uptake. Notably, their integrated life-cycle assessment identifies the compression-molding step itself as a dominant contributor to environmental impact, exceeding that of material selection. Chandradass et al. [239] further demonstrate that compression molding readily accommodates hybrid reinforcement architectures: in PLA composites reinforced with hemp/glass fiber mats, increasing glass-fiber content produces pronounced gains in tensile, flexural, and impact properties, accompanied by reduced moisture absorption. Using response-surface methodology, Grubb et al. [240] systematically quantify the relative influence of processing variables in wet-formed paper/PP composites, identifying fiber content as the primary determinant of stiffness, strength, impact resistance, and water uptake, followed by molding time, with pressure and temperature exerting secondary effects. These results demonstrate that compression-molding performance is governed less by nominal pressure or temperature than by charge architecture, fiber fraction, and time–temperature history, consistent with the trends summarized in Table 19.

From a manufacturing perspective, compression molding offers high productivity, dimensional stability, and compatibility with continuous or long natural fibers, making it attractive for medium- to high-volume semi-structural components. Typical applications include automotive interior panels, furniture elements, and large molded boards where surface quality, repeatability, and moderate mechanical performance are required [236]. In this processing regime, compression molding provides a practical balance between throughput, fiber preservation, and achievable structural performance.

5.7. Additive Manufacturing (FDM/FFF)

Fused deposition modeling (FDM), also referred to as fused filament fabrication (FFF), enables layer-by-layer construction of thermoplastic components with geometric freedom that is difficult to achieve using conventional molding processes. Natural-fiber-filled filaments based on PLA, PHB, recycled polypropylene, and related matrices are typically processed at 180–220 °C. However, their broader adoption is constrained by filament brittleness, moisture sensitivity, nozzle wear, and weak interlayer bonding [221]. Consequently, additive manufacturing of natural-fiber composites is presently most suitable for prototyping, architected structures, and low-load components rather than load-bearing structural applications.

Across recent experimental studies and focused reviews, a consistent conclusion emerges: in FFF-printed natural-fiber composites, achievable performance is governed primarily by feedstock engineering and printability constraints rather than by intrinsic reinforcement efficiency [221,241]. Unlike conventional molding routes, FFF imposes strict limits on fiber size, dispersion, and moisture content, such that filament integrity and interlayer bonding dominate the mechanical response of printed parts. Moisture-induced porosity, filament brittleness, and weak fiber–matrix adhesion are therefore repeatedly identified as the principal barriers to mechanical property retention during layer-by-layer deposition, underscoring the importance of surface modification, drying protocols, and processing-window optimization.

These trends are systematically illustrated by the experimental studies summarized in

Table 20. Additive manufacturing (fused filament fabrication) of NFRCs: feedstock strategies, printability constraints, and key structure–property responses.

Processing Strategy	Matrix/Fiber	Processing Conditions	Reported Property Trends	Ref.
Short-fiber filament design illustrating printability-limited reinforcement due to agglomeration and filament brittleness.	PLA/hemp powder (2.5, 5 wt%)	Loadings > 5 wt% cause unstable extrusion and nozzle clogging; printing parameters must be re-tuned with increasing fiber content.	Demonstrates a practical upper bound on natural-fiber loading and the dominance of filament integrity over reinforcement efficiency.	[242]
Particle-filled formulation (hemp shive) linking melt rheology and interlayer diffusion to print quality.	PLA, PBS, PLA/PBS (50/50)/hemp shive (3, 5 wt%)	Defect-free printing achieved only within a narrow thermal window; higher temperatures risk filler degradation.	Optimal window improves print quality while hemp increases stiffness but modestly reduces strength when adhesion is limited.	[243]
Continuous-fiber reinforcement via layer-paused FFF without special filament.	PLA/continuous hemp fiber	Manual fiber insertion into CAD-defined internal channels during printing.	Maximum tensile force increases by ≈30%, while displacement at peak load decreases by ≈9%; failure dominated by fiber pull-out and interfacial separation.	[244]
Process-window shift for biocomposite filaments relative to neat PLA.	PLA/date-palm fiber	Low-speed filament extrusion and higher effective melt temperatures required for stable printing relative to neat PLA.	Neat-PLA parameter sets are not transferable; fiber hygroscopicity increases water uptake and reduces tensile strength.	[245]
Fiber heat treatment to reduce hygroscopicity and suppress porosity-driven defects.	PLA/heat-treated wood fibers (20 wt%)	Pre-treatment modifies fiber chemistry and moisture response prior to compounding.	Crystallinity increases 23.4%→34.0–43.9%; 24 h water uptake decreases 3.9%→3.2%, improving dimensional stability of printed parts.	[246]

. In PLA systems reinforced with hemp powder, Srivastava et al. [242] intentionally limited fiber contents to 2.5 and 5 wt% because preliminary trials showed that loadings above 5 wt% caused unstable extrusion and nozzle clogging due to agglomeration and filament brittleness. Filaments were produced by single-screw extrusion at 170–185 °C, and printing parameters had to be re-tuned upward in nozzle temperature and downward in print speed as fiber content increased. These results demonstrate a practical upper bound on natural-fiber loading in FFF filaments and illustrate that conventional composite design heuristics based on maximizing fiber fraction are poorly suited to material-extrusion additive manufacturing.

In a complementary formulation-driven study, Garofalo et al. [243] investigated PLA, PBS, and a 50/50 wt% PLA/PBS blend filled with hemp shive particles (3 and 5 wt%). They demonstrated that print quality is governed by melt rheology and the resulting die-swell behavior at the nozzle exit, which controls bead geometry and interlayer diffusion. A narrow thermal window was identified in which defect-free deposition and adequate interlayer bonding could be achieved, with an optimal balance between print quality and mechanical response at ≈190 °C nozzle temperature and 70 mm s⁻¹ print speed. At these conditions, hemp shive increased flexural stiffness but modestly reduced strength, consistent with limited fiber–matrix adhesion, whereas the PLA/PBS blend provided enhanced ductility while maintaining printability.

Alternative reinforcement strategies attempt to decouple fiber incorporation from filament fabrication. Using a layer-paused FFF strategy with manual placement of continuous hemp fibers into CAD-defined internal channels, Karaca and Öztürk [244] increased the maximum tensile force of printed PLA specimens from ≈1.55 to 2.00 kN (≈30%), while reducing the displacement at peak load by 8.7%. Fracture analysis revealed fiber pull-out and interfacial separation as dominant failure modes, indicating that continuous-fiber placement enhances load capacity but remains limited by interfacial adhesion.

A further illustration of composition-specific process windows is provided by Mian et al. [245], who examined FDM printing of PLA/date-palm-fiber filaments. Filament extrusion required low screw speeds (5–6.5 rpm) and modified temperature profiles to prevent melt instability and nozzle blockage relative to neat PLA. Printing parameters optimized for neat PLA (210 °C nozzle) could not be directly transferred to the fiber-filled system, which required higher effective melt temperatures to maintain stable flow. Despite successful printing, the biocomposite exhibited increased water uptake (0.58% vs. 0.10% for neat PLA), attributed to the hydrophilic nature of date-palm fibers, and reduced tensile strength relative to unfilled PLA. These findings reinforce that in FFF of natural-fiber composites, processing windows and performance are dictated primarily by filament rheology and fiber hygroscopicity rather than by reinforcement efficiency.

Feedstock preconditioning offers a route to improve dimensional and moisture stability without increasing fiber content. Chien and Yang [246] heat-treated wood fibers at 180 °C for 2–6 h prior to compounding with PLA (PLA itself was only dried). Fiber heat treatment increased the crystallinity of the PLA matrix from 23.4% to 34.0–43.9%, reduced filament porosity, and lowered the water absorption of printed parts after 24 h from 3.9% to 3.2%. These effects were attributed to reduced fiber hygroscopicity and to enhanced nucleating activity of the thermally modified fiber surface, which promotes PLA crystallization during cooling. The results demonstrate that microstructural control of the filament via fiber preconditioning can improve dimensional stability and property retention in printed components without increasing fiber loading.

Successful additive manufacturing of natural-fiber composites, therefore, requires an integrated feedstock–process–interface design strategy. Performance gains are achieved more effectively by controlling moisture uptake, stabilizing filament rheology, and strengthening fiber–matrix interfaces than by increasing fiber content alone. Future progress will therefore depend on interfacial engineering and feedstock control strategies tailored to the thermal and flow constraints of fused filament fabrication.

5.8. Outlook and Processing Trade-Offs

Each fabrication route for natural-fiber–reinforced biocomposites involves inherent trade-offs among process control, achievable microstructural quality, and scalability. Low-temperature, solvent-based techniques offer fine control over fiber dispersion and interfacial interactions but are inherently limited in throughput and industrial viability. Thermomechanical processing routes enable scalable manufacturing but introduce fiber attrition, orientation effects, and process-induced anisotropy that complicate property control. Liquid composite molding provides superior impregnation and higher fiber volume fractions for continuous reinforcements, whereas additive manufacturing offers unmatched design freedom at the expense of mechanical efficiency and process robustness. Future progress will depend on hybrid processing strategies, advanced compatibilization and surface-modification approaches, and predictive process–structure–property frameworks that enable reliable deployment of natural fiber composites across structural, functional, and packaging applications.

6. Degradation, End-of-Life, and Circularity

Figure 9 summarizes how composite design choices govern degradation pathways and, in turn, determine recyclability, biodegradation, and circularity outcomes. The diagram emphasizes that end-of-life performance is not an intrinsic material property, but an emergent result of fiber chemistry, matrix selection, interfacial engineering, and processing history.

6.1. Circularity Framework and End-of-Life Pathways

From a systems-design perspective, sustainability, recyclability, and biodegradation are therefore treated here as performance-constraining design variables rather than as post-use considerations. The end-of-life (EoL) behavior of natural-fiber–reinforced composites arises from coupled degradation processes within the polymer matrix, the lignocellulosic reinforcement, and the fiber–matrix interphase. Natural fibers promote biodegradation through their hydrophilic cellulose and hemicellulose, which facilitate moisture uptake and microbial colonization; however, this same hydrophilicity can accelerate interphase swelling, hydrolysis, and debonding during service. Biocomposites must therefore balance durability during use with controlled degradation after disposal, a challenge that becomes increasingly significant as global plastic consumption continues to rise. Worldwide plastic use is projected to increase from 464 Mt in 2020 to 884 Mt in 2050, with cumulative environmental accumulation potentially reaching 4725 Mt under business-as-usual scenarios [247].

Mechanical recycling preserves polymer chains, in contrast to thermal recycling (combustion or pyrolysis), which breaks them down for energy recovery. It remains the most accessible circularity route for thermoplastic biocomposites, though repeated melt reprocessing can induce chain scission, thermo-oxidation, fiber attrition, and interfacial debonding. PLA/sisal composites, for instance, retain modulus over multiple cycles but exhibit a ≈21% reduction in tensile strength after three cycles and a decrease in $T_{\mathrm{g}}$ from 68.9 to 61.8 °C, indicating molecular-weight degradation and limited suitability for repeated recycling [248].

