Ultrahigh Tough, Super Clear and Highly Anisotropic Nanofibers

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Ultrahigh Tough, Super Clear and Highly Anisotropic Nanofibers-Structured Regenerated Cellulose Films Dongdong Ye, Xiaojuan Lei, Tian Li, Qiaoyun Cheng, Chunyu Chang, Liangbing Hu, and Lina Zhang ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.9b02081 • Publication Date (Web): 03 Apr 2019 Downloaded from http://pubs.acs.org on April 3, 2019

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Ultrahigh Tough, Super Clear and Highly Anisotropic Nanofibers-Structured Regenerated Cellulose Films Dongdong Ye †,‡, Xiaojuan Lei †, Tian Li §, Qiaoyun Cheng †, Chunyu Chang †, Liangbing Hu *,§ and Lina Zhang *,† †

College of Chemistry and Molecular Sciences, Wuhan University, Wuhan 430072, China

‡ School

of Textile Materials and Engineering, Wuyi University, Jiangmen 529020, China

§ Department

of Materials Science and Engineering, University of Maryland, College Park, Maryland

20742, USA

*E-mail: [email protected] (L. Hu) *E-mail: [email protected] (L. Zhang)

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ABSTRACT: Despite that the tremendous efforts have been dedicated to developing environmentallyfriendly films made from natural polymers and renewable resources, in particular, multifunctional films featuring extraordinary mechanical properties, optical performance and ordered nanostructure, challenges still remain in achieving all these characteristics in a single material via a scalable process. Here, we designed a green route to fabricate strong, super tough regenerated cellulose films featuring tightly stacked and long-range aligned cellulose nanofibers self-assembled from cellulose solution in alkali/urea aqueous systems. The well-aligned nanofibers were generated by directionally controlling the aggregation of cellulose chains in hydrogel state using a pre-orientation-assisted dual crosslinking approach, i.e. a physical crosslinking was rapidly introduced to permanently reserve the temporaryaligned nanostructure generated by pre-orienting the covalent cross-linked gels. After a structural densification in air-drying of hydrogel, high strength was achieved, and more importantly, a recordhigh toughness (41.1 MJ m-3) in anisotropic nanofibers-structured cellulose films (ACFs) was reached. Moreover, the densely packed and well-aligned cellulose nanofibers significantly decreased the interstices in the films to avoid light scattering, granting ACFs with high optical clarity (91%), low haze ( 91%), low haze (< 3%), water and thermal stability, as well as attractive light management capacity. This method might be very scalable in fabricating regenerated cellulose films featuring tightly stacked and long-range aligned cellulose nanofibers, and solving the conflict between strength and toughness of anisotropic cellulose films, which should be very useful for guiding the controllable construction of nanofibers structure in other polysaccharide-based film materials and broadening their applications in biodegradable packaging, flexible electronic and optoelectronic


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material. RESULTS AND DISCUSSION We prepared anisotropic cellulose films featuring tightly stacked and long-range aligned nanofibers using a pre-orientation-assisted dual crosslinking strategy followed by structural densification approach (Fig.1; Fig. S1). The chemical gels were synthesized by reacting high clarity cellulose solution (Fig. S2) with certain amounts of a chemical crosslinker, epichlorohydrin (EPI),42,43

Fig. 1. Schematic image of the bottom-up approach for fabricating ACFs. By directionally controlling the architecture of the cellulose chains with a facile pre-stretching strategy in the chemical cross-linked gel state followed by locking the aligned nanostructure with physical crosslinking (hydrogen bonds) among neighboring cellulose chains after destroying the alkali/urea solvent shells on the cellulose chains in acidic solution. The resulting anisotropic cellulose hydrogel combines a long-range aligned structure and dual-crosslinking networks. ACFs were constructed by evaporating water from the anisotropic hydrogels via air-drying.


