Current Challenges in Melt Extrusion of Cellulose-Based

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Chapter 7

Current Challenges in Melt Extrusion of Cellulose-Based Nanocomposites Alain Dufresne* Université́ of Grenoble Alpes, CNRS, Grenoble INP, LGP2, F-38000 Grenoble, France *E-mail: [email protected].

The demand for the use of biomass as the source of renewable energy and materials has exploded during the last years in response to socio-economic and environmental issues. Suitable chemical or mechanical treatments can be applied to produce materials with dimensions in the nanometer range from many naturally occurring sources of cellulose. These nanomaterials display interesting properties for their use as a biobased reinforcing nanofiller in polymer nanocomposite applications. Recent initiatives for the commercial exploitation of these cellulosic nanomaterials have emerged, and melt processing methods, such as extrusion and injection-molding, naturally appear as a step towards their larger scale use. However, identified difficulties are associated with the melt processing of cellulose based nanocomposites. Some strategies have been proposed in the literature to tackle at least part of these issues, even if significant scientific and technological challenges to take up remain.

© 2018 American Chemical Society

Introduction Broadly defined, nanomaterials can be categorized as those which have structured components with at least one dimension less than 100 nm. In one-dimensional nanomaterials (1D), one dimension is outside the nanoscale, as for nanotubes, nanorods, and nanowires. Unexpected and attractive properties can be observed when decreasing the size of a material down to the nanoscale compared to its bulk counterpart. Cellulose is no exception to the rule. The specific properties of nanomaterials result from increased relative surface area, and quantum effects. It can impact properties such as reactivity, mechanical strength and electrical characteristics. Natural fibers are inherently composite materials primarily consisting of cellulose, hemicelluloses, pectin and lignin. The individual content of these components varies with the different types of fibers. This composition as well as the structure of the fiber can also be affected by growing and harvesting conditions. Chemical treatments can be applied to natural fibers to remove most of the non-cellulosic components and extract cellulose from the biomass. Suitable chemical and/or mechanical treatments can then be used to convert cellulose into fibrous material with one or two dimensions in the nanometer range. The initial concept of the chemical extraction of cellulose nanomaterials through an acid hydrolysis process was first reported in 1947 (1), and the mechanical destructuration of cellulose fibers was pioneered in 1983 (2, 3). These strategies involve a top-down deconstructing approach and differ therefore from bacterial cellulose (BC), another cellulosic nanomaterial, produced using a bottom-up method. As is usually the case when a new field is developing, the terminology has been somewhat confused for a while, but the term "nanocellulose" is now used to cover the range of materials derived from cellulose with at least one dimension in the nanometer range. The potential of these nanomaterials has been proved for special functional nanomaterials (4) but it is as a biobased reinforcing nanofiller that it has attracted the most significant interest (5–7). It is mainly attributed to the low cost of the raw material, low density, high specific strength and modulus, and renewability. In addition, the nanoscale disintegration of natural fibers eliminates the macroscopic flaws therefore limiting the big variation of structure and huge scatter of mechanical plant fiber properties inherent to any natural products. However, as for any nanoparticle, the main issue is related to its homogeneous dispersion within a polymeric matrix. For a long time processing has been mainly limited to wet techniques, i.e. casting a liquid mixture of the cellulose nanomaterial and polymer matrix, followed by evaporation of the liquid. However, when high volume products are targeted this wet processing strategy appears difficult to scale up. More conventional melt processing techniques are therefore expected to be the key processing methods. This chapter describes the difficulties associated with melt processing of cellulose based nanocomposites and approaches proposed in the literature to tackle these issues.