Recycled polyolefin composites, produced by reinforcing recycled polyethylene or polypropylene matrices with natural fibers, remain essentially non-biodegradable because polyolefins are chemically inert and highly resistant to microbial attack. Their durability, however, is offset by excellent melt recyclability, providing a circularity pathway distinct from biological degradation and well suited to closed-loop material recovery. In LDPE–lignocellulosic composites, biodegradation originates almost exclusively from the filler rather than the matrix. Zykova et al. [249] showed that LDPE–flax-straw composites exhibit a strong morphology–degradation relationship: elongated, high-aspect-ratio flax-straw particles create a more porous and interconnected microstructure within the hydrophobic LDPE matrix, enhancing water ingress and microbial access. By contrast, LDPE–wood-flour composites, containing shorter and more equiaxed particles, show substantially lower mass loss under identical conditions. Across all systems, the LDPE matrix remains intact, confirming that biodegradation proceeds exclusively through the natural filler while the polyolefin phase supports recyclability rather than biological breakdown.

6.2. Life-Cycle Assessment and Circular Materials Selection

Circularity begins at the materials-selection stage. Agricultural residues such as flax straw and bagasse, recycled textiles including cotton and jute, and biochar increasingly serve as low-impact reinforcements or fillers, and their benefits are routinely quantified using life-cycle assessment and life-cycle engineering frameworks [250,251,252]. Recent cradle-to-gate analyses show that incorporating upcycled end-of-life textiles into natural-fiber laminates can both improve mechanical performance and reduce environmental burden. In particular, hybrid flax–recycled-polyester composites exhibit lower impacts across all ReCiPe midpoint categories and significantly higher toughness and interlaminar shear strength than flax-only laminates [253].

6.2.1. Environmental Burdens of Natural-Fiber Supply Chains

LCA studies of flax-fiber production show that environmental performance is governed by both agricultural practice and processing route. Detailed life-cycle inventories reveal that only ≈5% of harvested stem mass is converted into long fiber, with the remainder entering low-value waste streams, and that agrochemical manufacture accounts for more than 80% of agricultural energy demand, driven largely by nitrogen fertilizer production (66 GJ t⁻¹ of nutrient) [254]. Processing contributes additional variability: warm-water retting combined with no-till cultivation can yield sliver with an embodied energy of 59 GJ t⁻¹, whereas traditional ploughing and bioretting increase this to nearly 200 GJ t⁻¹. Spinning remains the most energy-intensive step, adding roughly 24 GJ t⁻¹ even under optimized conditions.

6.2.2. Allocation Methods and LCA Sensitivity

LCA results for flax depend strongly on how environmental burdens are divided between co-products. Because fiber and seed are produced from the same crop, impacts may be allocated to the high-value seed (economic allocation) or distributed by mass across all outputs. Economic allocation or system expansion assigns most burdens to the seed, making the fiber appear to have very low embodied energy, whereas mass-based allocation provides the fiber with a much larger share and substantially higher impacts. As a result, flax composites can appear either environmentally superior or inferior to synthetic fibers depending on the allocation method rather than any change in process efficiency [255].

These analyses make clear that the sustainability of natural-fiber composites is governed by both process realities (agrochemical intensity, retting route, spinning energy) and LCA methodological choices (allocation strategy, co-product treatment, system boundaries). Incorporating recycled polymer textiles offers an additional route to reducing impacts, as these materials carry minimal new environmental burdens. For example, flax–recycled-polyester laminates exhibit lower cradle-to-gate impacts than flax-only systems by partially substituting high-impact natural fibers with end-of-life textile layers whose embodied burdens have already been accounted for [253]. Robust comparison with glass or carbon fibers, however, requires consistent LCA assumptions (allocation rules, boundary definitions, and impact categories) since methodological differences alone can shift the calculated embodied energy of flax by more than a factor of four. Without such alignment, the circularity and sustainability potential of bio-based and recycled reinforcements cannot be reliably assessed.

6.3. Environmental Controls on Polymer Biodegradation

The biodegradation of biopolymers and natural fiber composites depends strongly on the disposal environment as well as on the chemistry of both the polymer matrix and the lignocellulosic reinforcement. Representative degradation behaviors across polymer families and fiber systems are summarized in

Table 21. Comparative summary of biodegradability and recyclability characteristics of selected natural-fiber–reinforced biocomposites. Matrix systems include PLA and PLA-based materials, PHB/PHBV biopolyesters, starch-based thermoplastics, polyolefins, and bio-based epoxies. Mechanical retention refers only to values explicitly reported after degradation or recycling; where mechanical testing was not performed or not reported, this is stated explicitly.

Matrix	Fiber	Degradation Rate	Mechanical Property Retention	Ref.
PLA	Flax	Hydrothermal aging in water (20–50 °C, 144 h): water uptake and swelling increase with fiber content; tensile strength loss increases with temperature; specimens aged at 50 °C fracture prematurely during testing.	Tensile strength retention of ≈47–80% after 144 h at 20–35 °C; mechanical retention not measurable after aging at 50 °C.	[256]
PLA	Flax (aligned vs. random)	Compost soil burial (120 days): mass loss of 19% (aligned) and 27% (random); surface erosion and fiber pull-out observed.	Residual flexural strength of ≈43% (aligned) and 20% (random) after 120 days.	[257]
PLA blends	None/CaCO3/lignin-coated CNC	Freshwater exposure and industrial composting (58 °C): neat PLA shows minimal freshwater degradation; PLA/ATBC/CaCO3 reaches 94% biodegradation after 60 days in compost.	Mechanical retention not reported; SEM, FTIR, and DSC show surface erosion, ester bond scission, and reduced $T_{\mathrm{g}}$.	[258]
Recycled PLA	Sisal (NaHCO3-treated)	Repeated melt reprocessing (up to 8 cycles): progressive chain scission; $T_{\mathrm{g}}$ decreases from 68.9 to 61.8 °C by 8 cycles.	Tensile and flexural strength decrease by ≈21% after 3 cycles; elastic modulus remains approximately stable up to 4 cycles.	[248]
PHBV	Sisal (propionylated)	Chemical modification reduces hydrophilicity and bacterial attachment (S. aureus, E. coli); thermal stability increases ($T_{\mathrm{d}}$: 265→274 °C).	Tensile strength increases from 21 to 26 MPa; elastic modulus increases from 1.69 to 1.92 GPa.	[259]
PHBV	Oil-palm empty fruit bunch fiber	Soil burial (16 weeks, 25–30 °C): mass loss increases from ≈69% (neat PHBV) to ≈99% with 10 wt% fiber; enhanced water uptake and microbial attack via fiber–matrix gaps.	Mechanical retention not reported; SEM shows interfacial debonding and fiber exposure during degradation.	[260]
PHB	Jute	Soil burial (45 days): ≈65% mass loss.	Flexural strength retention of ≈50% after 45 days.	[261]
Plasticized starch (potato, sweet potato, corn; glycerol/water/glycols)	Sisal/Jute/Cabuya	Compost exposure (30–40 days): initial moisture uptake during first ≈7 days followed by net mass loss of ≈14–17%; jute- and cabuya- reinforced composites show slightly higher loss than sisal; microbial colonization and matrix erosion observed by microscopy.	Mechanical retention not quantitatively reported; morphological damage includes surface cracking, matrix deterioration, and fiber protrusion, while fibers remain largely intact.	[262]
TPS (glycerol–water)	Flax/Date palm/Banana/Bagasse (50 wt%)	Soil burial (6 weeks): residual mass of 59% (flax), 47% (palm), 46% (banana), and 35% (bagasse); progressive matrix erosion and fiber exposure observed by SEM.	Tensile strength and modulus decrease by >50% within the first week for all composites; flax-reinforced TPS retains the highest relative mechanical integrity throughout burial.	[112]
Arrowroot starch (plasticized, solution-cast film)	Arrowroot fiber (2–10 wt%)	Soil burial (up to 12 days): rapid mass loss; neat TPAS loses 64% mass by day 8, while TPAS/AF-10 loses 78.7% by day 8 and fully disintegrates by day 12.	Mechanical properties after degradation were not reported; no post-burial tensile testing was performed.	[263]
TPS (cassava starch)	Milled brown algae	Accelerated weathering (24 h): mass loss of 51.4% for neat TPS and 20.7% for TPS with 10 wt% algae.	Tensile strength increases up to 2.22 MPa (10 wt% algae, sonicated); reduced embrittlement relative to neat TPS.	[264]
Recycled LDPE	Flax straw/wood flour	Soil burial: biodegradation rate correlates with filler particle size ($R^2=0.99$).	Elastic modulus exceeds 200 MPa; tensile strength decreases with agglomeration.	[249]
PP	Pine wood + propolis–silane	UV and fungal exposure: fungal mass loss reduced from 3.5% to 1.5%.	≈90% retention of tensile strength and modulus after combined UV and fungal exposure.	[261]
PP/starch resin/PCL blend	Bamboo fiber	Soil burial (90 days): mass loss up to ≈19% (23.1 wt% fiber); surface fading and pit formation. Microbial exposure (≤49 days): progressive mass loss and surface erosion.	Tensile strength decreases with fiber content; elongation at break decreases sharply during microbial exposure.	[265]
Cashew nut shell liquid (CNSL) epoxy	Banana fiber	Aerobic composting and mud exposure (60 days): mass loss up to ≈19 wt%.	Mechanical retention not reported.	[266]
Epoxy (NEMresin 1011)	Bacterial cellulose	Aerobic respirometric testing (33 days): CO2 evolution approximately twice that of glass–epoxy reference.	Mechanical retention not reported.	[267]
Epoxy, bio-epoxy, or acrylic resin	Flax fiber (UD laminate)	Water immersion (23–60 °C, 56 days): progressive moisture uptake and fiber swelling; matrix plasticization and interfacial debonding.	Flexural strength retention of ≈27–38% after 56 days; glass-fiber composites retain substantially higher strength.	[195]
Epoxy, bio-epoxy, or acrylic resin	Flax fiber (UD laminate)	Accelerated UV/condensation weathering (56 days): photo-oxidation, resin erosion, cracking, and fiber exposure.	Flexural strength decreases by ≈25–50%, depending on matrix chemistry; glass-fiber composites show superior retention.	[268]

.

Thermophilic industrial composting provides elevated temperatures, high microbial activity, and sufficient moisture to accelerate hydrolysis and microbial assimilation, whereas marine environments support far slower degradation owing to lower temperatures, reduced enzymatic activity, and limited nutrient availability [201]. Consequently, PLA undergoes negligible chain scission in seawater because its ester-hydrolysis pathways require temperatures well above ocean conditions, while polyhydroxyalkanoates such as PHB and PHBV degrade more rapidly due to the activity of marine microorganisms that express PHA depolymerases. Other biodegradable polyesters, including PBAT and PBS, exhibit only limited chain cleavage and slow surface erosion. The resulting trend in seawater degradability

$$\mathrm{PHAs}\gg \mathrm{PBAT}>\mathrm{PBS}\gg \mathrm{PLA}$$

demonstrates that materials certified as industrially compostable cannot be assumed to degrade effectively in marine environments.

Matrix-Dependent Biodegradation Pathways

Within natural fiber composites, degradation behavior reflects the coupled effects of matrix chemistry, fiber composition, environmental exposure, and microstructural pathways governing water and microbial ingress. Hemicellulose-rich fibers degrade more rapidly because of their amorphous and highly hydrophilic structure, whereas lignin-rich fibers exhibit slower mass loss due to their aromatic, cross-linked architecture [252]. These compositional effects, together with variations in matrix hydrophobicity and microvoid content, account for the wide range of durability outcomes summarized in Table 21.

