Greening Fused Deposition Modeling: A Critical Review of Plant Fiber-Reinforced PLA-Based 3D-Printed Biocomposites

Highlights

• Identification of gaps in FDM biocomposite research, post-critical analysis of scientific literature, and proposed targeted qualitative and quantitative approaches to accelerate progress in this area
• Proposal of a classification framework for PLA–cellulosic FDM 3DP biocomposites based on reinforcement form and biocomposite filament-processing equipment
• Compilation of diverse processing conditions and their effects on biocomposite filament and 3D-printed biocomposite structure–property relationships to guide future research

Abstract

Fused deposition modeling (FDM) 3D printing (3DP) of PLA biocomposites reinforced with plant-derived cellulosic fibrous materials, including spun yarn, microcrystalline, microfibrillar, nanofibrillar cellulose, and cellulose nanocrystals, offers an environmentally sustainable solution to the mechanical limitations of polymer-only printed materials. Micron- and submicron-scale cellulosic fibers are valued for their renewability, non-toxicity, high surface area, and favorable elastic and specific moduli; notably, micron-scale reinforcements are particularly attractive due to their ease of large-scale industrial production and commercial viability. Similarly, PLA benefits from large-scale production, contributes to CO₂ sequestration through its raw material precursors, and requires less energy for production than non-biodegradable petroleum-derived polymers. Incorporating these raw materials, each of which offers attractive performance properties, complementary commercial strengths, and environmental benefits, as constituent phases in FDM 3D-printed biocomposites (FDMPBs) can further enhance the environmental responsiveness of an already low-waste FDM 3DP technology. Inspired by these compelling advantages, this paper critically reviews research on FDMPB with cellulosic reinforcements in a PLA matrix, uniquely categorizing studies based on the form of cellulosic reinforcement and its impact on the biocomposite’s structure and mechanical performance. Additionally, the review covers biocomposite filament production methods and the equipment involved, presenting an alternative framework for cataloging FDMPB research. A comprehensive literature analysis reveals that the wide variation in feedstocks, fiber–matrix compounding methods, equipment, and processing parameters used in filament production and 3DP complicates the comparison of FDMPB mechanical properties across studies, often resulting in conflicting outcomes. Key processing parameters have been compiled to bridge this gap and offer a more nuanced understanding of the cause-and-effect relationships governing biocomposite properties. Finally, targeted recommendations for future research on developing FDMPB with a PLA matrix and micron-scale cellulosic reinforcements are provided, addressing the knowledge gaps and challenges highlighted in the peer-reviewed literature.

1. Introduction

Initially viewed as a rapid prototyping tool, three-dimensional printing (3DP) is now recognized for its potential in customized fabrication, offering unparalleled design flexibility by enabling the production of complex geometries unattainable with traditional subtractive methods [1]. Opinions diverge on its future role; while some view 3DP as ancillary, and not having a driving role in the technological milieu of the fourth industrial revolution [2], others foresee its exponential growth and widespread use in large-scale production. The current state of 3DP mirrors the early personal computer era, with increasing utilization despite processing limitations [3]. The democratization of 3DP technologies such as fused deposition modeling (FDM), spurred by lapsing patents and affordable open-source printers, is accelerating adoption among businesses and hobbyists alike [4,5], making it the world’s most prevalent 3DP technology. Additionally, recent geopolitical upheavals and health crises underscore the need to revive domestic manufacturing in developed economies [6], with 3DP offering solutions to high labor costs and supply chain challenges.

Despite challenges in large-scale production due to slow printing speeds, 3DP is gaining traction in the medical, automotive, aerospace, and defense industries. In bioengineering, for instance, 3DP enables faster and more customizable production of body implants than traditional methods, saving time and resources. However, due to interlayer porosities or voids from the sequential layering production methodology of 3DP, the inferior mechanical properties of 3D-printed objects remain a constraint on these objects’ widespread adoption in some functional applications [7]. To overcome this limitation, 3D-printed composites show improved mechanical performance compared to polymer-only prints, addressing some concerns [8,9]. However, while 3D-printed composites mitigate mechanical shortcomings, the reliance on non-biodegradable polymer feedstocks in 3D-printed composites poses challenges for end-of-life disposal [10,11]. Similarly problematic, biopolymeric matrices offer an environmentally friendly alternative, but their use as a printing mono-material in 3DP suffers from the same issues of inferior mechanical performance as non-biodegradable polymer-only prints, and they are also more expensive than petroleum-based thermoplastic matrices.

The competing demands of mechanical performance, sustainability, and cost can be reconciled by supplementing biopolymeric matrices with relatively inexpensive biorenewable fibrous reinforcement that not only offsets the resulting material costs compared to biopolymer-only printed objects but also enhances the mechanical properties of the final product, thereby making sustainable solutions more economically viable. This makes FDM 3D-printed biocomposites (FDMPBs) essential for integrating FDM 3DP more firmly within the framework of green technologies.

Like other composite materials, FDMPBs comprise matrix and reinforcement phases, with the distinction that these phases are derived exclusively from biorenewable materials and are amenable to degradation at the end of the product’s service life. In practice, the matrix and reinforcement are joined thoroughly under mechanical shear or agitation in one or more melt-compounding steps to prepare feedstocks for FDM 3DP. These feedstocks are then melted and deposited raster by raster, layer by layer during FDM 3DP to 3D-print the FDMPB. Figure 1 provides a schematic representation of these processes to aid the reader’s understanding.

With the manufacturing processes and sustainability challenges of FDM 3DP and its composites now outlined, we now turn to a focused examination of the ideal fibrous reinforcement and matrix materials for FDMPBs. PLA, derived from renewable biomass, is the most widely produced biopolymer and a popular material for FDM 3DP. While PLA boasts a high tensile strength and modulus, its low toughness necessitates strategies such as chain orientation, blending, or copolymerization to broaden its applications [12]. In contrast, cellulose, nature’s most abundant biopolymer, offers high stiffness, tensile strength, and crystallinity, making it an excellent renewable and sustainable reinforcing element for FDMPBs [13].

Moreover, PLA’s low glass transition temperature (approximately 60 °C) and melting temperature range (130–180 °C) enable processing at relatively low temperatures, which helps protect thermally sensitive cellulosic fibers (degradation temperature: 200–350 °C) and minimizes their degradation during melt compounding and/or filament extrusion and 3DP, also making it an attractive matrix material for FDMPBs.

PLA exhibits slow crystallization behavior, which can be altered by adding cellulosic fibers as crystal nucleation sites that speed up its crystallization [14]. Slow crystallization adversely impacts the printed material’s dimensional accuracy and stability due to delayed solidification of the deposited extrudate [15]. Slow crystallization of PLA and attached delay in printed PLA solidification would further reduce the 3D printing speeds of an inherently slow printing process. The required mechanical performance, contingent upon the optimum polymer crystallinity, might not be achievable in the case of slow-crystallizing polymers, as the optimum crystallinity levels would not be reached in the printed parts. This would necessitate post-processing steps like thermal aging to arrive at the optimum crystallinity levels and prolong processing times. Additionally, PLA’s slow degradation rate is a concern that blemishes its credentials as an environmentally friendly polymer. Incorporating cellulosic reinforcement accelerates PLA degradation in biocomposites [16].

Overall, the synergistic properties of PLA and cellulosic fibers make them an ideal pairing for FDMPBs. However, hydrophilic cellulose’s low compatibility with hydrophobic PLA is a major challenge. Increasing the contact surface of cellulosic reinforcement by going into micro- or nano-scales can enhance the interaction between PLA and the reinforcement thereby improving mechanical performance [17,18]. While nano-scale cellulose fibers are costly and mostly confined to lab-scale production, micron-scale cellulose fibers, which are already extensively used in the pharmaceutical industry [19,20], are economically viable to produce on an industrial scale.

Therefore, developing PLA/micron-scale cellulose fiber-reinforced 3D-printed biocomposites that combine these materials not only leverages the mechanical and processing advantages of PLA and cellulose but also fuels the demand for eco-friendly products and promotes the usage of sustainable materials.

Drawing inspiration from the points discussed above, this work examines advancements in FDMPB research, focusing exclusively on using plant-based cellulosic fibers, ranging from micron-scale fibers to macro-scale yarns as reinforcements, with PLA as the matrix material. Initially, the review details the structure and synthesis of micro-cellulosic materials, discussing cellulose’s molecular organization, its hierarchical assembly from elementary fibrils to macrofibrils, and the engineering processes used to extract or “engineer” micron-scale reinforcements from native cellulose. These subjects are covered to shed light on the native cellulosic structure and the properties that can be expected to arise from this structure. Furthermore, “engineering” specific cellulosic structures as reinforcement candidates in PLA-based FDMPBs requires an intimate understanding of which structural constituents must be retained, which ones must be eliminated, and which processing routes are needed to achieve the desired micro-cellulosic structures. Following this, the review examines PLA comprehensively, including its synthesis via ring-opening polymerization, structural features, and degradation behavior under various environmental conditions. After developing an understanding of cellulose, PLA chemical structures, and key chemical moieties in these structures, this work explores the chemical interactions between PLA and cellulosic fibers, emphasizing interphasic adhesion mechanisms such as hydrogen bonding and van der Waals forces that are critical to the structure and performance of FDMPBs. By integrating knowledge of the cellulosic native and engineered structures, PLA chemistry, and PLA–cellulose interfacial interactions, researchers can optimize the mechanical and thermal performance of FDMPBs. The intrinsic properties of cellulosic reinforcements, PLA matrix characteristics, and filament processing techniques are interdependent in fabricating robust PLA-based cellulosic fiber-reinforced FDMPB composites, emphasizing the need for a holistic approach in composite design. This integrated perspective informs the development of effective strategies to enhance interfacial bonding and overall FDMPB durability while overcoming challenges in process optimization, sustainability, and interface compatibility. The discussion then transitions to introducing the FDM technology and the centrality of filaments in the 3DP of FDMPBs. A novel framework for classifying the FDMPB research is offered based on the methods and equipment used to produce 3D-printable biocomposite filaments, explicitly linking these production processes to key challenges in FDMPB development, such as process optimization and sustainability. Furthermore, FDMPB research is systematically categorized by cellulosic reinforcement types and forms, evaluating their influence on the microstructure and mechanical properties of the biocomposites. Finally, this critical analysis highlights existing research gaps by synthesizing insights from diverse academic sources while suggesting future directions for developing PLA and cellulosic fiber-based FDMPBs.

2. Micro-Cellulosic Reinforcements and PLA Matrix: Structure, Synthesis, and Interactions

2.1. Cellulose Structure and Micron-Scale Cellulose Synthesis

2.1.1. Structure

Understanding cellulosic molecular and hierarchical structure is the key to predicting the properties that can be derived from it and the nature of its interactions with the PLA matrix in FDMPBs. Moreover, an intimate understanding of this structure also aids in engineering cellulosic materials and adopting the most impactful techniques for their engineering.

Different cellulosic allomorphic structures exhibit varying levels of crystallinity, which influences the reinforcement phase’s stiffening potential within the polymeric matrix. More amorphous celluloses have a higher affinity for water, which can potentially weaken their interactions with the non-polar PLA matrix in FDMPBs. Cellulose (structural formula: (C6H10O5)n), composed of linear polysaccharide chains made up of glucose monomers covalently linked by β(1→4) glycosidic bonds, forms two hydrogen bonds between neighboring glucose units [21], as demonstrated in Figure 2. With a non-reducing hydroxyl group at one end and a reducing hydroxyl group at the other, the latter hydroxyl group is amenable to chemical modifications for cellulose’s property adjustments [22]. Equatorial hydroxyl (-OH) groups extending radially from the glucopyranose ring are responsible for cellulose’s hydrophilic nature and intermolecular hydrogen bonding [23].

Elementary Fibrils, Microfibrils, and Macrofibrils

Glucan chains in cellulose self-organize into elementary fibrils, or microfibrils, through hydrogen bonding and van der Waals forces [25,26,27,28]. Due to their high surface energy, these microfibrils can aggregate into larger structures known as macrofibrils [29,30]. The structural organization of cellulose from the molecular level to elementary fibrils across different dimensional scales is illustrated in Figure 3.

A proposed model suggests that an elementary fibril consists of a crystalline cellulose core encased in sub-crystalline and non-crystalline layers [26]. A microfibril is a cellulosic structure comprising a single elementary fibril coated with polysaccharides such as hemicelluloses and pectin. Figure 4 accurately illustrates the elementary fibril, microfibril, and macrofibril structures according to this model.

During the expansion of the plant’s primary cell wall, elementary fibrils detach from the distal ends of macrofibrils and become coated by polysaccharides, forming microfibrils. Due to conformational differences, the crystallite sizes of elementary fibrils within macrofibrils are larger than those of detached microfibrils [26]. Some perspectives suggest that microfibrils comprise multiple elementary fibrils rather than a single crystalline core [31]. Resolving this discrepancy may involve recognizing the interchangeable terminology used for similar cellulosic structures. For instance, “microfibril” and “macrofibril” may refer to the same structure in different studies [26,31]. Similarly, “elementary fibrils” may be synonymous with “microfibrils”, while “macrofibrils” might describe bundles of microfibrils in some studies [32]. In the secondary plant cell wall, elementary or microfibrils align parallelly, individually, or in aggregated formations without splitting from macrofibrils, characteristic of the secondary cell wall after cell growth cessation [27].

2.1.2. Micro-Cellulosics Synthesis

A well-developed understanding of native cellulose structure guides the engineered micro-cellulose extraction process. In turn, the physical, mechanical, and thermal properties of micro-cellulosic reinforcements are deeply linked to their engineered structures. Therefore, investigating micro-cellulosic fiber extraction processes is important to predicting how these reinforcing media influence the performance of micron-scale cellulosic fiber-reinforced PLA-based FDMPBs.

The scientific literature needs a well-defined pathway delineating the progression from naturally occurring nanoscale fibrillar structures, such as elementary fibrils, microfibrils, and macrofibrils, to larger engineered or extracted micron-scale fibrillar structures. This gap in the literature may be attributed to the variability in the dimensions of these engineered structures, which makes their characterization challenging [33]. The following sections describe the processes for engineering or extracting micron-scale fibers with high cellulose crystallinity and low extra cellulosic content.

Microfibrillar Cellulose

The degree of nano-scale fibrillation in cellulosic materials depends on the extent of homogenization, which is largely influenced by financial and energy inputs. Historically, Turbak et al. and Herrick et al. laid the foundation for microfibrillar cellulose (MFC) development by demonstrating that passing a wood pulp–water suspension through narrow homogenizer orifices generates immense shear and turbulence, breaking cellulosic fibers into engineered microfibrillar structures [34].

The extraction processes for microfibrils invariably produce a heterogeneous mixture of particles within an interconnected network comprising nanofibrils and micron-scale fibrillar bundles [33]. A high degree of homogenization results in nanocellulose being the primary component of microfibrillar cellulose (MFC) [33]. Mechanical extraction methods such as grinding, cryo-crushing, and intense ultrasonic treatment can extract MFC [35] but yield a polydisperse product containing both micro- and nanofibrils [36]. These methods are energy-intensive, but combining them with chemical or enzymatic pretreatments can optimize energy use [34]. Thus, MFC extraction generally involves a two-step process: initial chemical or enzymatic preprocessing to dissolve lignin and hemicellulose, followed by mechanical fibrillation [37,38]. In fact, chemical pretreatment to remove or reduce extra-cellulosic components facilitates the breakdown of fiber diameters from micron scales down to nanometer scales even during the refining pretreatment. Conversely, the presence of high lignin content, as seen in untreated thermomechanical pulp fibers, can inhibit this diameter reduction, preventing fibers from reaching the nano scale even after homogenization [39].

Recent advances include TEMPO-mediated oxidation, which weakens intermolecular bonds between cellulosic chains to promote defibrillation and reduce energy expenditure during fibrillation. In this process, the treated cellulosic substrate undergoes cryogenic treatment and multiple crushing steps via twin-screw extrusion, resulting in an MFC product with a decline in crystallinity from 68.93% to 59.2%, as measured by X-ray diffraction [40].

Mechanical fibrillation itself generally consists of two steps. The first step pertains to refining the cellulosic pulps, during which the fiber size is decreased, some fibrous structure is broken into fibrils, and a consistent fiber slurry or water suspension is often formed. The second step involves subjecting the refined substrate to intense shear or impact stresses during cryo-crushing, grinding, or low-pressure homogenization. A pure two-step mechanical refining approach has also been reported in the literature, wherein the first step employs beater rolls and plates of a valley beater to weaken or sever the intermolecular connections in water-soaked bleached softwood kraft pulp, encouraging initial defibrillation. The second step uses a ball mill at a speed of 500 revolutions per minute for 2 h to dismantle the remaining hydrogen and van der Waals bonds, promoting greater microfibrillation and reducing the fibril diameter [41]. Moreover, the reduction in the size of the raw material used for MFC extraction can ease the individualization of the fibrillated fibers, as intermolecular forces make it difficult for large-sized raw materials to undergo complete individualization during the fibrillation process. For this purpose, microcrystalline cellulose (MCC) becomes an attractive raw material for MFC extraction through mechanical grinding. Utilizing never-dried MCC reduces the energy input required to extract MFC via mechanical grinding. Drying MCC causes hornification, which removes water from hydroxyl groups and forms additional hydrogen bonds between cellulosic chains. This ultimately makes cellulose fibrillation into single fibrils more challenging [42]. Typically, sodium hydroxide alkaline pretreatment disrupts the hydrogen bonding between cellulosic chains to aid mechanical fibrillation and reduce the energy expenditure of the fibrillation process. Recently, calcium hydroxide has been used to aid mechanical grinding for MFC production due to its low corrosiveness, which may limit damage to the crystalline cellulosic structure [43]. Similarly, water too can break hydrogen bonds and swell cellulosic structures while making a fiber–water suspension. For instance, one approach used a disintegrator machine for initial fiber size reduction, followed by mechanical shearing at up to 125,000 revolutions in a PFI refiner. The refining step fibrillated the cellulose but did not individualize the fibrils. Therefore, cryogenic crushing was carried out to separate microfibrils from the fiber cell wall. The output from the refining process was divided into two batches: one converted into a water suspension and treated in a disintegrator for 15 min and another freeze-dried at −50 °C, ground in a Wiley mill, and passed through a 60-mesh sieve [44].