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which can then be pre-stretched to form a temporarily oriented structure (Fig. 1; Fig. S3A, B). Subsequently, anisotropic cellulose hydrogels were obtained by immersing the pre-stretched chemical gel in a sulfuric acid aqueous solution to destroy the alkali/urea solvent shells around the cellulose chains and permanently fix the re-aligned structure with physical crosslinking (hydrogen bonds) between neighboring cellulose chains. Then, bright birefringence colors appeared in the highly transparent anisotropic hydrogel fabricated with above pre-orientation-assisted dual crosslinking strategy under polarized light due to permanently frozen nanofibers structure confirmed by AFM, compared to the isotropic hydrogel without pre-orientation (Fig. S3C, D). Importantly, anisotropic nanofibers-structured cellulose films, coded as ACFs, were constructed by evaporating water from the anisotropic hydrogels using a facile air-drying process. Subsequently, we in detail monitored the surface and cross-sectional morphologies of the anisotropic hydrogels before and after air-drying (Fig. 2). The cellulose hydrogel exhibited a highly aligned and porous structure along the pre-stretching direction in the XY and YZ planes in scanning electron microscopic images (Fig. 2A), and aligned hierarchical fibrous structures with sub-microscale hole-like aisles in the XZ plane, confirming the locking of the oriented structure. After the air-drying process, ACFs showed a densified and ordered nanofiber structure along both the surface and crosssection (Fig. 2B; Fig. S4). The magnified surface SEM image exhibited extremely low surface roughness, with a densely packed, interlocked, and well-aligned nanofibers structure along the prestretching direction in the XY plane (Fig. 2B; Fig. S4A). Furthermore, cross-sectional SEM images taken in parallel (YZ plane) and perpendicular (XZ plane) to the film orientation displayed an assembled and aligned nanofibers compact structure (Fig. 2B; Fig. S4B), endowing the anisotropic films with reduced internal defects, which might enhance the mechanical performance of the material. Besides, the surface morphologies of the ACFs with increased pre-stretching strains were further investigated by atomic force microscope (AFM) (Fig. 2C). Apparently, the ACF with 40% prestretching strain exhibited a slightly oriented nanostructure along the pre-stretching direction in the


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Fig. 2. Morphology Characterization of structural densification in air-drying of anisotropic hydrogel. (A, B) SEM images of the anisotropic hydrogel and film, showing the densification of the air-drying process and the well-aligned nanofibers in the resulting anisotropic film. The diameter distribution of the cellulose nanofibers is shown in the insert of the cross-sectional SEM image of the ACF’s XZ plane. (C) AFM images of the surface of the anisotropic cellulose films with various prestretching strains, including 40%, 80%, 120%, and 160%. magnified region, whereas as the strain increased to 120%, the nanofibers became closely aligned, leading to a tendency to assemble and stack with each other into an extremely compact nanostructure. Importantly, tightly stacked and well-oriented nanofibers with denser structures appeared in the 160% strain sample, as well as demonstrating an extremely low surface roughness (Fig. 2C) and extremely stable in water, even under ultrasonic treatment (Fig. S5A and Movie 1). What’s more, the submerged


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dual-cross-linked films still maintained the strength in order of Megapascal as long as three weeks (Fig. S5B, C), further confirming the feasibility of our rational design in fabricating the anisotropic film. The ordered structure of the ACFs was studied with both wide-angle (WAXS) and small-angle Xray scattering techniques (SAXS) in a synchrotron radiation facility (Fig. 3). For the film without prestretching, there was a nearly uniform diffraction pattern in the WAXS pattern at all azimuthal angles on the (110) scattering plane, and a ring with nearly uniform intensity in the SAXS pattern, revealing the isotropic structure of the film. In contrast, ACFs made using the pre-stretching process displayed clear equatorial arcs in the WAXS patterns and elongated longitudinal 2D SAXS patterns, suggesting

Fig. 3. Well-ordered nanostructure of ACFs confirmed by X-ray scattering technology. (A) 2D WAXS patterns of ACFs with different pre-stretching stains, including 0%, 40%, 80%, 120% and 160%. The arrows represent the pre-stretching direction. (B) 2D SAXS results of ACFs with different pre-stretching strains. (C) Azimuthal-integrated intensity distribution curves of the WAXS patterns, where 0° represents the perpendicular direction. (D) Effect of pre-stretching strains on Herman’s orientation parameter (fc) in WAXS. (E) Azimuthal-integrated intensity distribution curves of the SAXS patterns. (F) Effect of pre-stretching strains on orientation parameter in SAXS.