PLA-based composites are particularly vulnerable to degradation under hydrothermal and composting conditions. In PLA/flax biocomposites, hydrothermal ageing primarily compromises mechanical integrity rather than mass loss, with tensile strength decreasing by 20–53% after 144 h of immersion at 20–35 °C and approaching complete embrittlement at 50 °C [256]. Compost burial further induces surface erosion and fiber pull-out, leading to substantial reductions in flexural strength depending on fiber alignment and exposure time [257]. Plasticizers such as acetyl tributyl citrate (ATBC) and reduced crystallinity exacerbate these effects by promoting water diffusion and microbial colonization [258].

Fiber-induced microstructure strongly influences biodegradation kinetics in biopolyesters and starch-based systems. In PHBV composites, incorporating 10 wt% oil-palm empty-fruit-bunch fibers increases porosity and water uptake, accelerating mass loss from approximately 69% for neat PHBV to nearly 99% after 16 weeks of soil burial at 25–30 °C [260]. PHB/jute composites show a similar trend, undergoing rapid soil degradation while retaining measurable flexural load-bearing capacity during early exposure, indicating that substantial mass loss does not immediately translate into mechanical failure [261]. Comparable behavior is observed in flax-reinforced starch/glycerol thermoplastics, which lose more than 90% of their mass within 60 days under aerobic composting, although their high water uptake and low tensile strength (<20 MPa) limit structural applicability [112].

Thermoplastic starch systems exhibit comparable vulnerability unless stabilized by fillers that improve interfacial adhesion and restrict moisture transport. In starch–fiber composites, an initial moisture-driven mass gain typically occurs during the first week of composting, followed by 14–17% mass loss after 30–40 days; higher fiber or glycerol contents increase porosity and microbial accessibility, while UV exposure contributes only modest additional degradation [262]. In soil-burial tests, composites of thermoplastic arrowroot starch (TPAS) reinforced with arrowroot fiber exhibited significantly greater mass loss than TPAS films alone [263]. Mineral particulates or algae-derived additives can partially suppress embrittlement and delay hydrolytic attack by reducing interfacial void formation and limiting crack propagation [264].

Moisture-driven degradation dominates the long-term durability of flax-fiber composites. Hygrothermal immersion studies show that unidirectional flax laminates reinforced with petro-epoxy, bio-based recyclable epoxy, or liquid thermoplastic acrylic resins exhibit initial flexural strengths of 220–280 MPa but retain only 27–38% of their unaged strength after 56 days of water ageing as a result of extensive moisture uptake, fiber swelling, matrix plasticization, and interfacial debonding [195]. Accelerated UV/condensation weathering over the same duration produces moderate mass changes but substantial surface deterioration, including resin erosion, photo-oxidation, cracking, and fiber exposure, reducing flexural strength by 25–50% depending on matrix chemistry [268]. In both hygrothermal and UV-driven ageing, glass-fiber composites retain substantially more of their mechanical performance because their non-hygroscopic, inorganic reinforcement does not swell or promote capillary moisture transport and is resistant to UV-induced degradation. These results demonstrate that, even for recyclable or bio-based matrices, durability is largely governed by the hydrophilicity and UV sensitivity of flax fibers.

Surface modification of the reinforcement plays a critical role in mitigating such degradation. Untreated fibers contain surface waxes, extractives, and loosely bound hemicellulose that hinder wetting by the polymer matrix and weaken interfacial adhesion, producing microvoids and capillary pathways that facilitate water ingress and microbial colonization. Alkali treatment (2–10% NaOH) removes hemicellulose, pectin, and surface impurities while increasing surface roughness and exposing cellulose microfibrils. Although cellulose remains intrinsically hydrophilic, removal of amorphous polysaccharides and reduction of capillary channels lowers the density of accessible hydroxyl groups, thereby reducing apparent hydrophilicity and slowing moisture diffusion. Silane treatments further suppress water uptake by forming covalent siloxane linkages at the interface, improving adhesion and reducing moisture-accessible sites.

PHB-based matrices offer favorable biodegradability and mechanical performance for packaging, biomedical scaffolds, and short-service-life components. However, their high crystallinity leads to brittleness, so improved ductility typically requires copolymerization (e.g., PHBV) or blending with PLA or PCL to reduce crystallinity while maintaining toughness during environmental ageing.

6.4. Microbial Attack in Natural Fiber Composites

Natural-fiber–reinforced composites are inherently susceptible to microbial attack because both the lignocellulosic reinforcement and many bio-based polymer matrices contain chemically accessible, metabolizable groups. Cellulose, hemicellulose, and pectin are strongly hydrophilic, promoting moisture uptake that plasticizes the matrix, induces swelling stresses, and weakens the fiber–matrix interface. This moisture-assisted interfacial degradation creates favorable conditions for fungal and bacterial colonization and facilitates progressive penetration of microorganisms into the composite interior.

6.4.1. Microbial Degradation of Lignocellulosic Reinforcements

Microorganisms preferentially degrade the amorphous regions of lignocellulose, and white-rot and soft-rot fungi readily colonize composites with high natural-fiber content, particularly when fibers are poorly encapsulated by the polymer matrix. Systems such as bamboo/PLA and wood–plastic composites exhibit substantial mass loss and mechanical deterioration under fungal exposure. Degradation severity increases with increasing fiber loading and with the presence of readily hydrolyzable constituents, including soluble polysaccharides or proteinaceous additives. High-aspect-ratio fillers and nutrient-rich particulates (e.g., milled straw or hydrolyzed keratin) further intensify microbial activity by increasing accessible surface area and providing additional nutrient sources [269].

6.4.2. Bio-Fillers as Degradation Initiation Sites in Composite Matrices

The presence of lignocellulosic fillers can initiate degradation even in matrices that are otherwise resistant to microbial attack. In Bioplast–spruce composites (where Bioplast denotes a commercial starch-based bioplastic blend), increasing wood content leads to more severe fungal and termite damage as a result of enhanced moisture uptake and microbial accessibility, manifesting as surface microcracking, embrittlement, and localized failure originating in the fiber phase [270]. A related mechanism is observed in soil-buried LDPE/corn-flour composites, where measurable mass loss occurs only in the presence of the biodegradable starch filler; neat LDPE remains essentially unaffected. Microorganisms selectively consume the corn flour, producing pits, voids, and surface erosion that subsequently compromise the surrounding polyolefin matrix [271]. These examples illustrate how bio-fillers can act as preferential degradation pathways, enabling secondary damage in otherwise durable polymer systems.

Matrix chemistry further modulates microbial susceptibility. Hydrophobic polymers such as PE and PP limit bulk moisture ingress, but incomplete fiber encapsulation permits localized microbial attack at exposed lignocellulosic sites. In contrast, biopolyesters such as PLA and polyhydroxyalkanoates may undergo concurrent matrix and fiber degradation under composting or soil-burial conditions, accelerating structural deterioration. Across matrix classes, interfacial weakening, microcracking, and surface roughening enhance microbial adhesion and activity, driving progressive loss of mechanical integrity.

6.5. Durability Versus Degradability in Bio-Based and Recycled Composites

Biodegradation studies on epoxy-based natural fiber composites consistently show that cellulose-rich reinforcements strongly influence end-of-life behavior, largely independent of matrix origin. Bio-epoxy systems derived from CNSL and reinforced with Musa acuminata banana fibers exhibit compost-mediated mass loss approaching 19 wt% after 60 days under aerobic conditions [266]. Similarly, epoxy laminates containing bacterial cellulose release nearly twice as much CO₂ as glass–epoxy controls during 33 days of aerobic respirometric testing, indicating enhanced microbial activity associated with the degradable cellulose phase [267].

6.5.1. Hygrothermal and UV-Driven Aging of Flax Composites

Moisture exposure represents the dominant pathway for early degradation in epoxy composites reinforced with hydrophilic natural fibers. Flax/epoxy laminates absorb only ≈1 wt% moisture at low relative humidity but reach ≈6 wt% at 90% RH, while flax fibers swell by 12–20% in diameter [198]. This hygroscopic mismatch generates interfacial stresses that promote microcracking, debonding, and progressive interfacial weakening. Liquid-water immersion is substantially more damaging: flax/bio-epoxy composites absorb ≈ 11% water and suffer reductions of 64% in flexural strength, 70% in flexural modulus, and more than 50% in tensile modulus [272]. By contrast, samples conditioned in a warm–humid air environment (75% RH, 45 °C), where moisture uptake occurs via vapor diffusion rather than liquid ingress, absorb only ≈2.6% moisture and retain most of their mechanical properties.

6.5.2. Moisture-Driven Interfacial Failure in Recycled and Textile-Based Systems

Comparable moisture-driven interfacial degradation has been reported in epoxy composites reinforced with recycled cotton or jute textiles, where incomplete fiber encapsulation and discontinuous bonding accelerate damage evolution [273]. In such systems, swelling, sorption hysteresis, and the intrinsic incompatibility between hydrophobic epoxies and hydrophilic lignocellulosic fibers collectively dominate interface deterioration and increase susceptibility to microbial attack.

Durability can be improved through targeted chemical modification of the reinforcement. Propionylation of sisal fibers suppresses microbial degradation in PHBV composites primarily by reducing surface hydrophilicity and bacterial adhesion, while concomitantly improving thermal stability and tensile strength through enhanced fiber–matrix compatibility [259]. In wood–PP composites, hybrid propolis–silane treatments reduce mass loss of the composite due to fungal degradation by more than half, while preserving ≈ 90% of stiffness and strength after combined UV and fungal exposure [261]. These approaches strengthen interfacial bonding, delay microcrack formation, and mitigate biological deterioration while maintaining the compostability or recyclability of the underlying polymer matrix.

6.6. Engineering the Degradation Rate

Achieving predictable and tunable biodegradation in natural fiber composites requires coordinated control of chemistry, microstructure, and environmental interactions. Rather than a single optimal formulation, degradation behavior emerges from the coupled effects of interfacial chemistry, filler architecture, and matrix mobility, which together govern moisture transport, microbial accessibility, and enzymatic activity. Consequently, degradation can be deliberately slowed, delayed, or accelerated through targeted materials and processing strategies, depending on the intended service life and disposal environment.

6.6.1. Chemical and Interfacial Control of Biodegradation

Chemical modification of fibers and interfaces provides a direct route to regulating microbial activity and degradation kinetics. Bio-derived antifungal and antioxidant compounds, including tannins, essential oils, and propolis, suppress microbial colonization by inhibiting extracellular enzymatic activity and scavenging reactive oxygen species [274]. Hybrid propolis–silane treatments exemplify this approach, combining bio-based antimicrobial functionality with improved interfacial stability to delay degradation while preserving composite integrity [261]. By simultaneously reducing surface polarity and strengthening fiber–matrix bonding, such treatments limit moisture uptake and restrict microbial access to degradable phases.

6.6.2. Additive-Driven Stabilization and Environmental Shielding

Degradation resistance can also be enhanced through the incorporation of bio-derived fillers that modify local chemistry, radiation sensitivity, and interfacial packing. TPS composites reinforced with milled brown algae exhibit reduced photodegradation because polyphenolic compounds in the algae act as radical scavengers and UV stabilizers [264]. Beyond chemical effects, improved filler dispersion and reduced particle size promote denser interfacial packing, suppress crack initiation, limit moisture ingress, and slow degradation kinetics. In this sense, bio-derived fillers can serve dual roles as reinforcement agents and environmental stabilizers.