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Mechanical Properties of Cellulose at the Nanoscale The nanoscale structure of cellulose is intrinsically present in lignocellulosic fibers designed by nature. Most natural fibers are constituted of cellulose fibers, consisting of helically wound semi crystalline cellulose microfibrils, typically 10-30 nm in diameter, bound together by an amorphous lignin matrix. As the size of the cellulose crystals is reduced to nanoscale, they become insensitive to flaws with strength approaching the theoretical strength of atomic bonds. The optimized tensile strength and modulus of cellulose crystals thus allows a large amount of fracture energy to be dissipated in the matrix via shear deformation and consequently enhances the fracture toughness of these biocomposites. Evaluation of the tensile modulus and strength of the cellulose crystal is challenging because of the limitations in measuring the mechanical properties of nanomaterials. Theoretical calculations and indirect experimental measurements using atomic force microscopy (AFM), X-ray diffraction analysis, inelastic X-ray scattering, and Raman spectroscopy have been used to calculate the elastic properties of cellulose nanocrystal (8). The theoretical tensile strength of the cellulose crystal was found to be in the range of 7.5–7.7 GPa, which is much higher than that of steel wire and Kevlar-49 (8). A broad range of values, between 56 and 220 GPa, was reported for the tensile modulus of single cellulose I (native cellulose) crystal, with an average value of 130 GPa (9). The lower crystallinity of cellulose microfibrils results in a lower modulus, which average value is around 100 GPa (7). The specific tensile modulus, which is the ratio between the tensile modulus and the density (1.5-1.6 g.cm-3 for crystalline cellulose), was estimated around 85 and 65 J.g-1 for the cellulose crystal and cellulose microfibril, respectively, whereas it is around 25 J.g-1 for steel (10). These data justify the numerous research efforts that have been poured into the use of cellulosic nanomaterials as reinforcement for polymers.

Nanocellulose Two main strategies are commonly used to release nanoparticles from cellulose fibers, viz. mechanically-induced destructuration and chemically-assisted method. Both approaches result in morphologically different nanomaterials. The mechanically-induced destructuration method consists in submitting a diluted cellulosic fiber suspension to severe mechanical shearing actions in order to defibrillate the material and release more or less individually the constitutive microfibrils. Different shearing approaches such as high-pressure homogenization (2, 3), microfluidization (11), ultra-fine friction grinding (12), high-intensity ultrasonication (13), aqueous counter collision (14) or ball milling (15) can be used. The ensuing nanomaterial is usually called microfibrillated cellulose (MFC), nanofibrillated cellulose (NFC) or cellulose nanofibrils (CNF) and it is obtained as a much diluted aqueous suspension (typically 2 wt%) for viscosity issue. Indeed, even when diluted the suspensions displays a gel-like behavior (Figure 1a) (16). 139

Figure 1. (A) picture of a 2 wt% CNF suspension obtained from eucalyptus, enzymatically pretreated. Reproduced with permission from ref. (16). Copyright 2012 Elsevier. (B) TEM from a dilute suspension of CNF prepared from Opuntia ficus-indica. Reproduced with permission from ref. (20). Copyright 2003 Elsevier. (C) TEM from a dilute suspension of CNC prepared from capim dourado. Reproduced with permission from ref. (21). Copyright 2010 Springer. (D) picture of an aqueous suspension of CNC prepared from capim dourado (0.50 wt%) observed between cross nicols showing the formation of birefringent domains. Reproduced with permission from ref. (21). Copyright 2010 Springer. To decrease the high energy consumptions for fiber delamination associated to this mechanical treatment, different pretreatments of the cellulosic fiber raw material have been therefore proposed to facilitate this production route. It includes enzymatic pretreatment (17), or introduction of charged groups through carboxymethylation (18) or 2,2,6,6-tetramethylpiperidine-1-oxyl (TEMPO)-mediated oxidation (19). The resulting nanomaterial occurs as long high aspect ratio flexible entangled filaments made of alternating crystalline and amorphous cellulose domains (Figure 1b) (20). The width is generally in the range 3-100 nm depending on the source of cellulose, defibrillation process and pretreatment and the length is considered to be higher than 1 µm. The chemically-assisted method generally consists in applying a controlled strong acid hydrolysis treatment to cellulosic fibers. During this treatment, the hydronium ions can penetrate the cellulose chains in the amorphous domains, for which the hydrolysis kinetics is facilitate, promoting the hydrolytic cleavage of the glycosidic bonds and releasing individual nanocrystalline domains. Even if 140

different strong acids have been shown to successfully degrade cellulose fibers, sulfuric acid is generally preferred because it promotes the grafting of surface anionic sulfate ester groups favoring the dispersion of released nanoparticles in water. The ensuing nanoparticles are generally called cellulose nanocrystals (CNCs) and occur as spindle-like acicular nanoparticles (Figure 1c) (21). The average length is generally of the order of a few hundred nanometers and the width is of the order of a few nanometers, but these dimensions depend on the origin of the cellulose and hydrolysis conditions. When observed between cross nicols, the CNC suspension shows a birefringent typical feature (Figure 1d) originating from the structural form anisotropy, and flow anisotropy resulting from the alignment of the spindle-shaped nanoparticles generally operated before observation (21).