Microcrystalline Cellulose

• Overview

Micron-scale microcrystalline cellulose (MCC) is a highly crystalline cellulosic material with varying-sized hydrogen-bonded microfibrillar aggregates [45]. Its extraction from fiber feedstock typically involves a two-step process: pretreatment of the cellulosic fibers to improve accessibility to the MCC followed by acid hydrolysis. These steps aim to remove hemicellulose and lignin biopolymer matrices surrounding cellulose in the plant fiber [46,47]. Figure 5 displays the transformation of the cellulosic substrate after various processing steps to obtain MCC from different origins. Cellulose nanocrystals (CNCs) and MCC are extracted using a similar two-step methodology. The primary distinction between their extraction lies in isolating nano-scale cellulose crystalline regions from the supra-cellulosic structure by eliminating the amorphous cellulosic domains during acid hydrolysis in the case of CNCs [48,49].

Pretreatment Methods

Physical pretreatment methods enhance the accessibility of MCC, boost the substrate’s surface energy by expanding its surface area, remove extra-cellulosic materials (lignin, hemicellulose, pectin) from the substrate, and minimize particle sizes without using chemical agents. These methods include mechanical, hydrothermal, pressure-induced fiber explosion, and ultrasonic irradiation.

Ball milling is a mechanical pretreatment method in which the hydrogen bonds between cellulosic chains are broken through the shearing action of solid balls hitting the cellulosic material and the size of the cellulosic substrate is reduced. This process also disrupts the crystalline structure of the cellulose, making the overall structure more responsive to subsequent treatments [50].

The simplest forms of hydrothermal pretreatments are autocatalyzed and conducted at temperatures ranging between 150 and 230 °C. These autocatalyzed hydrothermal pretreatments are responsible for the partial hydrolysis of hemicellulose and fragmentation of lignin from plant material. Acidic catalysts lower the pH during hydrothermal pretreatment, leading to the comprehensive dissociation of hemicellulose into monosaccharides, whereas an alkaline pH dissolves lignin by protonating phenolic hydroxyl groups into phenolate ions, with some hemicellulose also being affected [51].

The most common pressure-induced pretreatment for lignocellulosic materials is steam explosion. The detailed mechanisms of steam explosion of lignocellulosic biomasses are poorly understood. The process can be divided into two stages. In the first stage, under high pressure (5–45 bar), the sub-critical water (160–260 °C) diffuses into the porous lignocellulosic biomass structure, causing hydrolysis, deacetylation, and depolymerization of hemicellulose, while cleaving structural bonds in lignin and reducing the degree of polymerization in cellulose. In the second stage, a sudden drop in pressure causes the absorbed water to revert to a gaseous state, converting thermal energy into mechanical energy and causing shear-induced fissures and fibrillation in the cellulosic structure [52].

Ultrasonic waves create regions of high and low pressure, leading to the formation of bubbles or cavities; when these collapse, they generate localized very high temperature and pressure areas. This process may loosen the lignocellulosic structure and enable penetration of de-lignifying agents like hydrogen peroxide depolymerizing lignin. However, the duration of such “ultrasonic–chemical” treatment must be optimized to prevent excessive hydrolysis of the cellulosic fraction [53].

Chemical pretreatment is the second method for modifying the lignocellulosic biomass to increase access to the cellulose for the next processing steps. Mainstream chemical pretreatments include acidic, alkaline, and ozonolysis methods. Acidic pretreatment is typically carried out in the presence of diluted or concentrated (0.2–2.5% w/w) inorganic mineral acids like sulfuric, hydrochloric, phosphoric, and nitric acids under elevated temperature conditions ranging from 130 °C to 210 °C [54]. Dilute acids, when coupled with high temperatures, are generally used for hemicellulose removal, while strong acids operate at lower temperatures, cleaving glycosidic bonds in hemicellulose and solubilizing, then condensing, lignin into larger fragments [55]. This pretreatment method is most effective for hemicellulose removal, but lignin removal rates depend on treatment conditions like temperature, acidic concentration, and pretreatment duration. The corrosive nature of treating necessitates reactor vessels, and the treated substrate requires extensive washing to neutralize the acidity and remove the condensed lignin, a process that poses sustainability challenges because of the large volumes of water used and contaminated post washing. Consequently, there is growing interest in using weak organic acids (formic, citric, maleic, oxalic), which are more sustainable and less corrosive, although they may exhibit slower hydrolysis rates due to the lower hydrogen ion concentrations [56]. Alkaline agents like sodium, potassium, calcium, and ammonium hydroxides sever the ester linkages connecting lignin and hemicellulose to cellulose in saponification [54,57]. This saponification process also effectively breaks the ester linkages between lignin and hemicellulose, rendering lignin soluble. Alkaline pretreatment is regarded as the most capable technique for lignin removal from lignocellulosic biomass [58], and it also hydrolyzes hemicellulose by breaking down its functional groups [59]. A study on rice straw lignocellulosic biomass processing has shown that 1% sodium hydroxide treatment for 3 h rationalized hemicellulose and lignin contents without damaging the cellulose [60]. Ozonolysis and ionic liquid chemical pretreatments are not widely adopted chemical pretreatment techniques. Electron-deficient ozone selectively attacks the lignin due to the presence of high-electron-density phenolic and aromatic compounds. This breaks these compounds into smaller, water-soluble compounds. Ozonolysis’s application as a chemical pretreatment is limited due to the high costs involved in running this process [61].

Bleaching chemical pretreatment is often carried out after other pretreatments to remove residual lignin from pretreated cellulose to boost access to cellulose during the MCC extraction step and increase the process yield. Cellulose is also discolored in this process. The most commonly used bleaching agents in the literature are sodium hypochlorite and hydrogen peroxide.

Acid Hydrolysis and MCC Extraction

The MCC extraction process can begin when cellulose has been pretreated and made more accessible for extraction. Different synthesis routes have been adopted in the literature, but acid hydrolysis is the most widely reported route. Recent research has focused on nitric acid hydrolysis for MCC extraction, as it provides a superior yield compared to sulfuric and hydrochloric acids, primarily due to its efficiency in hydrolyzing amorphous cellulose and producing fewer harmful byproducts. However, the percent crystallinity of the extracted MCC, as observed by X-ray diffraction, is lower than that obtained via hydrochloric acid [62]. For example, MCC was obtained by continuously stirring wheat straw cellulose pretreated with sodium hydroxide alkali and sodium hypochlorite bleach in 140 mL of 0.5 M nitric acid at 80 °C for 30 min. The hydrolyzed cellulose was repeatedly washed until neutralized, dried, and ground into a fine powder using a rotary ball mill [62]. Similarly, pedicle fibers were alkali pretreated and bleached to prepare them for MCC extraction. Four g of bleached fibers were combined with 10 mL of water in a flask and refluxed with 160 mL of 40% sulfuric acid at 85–95 °C for 40 to 50 min. The reaction was halted by removing the heat and adding distilled water when white cloudiness was observed. After washing, the hydrolyzed MCC was centrifuged, sonicated, and freeze-dried. X-ray diffraction analysis revealed that the crystallinity percentage increased from 40.43% after alkali pretreatment to 64.53% after acid hydrolysis to extract MCC [63]. Teff grass has also been investigated as a cellulosic source for MCC extraction. The pretreatment of teff grass was performed in three steps using multiple reagents. First, it was boiled with acetic and formic acids in a water bath at 90 °C for 90 min. Next, de-lignification was achieved by treating the grass with acetic acid, formic acid, and hydrogen peroxide, followed by bleaching with hydrogen peroxide for 1 h. Finally, the pretreated substrate was boiled for 30 min in the presence of 2.5 M hydrochloric acid (with a 1:20 fiber-to-liquor ratio) to hydrolyze the material and extract MCC. The MCC was then water-washed and treated with an ammonium hydroxide solution to neutralize all acidity. Three different MCC samples were subsequently prepared from the milled MCC: oven-dried, as a slurry in water, and spray-dried, with the oven-dried MCC exhibiting the highest crystallinity, at 85% [64]. MCC extraction from coffee husks consists of removing waxes and tannins as the first pretreatment step, followed by two-time sodium hydroxide alkali treatment for the removal of hemicellulose and some amount of lignin. Acid hydrolysis of the pretreated coffee hull substrates was performed in 7% sulfuric acid with a 1:20 fiber-to-liquid ratio (g/mL) under stirring at room temperature. In this case, bleaching with hydrogen peroxide for de-lignification was done after acid hydrolysis. The pH of the bleaching solution was increased to 11 by 8% sodium hydroxide inclusion. All treatment steps were repeated twice. The crystallinity of the untreated, dewaxed, alkali-treated, and bleached coffee husk samples after different treatments was 69.7%, 76.2%, 81.2%, and 89.9%, respectively [65]. Another study has compared the properties (microstructural, morphological, and thermal) of alkali-treated fibers, bleached fibers, and acid-hydrolyzed MCC, all derived from washingtonia plant fiber. Alkaline treatment entails treating 10 g of fiber in 600 mL of 2% sodium hydroxide alkali at 80 °C for 2 h to remove lignin. The bleaching treatment follows the alkali treatment, using 600 mL of 2% sodium hypochlorite acidified with 5 mL of acetic acid to remove residual lignin and hemicellulose at 80 °C for 2 h. Finally, the alkali-treated, bleached substrate is hydrolyzed by 2.5 M hydrochloric acid solution for 30 min at 80 °C. The SEM morphological analysis showed ruptured features for alkali-treated fibers due to fiber swelling. Bleached fibers had smooth surfaces due to the removal of lignin and hemicellulose. Acid-hydrolyzed fibers showed fibrillation due to partial depolymerization of amorphous cellulosic domains. Reinforcing the morphological findings was the X-ray diffraction analysis-derived crystallinity of fibers that underwent different treatments. Alkali-treated, bleached, and acid-hydrolyzed MCC had crystallinity values of 51.5%, 69.4%, and 72.6%, respectively. The higher crystallinity of the MCC made them the most thermally stable out of the three washingtonia fiber derivatives [66], which is also confirmed in another similar study [67].

Figure 5. (A1) Raw flax fibers; (A2) flax fibers post chemical pretreatment for removal of biopolymer matrix materials; (A3) MCC extracted from flax fibers; (B1) MCC extracted from cotton linter; (B2) MCC extracted from birch sulfite pulp; (B3) MCC extracted from poplar kraft pulp. The scale bar is 10 μm for all the figures. Reprinted with permission from Springer Nature [68].

Figure 5. (A1) Raw flax fibers; (A2) flax fibers post chemical pretreatment for removal of biopolymer matrix materials; (A3) MCC extracted from flax fibers; (B1) MCC extracted from cotton linter; (B2) MCC extracted from birch sulfite pulp; (B3) MCC extracted from poplar kraft pulp. The scale bar is 10 μm for all the figures. Reprinted with permission from Springer Nature [68].

2.2. PLA Matrix: Structure, Synthesis, and Degradation

Elucidating PLA’s molecular structure is crucial for understanding how it interacts with cellulosic reinforcements and how these interactions impact the mechanical properties and durability of FDMPBs. This understanding also provides insight into PLA’s inherent mechanical properties and degradation mechanisms.

2.2.1. PLA Synthesis

PLA is produced from lactic acid monomers obtained through bacterial fermentation of plant-derived carbohydrates. The type of bacteria used dictates the form of lactic acid, either as L- or D-lactic acid enantiomers [69]. These monomers are then converted into lactide stereoisomers L-lactide, D-lactide, or meso-lactide [70]. The predominant method for PLA production is ring-opening polymerization, which effectively polymerizes the lactide stereoisomers in the presence of organometallic catalytic compounds like stannous octoate to produce high-molecular-weight PLA [69,71,72]. During the ring-opening polymerization, stannous octoate, when added with alcohol as an initiator, forms a reactive complex with the alcohol’s hydroxyl group. This reactive complex interacts with the lactides and the alcohol’s oxygen, acting as a nucleophile, attacking the lactide’s ester group, resulting in double-bonded oxygen in the lactide losing its double bond. With the loss of double-bonded oxygen, the lactide ring is opened and the opened lactide units begin arranging themselves into a growing polylactic acid chain [71,72]. The chemical structures of lactic acid, lactides, and PLA, both before and after polymerization, are depicted in Figure 6, while Figure 7 illustrates the ring-opening polymerization of PLA via the coordination-insertion mechanism of lactide with an organometallic stannous octoate catalyst.

2.2.2. PLA Structure

PLA (structural formula: (C₃H₄O₂)n) is an aliphatic polyester characterized by monomeric units linked by ester bonds [69]. As is discernible in Figure 6a, PLA lacks any functional side-chain groups [73]. The PLA chain terminates with reactive carboxyl (-COOH) and hydroxyl (-OH) groups. Depending on the lactide used in polymerization, the microstructure of PLA can be tailored by adjusting the L- and D-contents; a higher L-content (>90%) produces a more semi-crystalline PLA, while increased D-content results in a more amorphous polymer and a lower melting temperature [74,75].

2.2.3. Breakdown of PLA Structure

PLA’s degradation process dictates its service life and reveals its vulnerabilities to specific environmental conditions, such as exposure to heat, pH, moisture, UV radiation, and microbial activity. The PLA disintegration kinetics also depend on material factors such as its molecular weight distribution, microstructure, and the surface area exposed to degradation processes. Understanding how PLA disintegrates helps predict the potential applications of PLA-based FDMPBs and informs waste management strategies at the end of their lifecycle.

Typically, PLA degradation follows a two-stage process, with UV, pH-led, or thermochemical hydrolysis serving as a precursor to the biodegradation of PLA. During hydrolysis, ester linkages in PLA’s backbone are cleaved, converting high-molecular-weight PLA fractions into low-molecular-weight ones that microbes can digest via second-stage enzymatic activity [76,77,78]. Crystalline domains in PLA exhibit better resistance to hydrolytic degradation compared to amorphous domains when exposed to hydrolytic agents like water [79]. Above the glass transition temperature, water diffusion into the thermally mobile polymer chains accelerates this process. High-humidity environments (relative humidity >60%) also increase PLA’s hydrolytic degradation. Thermal degradation raises the number of terminal hydrophilic -COOH groups, enhancing water uptake and further facilitating hydrolytic degradation [69,80].

2.3. PLA–Cellulose Interaction

The interaction between its two formative phases is the centerpiece predictor of how FDMPBs behave under various mechanical or thermal stimuli. Interphasic adhesion in PLA–cellulose biocomposites relies on bond strengths and the molecular surface roughness of polymeric chains, with increased roughness enhancing hydrogen bonding sites between PLA and cellulose [81]. Hydrogen bonds are the key to PLA–cellulose interactions. In PLA–cellulose biocomposites, polar carbonyl (C=O) and terminal -OH groups from PLA form hydrogen bonds with cellulose’s polar -OH groups [82]. Additionally, PLA’s double-bonded oxygen atom engages in weak van der Waals interactions with cellulose’s nonpolar (C-H) part. Sufficient van der Waals effects in cellulosic crystalline plains can yield adhesion quality comparable to that achieved through polar moiety interaction-originated hydrogen bonds between PLA and cellulose [81]. Molecular dynamics simulations reveal a low propensity for cellulose -OH groups to form hydrogen bonds with PLA’s C=O groups, a trend that intensifies with increased cellulosic content in the composite. Conversely, as the cellulose concentration rises, its polar -OH groups show a higher affinity for PLA’s terminal -OH group, leading to more frequent hydrogen bond formation [82]. Despite undergoing numerous melt extrusion cycles that typically reduce molecular weight, recycled PLA–cellulose biocomposites have demonstrated the preservation of specific mechanical properties. This observation may suggest an enhancement in hydrogen bonding interactions between the composite phases and increased -OH end groups within the PLA. However, this phenomenon has been primarily explained in terms of improved polymer crystallinity, cellulosic reinforcement disaggregation, and dispersibility in the literature [83,84,85].

3. Fused Deposition Modeling

3.1. Historical Development and Prospects

In 1989, Scot Crump pioneered FDM, co-founding Stratasys, which later patented and commercialized the technology [86]. The 2005 replicating rapid prototyper (RepRap) project helped democratize FDM by enabling the self-replication of 3D printers through open-source designs [87]. RepRap’s replicators employed the exact printing mechanism as FDM; however, due to active FDM patents, the open-source community and RepRap-inspired startups adopted the term “fused filament fabrication” (FFF) to describe the process. The lapse in Stratasys patents post 2005, coupled with the open-source community’s accumulated expertise in hardware design, catalyzed a boom in FDM/FFF 3D printer startups. As more patents expire, developers can access proprietary technologies to improve affordable open-source FDM 3DP hardware, further accelerating its adoption [88]. This trend is also underscored by FDM’s dominant market presence and its leading annual growth rate of 21.15% in the 3DP sector [89,90].