6.6.3. Designed Acceleration and Triggered Compostability

Conversely, accelerated or “on-demand” degradability can be engineered by deliberately enhancing hydrolysis and microbial accessibility. Plasticizers such as acetyl tributyl citrate (ATBC) lower the glass-transition temperature of PLA, increase chain mobility, and promote surface microcracking, thereby accelerating water diffusion and hydrolytic chain scission. Calcium carbonate further enhances biodegradation by buffering acidic hydrolysis products and sustaining microbial activity, enabling composting efficiencies approaching 94% within 60 days [258]. Lignin-coated cellulose nanocrystals provide a complementary, temporally programmed strategy: the lignin shell initially reduces surface hydrophilicity and delays degradation, but its gradual oxidative breakdown increases porosity and ultimately accelerates microbial access [258].

6.6.4. Microstructural Design and the Strength–Degradation Trade-Off

Across these strategies, degradation kinetics are ultimately governed by microstructural design. Higher fiber contents, larger aspect ratios, stronger interfacial adhesion, and increased crystallinity generally improve stiffness and strength but restrict moisture transport and microbial penetration, thereby slowing biodegradation. In contrast, high amorphous fractions, increased porosity, loosely bound interfaces, and poorly dispersed or low-aspect-ratio fillers facilitate enzymatic and microbial attack by promoting moisture ingress and colonization. Optimal performance, therefore, lies not at a universal maximum in reinforcement efficiency, but at a tailored balance between mechanical integrity and environmental degradability, defined by the intended service life and end-of-life pathway [36].

6.7. Thermal Decomposition of Lignocellulosic Fibers

Thermal decomposition of lignocellulosic fibers defines the upper processing limits of biocomposites, constrains mechanical recycling, and governs suitability for thermal recovery or energy conversion at end-of-life. Upon heating, these fibers decompose through a characteristic multistage sequence dictated by the chemistry and supramolecular organization of hemicellulose, cellulose, and lignin. Hemicellulose degrades first, typically between 200 and 350 °C, followed by a sharper cellulose decomposition event at 320–400 °C, while lignin undergoes slow, distributed degradation over a broad range extending from ≈200 to 500 °C and, in some systems, to even higher temperatures [212,275]. This staggered behavior reflects increasing molecular order, aromaticity, and thermal stability across the three constituents.

Hemicellulose decomposes readily due to its amorphous, highly branched architecture and the presence of labile acetyl and glycosidic linkages, which promote early volatilization and acid-catalyzed bond cleavage. Cellulose degradation is dominated by random scission of the $\beta$-1,4-glycosidic backbone within its microfibrillar framework, producing a pronounced mass-loss peak once sufficient thermal energy overcomes crystalline constraints. In contrast, lignin’s heterogeneous, cross-linked aromatic network decomposes gradually through parallel bond-cleavage pathways, yielding a wide distribution of activation energies and persistent char formation that strongly influences flame resistance and residual mass.

Kinetic analyses indicate a transition in degradation mechanisms as conversion progresses [275]. At low conversion ($x\lesssim 0.4$), decomposition is dominated by the volatilization of water and extractives, the breakdown of hemicellulose and lignin, and the degradation of amorphous cellulose domains, with heat and mass transfer occurring more readily in non-ordered regions. At higher conversion levels ($x\gtrsim 0.5$), degradation is increasingly governed by crystalline cellulose, whose highly ordered microfibrillar packing impedes heat transfer and shifts decomposition to higher temperatures. Fiber-specific variations arise from differences in cellulose crystallinity, extractive content, and lignin chemistry: higher crystallinity delays mass loss and raises activation energies, while hardwood lignins rich in syringyl units generally degrade more readily than guaiacyl-dominated softwood lignins [275].

This hierarchical decomposition sequence (hemicellulose preceding cellulose, followed by lignin) directly governs heat-release profiles, char yield, and flammability limits in natural-fiber composites. These thermal responses constrain processing windows, influence fire performance, and determine the viability of end-of-life pathways such as mechanical recycling, pyrolysis, or energy recovery, linking lignocellulosic chemistry and microstructure directly to biocomposite thermal design.

6.8. Trade-Offs and Future Directions

Biocomposite design requires explicit trade-offs between in-service durability and end-of-life degradability. Coupling agents, reversible bonding chemistries, and bio-derived antimicrobial barriers (e.g., propolis, chitosan) can extend service life while preserving pathways for controlled disassembly or biodegradation, but moisture-driven dimensional changes, such as the ≈20% radial swelling of flax fibers under high relative humidity [198], must be incorporated into predictive lifetime and reliability models. Fire performance imposes an additional constraint, as flame-retardant additives that improve safety often compromise recyclability or biodegradation. Addressing these competing requirements demands circular-design frameworks that integrate degradation kinetics, mechanical aging, fire safety, recyclability limits, and life-cycle assessment/engineering, enabling materials with predictable lifetimes and application-specific end-of-life pathways across automotive, construction, packaging, marine, and consumer sectors.

7. Applications of Natural-Fiber–Reinforced Composites

Figure 10 summarizes a design-driven pathway linking material selection, interfacial engineering, processing history, durability, and circularity considerations to the range of applications in which natural-fiber–reinforced composites (NFRCs) can be practically deployed. Rather than prescribing applications a priori, this framework emphasizes application screening based on the coupled constraints that define the achievable performance envelope.

7.1. Application-Level Design Constraints

Guided by the framework in Figure 10, the application of NFRCs is governed by coupled constraints on interfacial integrity, thermal stability, and resistance to moisture and environmental degradation, which together determine stiffness, toughness, and long-term reliability. Within these bounds, NFRCs have gained traction across automotive, construction, consumer products, biomedical devices, and packaging, driven by low density, renewability, and favorable specific mechanical performance. Processing route plays a central role in translating these material attributes into application-ready components: compression- and injection-molded NFRCs satisfy the dimensional tolerances and surface-quality requirements of automotive interior panels, trims, and housings, whereas extrusion-based systems enable scalable production of building elements and consumer goods. Additive manufacturing, although currently constrained by interlayer bonding and fiber printability, expands the design space by enabling complex geometries for biomedical scaffolds and customized packaging. Application viability therefore emerges from the interaction of processing history, interfacial quality, and durability requirements, explaining why NFRCs have become credible substitutes for petroleum-derived composites across diverse use scenarios.

7.1.1. Regulatory and Safety Constraints

From a fire-engineering perspective, the dominant variables governing the fire response of natural-fiber composites are heat-release rate, char formation, and the stability of the fiber–matrix interphase under thermal decomposition. Beyond intrinsic material performance, the practical deployment of NFRCs is shaped by regulatory approval, technological readiness, and scalability constraints [36,37,214]. In automotive and construction applications, NFRCs must satisfy fire, smoke, and toxicity (FST) requirements, crash and impact standards, and building codes, which currently limit their widespread use mainly to interior, trim, and semi-structural components rather than primary load-bearing elements [37,214,276]. In packaging and consumer products, food-contact and chemical-safety regulations restrict allowable matrices, additives, and fiber surface treatments [277], while biomedical and marine applications face additional barriers related to biostability, extractables and leachables, and long-term biocompatibility [278,279].

7.1.2. Technology Readiness and Qualification

From a technology-readiness perspective, NFRCs are already used at high TRL in automotive interiors, decking, molded consumer goods, and packaging, where performance, cost, and certification align with industrial requirements [36,37]. In contrast, structural automotive components, marine laminates, aerospace interiors, and biomedical devices remain at lower readiness levels owing to unresolved challenges in durability, moisture sensitivity, reproducibility, and regulatory qualification [37,213]. These gaps reflect not only material limitations but also the difficulty of demonstrating long-term reliability under realistic environmental and mechanical loading.

7.1.3. Scalability and Manufacturing Constraints

Scalability strongly differentiates viable application domains for NFRCs. Melt-processed NFRCs are compatible with established extrusion, injection-molding, and compression-molding infrastructure, enabling high-volume, cost-competitive production for interior, packaging, and consumer markets [214,216,217]. In contrast, resin-infusion routes, hybrid laminates, and additively manufactured NFRCs exhibit lower throughput, higher cost, and more stringent quality-control requirements, which limit near-term scalability. As a result, melt-processed NFRCs occupy higher technology- and manufacturing-readiness levels, whereas infusion-based and additively manufactured systems remain constrained to lower readiness levels despite favorable laboratory-scale performance.

7.1.4. Food-Contact and Packaging-Specific Constraints

In food-contact packaging, regulatory and scale-up constraints are particularly stringent. Reviews of advanced packaging systems emphasize compliance with the EU food-contact framework (EC 1935/2004 [280]) and plastics-specific regulations (EU 10/2011 [281]), with materials incorporating nanoparticles evaluated on a case-by-case basis prior to market entry [277]. Non-intentionally added substances (NIAS) arising from impurities, degradation, and side reactions further complicate safety qualification, since their identification and toxicological assessment remain analytically demanding [277]. Recent packaging reviews also highlight a persistent gap between laboratory-scale performance and industrial implementation, noting that many bio-based and natural-polymer systems lack validation under real processing and storage conditions [282]. End-of-life claims are likewise regulated, with inconsistent use of terms such as “biodegradable” motivating the need for clearer labeling standards and harmonized ISO/ASTM test protocols, as different environments can yield markedly different degradation outcomes for the same material [282].

These regulatory, technological, and scale-up constraints explain why NFRC adoption is most advanced in interior, packaging, and consumer applications, while structurally demanding, safety-critical, and biomedical uses remain longer-term targets despite strong sustainability drivers.

7.2. Packaging

Packaging represents one of the most mature and rapidly advancing application spaces for NFRCs, driven by regulatory pressure, high material turnover, and the urgent need to replace single-use petrochemical plastics. In this sector, relatively modest mechanical demands, short service lifetimes, and well-defined end-of-life pathways align closely with the strengths of bio-based and biodegradable composite systems.

Agro-industrial residues, including coconut shells, rice husks, walnut shells, bagasse, wool waste, fruit peels, and grain by-products, provide abundant, low-cost feedstocks for biodegradable packaging films and molded components [283]. When incorporated into biopolymer matrices such as PVA, PLA, starches, soy resins, or kraft pulp, these fillers enhance stiffness and strength, tailor gas and moisture barrier properties, and enable controlled biodegradation. In many cases, lignocellulosic chemistry and associated phytochemicals impart antimicrobial, antioxidant, or hydrophobic functionality, contributing directly to shelf-life extension and packaging performance. This valorization of agricultural waste streams strongly aligns with circular-economy objectives, transforming residues into functional packaging materials. Representative systems summarized in

Table 22. Representative biocomposite materials for sustainable packaging, summarizing composition, tensile strength, barrier behavior, and biodegradability.