Preparation of Nanocomposites Since cellulose nanomaterials are obtained as diluted suspensions, usually in water since it is the most convenient and cheapest polar liquid medium, the simplest processing method consists in mixing the cellulose nanomaterial dispersion with a polymer/ prepolymer/monomer dispersed itself in the same liquid medium or in a liquid miscible with the liquid in which the nanoparticles are suspended. This mixture can then be cast and the liquid evaporated, avoiding the aggregation of the nanomaterial because of the intercalated matrix material and preserving its individualization state. This approach can be extended to other liquids to cover a broader range of polymer matrices by suspending the cellulose nanoparticles in the suitable liquid, i.e. the solvent in which the polymer matrix can be solubilized. Solvent exchange can be conducted in this endeavor. A pretreatment of the nanomaterial consisting in surface functionalization can also be conducted. Coating the nanoparticles with a surfactant or covalent coupling of hydrophobic moieties directly on their surface have been extensively used to tune their surface chemistry and enhance the dispersion of the filler in a broad variety of nonpolar liquid media or polymers. The negatively charged surface groups setting at the surface of H2SO4-prepared CNC can be advantageously used to promote strong interactions with cationic surfactants and the omnipresent surface hydroxyl groups of cellulose nanomaterials can be the site of various chemical reactions because of their high reactivity. Among many others, the surface chemical modification of cellulosic nanoparticles by acetylation (22), esterification (23), and silanization/ silylation (24, 25) has been reported. Polymer surface modification based on the "grafting onto" strategy with different coupling agents (26, 27) or "grafting from" strategy with the radical polymerization or ring opening polymerization (ROP) (28), atom transfer radical polymerization (ATRP) (29) and single-electron transfer living radical polymerization (SET-LP) (30) have also been reported. The modified cellulose nanomaterials can be effectively dispersed in nonpolar liquid media and mixed with the corresponding polymer solution prior to be cast and evaporated. This wet processing method generally leads to high stiffness nanocomposites and from the earliest studies, it was show that this outstanding reinforcing effect cannot be attributed only to the high stiffness of cellulose nanoparticles (31). 141

It was shown that above the percolation threshold, which value depends on the aspect ratio of the nanoparticles, the cellulosic nanoparticles can connect and form a 3D continuous pathway through the nanocomposite film. The formation of this cellulose network was supposed to result from strong hydrogen bonding interactions between nanoparticles. Therefore, any factor that affects the formation of the percolating nanocrystal network or interferes with it changes the mechanical performances of the composite (32).

Melt Processing of Nanocomposites The industrial production of nanocellulose is running, and this wet processing strategy appears difficult to scale up when high volume products are targeted. Melt processing techniques, such as extrusion and injection-molding are therefore expected to be the key processing methods. These techniques are cheap, fast, industrially and economically viable, and solvent-free. The main issues associated to the efficient melt processing of CNF/CNC reinforced polymer nanocomposites have been reported in the literature and some strategies have been proposed to overcome them. The identified issues are 1) the aggregation of the nanofiller upon drying prior to melt processing, 2) the difficulty associated to its dispersion within the polymer melt, 3) its thermal stability, 4) its structural integrity, and 5) its orientation.

Drying of Cellulose Nanomaterials When drying the CNC/CNF suspension, irreversible aggregation is expected because of the highly reactive surface of the nanoparticles and strong hydrogen bonding forces. This effect can be limited by using adapted drying techniques that produce a porous weakly-bonded solid material keeping as most as possible the nanosize structure. Spray drying has been proposed as a suitable technique to dry nanocellulose suspensions and particle size ranging from nano to micron were obtained (33). However, another study showed that conventional spray drying resulted in a compact solid structure with very low porosity and spray freeze drying was suggested as a more suitable technique (34). Physical coating or functionalization of the nanoparticles can also be used to avoid this irreversible aggregation exploiting the presence of the intercalated additive material and preserving their individualization state. However, this additional material will be in the final nanocomposite "for better and for worse". The effective redispersion of dried cellulose nanomaterials is also a drag for its commercialization. It is generally necessary to market diluted aqueous dispersion and therefore transport huge amounts of water. Maltodextrin was recently suggested as a green, easily removable and non-toxic additive improving the redispersability of dried CNF in water (35).