3.2. FDM 3DP Process

In FDM, the nozzle deposits molten or softened thermoplastic extrudate onto the print bed as rasters, arranged side by side, forming a single printed layer. Individual layers are stacked to form a complete object. After each layer is deposited, the print bed lowers, or the printhead rises along the z-axis by the quantum of layer height specified in the slicing software, which generates the G-code (a language that the 3D printer can understand and that carries all instructions necessary to execute a print job). The layer-by-layer object construction results in mechanical anisotropy, favoring performance in the printing directions [91]. This anisotropy results from the limited diffusion of polymer chains between adjacent and stacked rasters, beads, or lines and the alignment of polymer chains in the printing direction. Furthermore, the alignment of reinforcing fibers along the printing direction further accentuates the anisotropic behavior in FDMPBs [92,93].

When rasters encounter the print bed or previously deposited rasters, they spread horizontally, influenced by the extrudate’s viscosity and surface tension. The interaction between the hot, newly extruded raster and the previously laid raster triggers a localized phase change at their interface, and inter-raster fusion occurs [94]. However, complete inter-raster fusion is challenging due to limited surface wetting and low segmental motion in polymeric chains, leading to suboptimal inter-raster linkages compared to intra-raster polymeric chains [95]. The subpar fusing of neighboring rasters produces inter-raster and, by extension, interlayer porosities [96,97]. In FDMPBs, intra-raster porosities can also form due to incompatibility between polar fibrous reinforcements and the nonpolar thermoplastic matrix, causing phase separation, fiber aggregation, and flow instabilities [92].

Heated print beds enhance first-layer adhesion in FDM by supplying the necessary thermal energy to the polymeric systems. Thermal stresses induced during polymer cooling can cause uneven shrinkage rates in semi-crystalline thermoplastics, leading to warpage and potential print failures. Heated print beds address this structural warpage effectively. Additionally, some 3D printers maintain a uniform thermal profile by regulating the temperature of the entire print chamber, thus enhancing the dimensional stability of the printed material [98]. Fibrous additives in the semi-crystalline polymer can also counteract shrinkage [99].

A 3D printer typically uses a rotatable filament spool to supply thermoplastic filament feedstock for FDM, while pallet-based FDM printers, like big area additive manufacturing (BAAM) 3D printers, use thermoplastic pellet feedstock. Drive wheels or gears pull the printing filament into the printhead, acting like a plunger, pushing its melted polymer part from the heating block into the printing nozzle to sustain extrusion [100]. In pellet-based FDM, the 3DP process utilizes a melt extrusion mechanism driven by screw rotation and forward pressure. Figure 8 labels parts of commercially available direct-drive FDM printers. Optimum filament stiffness is crucial to prevent bending and ensure consistent flow, avoiding under- or over-extrusion of polymeric melt during 3DP. Brittle filaments can break, causing print failures, while inconsistencies in filament diameter can lead to filament slippage in the filament feeding gears and polymer melt flow instability [101]. Thus, the filament-making process and its influence on the biocomposite filament’s structural and functional properties is vital for producing high-quality FDMPBs. The following discussion will critically examine various biocomposite filament fabrication techniques and their intended outcomes.

3.3. FDM Biocomposite Filament Production Methods

Some FDM printer manufacturers restrict end users to proprietary filaments and 3DP process settings suitable for these proprietary filaments. In contrast, FDM printers developed by startups inspired by the RepRap movement, leveraging open-source hardware designs and software, typically provide greater flexibility in filament feedstock choices and printing parameters. This openness encourages innovative experimentation with various cellulosic reinforcement types to develop and optimize new cellulosic fiber-reinforced PLA-based biocomposite filaments for FDM 3DP.

One of the key prerequisites for producing 3D-printable cellulosic fiber-reinforced PLA-based biocomposite filaments requires uniform distribution and dispersion of fibers within the thermoplastic PLA matrix before their melt extrusion as biocomposite 3DP filaments. This prevents fiber agglomeration in the molten biocomposite filament during printing, which can clog the 3DP nozzle. Moreover, consistent fiber distribution and dispersion in the biocomposite filament are also essential for maintaining stable rheological properties of the molten extrudate during the FDM 3DP process, preventing melt flow instability caused by uneven fiber distribution and dispersion within the polymer melt. Consequently, attaining homogeneous dispersion of the fibrous reinforcement in the extruded matrix during FDM 3DP largely depends on the quality of the fibrous reinforcement in the biocomposite filament feedstock. Distributive mixing preserves fiber integrity and ensures even distribution, while dispersive mixing breaks down fiber agglomerates [102] but may cause fiber dimensional attrition [103]. Single-stage and double-stage fiber–matrix mixing processes for developing 3D-printable biocomposite filaments can incorporate both distributive and dispersive mixing modes.

In the double-stage biocomposite filament-making process, solution casting is often the first step to enhancing fiber dispersion within a PLA polymeric matrix when the reinforcing fibers are highly prone to clustering due to their high surface energies or loading contents. In this process, a solvent dissolves PLA, disentangling its chains and lowering the system’s viscosity. Subsequent physical agitation allows uniform fiber dispersion in the low-viscosity PLA solution. After solvent evaporation, solid fiber-reinforced films form, which can be crushed and used as feedstock for melt extruded into 3D-printable filaments. Conversely, the single-stage melt compounding process uses heat and mechanical shear to transition the thermoplastic PLA matrix from solid to molten form, enabling the mixing of fiber reinforcement in polymeric matrices. However, the high viscosity of PLA melt complicates fiber dispersion in highly fiber-loaded biocomposites. Therefore, melt compounding is often combined with solution casting in a double-stage process or carried out in two melt extrusion cycles as a double-stage process to optimize fiber dispersion [104].

Table 1. Solution casting vs. melt compounding.

Solution Casting	Melt Compounding	Refs.
Room-temperature processing to produce biocomposite films is possible. Hence, thermal degradation of fibrous reinforcement does not occur.	Less thermally stable fibrous reinforcements are prone to degradation during compounding, especially when subjected to multiple compounding cycles. However, melt-compounded composites exhibit superior thermal stability than solution-cast ones.	[105,106,107]
Biocomposite crystallinity is retained for longer durations in composting conditions in composting conditions, which means these biocomposites are slow to biodegrade.	A quicker rate of decline takes place in both biocomposite crystallinity and biodegradation under composting conditions.	[105]
The mechanical degradation of the reinforcing fiber aspect ratio is prevented.	Reinforcing fibers can lose their higher aspect ratios in the shearing action associated with melt compounding.	[108]
Solution casting can effectively disperse micro and nano-scale cellulosic fibrous reinforcement.	Hydrogen bonding in molecules of cellulosic fibrous reinforcement causes their agglomeration in the polymer matrix during melt compounding processes.	[109,110]
Solvent reclamation is essential for reusing toxic, widely used, costly solvents, necessitating appropriate infrastructure.	No solvent use is required for typical melt compounding operations, making it a cleaner process.	[111]
The scalability of the solution casting process is challenging.	Melt compounding processes are industrially scalable and cheaper manufacturing routes compared to solution casting.	[104,112]
Re-agglomeration of nano-cellulosic reinforcement occurs when a melt processing step is added to produce the biocomposite.	Typically, multiple melt compounding cycles ameliorate the dispersion and distribution of fibrous reinforcement in the polymeric matrix.	[113]

contrasts melt compounding and solution casting techniques in greater detail.

3.4. FDM Filament Melt Processing Equipment

From a sustainability perspective, single-stage and double-stage melt compounding and extrusion should be considered leading manufacturing methods for producing 3D-printable cellulosic fiber-reinforced PLA-based biocomposite filaments, given their scalability and environmental advantages, as depicted in Table 1. The following sections elaborate on the equipment used in melt compounding the fiber–matrix phases and extruding fiber-reinforced PLA-based biocomposite filament feedstock for FDM.

3.4.1. Single-Screw Extruders

Single-screw extruders utilize a flighted screw with progressively shallower depths encased in a metallic barrel. These extruders can have multiple temperature zones where the heat is transmitted from the heating elements to the barrel wall. The extruder channels solid, plasticated, or molten materials forward under screw rotation and pressure build-up. These extruders build up and maintain pressure using shear forces that dominate when the thermoplastic polymer is in a solid state and polymer viscosity when the polymer is in a molten form to extrude the melt [114].

The screw in a single-screw extruder is divided into three zones based on flight depth: the feeding zone, the transition zone, and the metering zone. With the deepest flights, the feeding zone transports pellets or powder from the hopper to the transition zone. The increased flight depth in this zone minimizes excessive shear forces between the solid polymer compacted into a solid bed and the barrel wall, which would otherwise require high torque to overcome. The temperature of the solid bed rises due to friction-stimulated heating and heat transfer from the barrel wall. In the transition zone, flights become progressively shallower, accommodating the melt pool formed as the solid bed melts. With the shallowest flights, the metering zone breaks up any remaining unmolten particles [115,116]. As polymer viscosity decreases and shear forces weaken near the screw’s terminal end, minimal flight depth manages the pressure drop and maintains uniform pressure at the die. Inconsistent pressure at the die can lead to an extruded product lacking dimensional uniformity due to fluctuations in the extrudate output.

For fiber-loaded polymeric systems, the shear forces necessary for homogeneous fiber dispersion and deagglomeration arise from the shearing of high-viscosity melt and frictional interactions between the fiber-loaded polymer melt and the barrel wall. However, single-screw extruders offer limited opportunities for achieving high levels of fiber dispersion compared to twin-screw extruders [117].

3.4.2. Twin-Screw Extruders

Twin-screw extruders, housed within a metallic barrel, feature two screws that may be intermeshing, partially intermeshing, or non-intermeshing, and either co-rotating or counter-rotating [118]. In intermeshing counter-rotating twin-screw extruders, the opposite rotation of the screws allows one screw to seal the material circulating in the other screw’s channel, resulting in the material moving down and recirculating in a C-shaped section [119]. The distance between the screws determines the tightness of the seal. Unlike single-screw extruders, intermeshing counter-rotating twin-screw extruders do not primarily rely on the polymer’s shear and rheological properties for melt extrusion [119,120]. In the case of high-viscosity polymeric systems, such as fiber-loaded ones, even larger screw clearances can achieve positive melt conveyance. However, low inter-screw distances and tight seals require higher torque levels to overcome shear forces [118]. Non-intermeshing counter-rotating twin-screw extruders lack high dispersive capacity and self-cleaning features [121]. In a “self-wiping” intermeshing co-rotating twin-screw extruder, the material in one screw’s channel is expelled by the other screw’s flight, moving longitudinally around both screws in a pattern resembling the number 8, enhancing material conveyance and mixing [121]. In twin-screw extruders, the screw rotation speed influences the residence time of the material and the extent of screw flight channel filling.

Limited mixing is caused by the compression and expansion in alternating low- and high-pitch screw elements. Therefore, specialized shear mixing elements, such as kneading blocks and gear mixers, are introduced in the twin screws to enhance mixing. These elements provide superior distributive and dispersive mixing. Extrusion screws in single-screw extruders can also be designed to possess these specialized mixing components to add to the extruder’s distributive and dispersive mixing potential [122,123,124]. The material faces less shear force in narrow kneading blocks as it moves in and around them, leading to adequate distributive mixing. Wider kneading blocks offer broader contact areas and enhance exposure to planar shear, resulting in superior dispersive mixing compared to narrower blocks.

3.4.3. Torque and Capillary Rheometers

Rheological properties of polymeric systems, including viscosity and melt flow, are commonly characterized using torque and capillary rheometers under different shear and thermal conditions. Torque rheometers utilize blades to apply shear stresses, inducing phase changes in polymer solids at elevated temperatures and mixing added fibrous reinforcement in the molten polymeric matrix. In contrast, capillary rheometers shape the polymer melt by extruding it through a narrow capillary tube using a piston. Researchers have used these devices to produce small biocomposite filament batches for 3DP [125]. Additionally, plunger-type batch rheometers can produce biocomposite filaments with highly oriented polymeric chains and fibrous reinforcement [126,127].

3.4.4. Thermo-Kinetic Mixers

Thermo-kinetic mixers, equipped with high-speed rotating bladed rotors, generate heat by converting the kinetic energy of mobile solid polymer particles colliding with each other and the chamber walls, inducing a polymeric phase change. These mixers achieve high fiber dispersion within short time scales, thus minimizing thermal exposure and limiting the polymer’s and fiber’s thermal degradation [128]. However, the high rotational speeds can produce high mechanical shear rates, which may negatively impact the fiber aspect ratio in the resulting biocomposite filament and FDMPB [129].

3.5. Cellulosic Reinforcement-Based FDM Biocomposite Classification

Plant cellulose-reinforced PLA-based FDMPBs can be classified and subclassified based on the form of the reinforcement phase, the filament manufacturing technique, and the scale of the fibrous reinforcement used in biocomposite filament-making, as depicted in Figure 9.

3.5.1. Cellulosic Yarn-Reinforced PLA-Based FDMPBs

Composites reinforced with spun yarns or long fiber tows demonstrate superior flexural and tensile properties than those reinforced with nano- and micron-scale fibers due to the larger dimensional continuity in the reinforcement phase [135]. At the core, spun yarns or fiber tows are also fibrous assemblies. Consequently, there is growing interest in spun yarn-reinforced PLA-based FDMPBs. The pioneering research in this area involved saturating jute yarn with molten PLA in an FDM printer’s heating element. However, achieving uniform yarn placement was challenging due to the difficulty of maintaining yarn tension during 3DP [130]. A recently developed double-stage in situ impregnation method for yarn-reinforced FDMPBs improves yarn flattening and matrix infiltration into the yarn structure compared to single-stage impregnation. Flattened yarn allows for greater lateral spreading of the constituent fibers within the polymer matrix, reducing fiber-rich or matrix-deficient zones in the printed biocomposite. Additionally, reduced fiber pull-out in samples 3D printed from double-stage impregnated filaments indicates superior fiber–matrix adhesion compared to single-stage impregnated filaments [136].

Another method, called coating continuous fiber fabrication, produces pre-impregnated biocomposite filaments [132,137,138]. In FDMPBs using flax yarn/PLA pre-impregnated filaments, porosities are primarily located within the matrix-deficient yarn component of printed rasters rather than between printed layers. Increasing the printing layer height from 0.2 mm to 0.6 mm significantly raises the porosity due to the inadequate yarn compaction at higher layer heights. However, increasing the number of printing layers improves yarn compaction. The porosity of the biocomposite also increases with greater hatch spacing between reinforcement yarns in adjacent rasters [139]. The number of printing passes and hatch spacing influence the available space for lateral matrix spreading, thereby affecting the dimensional accuracy of the printed object [137,138].

The fiber weight fraction in pre-impregnated yarn-reinforced biocomposite filaments can be modulated by adjusting the matrix-impregnating mold’s diameter if the yarn’s linear density is constant. A study found that utilizing a 1.2 mm orifice mold led to matrix-rich, unreinforced areas caused by a thicker polymer layer surrounding the fibrous core. Conversely, using a 0.8 mm orifice mold resulted in fiber-rich regions characterized by a thinner resin coating around the fibrous core and an uneven matrix distribution [132]. In yarn-reinforced FDMPBs, fiber content is also influenced by layer height and hatch spacing. Lower layer heights increase fiber content regardless of hatch spacing, while wider hatch spacing at a fixed layer height reduces fiber content in the printed biocomposite [139].

Increasing fiber content in yarn-reinforced composites does not indefinitely improve the mechanical performance of FDMPBs. Beyond optimal thresholds, weaker interfaces between natural fibers and the polymeric matrix become prevalent, increasing porosity occurrence. Structural porosities, in turn, significantly impact the dynamic strength of biocomposites. Among the factors influencing porosity evolution in yarn-reinforced FDMPBs, layer height plays a more critical role than fiber content, and it is the most influential factor affecting dynamic strength, followed by hatch distance [139]. Table 2 summarizes the plant-based spun yarn-reinforced FDMPB research detailing yarn specifications, fiber volume fractions achieved, and techniques for fabricating yarn-reinforced biocomposite filaments. Additionally, it presents the mechanical performance data reported in the studies listed in the table.

Improving compatibility between hydrophilic natural fiber-based yarn reinforcement and hydrophobic polymer matrices in FDMPBs remains largely unexplored. Surface modification of flax yarns with silane couplers has been shown to improve matrix penetration into the yarn structure, reduce porosity fraction, decrease fiber pull-out, and increase fiber breakages during mechanical testing, indicating robust reinforcement–matrix interaction in the FDMPB [140]. Moisture absorption into biocomposite triggers natural fiber swelling, affecting dimensional stability and creating internal stresses [141]. These stresses can lead to bonding failures at the fiber–matrix interface and the formation of microcracks, which further accelerate moisture ingress into the FDMPB interior and fiber–matrix debonding [142]. Matrix stiffness influences this behavior, with stiffer matrices reducing flax fiber swelling, resulting in lower biocomposite porosity, water absorption, and diffusion rates, while softer matrices exhibit the opposite trend [143].

Yarn-reinforced FDMPBs can be effective analogs for woven fabrics. 3DP structures inspired by woven fabric featuring orthogonally interwoven ramie yarns have been successfully developed, demonstrating high bending rigidity per layer. These structures exhibit superior energy absorption and maximum penetration force compared to unidirectionally printed materials. The undulated geometry of spun yarn reinforcement increases energy absorption, prevents fiber damage, and alters the fracture path during indenter penetration [144]. The dual nozzle printing method, initially introduced by Markforged for synthetic fiber-reinforced composites, has been adapted for FDM 3DP of spun yarn-reinforced biocomposites, rationalizing printing times. This method has demonstrated its potential in replicating woven composite pipe structures [145]. Axial and radial composite pipe structures were printed, with the former showing enhanced tensile properties and the latter improved compressive properties, while both exhibited similar bending performance [146].