Matrix	Fiber/Filler	Tensile Strength (MPa)	Barrier Performance	Biodegradability	Ref.
Soy resin	Jute felt + 5% Cloisite 15A nanoclay	59.2	Water absorption: 41.66%; contact angle: 72.9°	56.8% after 60 days (compost)	[3]
Arrowroot starch (thermoplastic)	Arrowroot-derived fiber (10%)	Not reported	WVP †: $5.20\times {10}^{-10}$ (30% reduction)	100% after 12 days (compost)	[263]
PVA	Coconut-shell cellulose nanofibers + linseed and lemon oil (2% CNF, 1% oils)	6.72	Contact angle: 91.3°; swelling ratio: 3.19	87.3% after 45 days (soil burial)	[284]
Corn starch	Rice husk fiber (20%) + benzalkonium chloride (0.05%)	1.08	Contact angle increased to 64°; moderate transparency	100% after 30 days (compost)	[285]
PLA	Microwave-functionalized wood fibers (10–20%)	54.5	WVP: $4.8\times {10}^{-11}$; nearly 100% UV blocking	Not reported	[286]
PLA	Silane-treated walnut-shell powder (10%)	29.2	Contact angle decrease with higher loading	17–19% after 60 days (soil burial)	[287]
Kraft pulp	Hydrolyzed wool fibers	56.2 (tensile index *)	Not quantified	100% after 90 days (soil burial)	[288]
PE/PP blend	Alkali-treated hemp fiber	49.7	WVP: $2.69\times {10}^{-11}$; WVTR ‡: 51; contact angle: 119°	Not biodegradable; recyclable	[289]

* The tensile index is the tensile strength of a paper or fiber network normalized by its basis weight. It is defined as the maximum tensile force per unit width (N m⁻¹) divided by the grammage (g m⁻²), yielding a thickness-independent measure of network strength in N m g⁻¹. ^† WVP: water vapor permeability, the intrinsic permeability to water vapor, calculated by normalizing WVTR to film thickness and the vapor-pressure gradient (g m⁻¹ Pa⁻¹ s⁻¹). ^‡ WVTR: water vapor transmission rate, the steady-state mass of water vapor passing through a unit area of film per unit time (g m⁻² day⁻¹).

illustrate how matrix–fiber synergy governs mechanical performance, barrier behavior, and degradability in emerging bio-composite packaging technologies.

7.2.1. Food Packaging and Functional Films

In food packaging, biocomposites are increasingly engineered not only for containment but also for active and intelligent functionality. Cellulose-based coatings formulated with waxes, fats, and gelatins have been widely reported to offer superior gas-barrier performance compared to many petroleum-derived polymers [290]. Recent studies further integrate natural fibers, antimicrobial agents, and botanical extracts to synergistically enhance mechanical and barrier performance while maintaining high biodegradability.

Poly(vinyl alcohol) films reinforced with coconut-shell cellulose nanofibers (CNF) and modified with linseed and lemon oil exhibit a substantial increase in tensile strength (2.56 to 6.72 MPa), enhanced hydrophobicity (contact angle 91.3°), reduced swelling, and ≈87% biodegradation after 45 days of soil burial [284]. Corn-starch films reinforced with rice-husk fibers and benzalkonium chloride (a quaternary ammonium antimicrobial) similarly show a 61% increase in ultimate stress (1.08 MPa), improved hydrophobicity (contact angle up to 64.2°), enhanced optical transparency, and complete biodegradation within 30 days [285].

Thermoplastic arrowroot starch (TPAS) films containing 10 wt% arrowroot fibers reduce water-vapor permeability by ≈30% relative to neat TPAS while degrading completely within 12 days under composting conditions [263]. Beyond mechanical and barrier enhancements, intelligent packaging systems based on microcellulose or methylcellulose matrices doped with anthocyanins or tea extracts function as freshness indicators, undergoing visible color changes in response to volatile amines released during food spoilage [291,292]. Despite these advances, potential nanoparticle migration and long-term cytotoxicity remain critical concerns for food-contact applications when functionality is imparted using non-biogenic or non-food-safe nanofillers [293]. By contrast, systems based exclusively on food-safe biopolymers, cellulose derivatives, and botanical additives largely mitigate these toxicity risks and are more readily aligned with regulatory and sustainability objectives.

7.2.2. Consumer and Structural Packaging

Bio-composite structural packaging systems are increasingly approaching the stiffness, strength, and durability of petroleum-based plastics while retaining compostability or recyclability. Jute-felt/soy-resin composites reinforced with 5 wt% nanoclay (Cloisite 15A) achieve a tensile strength of 59.2 MPa, compared with 35.8 MPa for the clay-free control, while reducing water absorption from ≈52% to 41.7% and increasing surface hydrophobicity (contact angle 72.9°) [3]. Although nanoclay incorporation slightly slows biodegradation (56.8% mass loss after 60 days versus 69.2% for the control), the composites remain highly compostable, illustrating the trade-off between reinforcement and degradation rate.

PLA-based systems further demonstrate the potential of bio-derived matrices for structural and semi-structural packaging. PLA reinforced with microwave-functionalized wood fibers reaches tensile strengths of 54.5 MPa, exhibits near-complete UV blocking, and shows a 48% reduction in water-vapor permeability relative to neat PLA [286]. Silane-treated walnut-shell particulates enable uniform additive-manufacturing PLA filaments, with biodegradation levels of 17–19% after 60 days depending on surface chemistry [287]. Finite-element analyses of PLA containers reinforced with banana fibers predict elastic moduli comparable to those of commercial PET packaging, indicating suitability for mechanically demanding packaging applications [294].

Starch/sisal foams densified with aluminum hydroxide, Al(OH)₃, exhibit tunable compressive behavior suitable for protective and cushioning packaging [295]. Kraft-pulp composites reinforced with hydrolyzed wool fibers exhibit tensile indices below that of pure kraft pulp (56.2 N m g⁻¹) but degrade completely within 90 days of soil burial [288]. Polyolefin-based composites, while not biodegradable, remain relevant for mechanically recyclable packaging: PE/PP laminated films reinforced with ≈29 wt% alkali-treated hemp fibers achieve tensile strengths of 49.7 MPa, low water-vapor permeability ($2.69\times {10}^{-11}$ g m⁻¹ Pa⁻¹ s⁻¹), and high hydrophobicity (contact angle 119°), making them competitive with commercial packaging films [289]. Finally, mycelium-based foams grown directly in molds or lightly compression-molded convert agricultural residues into rigid, fully biodegradable packaging components, offering a distinct biofabrication route for disposable structural packaging [77,296].

7.2.3. Processing and Performance Optimization

Processing strategy decisively governs how formulation translates into composite performance. Solution casting produces highly uniform, low-porosity films with excellent barrier properties and is widely used for PVA/CNF nanocomposites [297,298,299] and TPS systems [100,101,102,103]. However, long solvent-evaporation times, thickness limitations, and solvent handling constrain its industrial scalability.

Hot pressing offers a scalable alternative, enabling rapid consolidation, efficient moisture removal, and dense interfacial packing, which together enhance stiffness, strength, and dimensional stability in soy-based films and starch/sisal composites. Improved mechanical integrity in hot-pressed soy/glycerol films has been attributed to reduced microvoid content and tighter molecular packing [3]. When combined with in situ foaming, hot pressing further enables controlled cellular architectures: moderate aluminum hydroxide (ATH) loadings (10–20 wt%) promote uniform pore formation and improved tensile and compressive properties, whereas excessive filler disrupts foam development and degrades performance [295].

Beyond consolidation methods, targeted process intensification expands the performance envelope of bio-composite packaging. Microwave-assisted fiber functionalization provides an energy-efficient route to surface modification and markedly improves interfacial adhesion in PLA/wood-fiber systems [286]. Additive manufacturing enables geometrically complex PLA/walnut-shell components, although performance remains limited by interlayer bonding and filament uniformity [287].

7.2.4. Outlook for Sustainable Bio-Composite Packaging Technologies

Recent advances indicate that bio-composite packaging has moved beyond proof-of-concept toward mechanically robust, functionally differentiated systems suited to short- and medium-lifetime applications. As summarized in Table 22, tensile strengths approaching those of commodity plastics (≈30–60 MPa) are now achieved in fiber-reinforced soy-, PLA-, and pulp-based systems, while starch- and PVA-based matrices prioritize rapid biodegradation at lower strength levels. Barrier performance enhancements are formulation-specific and include reduced water-vapor permeability, increased hydrophobicity, and, in some cases, near-complete UV blocking. End-of-life behavior ranges from rapid compostability (12–30 days for starch-based materials) to slower, partial degradation in PLA systems, highlighting the importance of aligning material design with intended service lifetime. Continued progress will require integration of interfacial engineering, solvent-free and scalable processing routes, benign additive chemistries, realistic durability and migration testing, and comprehensive life-cycle assessment to ensure that improvements in performance and manufacturability translate into genuine environmental benefit within circular packaging systems.

7.3. Automotive, Aerospace, Marine, and Construction Applications

NFRCs have progressed from laboratory-scale innovations to widespread deployment in transportation and construction, with growing interest in aerospace and marine components exposed to moisture, impact loading, and corrosive environments. Their appeal derives from low embodied energy (the total energy required to extract raw materials, process them, and manufacture the final product), renewability, and highly tunable mechanical, acoustic, and thermal properties. However, their inherent flammability, driven largely by volatile-rich polymer matrices and the wicking-assisted combustion of fiber–matrix interfaces, remains a critical limitation in transportation and aerospace applications, necessitating flame-retardant additives, interfacial treatments, or hybrid laminates to achieve regulatory compliance [300].

Fire performance of aerospace interior materials is regulated by FAR 25.853 and is demonstrated primarily through the vertical burn test (Appendix F, Part I), with heat release evaluated using the OSU calorimeter (Appendix F, Part IV) and smoke emission quantified by the specific optical density (Ds) measured according to ASTM E662 [301,302].

Recent advances in fiber treatments, hybrid laminates, aerogel architectures, and hierarchical microstructural engineering have expanded the application space of NFRCs, enabling them to meet increasingly stringent performance requirements across interior automotive panels, structural building elements, multifunctional aircraft interiors, and coastal or marine hardware.

7.3.1. Automotive

Table 23. Representative NFRCs for automotive applications, illustrating mechanical performance, functional attributes, and sustainability.

Application	Matrix	Fiber	Performance Attributes	Ref.
Automotive interior components	PP	Nettle (5–30 wt%)	Tensile strength decreases with fiber loading (32.9→15.3 MPa), while tensile modulus increases to a maximum of 2203 MPa at 15 wt% nettle fiber; flexural modulus reaches 2429 MPa at 30 wt%; impact strength peaks at 27.9 J m−1 (30 wt%); enhanced stiffness relative to neat PP supports suitability for injection- and compression-molded interior automotive components.	[303]
Automotive lightweight components	Epoxy	Sea purslane (NaOH + acrylic-acid treated)	Tensile strength 49.6 MPa; flexural strength 56 MPa; modulus 2.38 GPa; impact strength 9.86 kJ m−2; microhardness 24.2 HV; increased crystallinity and thermal stability (onset 280–450 °C); optimum performance at 20 wt% due to strong interfacial adhesion and low void content; treated fibers exhibit lower moisture uptake than untreated.	[304]
Automotive interior and exterior components	PP/MAPP (IM) or polyester/epoxy (RTM)	Kenaf (10–60 wt%)	IM: tensile strength 38 MPa; RTM: 55 MPa. Flexural and impact properties peak at 20–30 wt% fiber content; RTM provides superior wet-out and higher ductility. Treated fibers reduce moisture uptake, and RTM parts exhibit lower water absorption than injection-molded parts.	[305]
Automotive interior components	Epoxy	Jute–glass hybrid laminates (glass skins)	Glass outer plies enhance stiffness, strength, and moisture resistance; tensile strength ≈115 MPa and flexural strength ≈217 MPa; the optimal G–J–J–G stacking sequence maximizes impact strength (378 kJ m−2).	[306]

summarizes representative automotive applications of NFRCs, highlighting component-level performance and functional attributes. The automotive sector represents the most mature and widely adopted application domain for NFRCs, driven by the combined imperatives of lightweighting, cost efficiency, and sustainability. Consequently, natural fibers are now routinely integrated into commercial vehicle components, including door and instrument-panel substrates, center-console carriers, parcel shelves, seat-back structures, trunk liners, wheel-arch covers, and under-body shields. Nettle-fiber/PP composites exemplify this balance of performance and processability: although increasing fiber content progressively reduces tensile and flexural strength, fiber loadings of 15–30 wt% substantially increase tensile and flexural moduli and yield a steady improvement in impact resistance, producing materials well suited for injection- and compression-molded interior components [303].