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Homogeneous Dispersion of Cellulose Nanomaterials within the Polymer Melt The issue of uniform dispersion of cellulosic nanomaterials in a polymer melt results from the hydrophilicity of cellulose and hydrophobicity of most polymeric matrices. The situation is similar to mixing oil with water, and it is even magnified when considering colloidal particles since the surface energy increases when decreasing the particle size. A recent literature survey (36) showed that melt extrusion of cellulose nanomaterial reinforced polymer nanocomposites has been mainly implemented with polylactic acid (PLA), polyethylene (PE) and polypropylene (PP). Moreover, it was reported that melt processing was mainly developed with CNCs that are considered as more adapted (36). It was attributed to the issue of homogeneous dispersion of CNFs within the polymer melt which is more difficult to perform than for CNCs because of entanglements between high aspect ratio and flexible CNFs. In these conditions, it is necessary to match the surface properties of the cellulosic material and the matrix. Hydrophilization of the matrix can be used but the chosen strategy generally consists in hydrophobization of the nanoparticle. Therefore, physical coating of the nanoparticle with a surfactant, polymer or copolymer, or covalent grafting of hydrophobic moieties have been extensively reported. One would think that when using a polar matrix, such as poly(ethylene oxide) (PEO) or starch, extrusion can be performed directly with raw cellulose nanoparticles since strong filler-matrix interactions are expected. However, even in these conditions aggregation of CNC was observed upon melt extrusion with PEO (37). Then, functionalization of the nanofiller can be implemented (e.g. with poly(ethylene glycol) (PEG)) to prevent strong self-aggregation, while preserving its polar nature. Assessment of the dispersion of the cellulosic nanomaterials within the polymeric matrix is generally performed using microscopic observations. Atomic force microscopy (AFM) or scanning electron microscopy (SEM) can be used, but the resolution is generally insufficient to identify individual nanoparticles. Nevertheless, it can be used to detect the presence of aggregates as shown in Figure 2. For example, Figures 2A and 2B show AFM images of nanocomposites obtained by melt extrusion followed by injection-molding consisting of polycaprolactone (PCL) as matrix and CNC extracted from ramie fibers as reinforcing phase (38). In Figure 2A, unmodified CNC was used and the elongated nanoparticles appear as larges aggregates up to a few microns in size. This aggregation was attributed to strong CNC-CNC interactions, largely accentuated by the freeze-drying process used before melt processing. Melt-blending performed at high temperature and under strong shear conditions was not able to finely disperse the nanoparticles within the PCL matrix. In contrast, when using PCL-grafted CNC, a totally different morphology was reported and no aggregates were observed (Figure 2B).

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Figure 2. (A,B) AFM images (phase) of cryo-microtomed PCL reinforced with 8 wt% neat CNC (A), and 8 wt% PCL-grafted CNC (B) obtained by extrusion and injection-molding. Reproduced with permission from ref. (38). Copyright 2011 Elsevier. (C-F) SEM images of the cross-section for neat PLA (C), and nanocomposites reinforced with 5 wt% neat CNC (D), and 10 wt% CNC coated with PEG-b-PLLA (E), and 10 wt% CNC coated with Im-PLLA (F). Reproduced with permission from ref. (39). Copyright 2017 American Chemical Society. 144

SEM images of the cross-sectional areas of nanocomposites obtained by melt extrusion followed by injection-molding consisting of PLA as matrix and CNC extracted from ramie fibers as reinforcing phase are shown in Figures 2 C-F (39). The surface of the cross-section for the unfilled PLA matrix was quite uniform (Figure 2C). When adding 5 wt% unmodified CNC, well distinguishable micrometric CNC aggregates were observed (Figure 2D). Additionally, voids in the PLA matrix proximal to the aggregates are visible. Two different non-covalent compatibilization strategies were investigated, both based on PLA-based surfactants, expected to improve the compatibility with the PLA matrix but differing by the polar head consisting of either a PEG block (PEG-b-PLLA) or an imidazolium group (Im-PLLA). For compatibilized nanocomposites (Figures 2E and 2F), the surface was similar to that of the unfilled material, and no aggregates were observed. These were strong lines of evidence of improved CNC dispersion and filler/matrix adhesion. If strong aggregation occurs, a simple visual inspection with the naked eye is enough. An example is provided in Figure 3 which shows extruded neat low density polyethylene (LDPE) film (left) that is obviously translucent as any low thickness polymeric film with a relatively low degree of crystallinity (23). When extruding the polymer with 10 wt% CNC (middle), the film becomes dotted with black spots. These heterogeneities reveal the poor and inhomogeneous dispersion of the filler within the polymeric matrix as well as thermal degradation of the cellulosic nanomaterial. When using organic acid chloride-grafted CNC (right), the occurrence of these aggregates progressively vanishes and the appearance of the composite film becomes similar to the one of the unfilled film as a result of improved dispersion. A combination of Raman imaging with image analysis has also been suggested as a powerful and useful tool to quantify the degree of mixing of CNC with melt compounded high density polyethylene (HDPE) (40).