The area of spun yarn-reinforced FDMPBs offers substantial research potential. One of the key challenges is the infiltration of viscous molten thermoplastic into the twisted fibrous structure of spun yarns while avoiding fiber-rich regions prone to stress-induced failure [140]. Enhancing polymer permeation through more accessible fibrous forms can improve the fibrous reinforcement–matrix distribution balance in FDMPB structures. The high bending stiffness of spun yarns complicates their precise nozzle path tracking during 3DP, resulting in inconsistent reinforcement distribution and alignment, thus diminishing reinforcement effectiveness [137]. To address these issues, roving fibrous structures have been suggested to allow greater flexibility and accuracy in following the nozzle path during printing [147].

FDM printing nozzles with plate geometry are employed during 3DP to address porosity issues by compressing spun yarn reinforcements. This process reduces porosity by promoting matrix infusion into the fibrous structure and spreading the yarn laterally to achieve a more homogeneous reinforcement–matrix balance [137,138]. However, the large dimensions of spun yarn reinforcements can cause breakages between the print surface and nozzle [148], restricting the process’s ability to reduce layer height and improve printing resolution. Reducing layer height is critical to minimizing porosity and increasing fiber content in FDMPBs, both of which are essential for enhancing print quality and mechanical performances [138,139]. Finally, despite the potential impact of cohesive fibrous bundles like spun yarns on the rheological profile of the exiting molten polymer, no research has modeled their effect on the biocomposite filament’s plunging behavior in FDM, per our knowledge.

Spinning yarn from fibers is a lengthy process involving multiple steps, significant waste, energy consumption, and financial costs compared to directly using fibrous reinforcement in FDMPBs. Annually, at least 137,493,000 tons of plant fibers are produced globally, which can be used for yarn spinning [149]; micron- and nano-scale fibers present a unique challenge, as they cannot be spun using conventional methods. Instead, specialized techniques like electrospinning are required, but these often rely on non-eco-friendly solvents, produce non-durable core-spun yarns, and are difficult to scale up [150,151]. The scalability barrier in producing plant-based nano-scale fibers has also limited their use as reinforcements in FDM biocomposite filaments. In contrast, micron-scale cellulosic fiber-reinforced PLA-based biocomposite filaments are proliferating the FDM marketplace due to their simpler production methods.

Table 2. Extrusion of PLA biocomposite filaments reinforced with spun yarns and their printed biocomposites: summary of maximal reported fiber volume fractions, mechanical performance, and research goals.

Reinforcing Yarn Specifications	Filament Fabrication Technique	Biocomposite Filament FVF (%)	Mechanical Properties	Filament Diameter (mm)	Objective of Study	Refs.
500 Tex double-plied jute yarn	In-nozzle impregnation	6.1 (Scanning electron microscopy for image analysis and measurement)	At 0.05% to 0.25% strain rates: Tensile modulus: 5.11 ± 0.41 GPa, a 157% increase over pure PLA. Tensile strength: 57.1 ± 5.33 MPa, a 134% increase over pure PLA (JIS K 7162 test standard).	-	Exploration of yarn-reinforced thermoplastic PLA’s 3D printability	[130]
68 Tex flax yarn with 320 turns/meter twist	Coating continuous fabrication (CCF)	30.4 ± 0.8 (Observation/calculation method not stated)	Longitudinal modulus (GPa): 23.3 ± 1.8 at 0.05% to 0.01% strain rates and 13.6 ± 0.8 at strain rate >0.04%. Transverse modulus (GPa): 3.5 ± 0.45. Elongation at break (%): 1.67 ± 0.20 longitudinally and 0.45 ± 0.08 transversely. Tensile strengths (MPa): 253.7 ± 15.0 longitudinally and 10.8 ± 1.2 transversely (ISO 527-4 test standard).	0.482 ± 0.03	Evaluation of mechanical property of flax spun yarn-reinforced 3D-printed biocomposite	[137]
68 Tex flax yarn with 320 turns/meter twist	Coating continuous fabrication (CCF)	26.4 ± 1.1 (Observation/calculation method not stated)	Stiffness and tensile strength increased by 210% at 0.2 mm layer height, with tensile modulus and strength improving by over 50% and 70%, respectively, at 10 layers (ISO 527-4 test standard).	0.503 ± 0.047	Examining how slicing parameters like layer height, layer count, trip number, and inter-raster distance affect sample geometry, tensile properties, and biocomposite microstructure	[138]
68 Tex double-plied	Coating continuous fabrication (CCF)	-	Compared to pure PLA, tensile strength rose by 89%, tensile modulus by 73%, flexural strength by 211%, and flexural modulus by 224% (ASTM D4018 test standard).	0.8, 1, 1.2	Using curved path planning for G-code generation to 3D print yarn-reinforced biocomposites with a fixed-axis 3DP machine	[132]
24 Nm/2R double-plied ramie yarn with 400 turns/meter twist	In-nozzle impregnation	-	Dynamic strength:152.56 MPa at a 2552.61 s−¹ strain rate with 1 mm hatch spacing, 94.21 MPa at a 508.53 s−¹ strain rate with 0.8 mm hatch spacing, both using a 0.3 mm layer height (no test standard defined).	-	Investigating how parameters like layer height, hatch spacing, and strain rate impact the dynamic strength of FDMPBs using experiments and machine learning models	[139]
24 Nm/2R double-plied ramie yarn	Two-stage in situ impregnation	20.1 (Calculated using fiber and filament dimensions)	Two-stage impregnated filament prints: fracture force (N): 27.1 ± 1.7, one-stage prints: 23.0 ± 4.0. Compared to one-stage prints, two-stage prints show a 20.1% and 9.5% improvement in tensile modulus at 0 and ±45° raster orientations, respectively, and a 7.6% and 9.2% improvement in tensile strength at the same orientations (no test standard defined).	-	Enhancing the mechanical properties of yarn-reinforced FDMPBs using a two-stage in situ impregnation for filament production	[136]
67 Tex flax yarn with 380 twists/m	In-nozzle impregnation	25 (Image analysis)	Compared to untreated samples, impact strength rose by >17.5% (70.65 kJ/m2) (ASTM D6110 test standard), flexural strength: 22% (136 MPa) (ASTM D7264 test standard; three-point bending), and delamination initiation value increased: 59% (2.91 kJ/m2). Tensile strength (MPa): Untreated sample: 161, treated sample: 170 Tensile modulus (GPa): Untreated sample: 7.8, treated sample: 9 (ASTM D3039 test standard)	-	Boosting mechanical properties of FDMPBs by surface treating flax yarn for improved PLA matrix compatibility	[140]
68 Tex flax yarn with 320 turns/m twist	Coating continuous fabrication (CCF)	32.6 ± 0.5 (Microscopic image analysis)	Longitudinal tensile modulus (MPa): cFF/PLA: 15799 ± 2154 (ASTM D638, ISO 527-4, and ISO 14129 test standards)	-	Investigating the hydroexpansion of yarn-reinforced FDMPBs and the effect of matrix stiffness on this expansion	[143]
36 Nm/2R ramie yarn with 400 turns/m twist with 0.35 mm yarn diameter	In-nozzle impregnation	24.3 (Calculated using fiber and filament dimensions)	Maximum force and energy absorption were 49.1% and 58.7% higher than pure PLA, respectively. Energy absorption: >4 J and penetration energy: ~0.5 JMaximum force: 450 N at a support span-to-indenter diameter ratio of 5 (ASTMD-6264 test standard)	-	Examining the penetration behavior of yarn-reinforced FDMPBs inspired by woven fabrics	[144]
Double-plied flax yarn with 0.4 mm diameter	Coating continuous fabrication (CCF)	16 (Observation/Calculation method not stated)	The radially reinforced composite failed at 370 N in compression (no test standard defined) and 332.6 N in three-point bending tests (ISO 14125:1998 test standard). Axially reinforced composite failed at 844.5 N in tensile testing (ISO 6259-1:1997 test standard).	1	Exploring print path planning algorithms and developing a five-axis, dual-nozzle system for printing and reinforcing curved parts like pipes	[146]
68 Tex flax yarn	Coating continuous fabrication (CCF)	32.6 ± 0.5 (Observation/Calculation method not stated)	-	0.5	Addressing geometrical limitations of yarn-reinforced biocomposite 3DP by examining discrepancies between intended and actual print paths	[147]
36 Nm/2R ramie with 400 m/meter twist	In-nozzle impregnation	-	Maximum tensile strength: 86.4 MPa (no test standard defined), peeling strength: 20.9 N (ASTM-D3167 test standard), under printing conditions of 220 °C, 0.3 mm layer height, and 100 mm/min speed.	0.35	Characterization of mechanical properties and interfacial characteristics of the FDMPB	[148]

Table 2. Extrusion of PLA biocomposite filaments reinforced with spun yarns and their printed biocomposites: summary of maximal reported fiber volume fractions, mechanical performance, and research goals.

Reinforcing Yarn Specifications	Filament Fabrication Technique	Biocomposite Filament FVF (%)	Mechanical Properties	Filament Diameter (mm)	Objective of Study	Refs.
500 Tex double-plied jute yarn	In-nozzle impregnation	6.1 (Scanning electron microscopy for image analysis and measurement)	At 0.05% to 0.25% strain rates: Tensile modulus: 5.11 ± 0.41 GPa, a 157% increase over pure PLA. Tensile strength: 57.1 ± 5.33 MPa, a 134% increase over pure PLA (JIS K 7162 test standard).	-	Exploration of yarn-reinforced thermoplastic PLA’s 3D printability	[130]
68 Tex flax yarn with 320 turns/meter twist	Coating continuous fabrication (CCF)	30.4 ± 0.8 (Observation/calculation method not stated)	Longitudinal modulus (GPa): 23.3 ± 1.8 at 0.05% to 0.01% strain rates and 13.6 ± 0.8 at strain rate >0.04%. Transverse modulus (GPa): 3.5 ± 0.45. Elongation at break (%): 1.67 ± 0.20 longitudinally and 0.45 ± 0.08 transversely. Tensile strengths (MPa): 253.7 ± 15.0 longitudinally and 10.8 ± 1.2 transversely (ISO 527-4 test standard).	0.482 ± 0.03	Evaluation of mechanical property of flax spun yarn-reinforced 3D-printed biocomposite	[137]
68 Tex flax yarn with 320 turns/meter twist	Coating continuous fabrication (CCF)	26.4 ± 1.1 (Observation/calculation method not stated)	Stiffness and tensile strength increased by 210% at 0.2 mm layer height, with tensile modulus and strength improving by over 50% and 70%, respectively, at 10 layers (ISO 527-4 test standard).	0.503 ± 0.047	Examining how slicing parameters like layer height, layer count, trip number, and inter-raster distance affect sample geometry, tensile properties, and biocomposite microstructure	[138]
68 Tex double-plied	Coating continuous fabrication (CCF)	-	Compared to pure PLA, tensile strength rose by 89%, tensile modulus by 73%, flexural strength by 211%, and flexural modulus by 224% (ASTM D4018 test standard).	0.8, 1, 1.2	Using curved path planning for G-code generation to 3D print yarn-reinforced biocomposites with a fixed-axis 3DP machine	[132]
24 Nm/2R double-plied ramie yarn with 400 turns/meter twist	In-nozzle impregnation	-	Dynamic strength:152.56 MPa at a 2552.61 s−¹ strain rate with 1 mm hatch spacing, 94.21 MPa at a 508.53 s−¹ strain rate with 0.8 mm hatch spacing, both using a 0.3 mm layer height (no test standard defined).	-	Investigating how parameters like layer height, hatch spacing, and strain rate impact the dynamic strength of FDMPBs using experiments and machine learning models	[139]
24 Nm/2R double-plied ramie yarn	Two-stage in situ impregnation	20.1 (Calculated using fiber and filament dimensions)	Two-stage impregnated filament prints: fracture force (N): 27.1 ± 1.7, one-stage prints: 23.0 ± 4.0. Compared to one-stage prints, two-stage prints show a 20.1% and 9.5% improvement in tensile modulus at 0 and ±45° raster orientations, respectively, and a 7.6% and 9.2% improvement in tensile strength at the same orientations (no test standard defined).	-	Enhancing the mechanical properties of yarn-reinforced FDMPBs using a two-stage in situ impregnation for filament production	[136]
67 Tex flax yarn with 380 twists/m	In-nozzle impregnation	25 (Image analysis)	Compared to untreated samples, impact strength rose by >17.5% (70.65 kJ/m2) (ASTM D6110 test standard), flexural strength: 22% (136 MPa) (ASTM D7264 test standard; three-point bending), and delamination initiation value increased: 59% (2.91 kJ/m2). Tensile strength (MPa): Untreated sample: 161, treated sample: 170 Tensile modulus (GPa): Untreated sample: 7.8, treated sample: 9 (ASTM D3039 test standard)	-	Boosting mechanical properties of FDMPBs by surface treating flax yarn for improved PLA matrix compatibility	[140]
68 Tex flax yarn with 320 turns/m twist	Coating continuous fabrication (CCF)	32.6 ± 0.5 (Microscopic image analysis)	Longitudinal tensile modulus (MPa): cFF/PLA: 15799 ± 2154 (ASTM D638, ISO 527-4, and ISO 14129 test standards)	-	Investigating the hydroexpansion of yarn-reinforced FDMPBs and the effect of matrix stiffness on this expansion	[143]
36 Nm/2R ramie yarn with 400 turns/m twist with 0.35 mm yarn diameter	In-nozzle impregnation	24.3 (Calculated using fiber and filament dimensions)	Maximum force and energy absorption were 49.1% and 58.7% higher than pure PLA, respectively. Energy absorption: >4 J and penetration energy: ~0.5 JMaximum force: 450 N at a support span-to-indenter diameter ratio of 5 (ASTMD-6264 test standard)	-	Examining the penetration behavior of yarn-reinforced FDMPBs inspired by woven fabrics	[144]
Double-plied flax yarn with 0.4 mm diameter	Coating continuous fabrication (CCF)	16 (Observation/Calculation method not stated)	The radially reinforced composite failed at 370 N in compression (no test standard defined) and 332.6 N in three-point bending tests (ISO 14125:1998 test standard). Axially reinforced composite failed at 844.5 N in tensile testing (ISO 6259-1:1997 test standard).	1	Exploring print path planning algorithms and developing a five-axis, dual-nozzle system for printing and reinforcing curved parts like pipes	[146]
68 Tex flax yarn	Coating continuous fabrication (CCF)	32.6 ± 0.5 (Observation/Calculation method not stated)	-	0.5	Addressing geometrical limitations of yarn-reinforced biocomposite 3DP by examining discrepancies between intended and actual print paths	[147]
36 Nm/2R ramie with 400 m/meter twist	In-nozzle impregnation	-	Maximum tensile strength: 86.4 MPa (no test standard defined), peeling strength: 20.9 N (ASTM-D3167 test standard), under printing conditions of 220 °C, 0.3 mm layer height, and 100 mm/min speed.	0.35	Characterization of mechanical properties and interfacial characteristics of the FDMPB	[148]

3.5.2. Cellulosic Fiber-Reinforced PLA-Based FDMPBs

Cellulosic fiber-reinforced PLA-based FDMPBs refer to biocomposites that incorporate fibers with diameters in the nano to micron scale and lengths ranging from nano to supra-micron dimensions. These materials can be subclassified as either nano- or microfiber-reinforced FDMPBs.

Micron-Scale Cellulosic Fiber-Reinforced PLA-Based FDMPBs

Biocomposites based on micron-scale fibrous reinforcements have demonstrated superior tensile performance compared to pure PLA prints post aging at room temperature, indicating their potential to enhance the longevity of PLA-based FDMPBs [152].

Micron-scale fiber reinforcement with a higher aspect ratio (L/D) enhances the tensile modulus and strength of FDMPBs [153,154,155]. Fiber-aspect-ratio attrition is typically mitigated by introducing fibers downstream, where a lubricating layer of molten polymer is already present, rather than at the extruder feed zone with solid polymer pellets [153,156,157]. Conversely, high-aspect-ratio fibers tend to agglomerate within the matrix more readily, disrupting melt flow and causing uneven fiber and matrix distribution [158,159]. Obstruction in melt flow may also prolong the matrix’s and fiber’s exposure to elevated temperatures within the 3D printer nozzle, risking thermal degradation of composite phases [160]. Additionally, high-aspect-ratio fibers may bend during melt compounding, leading to random fiber orientation and reduced filament stiffness [153]. Therefore, cellulosic reinforcements are often milled by design to attrite their aspect ratios in FDMPB research. Milling aids in more homogenous fiber dispersion in the molten PLA matrix during the fiber–matrix compounding stage [161,162]. Comparison between flax fibers and shives illustrates that while maintaining FDMPB printability, milled flax shives, with lower aspect ratios, can be added at higher levels (up to 30 wt.%) compared to flax fibers in flax/PLA biocomposite filaments [163].

High shear forces in melt compounding can reduce the fiber aspect ratio and the lumen size of cellulosic fibers, limiting polymer penetration into these hollow structures and increasing biocomposite porosity [164]. High shear conditions during melt compounding can also mechanically defibrillate micron-scale cellulosic fibers into nano-scale fibrils—a 5 wt.% content of these higher-aspect-ratio reinforcements in PLA matrix has been seen to improve reinforcement–matrix interaction, resulting in biocomposites with mechanical properties comparable to those of pure PLA. At higher MCC concentrations, such as 15 wt.%, high shear processing of maleic anhydride (MAH)-grafted PLA (MAH-g-PLA) matrix achieves strength and toughness similar to that of pure PLA FDM 3D-printed materials [107]. These two conflicting research outcomes on high shear compounding of cellulosic fibers and PLA matrix illustrate the variable effects of processing conditions, underscoring the challenges in achieving a scientific consensus in FDMPB research.