Achieving automotive-grade performance depends critically on interfacial engineering. Sea purslane fibers sequentially treated with NaOH and acrylic acid exhibit strong fiber–matrix adhesion and the lowest reported void content (1.04%), enabling marked gains in tensile (49.6 MPa) and flexural strength (56 MPa) at 20 wt% fiber loading [304]. These improvements arise from increased surface roughness and enhanced availability of reactive functional groups, underscoring the central role of surface modification in translating fiber reinforcement into load-bearing efficiency.

Kenaf-based systems further illustrate how processing route governs achievable performance in semi-structural automotive applications. Resin-transfer-molded kenaf/epoxy laminates reach tensile strengths up to 55 MPa at 40 wt% fiber, exceeding those of injection-molded PP/MAPP counterparts (38 MPa at 30 wt%) owing to superior fiber alignment, improved wet-out, lower void content, and the more homogeneous microstructures produced by controlled thermoset curing [305]. These comparisons highlight the decisive influence of fiber orientation and impregnation quality on the structural limits of bio-based automotive laminates.

Hybridization strategies provide an additional lever for performance optimization. While early studies broadly suggested benefits from combining natural fibers with thin glass skins, recent systematic evaluations clarify that stacking sequence is the dominant factor. An assessment of five jute–glass laminate architectures shows that configurations with glass skins and jute cores (G + J + J + G) deliver the highest tensile modulus (10.6 GPa), flexural modulus (13.9 GPa), and impact strength, while simultaneously reducing water uptake [306]. In contrast, laminates with jute on the outer surfaces, particularly those incorporating short or discontinuous jute fibers, exhibit reduced flexural performance and increased susceptibility to moisture-induced degradation. Strategic placement of thin glass layers therefore enables substantial gains in stiffness, strength, and durability while preserving the lightweight and sustainability advantages of natural-fiber composites.

7.3.2. Aerospace

Natural-fiber composites are increasingly evaluated for aircraft interiors and secondary aerostructures in response to decarbonization mandates and the demand for lightweight, low-toxicity materials. Current aerospace applications are concentrated in moderate-load interior components, including cabin sidewalls, seat shells, tray tables, trim and lining elements, overhead-bin housings, and related structures. Environmental-exposure studies demonstrate that flax/epoxy laminates remain chemically stable in hydraulic fluids and jet fuel, while moisture and UV exposure primarily induce surface-level degradation, manifested as swelling, gloss loss, erosion, and microcracking, rather than catastrophic mechanical failure [307]. Correspondingly, mechanical property reductions remain limited when appropriate protective coatings are applied.

Table 24. Representative NFRCs for aerospace applications and EM interference mitigation, showing matrix–fiber systems and measured functional performance.

Application	Matrix	Fiber	Performance Attributes	Ref.
External aircraft components (environmental resistance)	Epoxy	Flax fabric (2/2 twill FFRC)	Chemically stable in jet fuel and hydraulic fluid; water exposure causes up to 16% uptake with swelling and delamination; UV induces photo-oxidation, gloss loss, erosion, and microcracking. Fossil and partially bio-based coatings markedly improve moisture and UV resistance, preserving hardness and adhesion; coated FFRC exhibits substantially enhanced durability relative to uncoated laminate.	[307]
Non-structural aircraft interior components	Epoxy + flame-retardant treatments (FR A, FR B)	Flax/epoxy composite	Vertical burn lengths are well below FAR 25.853 limits for all samples (untreated and FR-treated). However, smoke density exceeds the FAR limit (Ds > 200) for both FR A and FR B, and all panels exhibit high peak heat release rates (PHRR = 131–158 kW m−2), far above the 65 kW m−2 requirement. FR treatments reduce CO, NOx, and SO2 emissions relative to untreated panels but do not sufficiently suppress smoke or heat release.	[308]
Aircraft interior and secondary structural components	Epoxy	Banana–sisal hybrid	Tensile strength 18–31.5 MPa; flexural strength up to 46.5 MPa; impact strength 32–42 J m−2; hardness 40–54 HRB; the hybrid composite (12 wt% banana + 5 wt% sisal) exhibits the highest mechanical performance due to improved fiber load sharing and fiber–matrix bonding.	[309]
EM attenuation and thermal insulation	Cellulose–chitosan aerogel	PANI coating	Reflection loss up to $-54.8$ dB at 13.8 GHz; strong dielectric attenuation; ultralight porous network; effective absorption across the X–Ku bands (8–18 GHz).	[310]
EM interference mitigation	Carbon aerogel	CNT/cellulose (carbonized)	RLmin = $-43.6$ dB (3 mm); effective absorption bandwidth 7.42 GHz; hierarchical micro/mesoporous structure; improved dielectric loss and impedance matching from CNT incorporation; carbonized at 550 °C.	[201]

compiles representative aerospace applications of NFRCs, emphasizing environmental durability, fire behavior, and multifunctional performance.

Coating systems, therefore, play a decisive role in enabling aerospace durability. A conventional petrochemical aircraft coating (≈220 µm), consisting of a high-solids polyurethane base layer and an abrasion-resistant topcoat, provides the most effective barrier against UV radiation and moisture ingress [307]. This system preserves matrix integrity, suppresses surface embrittlement, and significantly limits moisture uptake. A thinner, partially biobased polyurethane coating (≈70 µm) offers partial protection but is less effective in mitigating UV-driven oxidation and moisture-induced degradation. These results confirm that flax/epoxy composites, when paired with suitably engineered coatings, can meet durability benchmarks for many interior aerospace applications while delivering sustainability benefits.

Fire, smoke, and toxicity requirements defined by FAR 25.853 impose additional constraints on aerospace adoption. Phosphorus-based flame-retardant systems enable flax–epoxy laminates to satisfy the FAR 25.853 vertical-burn criterion, but peak heat release rates (PHRR) remain excessive: measured values exceed 140 kW m⁻² compared with the allowable maximum of 65 kW m⁻² [308]. This performance gap highlights the need for next-generation flame-retardant chemistries capable of suppressing heat release and smoke evolution without compromising mechanical integrity.

Hybrid natural-fiber architectures offer further performance enhancements for interior components. Banana–sisal/epoxy hybrids exhibit substantial improvements in tensile strength (31.5 MPa vs. 18 MPa for neat epoxy), flexural strength (up to 46.5 MPa), impact resistance (up to 42 J m⁻²), and hardness (up to 54 HRB) when 12 wt% banana and 5 wt% sisal fibers are combined [309]. Hybridization also reduces moisture uptake relative to single-fiber laminates, reflecting complementary reinforcement mechanisms between the two fibers. These property gains support the suitability of banana–sisal hybrids for non-structural aerospace cabin components such as seat shells and interior panels.

Beyond structural applications, cellulose-derived aerogels are gaining attention as ultralight materials for electromagnetic (EM) absorption in avionics housings. Polyaniline-coated cellulose–chitosan aerogels exhibit deep reflection-loss minima of up to $-54.8$ dB at 13.8 GHz, indicating highly efficient microwave attenuation [187]. CNT–cellulose aerogels carbonized at 550 °C similarly achieve strong absorption, with reflection-loss values of $-43.6$ dB and an effective absorption bandwidth of 7.4 GHz, defined as the continuous frequency interval over which reflection loss remains below $-10$ dB [201]. Reflection loss quantifies the fraction of incident electromagnetic energy absorbed rather than reflected, whereas effective bandwidth describes the frequency span meeting a specified absorption threshold. These results establish that plant-derived aerogels can match the EM-absorption performance of carbon foams and metallic meshes while providing substantially lower density and fully renewable material sources.

7.3.3. Marine

Marine applications impose distinct challenges, particularly with respect to moisture uptake, prolonged seawater exposure, and biological fouling. Key application areas for natural-fiber composites in the marine sector include non-structural and semi-structural components such as boat skin elements, interior panels and linings, housings and covers, buoys and flotation devices, and selected superstructure and decking elements. Flax/epoxy and hemp/epoxy laminates exhibit measurable hygroscopicity in seawater, reaching saturation levels of 7.5% and 9.8%, respectively, consistent with the hydrophilic nature of the fibers [311]. This water uptake leads to substantial reductions in tensile and flexural moduli (26–74%), although it can concurrently enhance ductility and impact resistance; notably, flax/epoxy shows a 117% increase in impact strength after conditioning. Biofouling studies further reveal heavy organism settlement on unprotected composite surfaces, particularly for hemp/epoxy systems, whereas conventional copper-based antifouling coatings markedly suppress mass accumulation [311].

Table 25. Representative NFRCs for marine applications, focusing on water resistance, impact behavior, and environmental sustainability.

Application	Matrix	Fiber	Performance Attributes	Ref.
Marine structural components (hull materials)	Epoxy (28% bio-based carbon)	Flax fabric	Seawater uptake ≈ 7.5%; tensile strength 68.6→70.4 MPa (dry→wet); flexural strength decreases from 73.8→39.1 MPa; apparent interlaminar shear strength increases from 10.4→20.7 MPa after seawater conditioning; impact energy rises from 1.72→3.73 J. Flax/epoxy shows good retention of wet mechanical properties and substantially reduced biofouling mass when protected by antifouling coating.	[311]
Marine structural components (hull materials)	Epoxy (28% bio-based carbon)	Hemp fabric	Seawater uptake ≈ 9.8%; tensile strength 45.7→31.3 MPa (dry→wet); flexural strength 81.2→60.4 MPa; apparent interlaminar shear strength decreases from 9.27→6.0 MPa; impact resistance remains nearly unchanged (1.85→1.93 J). Hemp/epoxy exhibits moderate retention of wet mechanical properties and benefits substantially from antifouling coatings, although overall durability is lower than that of flax.	[311]
Shipbuilding components (yachts and motorboats; non-load-bearing structural and lining elements)	DCPD polyester (yacht resin)	Hemp fabric (unmodified vs. NaOH-treated)	Unmodified hemp: tensile strength 81–86 MPa and flexural strength 11–41 MPa; NaOH-treated hemp exhibits reduced strength but increased elongation (12% vs. 4%) and higher impact energy (17–34 J cm−2); flexural load capacity increases with the number of fabric layers.	[312]
Shipbuilding components (impact and fire performance)	DCPD polyester (yacht resin)	Hemp fabrics (unmodified or NaOH-modified)	Impact strength after 3-month immersion: 35.4 J cm−2 (unmodified, HFRP) and 53.4 J cm−2 (modified, HFRPm); stable performance in fresh and brackish water; high total heat release (239–252 MJ m−2); nearly complete combustion (HFRP 96%, HFRPm 83%) in the context of energy-recovery disposal.	[313]

presents representative marine applications of NFRCs, emphasizing seawater exposure, impact resistance, and fire performance.