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anisotropic properties as a result of the aligned nanofibers. Additionally, we calculated the Herman’s orientation parameter (fc), which can quantitatively indicate the degree of orientation from the azimuthal-integrated intensity distribution curves of the X-ray scattering patterns according to Equations (1-1) and (1-2) (Methods section). With an increase in the pre-stretching strain from 0% to 160%, fc of the ACFs gradually increased from 0.034 to 0.88 and 0.021 to 0.89 according to the WAXS and SAXS results, respectively. The high fc values of the ACFs under large pre-stretching strain further demonstrated the existence of a long-range aligned nanofiber architecture along the pre-stretching direction. These results strongly confirmed the feasibility of locking the structural orientation of the hydrogel state via the pre-orientation-assisted dual crosslinking strategy and structural densification in air-drying process for preserving the nanofiber structures in films. Notably, the ACF exhibited a uniformly oriented nanostructure (Fig. 4). Three different areas in a ribbon-shaped 160% pre-stretched ACF displayed similarly aligned nanostructures and roughness, with the tightly stacked cellulose nanofibers oriented in the pre-stretching direction, resulting in nearly identical 2D WAXS patterns and appropriate fc values at the respective regions. The above results demonstrated the advantage of pre-orientation-assisted dual crosslinking strategy in successfully translating the ordered nanoscale building blocks to long-range aligned macroscopic bulk materials. To clarify the impact of the pre-orientation-assisted dual crosslinking strategy and structural orientation on the mechanical properties of the cellulose films, we first measured the tensile stressstrain performance of the isotropic cellulose films, including the physically, chemically, and dualcross-linked cellulose films (Fig. 5A-C). Obviously, the tensile strength and Young’s modulus of the physically cross-linked films were 106 MPa and 4.31 GPa, respectively, which were significantly higher than those of the chemically cross-linked films (81.8 MPa and 3.47 GPa, respectively). However, both the ruptured strain and work of fracture (a parameter that expresses one aspect of toughness) of the physically cross-linked film were much lower than those of the chemically crosslinked material, changing from 12.4% to 23.5% and 9.8 MJ m-3 to 15.3 MJ m-3, respectively. These


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results demonstrated that the physical cross-linking appears to contribute to the enhancement of the strength and stiffness of the cellulose film, whereas chemical crosslinking contributes to the enhanced extensibility, resulting in a relatively high toughness. After introducing both physical and chemical crosslinking, the dual-cross-linked film exhibited the highest fracture strain (44.1%) and an extremely high work of fracture (36.1 MJ m-3) (Fig. 5C). These results demonstrate that the dual-cross-linking strategy indeed contributes to the toughness of the material.

Fig. 4. The unifromly oriented nanostruture of ACFs. Birefringence patterns of a ribbon-shaped ACF sample pre-stretched 160%, and the corresponding internal oriented nanofibrous structure monitored by AFM and WAXS. A: analyzer; P: polarizer. The azimuthal-integrated intensity distribution curve of the WAXS pattern and the calculated fc appear in the inset of the corresponding 2D WAXS pattern.