The effects of high shear rates extend beyond fiber modification, inducing shear thinning in the polymer melt. This phenomenon enhances the polymer wetting of reinforcing fibers and reduces biocomposite filament roughness. These improvements can also be achieved with polymeric rheology modifiers like plasticizers [153,160], as depicted in Figure 10. Fibers endure higher shear rates during filament fabrication than in the printer nozzle, resulting in greater fiber alignment in the filament’s skin region than in the core, as illustrated in Figure 11. This fiber alignment during biocomposite filament extrusion is attributed to the maximum shear rates at the die walls.

Fiber and polymer chain orientation in the printing direction enhances the mechanical performance of FDMPBs in that direction [93]. Raster angle significantly affects performance under uniaxial tensile stresses at either a 0° or 90° raster angle, with only half of the biocomposite rasters contributing to stress distribution in biocomposites printed at a 0°/90° raster angle [165]. This directional dependency of mechanical performance indicates that aligning polymer chains and fibers predominantly on the x–y plane can reduce their ability to resist bending stresses perpendicular to this plane, i.e., the z-axis during flexural tests in three-point bending tests [93]. In contrast, MCC/PLA FDMPBs exhibit improved flexural strength and modulus compared to compression-molded samples due to better MCC fiber orientation in the bending direction [107]. This finding challenges the rationale presented in Ref. [93]. In large-scale FDM-based BAAM systems, which use pellet feedstocks and vertical screw-based melt extruders, higher shear rates are expected than those in smaller FDM 3D printers. Nonetheless, a study showed that while some fiber alignment occurs in the printed raster direction of BAAM-printed biocomposites, complete alignment is not achieved [154]. Intriguingly, complete fiber alignment was observed in a benchtop FDM 3D-printed wood fiber/PLA/PHA biocomposite [166]. These findings raise questions about the relationship between shear rate and fiber alignment in FDM systems of different scales. There is a possibility that high shear rates in the BAAM extrusion die may cause a more substantial die swell in the extrudate, disrupting the orientation of fibrous reinforcements in the raster printing direction. Future research should investigate how shear rates in the extrusion dies or nozzles impact fiber alignment in small benchtop FDMPBs compared to BAAM-printed biocomposites, especially when using the same base feedstocks in filament or pellet form.

The layer-by-layer construction of FDMPBs inherently results in a porous structure. This intrinsic porosity is further exacerbated by factors such as the transfer of porosities from biocomposite filaments, FDM printer settings, fiber content, fiber agglomeration, and moisture volatilization.

The extrudate die swell phenomenon enlarges preexisting porosities within the polymer melt during the biocomposite filament-making process [167]. Although the filament drawing narrows these porosities, it elongates them lengthwise. Porosities in the biocomposite filaments also transfer into FDMPBs, which already form inherently less dense, more porous structures due to their layer-by-layer construction [167]. This “porosity reflectance” from biocomposite filament to FDMPB is illustrated in Figure 12A,B.

Higher printing-layer heights exacerbate porosity within the FDMP, increasing moisture absorption while worsening tensile and flexural performance [168]. Compared to higher printing-layer heights such as 0.15 and 0.2 mm, the maximum cocoa hull fiber/PLA FDMPB tensile strength was achieved at a 0.1 mm printing-layer height. Lower layer heights increase the contact surface between neighboring rasters and rationalize the porosity sizes [169]. On the other hand, a higher layer height can also reduce the frequency of void formation by requiring fewer layers to print an FDMPB [170], which can minimize the effect of these defects on the FDMPB’s mechanical performance. Adjusting raster overlap through lower z-offset values (negative air gaps) can promote better inter-raster wetting and reduce the size of interlayer porosities [171]. Nevertheless, a negative air gap may compromise shape fidelity, presenting a trade-off between mechanical performance and dimensional accuracy. Elevated printing temperatures also reduce melt viscosity and facilitate matrix encapsulation of the reinforcing fibers, resulting in a denser FDMPB [172]. Fractographic analysis reveals fiber breakages in biocomposite fracture surfaces, while scanning electron microscopy images (Figure 13) reveal smoother fracture surfaces due to the formation of diffuse interfaces between the composite phases with increasing temperatures. However, these fiber breakages may also stem from the thermal degradation-prompted weakening of fibrous reinforcement in the FDMPB at high printing temperatures [173]. Furthermore, excessive melt temperatures may overly reduce viscosity, causing matrix-phase separation at fiber tips and creating porosities in the FDMBP [126]. Printing speed can optimize the residence time of polymer melt in the FDM printer’s printhead and, by extension, its viscosity to influence inter-raster and interlayer porosities in the FDMPB structure. Tensile and flexural strengths remain unaffected by printing speeds; however, compression strength and modulus decline as speed increases from 30 mm/s to 70 mm/s [174] due to reduced polymer residence time in the FDM printer printhead’s heating element. This leads to suboptimal viscosity, poor wetting, and insufficient fusion of rasters, creating porosities and premature compressive failure. Well-fused structures with low porosity can better withstand transverse stresses, which suggests that prolonged exposure to optimal temperatures or high-temperature processing enhances compressive strength [173,174]. Thermal degradation of cellulosic reinforcements in the polymer melt increase at lower printing speeds and extrusion temperatures above 200 °C [174]. Therefore, optimizing printing speeds and temperatures is crucial for achieving optimal mechanical performance in FDMPBs.

Sugarcane bagasse fiber-reinforced PLA-based biocomposites show lower thermal stability than neat PLA. However, these biocomposites exhibit good thermal stability around 200 °C, making them viable for 3DP at this temperature without any negative ramifications for the FDMPB performance [93]. Thermogravimetric analysis (TGA) conducted in the temperature range of 30–700 °C at a heating rate set of 10 °C/min and in the presence of nitrogen purging gas shows that the grafting of PLA with glycidyl methacrylate (GMA) compatibilizer disrupts PLA’s crystal structure and makes it less thermally stable, initiating GMA-g-PLA’s degradation before neat PLA. However, adding bagasse fiber reinforcement to GMA-g-PLA makes the biocomposite more thermally stable because the fibrous reinforcement acts as a strong nucleating agent, increasing the biocomposite crystallinity. Therefore, strong nucleating behavior can be correlated with improved thermal performance in FDMPBs [175]. Alternatively, TGA conducted in the heating range of 40 to 600 °C at a scanning rate of 10 °C/min under a nitrogen atmosphere showed that the inclusion of MAH-g-PLA in PLA/polycaprolactone (PCL)/MCC ternary biocomposite strengthens the chemical interactions between the components of the ternary system, enhancing the biocomposite’s initial degradation temperature beyond 290 °C, and limits the chances of 3DP at a printing temperature of 205 °C [176]. An increment in biocomposite thermal stability has also been attributed to strong chemical bonding interactions with PLA that limit its polymeric chain mobility. Therefore, the quantum of thermal input required to break the intermolecular bonds is increased [175,177].

Figure 12. (A,B) Labels: 30 wt.% wood/PLA (TM30), 1 wt.% wood/PLA (TM30_1MCC), 3 wt.% wood/PLA (TM30_3MCC), 5 wt.% wood/PLA (TM30_5MCC). (A): Scanning electron microscopy (SEM) images of biocomposite filament cross-sections with increasing incidence of porosity showing increasing MCC content (constant wood content); image resolution: 500 µm. (B): SEM images of FDM-printed biocomposite cross-sections showing the increasing incidence of porosity with increasing MCC content (constant wood content); image resolution: 1 mm. Reprinted from Creative Commons Attribution (CC BY) [178]. (C) Magnified images encapsulating the effects of increasing wood content on the filament surface roughness; image magnification: 70×, wood agglomeration, FDM-printed biocomposite surface; image magnification: 20× and edge quality; image magnification: 40×. Reprinted with permission from Elsevier [167].

Figure 12. (A,B) Labels: 30 wt.% wood/PLA (TM30), 1 wt.% wood/PLA (TM30_1MCC), 3 wt.% wood/PLA (TM30_3MCC), 5 wt.% wood/PLA (TM30_5MCC). (A): Scanning electron microscopy (SEM) images of biocomposite filament cross-sections with increasing incidence of porosity showing increasing MCC content (constant wood content); image resolution: 500 µm. (B): SEM images of FDM-printed biocomposite cross-sections showing the increasing incidence of porosity with increasing MCC content (constant wood content); image resolution: 1 mm. Reprinted from Creative Commons Attribution (CC BY) [178]. (C) Magnified images encapsulating the effects of increasing wood content on the filament surface roughness; image magnification: 70×, wood agglomeration, FDM-printed biocomposite surface; image magnification: 20× and edge quality; image magnification: 40×. Reprinted with permission from Elsevier [167].

Micron-scale sugarcane bagasse fibers endow larger contact surfaces at the interface of neighboring rasters, reducing inter-raster void sizes in FDMPBs compared to pure PLA specimens. However, increased fiber content leads to more porosities due to incompatibility between the polar fibrous and nonpolar matrix biocomposite phases, reducing both flexural and tensile strengths [93]. Similar behavior has been observed in FDMPBs reinforced with fibers such as Opuntia ficus-indica and Posidonia oceanica [164]. Additionally, increasing kenaf fiber content in FDMPBs leads to porosity and brittleness, causing a reduction in impact strength [162]. Figure 12A,B illustrate this issue by showing that increasing the MCC concentration from the optimal 1 wt.% to 5 wt.% in wood/PLA biocomposites, where MCC acts as a cross-linker and reinforcer, leads to a rise in porosities in both the biocomposite filament and the FDMPB. As MCC fiber loading increases, these porosities become more pronounced due to wood fiber agglomeration in the biocomposite filament. This results in a progressively rougher filament surface, which compromises the printing quality and resolution by producing void-rich FDMPB surfaces and edges [178], as depicted in Figure 12C. Conversely, increasing MCC content, where it is the sole reinforcing element, generates higher shear stresses in the PLA matrix, reducing melt viscosity during 3DP. This rationalizes inter-raster and inter-layer porosities in the resulting FDMPB [107]. Declining PLA melt viscosity has also been linked to MCC agglomeration in the PLA matrix and potential MCC degradation [179]. Similarly, shear-induced breakage of soy hull fiber agglomerates during 3DP has been discussed as a potential source of lowered melt viscosity of thermoplastic copolyester matrix [180]. PLA melt rheology modification is also evident in FDMPB fracture surfaces, where lower MCC content results in raster separation due to weak inter-raster interactions. In comparison, higher MCC content leads to denser inter-raster and inter-layer interfaces and smoother fracture surfaces, with transverse raster splitting and no inter-raster debonding [181].

Moisture outgassing at elevated temperatures from hydrophilic fibers into the polymer melt is another major contributor to void formation within the biocomposite filament and the FDMPB [182]. As fiber content increases, moisture outgassing also rises, creating larger and more frequent porosities within the material [183]. These porosities can coalesce [125], weakening the overall FDMPB structure and resulting in premature failure under low tensile and cyclic stresses. Consequently, the printed object’s effective tensile and fatigue strengths decline due to porosities and, by extension, moisture in either of the biocomposite phases [91,125].

Figure 13. Fracture surfaces of wood fiber/PLA FDMPB 3D printed at 200 °C, 210 °C, 220 °C, and 230 °C temperatures. Reprinted with permission from Creative Commons Attribution (CC BY) [173].

Figure 13. Fracture surfaces of wood fiber/PLA FDMPB 3D printed at 200 °C, 210 °C, 220 °C, and 230 °C temperatures. Reprinted with permission from Creative Commons Attribution (CC BY) [173].

Due to their polar chemistry, cellulosic fibers are prone to clustering in nonpolar thermoplastic matrices [107], especially at elevated concentrations [178,184]. This clustering can disturb the fiber dispersion in biocomposites and even cause FDM printer nozzle blockages [155,185]. However, chemical treatments can alleviate this issue by reducing fiber–fiber interactions, enhancing fiber dispersion, and improving interphase compatibilization. In one case, sodium hydroxide treatment of milled cocoa husk reinforcement for a PLA-based FDMPB made the reinforcement phase porous due to the partial or complete removal of extra-cellulosic reinforcement components like hemicellulose, lignin, pectin, and waxes. As a result, the tensile strength of the chemically treated milled cocoa husk-reinforced biocomposite filament increased by 18% compared to the untreated milled cocoa husk-reinforced biocomposite filament [169]. Naturally occurring pores (lumens) in plant fibers facilitate PLA permeation and mechanical interlocking with the reinforcement phase, enhancing matrix–fiber interfacial adhesion. However, this lumen penetration by PLA diminishes when the fiber size is reduced through grinding, which increases void content, as shown by SEM micrographs from FDMPB fractographic analysis. This phenomenon was observed with Opuntia ficus-indica (OFI) fibers, where the limited PLA infiltration into the fiber lumens compared to Posidonia oceanica leaf (POL) fibers resulted in inferior mechanical properties. In contrast, the POL/PLA FDMPBs demonstrated superior tensile strength, tensile modulus, elongation at break, flexural modulus, flexural strength, and impact strength due to more effective matrix penetration and interfacial bonding [164]. Therefore, similar chemical processing-induced morphological changes to the reinforcement structure can present opportunities for fiber–matrix mechanical interlock and improved biocomposite mechanical output. Chemically treated biocomposites often exhibit increased density compared to their polymer-only or non-compatibilized counterparts, owing to reduced porosities and fiber agglomeration [181].

One approach to chemical processing involves using compatibilizers that facilitate the replacement of weak interphasic bonds between the reactive species of cellulosic fibers and PLA by bonding with these composite phases [175,184]. Compatibilizers also retard agglomeration and ameliorate fiber dispersibility, so stress transfer efficiency improves throughout the biocomposite structure [175]. As a result, biocomposites with more compatible phases demonstrate increased tensile strength and greater elongation at break, allowing them to endure higher stresses before failure [175,177]. Variability in the biocomposite mechanical behavior can stem from localized distortions in the incompatible fiber–matrix interface. Thus, fibrous reinforcement and matrix compatibilization also alleviate variability in the mechanical behavior of the FDMPB [186]. The efficacy of compatibilizers is contingent on their optimal concentrations. For instance, exceeding 1 wt.% KH-550 silane coupler for micro–nano-cellulose (MNC) modification causes re-agglomeration in the PLA matrix, weakening interfacial connectivity, as shown by SEM-aided fractographic analysis, while 0.5 wt.% KH-550 is insufficient to promote cross-connections between composite phases [160]. Another chemical processing method for modifying biocomposite systems involves PLA melt rheology-modifying plasticizers. These plasticizers, typically low in molecular weight, insert themselves between higher-molecular-weight polymer chains, increasing chain mobility. Plasticizing the polymer matrix helps optimize melt rheology, improving fiber dispersion during melt compounding. Still, plasticizer saturation in FDMPBs must be avoided, as it compromises tensile strength, toughness [177], and elongation at the break of the biocomposite [160]. In one particular case, the incorporation of polyethylene glycol (PEG) into KH-550 silane-coupled micro–nano-cellulose/PLA biocomposites resulted in a universal deterioration in both tensile and flexural strength, irrespective of the PEG content added [160]. Despite the effectiveness of chemical compatibilizers and plasticizers, these chemicals, and the raw materials and processes involved in their synthesis, often pose risks to environmental and human health, as detailed in

Table 3. Key solvents, plasticizers, and compatibilizers in cellulosic fiber-reinforced PLA-based FDMPBs: functions, health, and ecological effects.

Chemical	Function	Hazards	Refs.
Tetrahydrofuran (THF)	Utilized as a PLA solvent in solution casting. Aids in modifying cellulosic surfaces to improve compatibility with PLA in biocomposites [187].	Carcinogen	[188,189]
Dichloromethane (DCM)	Utilized as a PLA solvent in solution casting. Aids in modifying cellulosic surfaces to improve compatibility with PLA in biocomposites [160].	Linked to attack on the ozone layer, a possible carcinogen	[190]
Maleic anhydride (MAH)	Enhances cellulosic-PLA adhesion by grafting MAH onto PLA’s backbone.	A key radical initiator in the MAH grafting reaction, dicumyl peroxide is an asthmatic agent. MAH itself is produced from furfural acid, which, in turn, is produced via unsustainable means. Furthermore, MAH irritates the skin, eyes, and respiratory tract on short-term exposure, and long-term exposure causes chronic bronchitis and asthma-like symptoms.	[191,192,193]
Polyethylene glycol (PEG)	PLA plasticization, cellulosic reinforcement dispersant, improving interfacial adhesion between PLA matrix and cellulosic reinforcements [160].	The raw materials used to produce PEG are primarily derived from the petroleum crude cracking process.	[194]
Isocyanates	These can be grafted onto a cellulose surface, reducing its hydrophilic potential and aiding in linking cellulose with thermoplastic matrices [195].	Cause intense irritation to the mucous membranes of the eyes and the respiratory and gastrointestinal tracts. Possible carcinogens.	[196,197]
3-Aminopropyltriethoxysilane (KH-550), 3-aminopropyltrimethoxysilane	Two of the most widely used silane couplers improve the interfacial adhesion of cellulosic materials with thermoplastic polymers.	Severely corrosive to the skin and eyes	[198,199,200]

.