The pronounced biofouling observed on unprotected hemp/epoxy composites is consistent with established mechanisms of marine adhesion: polymeric surfaces immersed in seawater rapidly acquire conditioning films of dissolved organic matter that facilitate microbial attachment and subsequent biofilm development, while polar, high-surface-energy, and high-modulus matrices such as epoxies provide strong interfacial adhesion sites for bioadhesive proteins and polysaccharides secreted by fouling organisms [314]. Because effective fouling-release materials typically require low surface energy and a surface modulus below ≈10 MPa, conventional epoxy laminates lack the interfacial characteristics needed to shed settled organisms and therefore accumulate substantial biomass. By contrast, copper-based antifouling coatings inhibit settlement through the controlled release of biocidal copper ions, which suppress microbial colonization and algal or larval attachment, explaining the markedly reduced mass accumulation reported for copper-coated composite surfaces [311].

Nevertheless, with appropriate barrier coatings, natural-fiber composites can achieve the durability required for marine environments, especially for components experiencing dynamic loads rather than sustained bending or tension. Further evidence of marine applicability comes from hemp fiber–reinforced polymers (HFRPs) manufactured with a DCPD-modified unsaturated polyester yacht resin, a low-viscosity construction resin widely used in boatbuilding for its favorable wet-out, good hydrolytic stability, and relatively low water uptake [312,313]. Unmodified HFRP laminates produced with this resin outperform alkali-modified counterparts in tensile and flexural strength and provide impact resistance comparable to, or better than, glass-fiber composites in both fresh and brackish water [312,313]. Recommended applications include boat-skin elements with moderate durability requirements, interior panels, housings, buoys, floating structures, and selected superstructure components [312]. Their end-of-life performance is also notable: controlled combustion yields high energy recovery (239–252 MJ m⁻²) with exceptionally low ash residues (3.7–17.4%), compared with more than 50% for conventional GFRP [312,313]. The performance of these laminates and the resin’s processing advantages position HFRPs as promising materials for low- to medium-load marine components within a circular-economy framework.

Moisture-induced degradation, however, remains a concern for certain high-fiber-content systems. Oil-palm-fiber/acrylic composites, for example, lose more than 65% of their flexural strength under prolonged alkali warm-water immersion [169], emphasizing the importance of fiber treatment, matrix selection, and barrier coatings for long-term performance in humid or marine environments.

7.3.4. Construction

NFRCs are gaining prominence in construction, where low density, reduced embodied carbon, and multifunctional performance position them as attractive alternatives to mineral- and petroleum-based materials [315]. Current application spaces span load-bearing and semi-structural panels, façade and cladding systems, lightweight beams and trusses, interior partition boards, insulation layers, roofing elements, and prefabricated modular components. Within this landscape, aligned sisal/epoxy structural panels produced by vacuum infusion demonstrate tensile strengths of 80–220 MPa and elastic moduli of 6–15 GPa over fiber volume fractions of ≈0.08–0.26, directly reflecting the decisive roles of fiber alignment, porosity control, and impregnation quality in achieving predictable load-bearing performance [316]. Microstructural analysis of technical sisal fibers further reveals lumen contents of $23\pm 9$% and cell-wall thicknesses of $3.4\pm 0.7$ µm, indicating that lumen impregnation and interfacial densification are central to realizing effective fiber moduli competitive with jute, flax, and selected E-glass grades.

Table 26. Representative NFRCs for construction and related sectors.

Application	Matrix	Fiber	Performance Attributes	Ref.
Structural panels	Epoxy	Sisal (technical fibers)	Tensile strength 80–220 MPa; modulus 6–15 GPa; strain 1.4–2.0%; lumen impregnation reduces porosity and enhances interfacial bonding, enabling predictable stiffness scaling with fiber volume fraction.	[316]
Acoustic panels	Green epoxy	Coconut/cotton/ bagasse	Sound absorption up to 0.18 at 1600 Hz; cotton provides highest flexural and impact strength, coconut provides highest damping due to intrinsic porosity; thermally stable to ≈300 °C.	[317]
Acoustic panels	Acrylic resin	Oil-palm fiber	Porous lightweight panels with NRC up to 0.37; good impact and indentation resistance; stable under room-temperature water immersion; severe flexural degradation above 60 °C.	[169]
Acoustic panels	PVA binder	Sugarcane bagasse	NRC increases with decreasing particle size; finest particles (0.29–0.37 mm) yield highest absorption and flow resistivity (6750 N s m−4); low-density panels (200 kg m−3).	[318]
Thermal/acoustic insulation	Lignocellulosic substrate	Fungal mycelium	Lightweight bio-composites (density ≈ 90–180 kg m−3); thermal conductivity ≈ 0.04–0.06 W m−1K−1; compressive stiffness and water uptake governed primarily by substrate type and extent of mycelial colonization.	[319]
Thermal/acoustic insulation	Lignocellulosic substrate	Fungal mycelium	Processing-controlled bio-foams with tunable density and moisture resistance; heat pressing increases stiffness and reduces water uptake; mechanical response governed by fungal species, substrate, and consolidation history.	[320]
Structural/insulation	Cement–lime binder	Hemp shives	Density 380–1030 kg m−3; compressive strength up to 9.6 MPa at ≤15 vol% shiv content; thermal conductivity 0.08–0.16 W m−1 K−1; intrinsic antimicrobial resistance.	[321]
Insulation/envelope	Low-carbon mineral binders	Hemp shives	Alkali-activated systems show ≈4× higher compressive strength and ≈31% lower thermal conductivity than hydrated lime; $\mathsf{\mu }$CT attributes strength–insulation synergy to phase connectivity and shiv orientation.	[322]
Insulation/infill	Lime-based binder	Hemp shives	Compressive strength 0.22–0.35 MPa at densities 350–600 kg m−3; strength variability dominated by testing definitions; permanent-strain criterion improves reproducibility and design relevance.	[323]

reports representative construction and related applications of NFRCs, illustrating structure–property relationships and multifunctional roles.

Hybridization strategies further extend the performance envelope of construction NFRCs by enabling targeted combinations of stiffness, durability, and functional response. Jute–glass laminates with glass outer plies exploit the barrier and stiffness contribution of glass to enhance flexural strength, surface hardness, and moisture resistance while retaining the low density and damping capacity of jute cores [306]. Beyond structural performance, acoustic insulation represents another rapidly expanding construction domain for NFRCs. Agricultural waste fibers such as sugarcane bagasse exhibit high sound absorption when finely milled, with absorption coefficients reaching $\alpha \approx 0.63$ and airflow resistivity values consistent with porous-media acoustic behavior [318]. In green epoxy matrices, coconut, cotton, and bagasse fibers provide complementary multifunctionality: coconut fibers yield superior acoustic damping due to high intrinsic porosity, whereas cotton fibers deliver higher flexural strength owing to greater fiber length and tenacity [317].

Hemp/lime composites (hempcrete) represent a distinct class of construction NFRCs in which thermal, hygrothermal, and environmental functions dominate over load-bearing capacity. Reported compressive strengths of hemp/lime composites typically fall in the range of 0.2–0.4 MPa at one year, with mechanical response governed primarily by binder hydraulicity, carbonation kinetics, and interfacial mineralization rather than by fiber reinforcement efficiency alone [324]. Increasing binder hydraulicity accelerates early-age strength development, while long-term strength converges across binder chemistries as carbonation progresses; freeze–thaw resistance and salt durability are similarly controlled by pore structure and binder hydration state. Within this performance envelope, hempcrete is best understood as a multifunctional building material whose value derives from carbon negativity, thermal inertia, and moisture-buffering capacity, rather than from load-bearing capability, supporting its widespread use in envelope and insulation applications [325].

Recent work has clarified structure–property relationships in hemp–lime composites with increasing precision. Low-carbon and alkali-activated binders have been shown to simultaneously increase compressive strength (by up to a factor of four relative to hydrated-lime reference systems) while reducing thermal conductivity, with micro-computed tomography revealing that improved binder-network connectivity, favorable shiv orientation, and enhanced hemp–binder interfacial bonding enable higher strength without sacrificing pore connectivity, allowing mechanical performance and thermal insulation to improve concurrently [322]. Variability in reported compressive strength has been traced largely to inconsistent testing definitions, motivating the adoption of permanent-strain-based criteria that improve reproducibility and design relevance for engineering applications [323]. At higher composite densities, cement–lime systems incorporating hemp shives achieve compressive strengths approaching 5–10 MPa while retaining low thermal conductivity and intrinsic antimicrobial resistance, delineating a continuum between insulation-grade hempcrete and lightweight structural bio-concretes [321].

Emerging bio-based construction materials further expand the NFRC landscape. Mycelium-based composites grown on lignocellulosic substrates form highly porous, low-density networks in which fungal hyphae act as a natural binder, enabling thermal insulation, acoustic damping, and lightweight panel applications with minimal processing energy [326]. Detailed experimental studies demonstrate that mechanical stiffness, moisture uptake, and durability in mycelium composites are governed primarily by substrate type, fungal species, and post-growth consolidation rather than by fiber chemistry alone, with heat pressing shown to significantly increase stiffness and reduce water absorption [319,320]. Together, these systems illustrate that construction NFRCs should be evaluated within application-specific performance envelopes that balance mechanical capacity, thermal and acoustic functionality, moisture regulation, durability, and environmental impact, rather than against a single structural metric.

7.4. Biomedical Applications

NFRCs have gained significant traction in biomedical engineering owing to their biocompatibility, low density, and tunable biodegradability. Their chemical affinity to human tissues and general lack of cytotoxicity could promote favorable host–material interactions without triggering inflammatory responses [327]. When combined with targeted fiber functionalization and interfacial engineering, NFRCs can deliver the mechanical robustness, sterilization compatibility, and bio-integration required for medical-grade applications.

Recent studies highlight diverse strategies for achieving these performance targets. Flax/epoxy laminates containing 3 wt% nano-silica exhibit improved fracture toughness and broad-spectrum antibacterial activity against both Gram-positive and Gram-negative bacteria, effects attributed by the researchers to membrane-disruptive interactions of SiO₂ nanoparticles [328]. This interpretation is noteworthy because membrane disruption typically requires cationic surface functionality [329], and although antimicrobial behavior is well documented for silver-coated nanoparticles [330], unmodified SiO₂ is not generally expected to act through such mechanisms. Tasar silk–fiber/epoxy composites reinforced up to 30 wt% achieve tensile and flexural strengths of 70.9 MPa and 43.5 MPa, respectively, while showing reduced water uptake and stable interfacial bonding, supporting their potential in orthopedic supports and medical housings [331].

Hybridization further broadens the design space by pairing complementary reinforcements to achieve balanced performance. Composites incorporating coconut tree primary flower leaf stalk fiber (extracted from the fiber-rich stalk supporting the coconut inflorescence) and short glass fibers in an unsaturated polyester matrix show well-balanced tensile, flexural, and impact properties alongside low moisture uptake, making them suitable for lightweight external braces and protective shells [332]. Similarly, kenaf–hemp systems modified with 0.5–1 wt% MWCNTs exhibit higher tensile (42.3 MPa) and flexural (59.7 MPa) strengths due to reduced porosity and improved interfacial continuity, supporting their use in orthopedic braces and fixation components [333].