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The effect of the chemical crosslinker dosage on the mechanical properties of dual-cross-linked cellulose films were also investigated, demonstrating that the dual-cross-linked film with an EPI/ anhydroglucose unit (AGU) molar ratio of 1:3.09 had the best ductility (Fig. S6). Additionally, we studied mechanical properties of the physically, chemically, and dual-cross-linked films in terms of how their respective crosslinking strategies affected their crystalline structure and film morphologies (Fig. S7, 8). Obvious changes in the characteristic NMR peaks (blue region: 89 ppm and 63 ppm) for the crystalline regions of the three films were observed, revealing the reduced crystallinity of the chemically cross-linked film due to a holistic weakening of the intramolecular hydrogen bonds between cellulose (Fig. S7). Consequently, dual crosslinking strategy maintains small amounts of crystalline regions among large amorphous regions, which synergistically enhance the toughness of the film. Besides, the crosslinking strategies greatly affected the normal morphologies and fractured morphologies of cellulose films. The surface morphology of isotropic cellulose films were shown in Fig. S8, where all specimens exhibited homogeneous structure assembled by nanofibers but showing different density of aggregation and defects, resulting in their mechanical difference and distinguishing fractured morphologies. Physical cross-linking film with the relatively dense nanofibers structure and the most crystallinity could withstand large ruptured strength but small fracture strain, whereas chemical cross-linking film fabricated from loosely network structured hydrogels (Fig. S9), featuring a relatively loose nanostructure and weak mechanical properties, barely resisted the large external force due to more defects in films, as shown in Fig. S8. After tensile tests, both the physical and chemical films displayed deep cracks (Fig. S8), suggesting their easy fracture. However, the dual crosslinked film containing nanofibers and chemical cross-linking networks exhibited relatively thin cracks, indicating a larger fracture strain, consistent with good toughness, as shown in Fig. 5B. In view of the above results, the dual-cross-linking structure contributed the both dissipating mechanical energy and the enhancing extensibility, leading to the good stress dispersion, thus the strain was improved while strength was preserved well. So the dual crosslinking strategy accompany with well-chosen EPI/AGU


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molar ratio (EPI: AGU=1:3.09) were chosen to fabricate the highly aligned cellulose films. Importantly, the structural alignment of the cellulose nanofibers could endow the dual-crosslinked films with improved tensile strength and toughness (Fig. 5D, E). Specifically, the tensile strength of the ACFs greatly increased to 253 MPa with an increase from pre-stretching values by 120%, which is 2.6-times higher than the isotropic dual-cross-linked film. This is attributed to the compact structure assembled by the ordered nanofibers in ACFs for stress transfer. Thus, the toughness of the cellulose films reached 41.1 MJ m-3, exceeding the pre-stretching strain by 40% and thereby revealing that both the strength and toughness of the cellulose films can be simultaneously enhanced by the synergetic effect of dual-crosslinking and the nanostructured orientation. Furthermore, the intrinsically light weight of cellulose resulted in a specific strength of the ACFs (217±15 MPa g-1 cm3, mean ± standard deviation), which was comparable to that of lightweight titanium alloy44 (nearly 244 MPa g-1 cm3), although no more than the densified wood5 and aligned BC film19 (Fig. S10). As summarized in Fig. 5F, the ductility of the anisotropic dual-cross-linked films in this work was higher than all other reported neat or modified isotropic and anisotropic cellulose films.5,19,20,22,23,25-29,36,45,46 In view of these results, the ultrahigh toughness, strong, and clarity of the ACF films were relative closely to the combination of the aligned nanofibril structure and the dual-crosslinking strategy, leading to the simultaneous strengthening and toughening in a regenerated cellulose film. The above tensile stress-strain tests were carried out with professional equipment and reasonable measuring parameters in order to avoid slipping or fracture in the clamp (Fig. S11).


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Fig. 5. Mechanical properties of the isotropic and anisotropic cellulose films. (A) Schematic illustration of the fabrication of the isotropic cellulose films, including those that had been chemically, physically, and dual-cross-linked. (B) Tensile stress-strain curves of the isotropic cellulose films. (C) The toughness of the isotropic physically, chemically, and dual cross-linked films, demonstrating that the dual-crosslinking strategy indeed contributes to the toughness of the dual-cross-linked film compared to physically and chemically cross-linked films. (D) Stress-strain curves of the dual-crosslinked films with various pre-stretching strains (0%, 40%, 80% and 120%). (E) The ruptured stress and toughness of the ACFs. (F) The mechanical properties of the ACFs and other cellulose-based films, including densified wood,5 anisotropic and isotropic wood-based cellulose film,20 cellulose nanofiber (CNF) film,36,45 anisotropic bacteria cellulose (BC) film,19 cellulose nanocrystal (CNC) film,22,46 microfibrillated cellulose (MFC) film,23 and commercial cellophane film25,28 and regenerated cellulose film26-29.