In response to the abovementioned concerns, FDMPB research has also explored eco-friendly compatibilization methods to address environmental sustainability. For instance, superheated steam (SHS) treatment reduces the hydrophilicity of natural fibers, enhancing their compatibility with PLA in FDMPBs. However, SHS-treated kenaf fiber-reinforced biocomposites exhibit degraded tensile properties due to increased porosity and thermal degradation of cellulosic fibers, undermining their reinforcing function [161]. Hemicellulose degradation during SHS treatment generates acidic by-products that further depolymerize cellulose, reducing tensile and flexural performance [173]. Thermal degradation-induced deterioration in the mechanical performance of cellulosic fiber-reinforced PLA-based FDMPBs has also been observed elsewhere [172,173]. These drawbacks highlight the need for alternative environmentally responsive fiber or PLA treatments that do not erode FDMPB thermal stability. Irradiative methods like vacuum gamma irradiation of MCC/PLA biocomposites enhance tensile strength and elongation by promoting interphasic chemical cross-linking, addressing the issue of thermal degradation. However, oxidative degradation and PLA chain scission occur in the presence of oxygen, compromising tensile performance [201]. In situ ultrasonic vibration assistance during FDM 3DP of hydroxyapatite (HA)/PLA composite has shown potential for optimizing the dispersion of HA filler within the PLA matrix. This technique breaks up HA agglomerates, ensures more compact inter-raster fusion, and minimizes porosities. As a result, the tensile and flexural strengths of the printed composites nearly match those of PLA-only prints [202]. Biodegradable polybutadiene succinate thermoplastic polymer is an alternative to hazardous chemical compatibilizers and irradiative methods [201], enhancing fiber dispersion in the PLA matrix and reducing water absorption by engaging cellulose’s hydroxyl groups [152]. Similarly, a soy hull fiber (SHF) and soy protein isolate (SPI) bi-filler system has been observed in SEM analysis of a biocomposite filament cross-section to improve interfacial adhesion between SHF and the PLA matrix [203]. Adding poly(2-ethyl-2-oxazoline) (PEOX) to PLA increases its hydrophilic potential to compatibilize it with hydrophilic composite constituents like SHF and SPI in the case of PEOX-compatibilized SHF/SPI/PLA FDMPB plant pots [204]. Flexible matrices are ideal for high-fiber-content biocomposite filaments, preventing excessive embrittlement [125,175,180,182], which can complicate winding and cause filament breakage during 3DP [145]. Although PLA is intrinsically brittle, it remains the most available and economical biomatrix material for FDMPBs due to the scale of its production. Alternatives like toughening PLA with polycaprolactone and reinforcing it with MCC have been explored. However, MCC and PLA incompatibility leads to phase separation, negatively affecting tensile properties. Introducing MAH as a compatibilizer significantly improves interfacial bonding, increasing tensile strength by 60%, elongation at break by 211%, and tensile modulus by 8%, which is also validated by image comparison from SEM analysis of MAH-grafted PLA/PCL/MCC and PLA/PCL/MCC composite [176].

The thermal behavior of biocomposites explicates interactions between biocomposite constituents and vice versa, predicting mechanical performance trends. For example, a robust interphasic adhesion between steam-exploded coconut fiber reinforcement and the bi-polymeric PLA/PBS matrix system supports effective coconut fiber dispersion, which in turn increases glass transition, melting, and cold crystallization temperatures of the coconut fiber/PLA/PBS biocomposite [152]. In MCC/PLA biocomposites, up to 1 wt.% MAH grafting of PLA enhances chain mobility and crystallinity, though it lowers the crystallization temperature [181]. Similarly, incorporating tributyl citrate plasticizer into an MCC/PLA/PCL biocomposite enhances PLA chain mobility through plasticization. The plasticizer amplifies biocomposite crystallinity by providing well-dispersed MCC crystal nucleation sites in the matrix while lowering the cold crystallization temperature [177]. In the case of titanate-coupled MCC/PLA biocomposites, crystallinity increases, and cold crystallization temperature decreases at lower titanate-coupled MCC contents (1 and 3 wt.%). However, when the content reaches 5 wt.%, these properties diminish due to intensified PLA and titanate-coupled MCC interactions, which curtail PLA chain mobility [187]. A similar trend is observed in wood fiber-reinforced PLA biocomposite filaments. Initially, increasing MCC content restricts polymeric chain mobility, lowering crystallinity. Beyond 5 wt.% content, MCC promotes PLA crystal nucleation through fiber agglomeration, leading to higher crystallization onset temperature and crystallinity. Despite the increment in biocomposite filament crystallinity, the attached tensile strength or elongation is unimproved due to stress distribution issues [167]. In general, as observed in

Table 4. Effect of PLA grade, reinforcing cellulosic fiber content, compatibilization or lack thereof, and FDM parameters on the mechanical properties of FDMPBs in the literature.

PLA Grade	Fiber Type, Sizes (µm), and Maximum Fiber Load (wt.%)	Treatment	Process and Equipment	Printing Layer Height (mm)	Printing Raster Angle (°)	Infill (%)	Printer Nozzle Size (mm)	Tensile Properties of Most Highly Fiber-Loaded FDMPBs and PLA	Impact Properties of Most Highly Fiber-Loaded FDMPBs and PLA	Compression Properties of Most Highly Fiber-Loaded FDMPBs and PLA	Flexural Properties of Most Highly Fiber-Loaded FDMPBs and PLA	Refs.
2003D	Kenaf particle size: 250 µm, 7 wt.%	-	Twin-screw compounding, then single-screw filament extrusion	-	±45	100	-	Fatigue ultimate tensile strength (MPa): pure PLA: 56.36, 7 wt.% kenaf/PLA composite: 11.82 (ASTM D638 test standard)	Izod impact strength: pure (kJ/m2) PLA: 3.04 7 wt.% kenaf/PLA composite: 2.1 (ASTM D256 test standard)	-	-	[162]
4032D	Sugarcane bagasse fiber (SCBF), raw sugarcane bagasse fiber (RSCB), 15 wt.%	Mercerization, bleaching to obtain SCB	Double-stage twin-screw extrusion	0.1	±45 (D), 0/90 (C), 0/0 (P), 90/90 Vertical	100	0.6	Tensile strength (MPa): pure PLA: 61.4, 15 wt.%; RSCB/PLA composite in D setting: ~35, 15 wt.%; SCBF/PLA composite in D setting: ~4 (ISO 527 test standard)	-	-	Flexural strength and modulus (MPa): pure PLA: ~103, ~3050, 15 wt.% RSCB/PLA composite in D orientation: ~50, ~3600 15 wt.% SCBF/PLA composite in D orientation: ~63, ~3800 (ISO 178 test standard; three-point bending)	[93]
Virgin PLA (unknown grade), ColorFabb Bamboofil 20% bamboo/PLA composite commercial filament	Bamboo, flax, median fiber diameter: 47–254 µm, median fiber length: 124–4849 µm, 15 wt.%	Plasticizer	Direct twin-screw extrusion (36 L/D) for bamboo composite filament, compression molding followed by twin-screw extrusion for flax composite	-	-	-	1	Printed parts not mechanically characterized	Printed parts not mechanically characterized	Printed parts not mechanically characterized	Printed parts not mechanically characterized	[153]
3052D	Hemp and harakeke fibers, diameter: 28.3 (±8.3) μm and 12.3 (±1.7) μm, respectively, 30 wt.%	Alkaline treatment	Intensive mixer compounding followed by twin-screw extrusion	1	-	-	1	Tensile strength (MPa): hemp/PLA FDMPB: ~28 tensile modulus (GPA): hemp/PLA FDMPB: ~4.2 (no test standard defined)	-	-	-	[155]
3251D	Agave fiber (AF), diameter: 37.7 ± 16.6 µm, length: 255 ± 108 µm, 10 wt.%	-	Single-stage twin-screw extrusion	0.3	±45, 0/90	100	0.3	Tensile strengths of PLA and 10 wt.% AF/PLA composite printed at ±45°, 0°/90° (MPa): ~51, 34, and 28, respectively. Tensile moduli of pure PLA and 10 wt.% AF/PLA composite printed at ±45°, 0°/90° (MPa): ~1100, ~1090, ~880, and 840, respectively (ASTM D638-03 test standard)	Charpy impact strength (J/m): pure PLA printed at ±45°, 0°/90°: 30, ~29 10 wt.% AF/PLA composite printed at ±45°, 0°/90°: 26, 25 (ASTM D6110-04 test standard)	-	Flexural strengths of pure PLA and 10 wt.% AF/PLA composite printed at ±45°, 0°/90° (MPa): 87, 82, ~56, and 48, respectively. The flexural moduli of the same (MPa): 3280, ~2700, and ~2500, respectively (ASTM D790-03 test standard; three-point bending)	[165]
4032D	Lemongrass fiber (LF), average fiber size: 65.6 µm, 10 wt.%	MAH grafting of PLA (0, 2, 5, 8, 10, 20%)	Double-stage twin-screw extrusion	0.1	±45	100	-	Tensile strengths of untreated pure PLA, untreated 10 wt.% LF/PLA composite, and 10 wt.% LF/5% MAH-g-PLA composite (MPa): 59.6 ± 0.8, 36.5 ± 1.3, and 54.0 ± 0.9, respectively (ISO 527 test standard)	Notched impact strengths of untreated pure PLA, untreated 10 wt.% LF/PLA composite, and 10 wt.% LF/5% MAH-g-PLA composite (kJ/m2): 2.6 ± 0.2, 1.9 ± 0.2, and 2.5 ± 0.2 (no test standard defined)	-	Flexural strengths of untreated pure PLA, untreated 10 wt.% LF/PLA composite, and 10 wt.% LF/5% MAH-g-PLA composite (MPa): 98.3 ± 4.4, 60.3 ± 5.5, and 96.2 ± 1.1, respectively. Flexural moduli of untreated, pure PLA, and 10 wt.% LF/5% MAH-g-PLA composites (MPa): 2220 ± 182, 2740 ± 113 and 3330 ± 25, respectively (ISO 178 test standard; three-point bending)	[184]
4043D	Pineapple leaf fiber (PALF), 5 wt.%	Alkali-treated pineapple fiber (APALF)	Single-stage single-screw extrusion	-	0/90	-	1.5	The tensile strengths of pure PLA, 5 wt.% PALF/PLA composite, and 5 wt.% APALF/PLA composite (MPa): 29.5, 40.1, and 41, respectively. The tensile moduli of the same (MPa): 879.4, 1152, and 1242.1, respectively (ASTM D638 test standard)	-	-	Flexural strengths of pure PLA, 5 wt.% PALF/PLA composite, and 5 wt.% APALF/PLA composite (MPa): 32.2, 44., and 49.4, respectively. The flexural moduli of the same materials (MPa): 1027.4, 1431.4, and 1481.5, respectively (ASTM D790 test standard; three-point bending)	[205]
4043D	Coconut fiber (CF), 10 wt.%	PBS added, no chemical modification	Double-stage twin-screw extrusion	0.1	45	100	1	Tensile strengths of pure PLA, 10 wt.% CF/PLA composite, and 10 wt.% CF/PLA/PBS composite (MPa): 3.31 ± 0.14, 4.31 ± 0.15, and 4.66 ± 0.38, respectively. The tensile moduli of the same (GPa): 0.90 ± 1.90, 1.13 ± 6.68, and 1.12 ± 4.77, respectively. These materials’ elongation at break (EOB) (%): 9.35 ± 0.01, 6.88 ± 0.03, and 5.73, respectively (ASTM D638 test standard)	Notched impact strength (kJ/m2): pure PLA: 4.10 ± 0.20 10 wt.% CF/PLA composite: 3.15 ± 0.25 10 wt.% CF/PLA/PBS composite: 4.30 ± 0.80 (ASTM D256 test standard)	The compressive strengths of pure PLA, 10 wt.% CF/PLA composite, and 10 wt.% CF/PLA/PBS composite (MPa): 86.29 ± 0.93, 44.80 ± 0.05, and 20.71 ± 0.53, respectively. The compressive moduli of the same (MPa): 1.65 ± 3.68, 0.97 ± 5.04, and 0.43 ± 5.43, respectively (ASTM D7336M-12 test standard)	Flexural strength (GPa): pure PLA: 2.56 ± 7.63 10 wt.% CF/PLA composite: 3.22 ± 6.65 10 wt.% CF/PLA/PBS composite: 3.13 ± 7.89 (D790-17 test standard; three-point bending)	[152]
4032D	Bagasse cellulose (BC) fiber, BCA, BCB, BCC (sizes: 80, 120, 200 mesh), 50 wt.%	Mercerization and bleaching of BC fiber, glycidyl methacrylate grafted PLA (PLA–GMA)	Single-stage twin-screw extrusion	-	-	-	-	Tensile strength (MPa): PLA, PLA–GMA, 50BCB: 45.44, 15.94, 47.55; tensile modulus (MPa): PLA, PLA–GMA, 50BCB: 2263.28, 969.01, 3233.29 (no test standard defined)	-	-	-	[175]
ColorFabb woodfil PLA/PHA composite commercial filament	Recycled pine wood fiber (WF), 30 wt.%	-	-	0.2	±45	100	0.4	Tensile strengths of 30 wt.% WF/PLA/PHA composite at 210°, 220°, 230°, 240°, and 250° (MPa): 19.2 ± 0.1, 20.3 ± 0.1, 20.8 ± 0.1, 19.8 ± 1.2, and 20.5 ± 3.2, respectively. Tensile moduli of the same at these temperatures (MPa): 426 ± 11, 453 ± 19, 446 ± 20, 438 ± 10, and 416 ± 76, respectively. Elongation at break (EOB) at these temperatures (%): 6.03 ± 0.14, 6.29 ± 0.72, 6.70 ± 0.01, 6.15 ± 0.03, and 7.06 ± 1.21, respectively (ISO 527-1/2 test standard)	-	-	-	[172]
Unknown PLA grade	MCC, average particle size of 38 µm, 3 wt.%	Gamma ray irradiation in vacuum (10, 30, 50 kGy)	Single-stage single-screw extrusion	-	-	-	-	Tensile strength (MPa): 50 kGy gamma ray-irradiated pure PLA: 64, 50 kGy gamma ray-irradiated MCC/PLA composite: ~66 (ASTM D638 test standard)	-	-	-	[201]
Easy WoodTM commercial composite filament	Cedar fiber (CF), 40 wt.%	-	-	0.2	0/0	-	0.4	Tensile strengths of 40 wt.% CF/PLA composite at 200°, 210°, 220°, and 230° (MPa): 20.0 ± 0.5, 19.5 ± 1.0, 18.1 ± 0.4, and 18 ± 0.1, respectively. The tensile moduli of the same at these temperatures (MPa): 1802 ± 32, 1717 ± 63, 1711 ± 39, and 1713 ± 15, respectively (ASTM D638 test standard)	-	Compressive strength (MPa): 40 wt.% CF/PLA composite at 200°, 210°, 220°, and 230°: 28.5 ± 0.4, 31.2 ± 0.6, 30.4 ± 0.5, and 32.8 ± 0.5 (ASTM D695 test standard)	Moments of rupture (MOR) of 40 wt.% CF/PLA composite at 200°, 210°, 220°, and 230° (MPa): 35.2 ± 1.0, 33.7 ± 1.6, 32.2 ± 1.4, and 32.8 ± 1.4, respectively. The moments of elasticity (MOE) of the same at these temperatures (MPa): 1928 ± 66, 1699 ± 84, 1806 ± 75, and 1557 ± 128, respectively (ASTM D790 test standard; three-point bending)	[173]
4043D	Cellulose microfiber (CMF) width: 20–30 µm, length: 700 µm, 10 wt.%	3-aminopropyltriethoxysilane	Torque rheometer compounding followed by plunger-type batch extrusion unit	0.4	-	-	0.8	Tensile strength (MPa): untreated CMF/PLA composite: ~26, treated CMF/PLA composite: ~28, tensile modulus (GPa): untreated CMF/PLA composite: ~4, treated CMF/PLA composite: ~3.4 (ASTM D638 test standard)	-	-	-	[186]
4043D, PLA waste (rPLA)	MCC, 5 wt.%	Joncryl chain extender (CE)	Single-stage twin-screw extrusion (L/D: 48:1), 27 mm screw diameter	0.38	±45	100	-	Tensile strengths of PLA/rPLA, PLA/rPLA/CE, and 5 wt.% MCC/PLA/CE (MPa): 44 ± 2, 20 ± 2, and 45 ± 2, respectively. The tensile moduli of the same materials (MPa): 2868 ± 140, 1662 ± 367, and 3051 ± 235, respectively. Elongation at break (EOB) for these materials (%): 2.6 ± 0.2, 2.0 ± 0.4, and 3.1 ± 0.5, respectively (ASTM D638 test standard)	Izod impact strength (J/m): PLA/rPLA: 28 ± 3, PLA/rPLA/CE:29 ± 3, 5 wt.% MCC/PLA/CE: 31 ± 3 (ASTM D256 test standard)	-	Flexural strengths of PLA/rPLA, PLA/rPLA/CE, and 5 wt.% MCC/PLA/CE (MPa): 71 ± 3, 83 ± 11, and 81 ± 6, respectively. The flexural moduli of the same (MPa): 2320 ± 121, 2571 ± 351, and 2731 ± 182, respectively (ASTM D790 test standard; three-point bending)	[179]
Commercial composite filament (unknown brand)	Wood fiber (WF), 30 wt.%	-	-	0.05, 0.1, 0.2, 0.3	-	-	0.4	Tensile strengths of 30 wt.% WF/PLA composite at 0.05, 0.1, 0.2, and 0.3 mm layer thickness (MPa): 35.5, 33.9, 28.7, and 20.5, respectively. Tensile moduli of the same at these thicknesses (MPa): 3642, 3410, 3115, and 2567, respectively (ISO 527 test standard)	-	-	Flexural strengths of 30 wt.% WF/PLA composite at 0.05, 0.1, 0.2, and 0.3 mm layer thickness (MPa): 128.3, 121.7, 113.6, and 84.3, respectively. Flexural moduli of the same at these thicknesses (MPa): 4887, 4350, 4125, and 3580, respectively (ISO 178 test standard; three-point bending)	[168]
4043D	Wood fiber (WF), 44.6% of fibers with lengths between 100 and 200 µm, 15 wt.%	-	Single-stage single-screw extrusion	-	0/90	100	0.6	-	-	-	Flexural strengths of pure PLA and 15 wt.% WF/PLA composite (MPa): 59.52 ± 8.25 and 80.14 ± 4.84, respectively. Flexural moduli of the same materials (GPa): 2.30 ± 0.39 and 3.67 ± 0.36, respectively (ASTM D790 test standard; three-point bending)	[183]
4032D	MCC, 5 wt.%	2, 4, 6, 8 wt.% TBC, 19 wt.% PCL	Single-stage twin-screw extrusion	0.2	-	-	-	Tensile strengths of 5 wt.% MCC/PLA//PCL (0 wt.% TBC) and 5 wt.% MCC/PLA//PCL(4 wt.% TBC): 28.2 and 39.4, respectively. Tensile moduli of the same (GPa): 1.29 and 1.28, respectively. Elongation at break (EOB) for these (%): 2.9 and 20.1, respectively (no test standard defined)	-	-	-	[177]
4043D	Cellulose microfiber (CMF), width: 20–30 µm, length: 550 µm, 20 wt.%	-	Torque rheometer compounding followed by plunger-type batch extrusion unit	200	-	-	-	-	-	-	-	[154]
3001D	MCC, 5 wt.%	Titanate coupling agent	Solution casting followed by twin-screw extrusion	-	-	-	-	-	-	-	-	[187]
Unknown PLA grade	Soy hull fiber (SHF), SPI, 7.5, 10 wt.%	PEOX	Single-stage twin-screw extrusion	0.2	-	95	0.4	Tensile strengths of pure PLA, 7.5 wt.% SHF/PLA PEOX compatibilized composite, 7.5 wt.% SHF/2.5 wt.% SPI/PLA PEOX compatibilized composite, and 10 wt.% SHF/PLA PEOX compatibilized composite (MPa): 36.384, 30.592, 30.938, and ~22.5, respectively (ASTM D638 test standard)	-	-	-	[203]
2002D	Opuntia ficus-indica (OFI), Posidonia oceanica leaves (POL), 20 wt.%	-	Single-stage single-screw extrusion (L/D = 25)	-	0	80	-	Tensile strengths of pure PLA, 20 wt.% OFI/PLA composite, and 20 wt.% POL/PLA composite (MPa): 60 ± 2, 32 ± 5, and 38 ± 7, respectively. Tensile moduli of the same (MPa): 2810 ± 40, 2610 ± 120, and 2560 ± 72, respectively. Elongation at break (EOB) for these (%): 3.7 ± 0.6, 1.8 ± 0.4, and 2.3 ± 0.5, respectively (no test standard defined)	Impact strength (kJ/m2): pure PLA: 21.2 ± 0.5 20 wt.% OFI/PLA composite: 11.3 ± 0.4 20 wt.% POL/PLA composite: 15.9 ± 0.8 (ASTM D6110 test standard)	-	Flexural strengths of pure PLA, 20 wt.% OFI/PLA composite, and 20 wt.% POL/PLA composite (MPa): 98 ± 6, 46 ± 2, and 60 ± 4, respectively. The flexural moduli of the same (MPa): 550 ± 11, 353 ± 22, and 350 ± 10, respectively (ASTM D790 test standard; three-point bending)	[164]
2003D	Wood fiber, particles that pass 237 µm mesh, 50 wt.%	-	Melt compounding cycle (unknown equipment) followed by single-screw extrusion	0.19	-	100	0.4	Printed samples not mechanically characterized	Printed samples not mechanically characterized	Printed samples not mechanically characterized	Printed samples not mechanically characterized	[178]
4032D	MCC, length: 24 µm, diameter: 10 µm, 10 wt.%	MAH	Melt mixing in capillary rheometer followed by filament extrusion on a single-screw extruder	-	-	-	-	Breaking stress, breaking strain, and tensile modulus of 10 wt.% MCC/MAH-g-PLA/composite: ~32 MPa, ~0.03%, and ~2700 MPa, respectively (no test standard defined)	-	-	-	[206]
LX175	MCC, average particle size: 51 µm, 25 wt.%	-	Twin-screw extrusion, L/D: 25	0.1	-	100	0.4	Tensile strength (MPa): pure PLA: ~57 25 wt.% MCC/PLA: ~34, tensile modulus (MPa): pure PLA: ~3500 25 wt.% MCC/PLA: ~3600, elongation at break (%): pure PLA: ~1.65 25 wt.% MCC/PLA: ~0.8 (ISO 527 test standard)	-	-	Flexural strength (MPa): pure PLA: ~120 25 wt.% MCC/PLA: ~62, flexural modulus (MPa): pure PLA: ~3400 25 wt.% MCC/PLA: ~3650, elongation at break (%): pure PLA: ~4.9 25 wt.% MCC/PLA: ~1.8 (ISO 178 test standard; three-point bending)	[107]
Formfutura BV commercial wood/PLA biocomposite filament (WPC)	Wood fiber, 40 wt.%	-	-	0.2	-	-	0.4	Tensile strength (MPa): WPC at 30 mm/s, 50 mm/s, and 70 mm/s: 19.8 ± 0.8, 19.2 ± 0.7, and 19.8 ± 0.3; tensile modulus (MPa): WPC at 30 mm/s, 50 mm/s, and 70 mm/s: 1731 ± 60, 1650 ± 60, and 1682 ± 27 (ASTM D638 test standard)	-	Compressive strength (MPa): WPC at 30 mm/s, 50 mm/s, and 70 mm/s: 31.8 ± 0.6, 27.7 ± 0.9, and 20.9 ± 1.8; compressive modulus (MPa): WPC at 30 mm/s, 50 mm/s, and 70 mm/s: 852 ± 8, 825 ± 11, and 728 ± 28 (ASTM D695 test standard)	Flexural strength (MPa): WPC at 30 mm/s, 50 mm/s, and 70 mm/s: 34.0 ± 1.5, 33.1 ± 0.5, 32.3 ± 0.4; flexural modulus (MPa): WPC at 30 mm/s, 50 mm/s, and 70 mm/s: 1680 ± 78, 1620 ± 70, and 1575 ± 34 (ASTM D790 test standard; three-point bending)	[174]
2003D	MCC, average particle size: 6–12 μm, wood fiber (WF), sieved through 237 µm sieve, 5 wt.%	-	Double-stage twin-screw extrusion	0.5	±45	100	1	Tensile strength (MPa): 30 wt.% WF composite/PLA: 28.31 30 wt.% WF/5 wt.% MCC/PLA composite: 19.38 (ISO 527-2:1996 test standard)	-	-	Flexural strength (MPa): 30 wt.% WF, composite/PLA: 24.35 30 wt.% WF/5 wt.% MCC/PLA: 24.94, flexural modulus (GPa): 30% WF/PLA: 2.97 30 wt.% WF/5 wt.% MCC/PLA: 1.93 (ISO 178 test standard; three-point bending)	[167]
4032D	MCC, diameter: 5–10 µm, 5 wt.%	Maleic anhydride compatibilizer, PCL toughener	Twin-screw extrusion, followed by single-screw extrusion	0.2	-	-	-	Tensile strengths of PLA and 5MCC/PLA/PLA-g-MAH/PCL (MPa): 46.67 and 49.4, elastic modulus (GPa): 1.5 and 1.39, EOB (%): 3.79 and 9.08 (ASTM D638 test standard)	-	-	-	[176]