Biodegradable biomedical composites for orthopedic and regenerative applications increasingly draw on fully biosourced reinforcements to achieve tunable mechanical and degradation behavior. Although not a fiber-reinforced system, hydroxyapatite obtained from renewable mussel shells can be incorporated into PCL or PLA to adjust stiffness, thermal transitions, and biodegradation rates in ways consistent with bone-scaffold design requirements [334]. PCL strengthened with natural lignocellulosic fibers provides a combination of strength, ductility, and energy absorption suitable for temporary musculoskeletal supports [335]. PLA reinforced with cellulose crystals derived from jute exhibits increased crystallinity, enhanced modulus and hardness, and mild antibacterial activity, indicating potential for sutures, resorbable fixation components, and other temporary biomedical devices [336].

Biofunctionalization strategies are rapidly advancing NFRC capabilities. Antibiotic, peptide, and nanoparticle-functionalized fibers (e.g., Ag, ZnO) impart antibacterial behavior and promote cell adhesion, enabling wound-contact materials and antimicrobial surface coatings [54]. Incorporation of essential oils and plant-derived flavonoids into starch- and cellulose-based films provides non-toxic antimicrobial functionality for wound dressings and temporary implants [337].

A summary of representative NFRC systems for biomedical applications, including matrix–fiber combinations, dominant property enhancements, and clinical relevance, is presented in

Table 27. Representative NFRCs investigated for biomedical applications. Systems are grouped by application class, highlighting matrix–filler combinations, dominant property gains, and intended biomedical relevance. Italicized entries denote application-class group headings.

Matrix	Filler	Key Property Enhancements	Biomedical Relevance	Ref.
External supports and housings
Unsaturated polyester (UP)	Coconut flower leaf-stalk fiber + glass fiber (hybrid)	Glass-rich hybrids show higher tensile, flexural, and impact strength; alkali treatment and higher glass fraction reduce water uptake, while increased coconut fiber raises moisture absorption.	Non-implant external supports and orthotic housings requiring impact resistance and moderate moisture tolerance.	[332]
Epoxy	Waste flax fiber + nano-SiO2 (0–5 wt%)	Mode I interfacial fracture toughness increases (max at 3 wt% nano-SiO2); antibacterial activity observed against Gram-positive and Gram-negative strains via inhibition-zone testing.	Damage-tolerant, contact-safe external housings with antibacterial surface functionality (biocompatibility not fully assessed).	[328]
Epoxy	Kenaf + hemp fibers with MWCNT (0.5–1 wt%)	Hybridization with low MWCNT loading increases tensile (to ≈42 MPa), flexural (to ≈56–60 MPa), and impact strength, with reduced porosity across formulations.	External fixation and support components where stiffness and impact resistance are critical, subject to nanoparticle safety evaluation.	[333]
Resorbable and orthopedic biomaterial systems
PLA/PCL	Hydroxyapatite particles (5–20 wt%)	HA increases flexural strength and modulus and shifts density toward cortical-bone values; higher HA contents accelerate hydrolytic degradation during saline immersion.	Resorbable bone-fixation and scaffold systems with stiffness and degradation rate tunable via HA loading.	[334]
Biofunctional and antimicrobial strategies
PCL	Lignocellulosic fibers (sisal, hemp, coir; 10–30 wt%)	At 30 wt% sisal, tensile strength and modulus increase substantially, while hardness remains comparable to commercial splinting sheets; elongation at break decreases markedly.	Thermoformable orthotic and splinting components combining bio-based reinforcement with clinically relevant stiffness.	[335]
PLA	Jute-derived crystalline cellulose (micro-/ nanocrystals; 3–15 wt%)	Cellulose acts as a nucleating agent, increasing crystallinity, hardness, and yield strength; antibacterial activity observed only at high loading, with limited interfacial adhesion.	Resorbable plates or thin biomedical components where stiffness and surface hardness are prioritized over ductility.	[336]
PLA (film)	Ag-nanoparticle-modified hazelnut husk flour (10–40 wt%)	Strong antibacterial activity against S. aureus, increasing monotonically with filler content (up to ≈100% reduction); Ag loading quantified by ICP-OES.	Antimicrobial bio-based films for hygiene-sensitive applications such as packaging or contact surfaces.	[338]
PLA (hot-pressed)	Sida hermaphrodita fibers with ZnO/cork/MMT additives	Significant bacterial log-reduction against S. aureus and K. pneumoniae after 24 h contact; antimicrobial effect observed across multiple composite variants.	Rigid antimicrobial biocomposites for non-implant contact surfaces and housings.	[339]
Epoxy	Kapok fiber + rice-bran particulate filler	Antibacterial inhibition against K. pneumoniae demonstrated by agar diffusion testing; tensile strength varies with formulation (21.7–34.0 MPa).	Hygiene-relevant composite concepts for contact-use environments requiring both mechanical integrity and antimicrobial screening.	[340]

. These systems illustrate how interfacial engineering, nanofiller incorporation, and controlled biofunctionalization can be strategically combined to achieve the necessary balance of structural integrity, sterilization compatibility, and biological safety for clinical use. Continued progress will depend on rigorous long-term cytotoxicity assessments, quantitative modeling of degradation and resorption kinetics, and regulatory validation for both load-bearing and fully resorbable biomedical devices.

7.5. Environmental Remediation and Cleanup Applications

For industrial water treatment and environmental remediation, several natural-fiber and biopolymer composites have demonstrated strong potential as sustainable adsorbents and sorbents. A composite based on chitosan, hydroxyapatite, and hide-derived collagenous proteins [341] incorporates hydroxyapatite particles within a collagen–chitosan matrix and effectively removes both methylene blue (MB) and sunset yellow (SY) dyes from aqueous solutions. MB uptake is maximized in strongly basic media, consistent with its cationic character, whereas SY uptake peaks near pH 3 due to its anionic nature. When applied to real tannery effluent, 100 mg of the composite was sufficient to decolorize a 100 mL batch, and regeneration with an ethanol–water mixture caused no significant loss in adsorption efficiency. Its broad pH versatility, durability, and reliance on readily available biomaterials highlight the viability of such biocomposites for remediation applications that demand both selectivity and sustainability.

Other biopolymeric systems have shown similar promise. Carboxymethylcellulose–citric-acid hydrogels and clay-impregnated cellulose networks have exhibited high organic-effluent uptake [292]. Chitosan has also served as an effective substrate for zinc oxide (ZnO) photocatalysts, enhancing photodegradation efficiency for organic contaminants [342]. Recent studies confirm the versatility of lignocellulosic and protein-based adsorbents across both cationic and anionic dye classes [343,344].

In carbon-management applications, biocomposite films generally exhibit superior CO₂ barrier and sorption behavior compared with traditional synthetic films [290]. A recently reported hybrid carbon–TiO₂–CaCO₃ biocomposite demonstrated CO₂ uptake capacities up to 1.76 mmol g⁻¹ at 273 K [345], suggesting potential for bio-based carbon-capture systems that, when integrated with carbon-negative construction materials, could contribute to global climate-mitigation goals.

For oil and hydrocarbon spill remediation, cross-linked cellulose–PVA networks have been shown to effectively break oil–water emulsions due to their high cross-link density, specific surface area, and low density [292]. Mats composed of plant fibers and animal hair have also demonstrated efficiency in aggregating oil droplets for subsequent removal [346]. Bio-based blends incorporating zeolites, hydrocarbon-oxidizing bacteria, and organic sorbents achieved oil-biodegradation efficiencies between 75 and 98% [347]. Irtiseva et al. [348] reported a vulcanized rubber–peat–cenosphere composite capable of absorbing up to 1.55 g g⁻¹ of diesel fuel. Because cenospheres are fossil-derived by-products, future formulations could incorporate biosourced hydrophobic microspheres that retain similar buoyancy and sorption performance.

For oil and hydrocarbon spill remediation, cross-linked cellulose–PVA networks have been shown to effectively destabilize oil–water emulsions, aided by their high cross-link density, large specific surface area, and low density [292]. Mats composed of plant fibers and animal hair similarly facilitate oil droplet aggregation and removal [346]. Bio-based blends incorporating zeolites, hydrocarbon-oxidizing bacteria, and organic sorbents achieve biodegradation efficiencies of 75–98% [347]. A vulcanized rubber–peat–cenosphere composite has also been reported to absorb up to 1.55 g g⁻¹ of diesel fuel [348]. Because cenospheres are fossil-derived by-products, future formulations may instead employ biosourced hydrophobic microspheres that provide comparable buoyancy and sorption capacity.

8. Conclusions and Outlook

This review demonstrates that, across material classes and application domains, the fiber–matrix interface functions as a central control point through which mechanical performance, environmental durability, processing robustness, and end-of-life behavior are intrinsically coupled. As understanding of these interdependencies has matured, natural-fiber–reinforced composites have progressed from niche materials toward credible alternatives for structural and semi-structural applications in transportation, construction, packaging, and selected aerospace and marine contexts. Their appeal lies in the combination of renewable feedstocks, low embodied energy, and functional mechanical performance within material systems increasingly aligned with circular-economy objectives. Over the past decade, substantial advances have clarified the structure–property relationships governing stiffness, toughness, moisture resistance, and long-term stability, with fiber chemistry, interfacial modification, and processing pathway acting in concert. While plant fibers continue to dominate due to their high cellulose content and favorable stiffness-to-weight ratios, animal- and fungal-derived reinforcements are expanding the design space through distinctive elasticity, bioactivity, and self-bonding architectures.

Despite this progress, several barriers continue to limit widespread industrial adoption. Reliable interfacial bonding under fluctuating humidity and temperature remains difficult to achieve, particularly in high-performance bio-based matrices with narrow thermal processing windows. At scale, variability in natural fiber quality, composition, and moisture content introduces property scatter that complicates process control, certification, and design allowables. These challenges are compounded by incomplete lifetime prediction under coupled mechanical and environmental loading, limited compatibility with high-throughput manufacturing and automation, and unresolved trade-offs between durability, recyclability, and controlled biodegradation. Nevertheless, hybrid composite concepts, hierarchical reinforcement strategies, and advances in solvent-free and additive manufacturing are beginning to bridge natural microstructures with engineered performance, offering pathways toward materials with improved predictability and environmental resilience.

Looking forward, the central challenge for next-generation natural-fiber biocomposites is the deliberate co-design of mechanical performance, environmental durability, and end-of-life behavior rather than treating these attributes as competing objectives. Achieving this integration will require closer alignment between molecular-level interfacial understanding, scalable processing science, and life-cycle assessment methodologies that inform material selection and design decisions at early stages. Continued progress will depend on predictive models of fiber–matrix interactions, standardized durability and biodegradation testing protocols, and harmonized reporting metrics that enable meaningful comparison across material systems. As advances in fiber preprocessing, moisture management, interfacial engineering, and digital process control converge with industrial-scale extrusion, compression molding, and additive manufacturing, natural-fiber composites are poised to transition from laboratory innovation to mainstream engineering solutions, enabling high-performance, low-impact materials tailored for realistic, near-term engineering deployment.