Figure 3. Photographs of neat LDPE film (left) and extruded nanocomposites films reinforced with 10 wt% unmodified CNC (middle) and 10 wt% acid chloride-grafted CNC (right). Reproduced with permission from ref. (23). Copyright 2009 Elsevier. Thermal Stability of Cellulose Nanomaterials Melt processing of polymers and composites involves heat and as such it is important to avoid the thermal degradation of the processed material. The thermal stability is quite limited, and this is exacerbated for H2SO4-hydrolyzed 145

CNC because the acid hydrolysis process introduces less thermally stable sulfate groups on the surface (41). Moreover, melt processing has been mostly developed for CNC that is considered as more adapted compared to CNF (36). The thermal degradation of H2SO4-hydrolyzed CNC was shown to occur as a two-step process, viz. a low temperature process involving the degradation of most accessible amorphous regions which are also highly sulfated, and a high temperature process attributed to the degradation of less accessible interior crystalline regions that are comparatively less sulfated. To tackle this issue, another acid can be considered for hydrolyzing cellulose. Hydrochloric acid has been used to prepare CNC and it was shown that ensuing nanoparticles display a higher thermal stability compared to H2SO4-hydrolyzed CNC. However, because of preserved surface hydroxyl groups and lack of charged moieties, HCl-prepared CNC tend to aggregate easily leading to poor stability of the aqueous dispersion. It was shown that on account of ionic repulsion between charged surface groups, slightly phosphorylated CNC, prepared by controlled hydrolysis with phosphoric acid, can be readily dispersible and can form stable dispersions in polar solvents. The H3PO4-prepared CNC were found to exhibit a much higher thermal stability than H2SO4-prepared CNC (42). When using H2SO4-hydrolyzed CNC improved thermal stability can be achieved by neutralizing the CNC suspension with a 1 wt% NaOH solution (28). After neutralization, the thermal stability of H2SO4-catalyzed CNCs is rather similar to that of HCl-catalyzed CNCs, but for the former, however, steric stabilization between the nanoparticles is maintained. Another strategy consists in shielding the thermally sensitive sulfate groups by coating the cellulose nanomaterial through either chemical grafting or physical wrapping with a surfactant or long chains (masterbatch approach) before melt processing. This method was shown to impart improved thermal stability to the nanoparticles (43–46).

Structural Integrity of Cellulose Nanomaterials Melt processing techniques, such as extrusion and injection-molding, not only involves heat but also high shear rates that can impact the structural integrity of the cellulosic nanomaterial. A very limited number of studies reported this mechanical degradation effect of cellulose nanomaterials upon melt processing, which obviously depend on extrusion conditions and viscosity of the polymer melt. The dimensions (length and diameter) of CNC extracted from ramie fibers before and after extrusion with PEO were determined through microscopic observations and compared (37). As shown in Figures 4A-C, the length was found to decrease from about 200 nm to 120 nm, and the diameter reduced from 2 nm to 5 nm, resulting in a limited change in the aspect ratio of CNCs from 28 to 24. In addition, a significant narrowing of the length distribution was reported, showing that the mechanical degradation process mainly concerned the longer nanoparticles.

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Figure 4. (A,B) TEM of CNC (whiskers) extracted from ramie fibers before extrusion (A) and after extrusion (B) with PEO, and (C) their length distributions. Reproduced with permission from ref. (37). Copyright 2011 Springer. (D) length distribution for parent CNCs and CNCs that had been extracted from PVAc/CNC nanocomposites containing 8.3 vol% CNCs prepared by solution casting (SC), solution casting and reprocessing in a roller blade mixer (SC+RBM), or solution casting and reprocessing in a twin-screw extruder (SC+TSE). Reproduced with permission from ref. (47). Copyright 2015 Wiley. 147

In another study, nanocomposite materials have been prepared from poly(vinyl acetate) (PVAc) and CNC using different processing methods that involved different shear rates (47). The impact on the morphology of the nanoparticles was investigated as shown in Figure 4D. PVAc/CNC nanocomposites with various CNC contents were first prepared by solution casting as reference materials. These materials were post-processed by mixing in a roller blade mixer or a twin-screw extruder and subsequent compression molding. When solution-cast materials were re-processed under low-shear mixing conditions in a roller blade mixer, a mechanical reinforcing effect similar to the one observed for the reference material was reported, suggesting that when pre-dispersing CNC within the matrix, the morphology can be maintained during re-processing. On the contrary, twin-screw extrusion proved responsible for mechanical degradation due to high shear melt-mixing environment, reducing the CNC length.