The thermo-mechanical properties of the ACFs were also evaluated (Fig. S12) and the tensile storage modulus (E’) as a function of temperature were determined by dynamic mechanical analysis (DMA). Apparently, the anisotropic cellulose films displayed significant mechanical reinforcement, for example, the E’ value of ACF-160% film was 23.3 GPa at 25 °C, which is 3.4-times that of the isotropic dual-cross-linked cellulose films (6.94 GPa). Additionally, the E’ of the isotropic film and


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anisotropic films with pre-stretching strain of 160% decreased from 14.7 GPa and 44.7 GPa at -100 °C to 2.9 GPa and 10.5 GPa at 250 °C, respectively, demonstrating that the oriented nanofibers dramatically enhanced the thermal and mechanical stability. Moreover, thermogravimentric analysis demonstrated that the ACFs had a relatively high decomposition temperature in the range of 300– 350 °C, indicating an excellent thermal stability (Fig. S13). Tan δ, i.e. the ratio of loss modulus to storage modulus, was measured to reflect a loss of energy due to the nature relaxation processes of films. Actually, the transition temperature of an isotropic film was -46.0 °C and increases to -41.6 °C in ACF-160% (Fig. S13), further confirming that the densely packed and intertwined nanofibers enhanced the attractive interaction between the cellulose, resulting in the transition of the aggregated structure and enhanced crystallinity after increasing the orientation degree, supported by X-ray diffraction (XRD) (Fig. S14). Furthermore, the tightly stacked and long-range aligned cellulose nanofibers in ACFs cause less light scattering to occur within films and allow more light to pass through it, which not only endows ACFs with a high transparency of 91.2% (Fig. 6A and Fig. S15), which is even higher than polyethylene (PE, < 88.6%), polyethylene terephthalate (PET, < 84.2%), the corresponding isotropic film (90.5%), and anisotropic cellulose hydrogel (87.3%), but also endowed the anisotropic films with an extremely low haze (2.1% at 800 nm), as shown in Fig.6B, similar to the aligned BC films19, aligned wood films20 and regenerated cellulose films26,28,29 (Fig. S15). Importantly, the haze of the resultant ACF was significantly lower than other cellulose-based composites (Fig. 6C).18,19,21,47, 48 When the ACF film was placed between crossed polarizers and rotated from 0–90° (Fig. 6D), a polarized He-Ne laser beam that passed through the ACFs having pre-stretching strains of 40%, 80%, 120%, and 160%, respectively. Their birefringence initially increased to reach a maximum power at 45° and then gradually decreased to a minimum at 90° (Fig. 6E), whereas that of the isotropic film hardly changed from around 0 mW. This reveals that the ACFs with oriented nanostructures could depolarize the beam so as to make the polarized laser pass through the analyzer and be collected by


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Fig. 6. Anisotropic cellulose films for light management. (A) The transmittance of the ACF, isotropic film, anisotropic hydrogel, and PE and PET plastics. (B) The haze of the isotropic and anisotropic films. Inset: photograph of ACF. (C) The haze of the ACF film in this work is shown to be lower than that of paper,48 anisotropic wood-based nanopaper,18 wood composites,21 anisotropic and isotropic BC film,19 nanopaper,47 and is similar to that of glass.47 (D) Schematic illustration of the angle variation of the anisotropic films between the crossed polarizers. (E) Depolarized laser power as a function of rotation angle (0–90°) of the anisotropic films with various pre-stretching strains from 0% to 160%. (F) Representative photographs showing the differences of the birefringence patterns with the rotation of the anisotropic film with different stacking densities of the cellulose nanofibers under the crossed polarizers. A: analyzer; P: polarizer. (G) Birefringence and polarization features of the anisotropic films observed between the crossed polarizers, indicating the ACFs feature a light management ability in depolarizing the polarized beam to successfully display the overlaid pattern in crossed polarizer. the detector. Moreover, the laser power enhanced with an increase in the nanostructural orientation under the same rotation angle (Fig. 6E). For example, the polarized laser could penetrate the anisotropic film with 160% pre-stretching strain and analyzer to reach 3.95 mW at 45°, which was nearly 4.43–times that of the anisotropic film with 40% pre-stretching at the same angle, demonstrating