, integrating fibrous reinforcement into PLA without adding compatibilizers or dispersion-aiding secondary polymers [152] typically reduces FDMPB tensile strength compared to pure PLA due to matrix discontinuities. However, exceptions exist where adequate fiber dispersion can be achieved [152,205]. Although increased crystallinity generally augments biocomposite stiffness [181], agave fiber-reinforced biocomposite filaments exhibit lower tensile moduli than pure PLA due to porosities introduced by fiber addition even after a rise in the biocomposite crystallinity [165], a phenomenon reported here as well [183,184] and ascribed to poor fiber dispersion in the matrix.

4. Prospects and Challenges

The increasing impacts of climate change are set to drive a global shift towards environmentally sustainable production technologies and materials through this century. This necessity positions FDM and FDMPBs as contributors to the 21st-century sustainable economy. Despite the challenges inherent to the early development phase, the prospects for FDMPBs remain promising. These prospects and challenges are outlined below.

4.1. FDMPB Degradation

Micron-scale reinforcing fibers are more effective at promoting biocomposite biodegradation than their nano-scale counterparts, though combining both types of fibers yields a synergistic effect that enhances overall degradability [207]. However, due to their relatively smaller surface area, the interface between micron-scale cellulosic fibers and PLA is less diffuse than that of nano-scale fibers, resulting in a higher propensity for interphasic voids/porosities that serve as conduits for moisture uptake. Furthermore, nano-scale reinforcements generally contain a lower fraction of hydrophilic amorphous cellulose than micron-scale fibers, whose amorphous content may only be partially hydrolyzed. Cellulosic fibers in the PLA matrix promote water diffusion, accelerating PLA hydrolysis. The moisture absorbed by these hydrophilic fibers causes swelling and cracks at the fiber–matrix interface, which are exacerbated at higher fiber loads, leading to interconnected cracks that allow more moisture to penetrate the biocomposite’s interior [125,208,209]. Additionally, interlayer and inter-raster porosities inside the FDMPB structure can trap surface infiltrating moisture, further facilitating structural breakdown. These moisture pathways, initially formed during reinforcement–matrix de-cohesion, are later exploited by microorganisms during biodegradation [210].

While cellulosic reinforcements enhance moisture-induced degradation, they can also delay UV degradation, preserving mechanical performance during the service cycle but slowing biodegradation at the end of life, as evidenced by UV testing conducted under a radiation exposure of 62 W/m² for four hours with forced airflow to prevent ozone formation [211]. Furthermore, cellulosic reinforcers can increase crystallinity by promoting polymeric crystal nucleation, which can impede degradation. Quenching extruded rasters can depress biocomposite crystallinity [212], but this may weaken inter-raster and interlayer fusion. Interestingly, hydrolytic degradation in high-crystallinity biocomposites accelerates beyond the fiber percolation threshold despite their higher crystallinity, as shown in degradation testing conducted in freshwater (pH ≈ 7.5, 45 °C) over 50 days [213].

In addition to the advantageous role of cellulosic fiber reinforcements in enhancing the degradability of PLA-based FDMPBs at the end of their service life, the potential of these reinforcements to cause undesirable molecular-weight degradation of the polymer matrix during filament production and the 3DP also requires closer scrutiny. High-molecular-weight PLA delays biodegradation but improves the tensile modulus [214]. A 22–25% drop in PLA molecular weight has been observed from extrusion to 3DP, with the most significant degradation occurring during filament production due to thermomechanical stress when PLA molecular-weight modification was assessed via the intrinsic viscosity method. Higher fiber loads increase melt flow resistance, requiring greater mechanical work and further degrading molecular weight [164]. Despite this, cellulose-reinforced FDMPBs exhibit greater stiffness than pure PLA while accelerating biodegradation. Reducing PLA molecular-weight degradation can be achieved by minimizing residence time in melt compounders, which is possible with higher screw speeds and moderate temperatures [215]. Cellulosic fibers also augment polymeric shear thinning [180], lowering viscosity and shortening PLA residence time in the equipment, thus limiting thermo-oxidative degradation. Future research should investigate the influence of various cellulosic fibers and their concentrations on the rheological properties, crystallinity, and molecular weight of polylactic acid (PLA) during filament production and three-dimensional printing (3DP). Such studies would provide critical insights into the effects of these factors on the biodegradability of FDMPBs under ambient conditions. Lastly, a PLA-specific research focus is needed to investigate the impacts of FDMPB feedstock fabrication techniques like melt compounding extrusion and solution casting on different PLA-based cellulose fiber-reinforced FDMPB degradation modes. Specifically, it is important to determine which method produces FDMPBs that degrade more readily under environmental conditions, as this has implications for sustainability. Since the biocomposite microstructure strongly affects its degradability, it is important to determine which FDMPB production method can lead to higher biocomposite crystallinity and higher void creation. Biocomposite crystallinity and void fractions can be used as proxies to explain the degradation results. Therefore, microstructure-characterizing techniques like differential scanning calorimetry, X-ray diffraction, and micro-computed tomography can be used to characterize the respective FDMPB crystallinity and void fractions. Subsequently, degradation kinetics can be evaluated under accelerated weathering conditions, following the relevant ASTM standards for thermal, photolytic, hydrolytic, and biodegradability testing.

4.2. FDMPB Compostability

In the green economy, compostability into plant nutrients is preferred over mere degradability. PLA, a compostable biopolymer, breaks down into water and carbon dioxide. However, its composting in soil is slow due to the scarcity of microbes capable of digesting it [216] and its high carbon-to-nitrogen ratio, which requires a nitrogen-rich composting substrate. Therefore, PLA often needs regulated high temperatures and humidity to decompose under the action of thermophilic bacteria [217,218]. While mesophilic bacteria can initiate PLA breakdown at ambient conditions, reducing the energy required for thermophilic composting [219,220], concerns remain about incomplete decomposition under mesophilic conditions [221]. This incomplete breakdown can result in microplastic residues, compromising compost quality and contaminating the food supply [222,223].

Integrating hydrophilic cellulosic fibers into the PLA matrix can accelerate composting, converting printed biocomposites into plant nutrients more rapidly than pure PLA [210,224]. To our knowledge, no studies have directly compared the generation of microplastic residues between pure PLA and PLA–cellulose fiber biocomposites under various composting conditions. However, the biocomposite’s increased degradation rate and reduced PLA content suggest a lower residual microplastic footprint. This raises a critical research question: Do FDMPB structures significantly shed less microplastic than pure PLA 3D-printed materials under real-world composting scenarios? Researchers can answer this question by utilizing pyrolysis–gas chromatography–mass spectroscopy (Py–GC–MS) to quantify and compare the microplastic content between PLA-only 3D prints and FDMPBs, while particle size analysis to keep track of microplastic shedding from the composting polymer can be conducted using scanning electron microscopy. However, since Py–GC–MS is a destructive technique, a new specimen would be required to quantify the microplastic content after a fixed interval. Therefore, multiple samples would be prepared and exposed to the same composting conditions, each undergoing Py–GC–MS analysis only once after a fixed interval.

Additionally, PLA biocomposites may reduce the energy needed for industrial thermophilic composting, which is typically energy-intensive due to the controlled temperature and humidity requirements. There are also no direct comparisons in the literature regarding the energy inputs required for industrial-scale composting of pure PLA versus PLA–cellulose fiber biocomposites.

4.3. FDMPB Durability

FDMPB durability can be defined as the ability of these materials to meet the design expectations in their application area during their defined service life. For instance, the water affinity of SHF-, SPI-, and PEOX-compatibilized PLA can accelerate the disintegration of PEOX-compatibilized SHF/SPI/PLA FDMPB plant pots, which benefits their fast degradability [204]. However, while functional in soil, such FDMPB plant pots risk premature structural compromise above ground, potentially causing soil and nutrient loss. Thus, characterizing their degradation behavior is essential to ensure durable above-ground service life [225]. Comparative literature points towards a more rapid moisture absorption under accelerated water immersion and freeze–thaw aging in wood fiber/PLA large-format 3D-printed biocomposites than their synthetic fiber-reinforced 3D-printed composite counterparts. These testing conditions simulate the environmental impacts on FDMPBs in outdoor applications by eroding their flexural strength and modulus to a larger extent than the baseline unaged wood fiber/PLA FDMPB and carbon fiber-reinforced acrylonitrile butadiene styrene (ABS) 3D-printed composites [226]. The hydrophilic character of wood fibers leads to fiber swelling post water absorption, de-cohesion from the matrix, matrix cracking, and moisture capillarity to the biocomposite interior, expediting the hydrolysis of the PLA polymeric chain. Therefore, emergent points of stress concentration and shortened polymeric chains collectively cause a decline in material strength, the modulus of elasticity, and toughness [226,227]. Higher-than-optimal printing speeds and/or lower printing temperatures can increase the extrudate viscosity and compromise robust inter-raster and, by extension, inter-layer adhesion during printing. This leads to the formation of voids, particularly at the surface layers or edges of FDMPBs, which serve as pathways for moisture ingress [228] and thereby accelerate hydrolytic degradation, ultimately limiting the material’s durability. Similarly, simulations of the combined impact of photo and thermo-oxidative degradation on wood flour/PLA biocomposite 3DP filaments showed that exposure to high temperatures and UV radiation resulted in discoloration, increased internal structural porosity due to composite phase separation, and increased moisture absorption in the aged filaments. In these tests, a fluorescent light apparatus was selected as the UV source per the ASTM G154-06 (2006) method, and three groups of samples were subjected to accelerated aging at 40, 50, and 60 °C. Their properties were evaluated at 0 h, 20 h, 40 h, 60 h, and 80 h intervals to monitor progressive degradation [229]. Hence, property tuning of such FDMPBs for outdoor applications becomes necessary to ensure durability. Copper metallic coating of ABS 3D-printed specimens has shown promise by protecting the coated specimen from UV and hygroscopic degradation. This insight was gained when environmental conditions were simulated in a controlled-environment chamber, where the 3DP samples were exposed to UV radiation at an intensity of 0.89 W/m² from 340 nm-wavelength UVA lights under a chamber temperature of 60 °C. A humidity level of 100% was maintained at 50 °C. The exposure cycle followed the ASTM G154 standard, consisting of sequential 8 h UV irradiation followed by a 4 h condensation phase [230]. Extending similar concepts to cellulose fiber-reinforced PLA-based FDMPBs can make these materials more resilient in their functional environments. A recent study has noted how the surface roughness of 3D-printed substrates can be tuned to trap air pockets that reduce the wettability of these substrates by controlling parameters like metallic deposition rates and substrate temperatures [231]. Other modes of introducing surface roughness to create hydrophobic materials like plasma etching can also be found in the literature [232]. The addition of highly crystalline fillers like cellulosic nanowhiskers to PLA is another strategy for creating durable materials, retarding water absorption, and, by extension, retarding the hydrolytic degradation of the PLA matrix in the biocomposite. This conclusion was drawn from a hydrolytic degradation test conducted in 10 mL of phosphate-buffered solution (pH = 7.4, temperature = 37 °C), where samples were removed at fixed intervals of one week, two weeks, one month, two and a half months, and three months; desiccated; and their degraded weights characterized [233].