Orientation of Cellulose Nanomaterials The high shear rates which are involved during melt processing of composites made of elongated particles can induce their orientation in the flow direction. Elongated cellulose nanomaterials are no exception to the rule. In itself this is not a problem as it can often lead to anisotropic composites with improved properties in one direction. However, it is worth noting that the outstanding mechanical properties observed for nanocellulose based nanocomposites are not only attributed to the high stiffness of cellulose nanoparticles, but mainly to the formation of a 3D cellulose network resulting from strong hydrogen bonding interactions between nanoparticles. Obviously, the orientation of cellulose nanomaterials in the flow direction limits the formation of this percolating network. This possible orientation obviously depends on the processing conditions and viscosity of the polymer. To evidence the orientation of CNC during melt extrusion with LDPE, small angle X-ray scattering (SAXS) experiments were performed (45). In this study, a triblock copolymer PEO-PPO-PEO was used as processing aid to limit the aggregation of the nanoparticles during drying, protect them from thermal degradation and compatibilize them with the non-polar polymeric matrix. The SAXS patterns obtained for extruded CNC/LDPE nanocomposites are shown in Figure 5A. An anisotropic shape of the SAXS patterns was observed with a higher intensity in the horizontal direction corresponding to a preferential orientation of CNC in the flow direction during the extrusion process. It can be seen that when increasing the CNC content from 1 wt% to 10 wt%, the anisotropic level was amplified which could be attributed to a higher orientation of the CNCs or to an increasing quantity of spindle-shaped nanoparticles orientated in the flow direction during the extrusion process.

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Figure 5. (A) 2D SAXS patterns for LDPE nanocomposites with different CNC contents. Reproduced with permission from ref. (45). Copyright 2016 Royal Society of Chemistry. (B) Short time-creep experiments (τ = 0.25 μN·m, T = 170 °C) for injection-molded CNC/PBAT nanocomposites: neat PBAT (■), PBAT reinforced with 1.8 wt% CNC (◆), and PBAT reinforced with 1.8 wt% CNC after thermal annealing for 30 min at 170°C (◇). Reproduced with permission from ref. (48). Copyright 2016 American Chemical Society. Orientation of CNC for injection-molded CNC reinforced polybutyrate adipate terephthalate (PBAT) nanocomposites was also reported using 2D-small amplitude oscillary shear (2D-SAOS) experiments (48). In these experiments, the expected result for a completely isotropic sample is symmetry and superposition of the normal and angular stress. It was observed that adding CNC caused a gradual distortion of the properties of the material by creating increasingly anisotropic samples due to CNC orientation. The absence of formation of a percolating CNC network was accessed by monitoring the response of the material under constant stress (creep experiment) (48). Figure 5B shows the evolution of the strain as a function of time when applying a constant stress to CNC/PBAT samples above the melting point of the PBAT matrix. As expected, the neat polymer 149

displayed a linear liquid-like behavior, with strain increasing regularly with the duration of the applied stress. The composite reinforced with 1.8 wt% CNC (corresponding to the percolation threshold) showed very distinct characteristics under the same testing conditions. The presence of CNC clearly modified the way the material can withstand the stress. Over the whole experimental time range, the experienced strain was clearly lower for the composite, the melt-processed material still presented almost linear behavior during the experiment. Besides its lower compliance value compared to that of the neat matrix, it corroborated that no percolating CNC network was present in the material. The CNC/PBAT composite was submitted to an annealing treatment at 170°C for 30 min before experiencing the creep experiment to induce a possible auto-reorganization of the nanofiller as a first step to create an isotropic material, the initial condition to induce 3D-network formation under mild conditions. Even if the high viscosity of the polymer melt limited the movement of the nanoparticles and hindered the formation of the desired percolating network, a spatial reorganization of the spindle-shaped nanoparticles was observed after short conditioning time as shown in Figure 5B. It suggested that it should be possible to partially alter the organization of the particle within the polymeric matrix imposed by the injection-molding process.