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that the orientation degree of the nanofibers in the cellulose films had a great influence on the depolarization effect. Furthermore, a wavelength selectivity from yellow, blue, to red-purple was observed with the film made with different stacking densities under the crossed polarizers, with a mass density of 0.684 ± 0.019 g/cm3, 1.087 ± 0.011g/cm3 and 1.245 ± 0.043 g/cm3. To visually display anisotropic films between the crossed polarizer, ACFs with increased mass density exhibited similar angular dependence and differential iridescent birefringence patterns, resulting from the differential in degree of orientation (Fig. 6F; Movie 2). As shown in Fig.6E, the depolarized light intensity reached a maximum at 45°, consistent with the above-mentioned result. Given the highly oriented nanofibers, the film exhibited bright iridescent color under polarized light with unrepentantly achieved wavelength selection effect with cellulose. As shown in Fig. 6G, the overlaid pattern in the crossed polarizer displayed with the assistance of the ACFs further confirmed that the oriented nanostructure in the ACFs endowed the cellulose films with potential applications in light management. In view of these results mentioned above, the self-assembled cellulose nanofibers played an important role in outstanding mechanical strength and extraordinary optical clarity of ACFs. Moreover, cellulose as carbohydrates can be used as raw materials for manufacture various sustainable polymers materials.49 More importantly, the regenerated cellulose films having excellent mechanical properties and transparency can be complete biodegradable, and their degradation products are water and carbon dioxide on the whole.29,50 Therefore, this work not only greatly broadened the cellulose applications from next-generation packaging to emerging flexible electronic and optoelectronic materials, and also achieved a virtuous cycle in nature due to the biodegradability in soil and at sea. This high-performance anisotropic cellulose film is a next generation of environmentally friendly material that will lead to the implementation of a future sustainable society. CONCLUSION In this work, construction of strong, ultrahigh tough, super clear, and biodegradable regenerated cellulose films was realized from the cellulose solution via a bottom-up route by combining a


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chemically and physically dual-cross-linked structure with highly ordered nanofibers. The anisotropic nanofibers-structured cellulose films (ACFs) were fabricated successfully from transparent cellulose solution in an alkali/urea aqueous systems, based on directionally controlling and freezing the nanofibrous architecture formed from cellulose chains in the hydrogel state via pre-orientation-assisted dual-crosslinking, followed by a structural densification during air-drying. The resultant ACFs exhibited record-high toughness (41.1 MJ m-3), high transparency (> 91%) and low haze (< 3%), polarization effect, water and thermal stability. In our findings, the rational combination of the nanofibers, and dual-crosslinking networks contributed the both dissipating mechanical energy and the enhancing extensibility, leading to the good stress dispersion, thus their strain was improved while strength was preserved well. The densely packed cellulose nanofibers significantly decreased the interstices in the film to avoid light scattering, leading to the high-performance optical properties in both of transmittance and haze. Moreover, the oriented nanostructure endowed the film with depolarizing a polarized beam and wavelength selectivity for optical polarizer. Our work shows an approach to fabricate extreme tough and super clear cellulose films with long-range ordered nanofibers, and greatly broadens the cellulose applications from next-generation packaging to emerging flexible electronic and optoelectronic materials, and as environmentally friendly material also will lead a virtuous cycle in nature. METHODS Materials. Cellulose (cotton linter pulp) with an α-cellulose content of > 95% was provided by Hubei Chemical Fiber Group Ltd. (Xiangfan, China). The cotton linter pulp was used after complete drying under vacuum at 60 °C and without further purification. The viscosity-average molecular weight of the cotton linter pulp in cadoxen was determined to be 9.6104 (degree of polymerization, DP=600) according to the Mark-Houwink equation [η] (mL g-1) = 3.8510-2 (Mw)0.76 and using an Ubbelohde viscometer at 25 °C. LiOH·H2O, urea, concentrated sulfuric acid, and epichlorohydrin (EPI) were purchased from Sinopharm Chemical Reagent Co., Ltd. (Shanghai, China). All of the reagents were