4.4. FDMPB Recyclability

Mechanical recycling routes for PLA-based FDMPBs comprise cutting, shredding, or grinding FDMPB printed objects and 3D-printable filament feedstocks [234], followed by their re-extrusion as 3D-printable recycled biocomposite filaments. Modes of loss in the mechanical performance of FDMPBs in mechanical recycling include degradation of the polymeric weight and fiber aspect ratio due to the high shear forces required to make the FDMPB size suited for re-extrusion. During re-extrusion, thermomechanical stresses can further degrade the polymer’s molecular weight and reinforcing fiber dimensions.

The studies covered in our work indicate mixed results regarding the ability of cellulosic fibers to preserve the mechanical performance of PLA/cellulosic fiber-reinforced PLA-based FDMPBs across multiple recycling cycles. For instance, upon recycling, micro-nano cellulose-loaded FDMPBs experience a reduction in thermal stability due to a decay in polymer molecular weight. However, this process enhances the crystallinity of the recycled biocomposite compared to recycled PLA, thereby increasing its tensile and storage moduli [235]. In another case, the tensile performance of recycled PLA–wood fiber biocomposite filaments and FDMPBs suffer despite the emergence of a more diffuse fiber–matrix interface after multiple recycling cycles and higher crystallinity compared to unrecycled composites. This indicates that the degradation of polymeric molecular weight or the average size and weight of polymeric chains can have a more pronounced effect on the mechanical performance of recycled biocomposites than macro- and microstructural characteristics, such as reduced composite porosity, delamination, and increased crystallinity [236].

Unlike injection molding of recycled PLA, which exposes it to high shear and attached thermal stresses, the low shear processing involved in FDM 3DP of cellulose-recycled PLA biocomposites prevents shear-instigated thermal degradation, thus preserving crucial rheological properties such as storage modulus, loss modulus, and complex viscosity of the resulting cellulosic fiber-reinforced recycled PLA-based FDMPB [179]. Additionally, the preservation of the reinforcing fiber length in the FDMPB is more evident than in injection-molded specimens after being subjected to multiple thermomechanical recycling cycles due to low shear processing conditions in 3DP [237].

Besides thermomechanical recycling, chemical recycling of PLA involves hydrolyzing it into lactic acid monomers for re-polymerization. The sustainability credentials of PLA are further burnished by the carbon dioxide-sequestering raw materials used for manufacturing PLA polymer precursors and the lower energy requirements for PLA production compared to other polymers. However, when PLA is combined with cellulosic fibrous reinforcements in FDMPB filaments and printed objects, the resulting material heterogeneity complicates the segregation of pure PLA, posing a challenge to its chemical recycling in an environmentally responsive and commercially viable manner [238]. Therefore, there is scope for innovative interventions at the process and equipment levels to overcome the material heterogeneity quandary and extract PLA as a mono-material for chemical recycling.

4.5. PLA–Cellulose Interactions

Table 4 demonstrates a broad and consistent research interest in micron-scale cellulose-reinforced PLA-based FDMPBs. In addition to FDMPB research, numerous studies have focused on traditional cellulose-reinforced PLA biocomposites, which are not covered in this analysis. However, it is perplexing that more concerted experimental efforts have not been directed toward characterizing the PLA and cellulose interfacial interactions at the molecular scale. A more refined understanding of these interactions could enable researchers to develop strategies, including environmentally sustainable ones, to enhance interfacial adhesion between these two incompatible composite phases, thereby addressing the prominent incompatibility issue that constrains the mechanical performance and applicability of biodegradable composites composed of PLA and cellulose.

4.6. FDMPB Filament Development

Despite the critical need for optimization in biocomposite compounding and 3DP filament extrusion, most research focuses on enhancing fiber dispersion and interphase compatibility through chemical modifications. However, parametric optimization like screw profiles, processing temperatures, die geometries, and screw speed remains underexplored, offering the potential for chemical-free improvements. A high degree of fibrous reinforcement dispersion in polymeric matrices can result in higher composite crystallinity [239], which in turn is responsible for improved composite mechanical and thermal properties [240]. Conversely, fiber agglomeration can immobilize polymer chains, preventing their rearrangement into more ordered domains, thereby reducing composite crystallinity.

Addressing the relationship between biocomposite filament processing parameters and fiber dispersion requires the execution of an experimental design aimed at identifying the statistically significant parameters that influence fiber dispersion. Coupling this approach with nondestructive micro-computed tomography enables precise mapping of fiber dispersion and agglomeration, helping to establish clear cause-and-effect relationships. Furthermore, well-dispersed fibrous reinforcement in the biocomposite filament also results in the fabrication of FDMPBs with homogenously dispersed fibrous reinforcement. This ensures less variability and consistency in mechanical performance in different FDMPB regions. Ultimately, understanding how these parameters shape the properties of the biocomposite filament can offer valuable insights into improving the overall performance of the FDMPB constructed from these filaments.

4.7. Brittle Cellulosic Fiber-Reinforced PLA-Based FDMPB Feedstock

PLA’s brittleness, due to its glass transition temperature above room temperature [241], is a significant challenge in producing highly loaded cellulosic fiber-reinforced PLA-based biocomposite filaments through melt compounding and extrusion. Brittle filaments are unsuitable for 3DP, as they tend to break during print jobs, wasting time, energy, and resources. In large-scale FDM 3D-printing systems like BAAM, which do not require biocomposite filament feed, the highly fiber-loaded FDMPB products produced may still face limitations in their applicability due to inherent material brittleness. PLA’s low elongation at break, typically around 10%, worsens with the addition of cellulosic fibers in FDMPBs, limiting their use in high-impact or load-bearing applications [242]. Biocomposite toughness can be improved by incorporating environmentally friendly plasticizers like TBC. Future research should explore more sustainable plasticizers to enhance the printability of cellulosic fiber-reinforced PLA-based biocomposite filaments and produce tougher FDMPBs for broader applications.

4.8. Highly Loaded Cellulosic Fiber-Reinforced PLA-Based FDMPBs

High natural fiber reinforcement content in composite structures boosts their environmental sustainability and cost-effectiveness by reducing reliance on expensive polymer matrices. However, maintaining the biocomposite’s mechanical performance with increased fiber content remains a considerable problem. High performance may only sometimes be feasible in highly fiber-loaded biocomposite systems, particularly those lacking chemical modifications to promote fiber dispersion or bonding with the matrix. Therefore, it is essential to identify target applications for these materials that do not require demanding mechanical properties. While this limitation is significant in filament-based FDM processes, it becomes less of a concern in large-scale additive manufacturing methods such as BAAM. BAAM does not rely on biocomposite filament feedstocks, as it 3D prints by melt-extruding biocomposite pallets. Hence, fiber-loading limitations like excessive and unsustainable filament embrittlement are not hindering factors for BAAM. Consequently, printed biocomposite parts can achieve high fiber-loading contents for specific, less mechanistically demanding applications. The high fiber content in FDMPBs can also have environmental benefits by expediting their biodegradation per the mechanisms discussed earlier.

4.9. FDMPB Void Defect Analysis and Quantification

FDMPBs show reduced inter-layer porosities at optimal fiber loads due to shear thinning-driven inter-layer wetting. However, there is a lack of peer-reviewed studies quantifying the frequency of void occurrence within rasters and between layers across different fiber loads. This raises critical questions: Are voids more prevalent within the rasters or at the inter-layer interfaces? How do intra-raster and inter-layer voids individually detrimentally affect the mechanical performance of FDMPBs? Image segmentation of the output from microscopy can construct a picture of the intra-raster and inter-layer voids, characterizing their void fractions to establish the frequency of void creation and distribution. Heat maps derived from these data showcasing the void defect distribution in the intra-raster and inter-layer spatial domains and their correlation with digital image correlation during mechanical testing (tensile, bending, impact) can shed light on how void location influences strain localization and failure modes in FDMPBs. Apart from heat maps, quantitative statistical analysis can provide a strong framework for correlating defect density with FDMPB mechanical performance.

Identifying where most defects concentrate under various processing conditions could help researchers propose more targeted solutions to lessen these defects in high-frequency regions, thereby enhancing the mechanical performance of FDMPBs. Therefore, it is essential to determine which processing conditions (e.g., printing speed, printing temperature, raster angle, nozzle geometry, cooling rate, layer height) contribute most significantly to the incidence of void frequency and their spatial distribution. FDMPB density characterization can correlate the influence of processing parameters on the incidence of void creation, while high-resolution synchrotron X-ray imaging can dynamically track the void formation and its spatial distribution within rasters and between printed layers during 3DP under different processing conditions. Defect hotspot maps generated from these data can then be mated with digital image correlation during mechanical testing of FDMPB samples to investigate how void location influences strain localization and failure modes. Furthermore, employing Taguchi methodology and statistical tools such as ANOVA and response surface methodology will help quantify the significance of parameters and their interactions.

4.10. FDMPB Parametric Studies

Extensive research has been conducted on the effects of printing parameters on polymer-only materials. However, limited studies explore the interaction of parameters such as printing speed, layer height, nozzle size, and printing temperature on FDMBPs. Design of experiments under the Taguchi methodology and subjecting the data to statistical analysis tools like ANOVA and response surface methodology can identify the critical interactions and optimize processing conditions for printing high-quality, defect-free FDMPBs. Moreover, Table 4 illustrates the varied printing process variables and testing standards used across different studies, complicating direct comparisons due to disparate processing conditions.

The impact of printing speed on shear-induced fiber orientation in FDMPBs remains an unexplored area of research as well. Questions such as, “Does the printing speed generate sufficient shear to force the fiber orientation in the deposited rasters to the extent where the emergent FDMPB mechanical properties are significantly altered?” deserve investigation. Capillary rheometry can simulate shear rates generated at high printing speeds, and various printing speeds and temperatures can be used in the process. In situ, dynamic visualization techniques like synchrotron X-ray imaging and high-speed hot-stage microscopy can be utilized to observe the fiber orientation in real time. Post-melt solidification visualization techniques like electron/polarized light microscopy and micro-computed tomography can illuminate the fiber orientation state. Then, mechanical characterization for tensile, flexural, impact, and compression properties of FDMPB samples printed at the conditions obtained from the rheological and fiber orientation state experiments can be carried out to correlate fiber orientation with mechanical performance. While the effects of varying raster angles and fiber contents on the tensile characteristics of FDMPBs have been examined in several studies, there is a significant gap in the literature regarding the compressive, flexural, and impact performance of FDMPBs across different raster angles. To fill this gap and answer the question about how varying printing raster angles influence the compressive, flexural, and impact properties of FDMPBs, experiments could be designed using a void mapping, nondestructive micro-computed tomography technique to characterize the void fraction; similarly, FDMPB density measurement following the Archimedes method could also be carried out. These void and density thresholds can then be correlated with the compressive, flexural, and impact performance of the FDMPB.

4.11. Scaling Up FDMPBs: Challenges and Research Directions

Until now, small-scale, specialized application areas have been the focus of attention for researchers to find application areas for FDMPBs. With the advent of BAAM, biocomposite research conducted by printing small-scale testing coupons or small objects on benchtop printers has attained the opportunity to transition to the production floor with these large-scale printers to print large-sized objects such as customizable wood-alternative products, including furniture pieces. While promising, defects such as incomplete wetting of neighboring rasters or partial necking may be amplified in BAAM, resulting in macro-porosities in larger structures. Print failures due to defects or unstable extrusion can lead to significant resource and time wastage. Additionally, loss of print resolution and layer stability to large layer heights can also be potential challenges in BAAM of biocomposites. Therefore, scientific research into mitigating defects in the large-scale printing of biocomposites is necessary to fully realize the potential and applicability of large-scale FDM for biocomposites.

Future studies could investigate how and whether BAAM printer design, processing parameters (e.g., printing temperature, die and screw profiles, layer height, extrudate output, flow rate), and feedstock characteristics (e.g., fiber content) influence macro void formation, printed layer stability, and attached print resolution. For example, a higher printing temperature, high printing speed, or modified die and screw profiles may slow the cooling rate, thereby affecting the cooling threshold required for optimal inter-raster fusion, while variations in layer height or extrudate flow rate stemming from these parameters can directly influence the geometry of deposited rasters. Similarly, differences in fiber content can alter the melt viscosity, surface tension of deposited melt, and heat transfer properties, thus impacting both the cooling behavior and the resultant void formation.

Thermal imaging can help assess how these variables affect the deposited raster’s cooling rate and solidification behavior. Mating thermal imaging to microscopic analysis could also provide insights into the balance between the optimum melt cooling rate and inter-raster fusion, as incomplete raster fusion can create voids in the FDMPB structure, while combining high-speed imaging with thermal imaging can reveal variations in raster spread and height that impact shape fidelity at different cooling rates. Dimensional metrology methods like structured light scanning can also quantify shape distortions arising from raster instability. Similarly, computed tomography can be employed to evaluate void geometries (qualitatively and quantitatively), distribution patterns (qualitatively), and fractions (quantitatively) in printed parts.

Statistical tools like ANOVA and response surface methodology can then be used to analyze data obtained from these techniques, identifying statistically significant main effects and interactions to guide the optimization of process parameters for defect mitigation.

4.12. FDMPB Mechanical Characterization

Extensive research has examined impact characteristics under pendulum impact stresses for cellulosic fiber-reinforced PLA-based FDMPBs. However, short, natural fiber-reinforced variants are seldom considered for drop-weight impact characterization. Evaluating drop-weight impact performance is essential for cellulosic fiber-reinforced biocomposites to become viable alternatives in high-volume applications such as wooden furniture, which aligns well with the customized manufacturing capabilities of FDM. The peer-reviewed literature also lacks studies on the compressive strength characterization of FDMPBs reinforced with high-cellulose-content micron-scale fibers, as highlighted in Table 4. Future research can focus on the effect of key processing parameters during biocomposite filament-making and 3DP stages on the material properties (reinforcing fiber aspect ratio and loading content) and their drop-weight impact and compression behavior. Suppose factors like void evolution, fiber agglomeration, printed raster, and layer de-cohesion influence material failure modes under drop-weight impact or compressive loading conditions. In that case, the way in which the optimization of processing parameters at the filament-making and 3DP stages can affect these factors can be explored. Statistical analysis can help establish the key processing parameters and interactions influencing the drop-weight impact performance and compression strength responses, which can then inform the optimization efforts. Nondestructive digital image correlation and micro-computed tomography can be used to quantify the frequency and scale of raster/layer de-cohesion in failed specimens and void fractions in both pre- and post-tested specimens to validate the effect of processing parameters on FDMPB drop-weight impact and compressive performance.

4.13. Parametric Optimization and Artificial Intelligence

Optimizing FDM biocomposite filament production and 3DP processes to achieve optimal mechanical properties in FDMPBs can be complex, costly, and time-intensive. This challenge arises from the multitude of parameters at each stage of the process and their interactions, which collectively influence the properties of the final FDMPB product. Emerging niche technologies, such as machine learning within the broader field of artificial intelligence, are waiting to be leveraged in research settings to address these complexities and enable the development of mechanically robust FDMPBs. However, to date, no peer-reviewed studies have explored the application of such artificial intelligence tools for simultaneously optimizing both cellulosic fiber-reinforced PLA-based biocomposite filament production and the 3DP stages. The quality of predictive machine learning models used to identify key parameters during biocomposite filament-making (such as processing temperatures, screw design and speeds, die design, and green chemistry-based interventions for biocomposite interphasic compatibilization) and during 3DP stages (such as printing speed, printing temperature, print bed temperature, nozzle geometry, and printing environmental conditions) is highly dependent on the quality of the experimental datasets used to train them. These datasets must be robust enough to reveal underlying patterns and ensure defect-free FDMPBs. The traditional design of the experiment approach is time-tested in generating robust and detailed datasets. Nevertheless, developing such datasets can be an expensive and time-consuming undertaking. Yet, high-quality predictive machine learning models capable of detecting subtle patterns beyond human observation have the potential to leapfrog FDMPB research. To validate the parameter optimization based on ML model predictions, qualitative and quantitative techniques such as micro-computed tomography (micro-CT) and optical microscopy can be employed to observe key morphological features, including voids, fiber-aspect-ratio degradation, and fiber agglomeration.