Conclusion Cellulose nanomaterials can be prepared in the form of cellulose nanofibrils (CNF) or cellulose nanocrystals (CNC). These materials exhibit interesting properties such as high stiffness and specific surface area, and low density for the processing of advanced polymer nanocomposites. In the last decade, initiatives for the commercial exploitation of nanocellulose have emerged. Melt processing methods, such as extrusion and injection-molding, are a step towards a larger scale use of cellulose nanomaterials. The main issues to overcome for an efficient melt processing of CNF/CNC reinforced polymer nanocomposites are the irreversible aggregation of the nanofiller upon drying prior to melt processing, its non-uniform dispersion within the polymer melt, its thermal stability, its structural integrity, and its orientation. Different strategies have been proposed to overcome these issues, but the fact remains that the final mechanical properties cannot compete with those obtained for nanocomposites prepared by casting/evaporation. This is mainly attributed to the proposed strategies that limit at the same time the formation of a 3D percolating nanoparticle network which is the basis of the outstanding reinforcing effect of cellulose nanomaterials observed when using wet processing techniques. One might think that simply increasing the CNC content would be enough to improve the mechanical properties of melt-processed nanocomposites. However, high loading level is impractical, for many reasons, among them cost, and viscosity changes. In addition, when using a semicrystalline polymer matrix, its crystallinity is often increased with the presence of the nanofiller and the impact of this crystallinity change from the exact role of the cellulosic nanomaterial is difficult to discriminate. 150

References 1. 2. 3. 4. 5. 6.

7.

8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25. 26.

Nickerson, R. F.; Habrle, J. A. Ind. Eng. Chem. 1947, 39, 1507–1512. Herrick, F. W.; Casebier, R. L.; Hamilton, J. K.; Sandberg, K. R. J. Appl. Polym. Sci. Polym. Symp. 1983, 37, 797–813. Turbak, A. F.; Snyder, F. W.; Sandberg, K. R. J. Appl. Polym. Sci. Polym. Symp. 1983, 37, 815–827. Lin, N.; Huang, J.; Dufresne, A. Nanoscale 2012, 4, 3274–3294. Azizi Samir, M. A. S.; Alloin, F.; Dufresne, A. Biomacromolecules 2005, 6, 612–626. Eichhorn, S. J.; Dufresne, A.; Aranguren, M.; Marcovich, N. E.; Capadona, J. R.; Rowan, S. J.; Weder, C.; Thielemans, W.; Roman, M.; Renneckar, S.; Gindl, W.; Veigel, S.; Yano, H.; Abe, K.; Nogi, M.; Nakagaito, A. N.; Mangalam, A.; Simonsen, J.; Benight, A. S.; Bismarck, A.; Berglund, L. A.; Peijs, T. J. Mater. Sci. 2010, 45, 1–33. Dufresne, A. Nanocellulose: From Nature to High Performance Tailored Materials, 2nd ed.; Walter de Gruyter GmbH & Co. KG: Berlin/Boston, 2017. Moon, R. J.; Martini, A.; Nairn, J.; Simonsen, J.; Youngblood, J. Chem. Soc. Rev. 2011, 40, 3941–3994. Dufresne, A. Curr. Opin. Colloid Interface Sci. 2017, 29, 1–8. Dufresne, A. Mater. Today 2013, 16, 220–227. Taheri, H.; Samyn, P. Cellulose 2016, 23, 1221–1238. Taniguchi, T.; Okamura, K. Polym. Inter. 1998, 47, 291–294. Zhao, H. P.; Feng, X. Q.; Gao, H. Appl. Phys. Lett. 2007, 90, 073112. Kondo, T.; Koseb, R.; Naito, H.; Kasai, W. Carbohydr. Polym. 2014, 112, 284–290. Zhang, L.; Tsuzuki, T.; Wang, X. Cellulose 2015, 22, 1729–1741. Lavoine, N.; Desloges, I.; Dufresne, A.; Bras, J. Carbohydr. Polym. 2012, 90, 735–764. Henriksson, M.; Henriksson, G.; Berglund, L. A.; Lindström, T. Eur. Polym. J. 2007, 43, 3434–3441. Wågberg, L.; Decher, G.; Norgren, M.; Lindström, T.; Ankerfors, M.; Axnäs, K. Langmuir 2008, 24, 784–795. Saito, T.; Kimura, S.; Nishiyama, Y.; Isogai, A. Biomacromolecules 2007, 8, 2485–2491. Malainine, M. E.; Dufresne, A.; Dupeyre, D.; Mahrouz, M.; Vuong, R.; Vignon, M. R. Carbohydr. Polym. 2003, 51, 77–83. Siqueira, G.; Abdillahi, H.; Bras, J.; Dufresne, A. Cellulose 2010, 17, 289–298. Sassi, J. F.; Chanzy, H. Cellulose 1995, 2, 111–127. de Menezes, A. J.; Siqueira, G.; Curvelo, A. A. S.; Dufresne, A. Polymer 2009, 50, 4552–4563. Grunert, M.; Winter, W. T. J. Polym. Environ. 2002, 10, 27–30. Goussé, C.; Chanzy, H.; Excoffier, G.; Soubeyrand, L.; Fleury, E. Polymer 2002, 43, 2645–2651. Kloser, E.; Gray, D. G. Langmuir 2010, 26, 13450–13456. 151