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used as received unless otherwise noted. Fabrication of the Isotropic Cellulose Hydrogels and Films. Isotropic cellulose films, including the chemically, physically, and dual-cross-linked samples, were fabricated through air-drying of chemically cross-linked hydrogel, physically cross-linked hydrogel and dual-cross-linked hydrogels without other treatment as Figure. S5A. The chemically cross-linked cellulose hydrogel was fabricated as reported in our previous work42: cellulose solution (6 wt%) was prepared by dissolving cotton linter pulp in a precooled 4.5 wt% LiOH/15 wt% urea aqueous solution (-12°C) within 2 min, to which EPI was added dropwise to a ratio of EPI/AGU (anhydroglucose unit) =1:1.33 at 0 °C for homogeneous mixing (2 h). The resultant solution was transferred into hand-made cuboid-shaped molds and kept at 5 °C for 8 h to form the chemical gels. After washing by distilled water, the colorless and transparent chemical hydrogels were prepared. The physically cross-linked cellulose hydrogel was fabricated as in our previous work28: cellulose solution (6 wt%) was prepared by dissolving cotton linter pulp in a 4.5 wt% LiOH/15 wt% urea aqueous solution (-12°C) within 2 min. After removing air bubbles via centrifugation, cellulose solution was directly cast onto a glass plate to give a thickness of 0.20–0.30 mm and then immediately coagulated in 5 wt% aqueous sulfuric acid for 1 min to obtain the physical hydrogels. Isotropic dual-cross-linked cellulose hydrogels were fabricated as follow: different amounts of EPI corresponding to EPI:AGU ratios of 1:9.28, 1:3.09, 1:1.86, and 1:1.33 were added to the cellulose solution (6 wt%) while stirring at 0 °C for 2 h. Then the resulting solution was poured into molds and kept at 5 °C for 8 h to gel. Subsequently, the as-prepared chemically cross-linked gels were directly immersed in 5 wt% aqueous sulfuric acid to remove the alkali/urea solvent sheath around the cellulose chains. After thoroughly washing with water, the dual-cross-linked isotropic cellulose hydrogels were obtained. Fabrication of Anisotropic Dual-cross-linked Cellulose Hydrogels and Films. Anisotropic dualcross-linked cellulose hydrogels were fabricated as shown in Fig. S1. EPI (EPI:AGU=1:3.09) was added into the bubble-free 6 wt% cellulose solution mentioned above with stirring at 0 °C for


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homogeneous mixing (2 h). Then, the resulted cellulose solution was poured into molds, and kept at 5 °C for 8 h for gelation. Subsequently, the as-prepared chemical cross-linked gels (containing alkaliurea solvent) with the elastic network were pre-stretched to specific strains and then immersed in the 5 wt% aqueous sulfuric acid to remove the alkali/urea solvent shell around the cellulose chains, resulting in permanent “freezing” of the highly oriented structure in the as-prepared hydrogels through the directional control of the hydrogen bonding. Ultimately, anisotropic cellulose hydrogels, with specific pre-stretching strains, were obtained after thoroughly washing the samples with deionized water. Anisotropic dual-cross-linked cellulose films were obtained by air-drying the hydrogel materials. Characterization. AFM images were recorded on a CypherTM S (Asylum Research) using a silicon nitride probe (RTESP-300, BRUKER) with a tip radius of 2 nm, a spring constant of 40 N/m, and a high resonance frequency of 300 kHz. All imaging results were analyzed using the AFM accessory software (Gwyddion). SEM images of hydrogels and films were taken with an field emission scanning electron microscopy (FESEM, Zeiss, SIGMA, Germany) using an accelerating voltage of 5 kV. Before testing, cellulose hydrogels were frozen in liquid nitrogen, snapped immediately, and freeze-dried to remove the water. Both the surfaces and cross-sections of the dried hydrogels and films were coated with gold vapor prior to imaging. DMA temperature sweeps under oscillatory stress were performed on the films using a DMA Q850 (TA Instruments, USA) in tension mode at a heating rate of 3 °C min-1 over a temperature range of -100 to 250 °C with a frequency of 1 Hz. 2D WAXD measurements were performed on a WAXD diffractometer (D/MAX-1200, Rigaku Denki, Japan). 2D SAXS profiles were obtained at the beamlike BL16B1 of the Shanghai Synchronization Radiation Facility. The wavelength used was 0.124 nm, and a CCD X-ray detector (MAR CCD 165, 2048 2048 pixels with pixel size of 80.5 m) was employed to collect data. The sample-to-detector distances were calibrated to 2540 mm by beef tendon. The data acquisition time was 30 s, and the 2D scattering images were analyzed with Fit2D software from the European Synchronization Radiation Facility. Herman’s orientation


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parameter (fc) was calculated from the azimuthal-integrated intensity distribution curves of the X-ray scattering patterns with the equations of

𝑓c =

3 < cos2𝝓 > ―1 2

(1-1) and

< cos2𝝓 > =

π/2

∫0 I(𝝓)cos2𝝓sin𝝓d𝝓 π/2

∫0 I(𝝓)sin𝝓d𝝓

𝝓 (1-2), where 𝝓 is the azimuthal angle, and 𝑰 (𝝓) is the 1-D intensity distribution

along with 𝝓. is calculated by integrating the intensity of the specific 2θ diffraction peak along 𝝓. Solid-state

13C

cross-polarization magic angle spinning (13C CPMAS) nuclear magnetic

resonance (NMR) experiments were carried out on a Bruker AVANCE III WB spectrometer (Rheinstetten, Germany). X-ray diffraction (XRD) spectra were recorded on a Riguka smartlab 9K diffractometer operated at 45 kV, 200 mA in reflection mode for Cu K radiation ( = 0.154 nm) with a scan speed of 2° min-1 and a step size of 0.05° in 2. Tensile tests of the cellulose films were performed on an INSTRON instrument (Model 5576, USA) equipped with pneumatic clamps so as to avoid slippage or fracture of the samples within the sample holder (Fig. S11). The rectangular cellulose films (0.6–1 mm thick and 10 mm wide) were stretched at a speed of 3 mm min-1. The light transmittance of the cellulose films was measured by a double-beam UV-visible spectrophotometer (Mapada UV-P7, China) with a wavelength ranging from 200 to 800 nm. The angular dependence between the ACFs and the depolarized light was measured on a liquid crystal device parameter tester (LCT-5160, the north LCD engineering R&D center, China). ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI:xxxx. Fabrication of ACFs from cotton linter pulp; image of the highly transparent cellulose solution; images of the chemical gel, pre-oriented chemical gel, isotropic and the anisotropic dual-cross-linked cellulose hydrogels under bright and polarized light; AFM images of the surface of the isotropic and anisotropic cellulose hydrogel; stability of the ACFs in water and ultrasonic treatment; the mechanical properties


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of the submerged ACFs as a function of times; surface and cross-section morphology of ACF–160% at different magnifications; structural morphology of chemically, physically, and dual-cross-linked cellulose films before and after tensile tests; 13C NMR data of chemically, physically, and dual-crosslinked cellulose films; image, structure and mechanical property of chemical hydrogel; comparasion the specific tensile strength of the ACFs with wood, densified wood, aligned BC film and titanium alloy; mechanical properties of the isotropic dual-cross-linked cellulose films with chemical crosslinking density; mechanical test instrument equipped with pneumatic zigzag clamping, sawtoothlike clamp and controllable air pressure for measuring the ACFs; dynamic thermal mechanical analysis test of the isotropic and anisotropic cellulose films with pre-stretching strain of 40%, 80%, 120% and 160%; thermalstability of ACFs; X-ray diffraction profiles of ACFs; comparasion the transparency of the ACFs with other cellulose-based films (PDF) Movie 1. Stability of anisotropic nanocellulose films in water under ultrasonic treatment (WMV) Movie 2. Highly transparent anisotropic nanocellulose films for light management (WMV) AUTHOR INFORMATION Corresponding Authors *[email protected] * [email protected] ACKNOWLEDGMENTS D. Ye, X. Lei and T. Li contributed equally to this work. This work was supported by the Major International (Regional) Joint Research Project of National Natural Science Foundation of China (21620102004) and scientific research startup funds for high-level talents of Wuyi University (AL2018010). REFERENCES


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