27. Habibi, Y.; Dufresne, A. Biomacromolecules 2008, 9, 1974–1980. 28. Habibi, Y.; Goffin, A. L.; Schiltz, N.; Duquesne, E.; Dubois, P.; Dufresne, A. J. Mat. Chem. 2008, 18, 5002–5010. 29. Yi, J.; Xu, Q.; Zhang, X.; Zhang, H. Polymer 2008, 49, 4406–4412. 30. Zoppe, J. O.; Habibi, Y.; Rojas, O. J.; Venditti, R. A.; Johansson, L. S.; Efimenko, K.; Österberg, M.; Laine, J. Biomacromolecules 2010, 11, 2683–2691. 31. Favier, V.; Canova, G. R.; Cavaillé, J. Y.; Chanzy, H.; Dufresne, A.; Gauthier, C. Polym. Adv. Technol. 1995, 6, 351–355. 32. Dufresne, A. J. Nanosci. Nanotechnol. 2006, 6, 322–330. 33. Peng, Y.; Gardner, D. J.; Han, Y. Cellulose 2012, 19, 91–102. 34. Kamal, M. R.; Khoshkava, V. Carbohydr. Polym. 2015, 123, 105–114. 35. Velásquez-Cock, J.; Gañan, P.; Gómez, H.; Posada, P.; Castro, C.; Dufresne, A.; Zuluaga, R. Food Hydrocoll. 2018, 79, 30–39. 36. Dufresne, A. Philos. Trans. Royal Soc. A 2018, 376, 20170040. 37. Alloin, F.; D’Aprea, A.; Dufresne, A.; El Kissi, N.; Bossard, F. Cellulose 2011, 18, 957–973. 38. Goffin, A.-L.; Raquez, J.-M.; Duquesne, E.; Siqueira, G.; Habibi, Y.; Dufresne, A.; Dubois, P. Polymer 2011, 52, 1532–1538. 39. Mariano, M.; Pilate, F.; Borges de Oliveira, F.; Khelifa, F.; Dubois, P.; raquez, J.-M.; Dufresne, A. ACS Omega 2017, 2, 2678–2688. 40. Lewandowska, A. E.; Eichhorn, S. J. J. Raman Spectrosc. 2016, 47, 1337–1342. 41. Lin, N.; Dufresne, A. Nanoscale 2014, 6, 5384–5393. 42. Camarero Espinosa, S.; Kuhnt, T.; Foster, E. J.; Weder, C. Biomacromolecules 2013, 14, 1223–1230. 43. Ben Azouz, K.; Ramires, E. C.; Van den Fonteyne, W.; El Kissi, N.; Dufresne, A. ACS Macro Lett. 2012, 1, 236–240. 44. Pereda, M.; El Kissi, N.; Dufresne, A. ACS Appl. Mater. Interfaces 2014, 6, 9365–9375. 45. Nagalakshmaiah, M.; Pignon, F.; El Kissi, N.; Dufresne, A. RSC Adv. 2016, 6, 66224–66232. 46. Nagalakshmaiah, M.; El Kissi, N.; Dufresne, A. ACS Appl. Mater. Interfaces 2016, 8, 8755–8764. 47. Sapkota, J.; Kumar, S.; Weder, C.; Foster, E. J. Macromol. Mater. Eng. 2015, 300, 562–571. 48. Mariano, M.; El Kissi, N.; Dufresne, A. Langmuir 2016, 32, 10093–10103.

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