New Approach for the Fabrication of Carboxymethyl Cellulose

Research Article Cite This: ACS Sustainable Chem. Eng. 2019, 7, 11850−11860

pubs.acs.org/journal/ascecg

New Approach for the Fabrication of Carboxymethyl Cellulose Nanofibrils and the Reinforcement Effect in Water-Borne Polyurethane Dan Cheng,†,§ Pingdong Wei,†,§ Lina Zhang,†,§ and Jie Cai*,†,‡,§ †

College of Chemistry and Molecular Sciences, Wuhan University, Wuhan 430072, China Shenzhen Research Institute, Wuhan University, Shenzhen 518057, China § Hubei Engineering Research Center of Natural Polymers-based Medical Materials, Wuhan University, Wuhan 430072, China Downloaded via UNIV OF SOUTHERN INDIANA on July 28, 2019 at 01:37:01 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

‡

S Supporting Information *

ABSTRACT: Nanocellulose has attracted significant attention due to its fascinating properties and great potential in the preparation of highperformance functional materials for specific end-use applications. Most cellulose nanocrystals and cellulose nanofibrils are in the cellulose I crystal structure; however, attempts to fabricate cellulose nanofibrils with the cellulose II crystal structure by self-assembly from cellulose solutions are rarely reported. Here, we demonstrate a new approach to the fabrication of carboxymethylated cellulose nanofibrils (CMCNFs) by homogeneous carboxymethylation and the self-assembly of cellulose chains from cellulose solutions using sequential “top-down” and “bottom-up” strategies. Furthermore, CMCNFs were used as reinforcing nanofillers to enhance the properties of water-borne polyurethane (WPU) and were evenly distributed in the WPU matrix. Dynamic thermomechanical analysis on the WPU/CMCNF nanocomposites demonstrated a significant reinforcement effect based on the tensile storage modulus above the glass transition temperature of the soft segments of the WPU, and the results were well-fit by the percolation model. The sequential top-down and bottom-up strategies developed here produce CMCNFs and should contribute to the reinforcement of polymer matrices, improving the practical exploitation of CMCNFs. KEYWORDS: Cellulose solution, Carboxymethyl cellulose nanofibrils, Percolation model, Water-borne polyurethane, Nanocomposites

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improve the nanofibrillation of cellulose.31−37 However, all of these CNCs and CNFs are in the cellulose I crystal structure. Additionally, CNFs with the cellulose II crystal structure were isolated from cellulose microfibrils from a combination of mercerization and mechanical treatment.38−40 However, attempts to fabricate cellulose nanofibrils with the cellulose II crystal structure by self-assembly from cellulose solutions are rarely reported. Here, we demonstrate a new approach to fabricate carboxymethylated cellulose nanofibrils (CMCNFs) by homogeneous carboxymethylation and self-assembly of cellulose chains from a cellulose solution using sequential “top-down” and “bottom-up” strategies. We previously found that cellulose can be dissolved in aqueous LiOH/urea solution by destroying the intra- and intermolecular hydrogen bonding interactions between the cellulose microfibrils.41,42 This top-down approach, the dissolution of cellulose, leads to the formation of semistiff cellulose chains in the cellulose solution. Further homogeneous carboxymethylation of the cellulose can be conducted by partial etherification of the hydroxyl groups of

INTRODUCTION With the progress and development of human society, research and advances in renewable and sustainable biomass-based materials benefit economies and the environment.1 Nanocellulose has attracted significant attention due to its fascinating properties, such as high specific strength and modulus, high specific surface area, high chemical reactivity, low coefficient of thermal expansion, and excellent biocompatibility and biodegradability.2−6 Nanocellulose has great potential in the preparation of high-performance functional materials for specific end-use applications, including separation membranes, packing materials, flexible electronic devices, biological platforms, polymer nanocomposites, and energy storage materials.7−16 The conversion of micrometer-scale cellulose into cellulose nanocrystals (CNCs) and cellulose nanofibrils (CNFs) with or without modification is achieved primarily through physical and chemical approaches, including strong acid hydrolysis,17−19 high-pressure homogenization or microfluidization,20−22 (2,2,6,6-tetramethylpiperidin-1-yl)oxyl (TEMPO)catalyzed oxidation,23,24 mechanochemical esterification,25,26 carboxymethylation,27,28 and nitro-oxidation.29,30 The pretreatment of cellulose microfibrils with solvent, including ionic liquids and deep eutectic solvents, was also introduced to © 2019 American Chemical Society

Received: May 1, 2019 Revised: May 30, 2019 Published: June 7, 2019 11850

DOI: 10.1021/acssuschemeng.9b02424 ACS Sustainable Chem. Eng. 2019, 7, 11850−11860

Research Article

ACS Sustainable Chemistry & Engineering

Figure 1. (a) Schematic illustration of the preparation of CMCNFs. (i) Dissolution of cellulose microfibrils in an aqueous LiOH/urea solution. (ii) Homogeneous carboxymethylation of cellulose chains in a cellulose solution. (iii) Self-assembly of carboxymethylated cellulose chains to give carboxymethylated cellulose nanofibrils. Photographs of a CMCNF suspension at 0.1 wt % (b) and concentrated CMCNFs at 0.4 wt % concentration (c). TEM image of isolated CMCNFs (d), and SEM image of a CMCNF aerogel obtained by freeze-drying the CMCNF suspension (e; inset, photograph of the CMCNF aerogel). Transparent CMCNF film prepared from the CMCNF suspension (f). from 3 to 48 h. The reaction temperature was fixed at 55 °C. Then, an aqueous 0.5 wt % NaCl solution was added to the reaction mixture under continuous stirring at room temperature.43 The mixture was subsequently transferred to a dialysis bag and equilibrated in deionized water for several days to remove residual chemical reagents to give a CMCNF suspension. Finally, the resulting CMCNF suspension was concentrated by reverse dialysis to 0.4 wt % solid content using 20 wt % PEG2000 solution, and the suspension was used without drying to avoid irreversible aggregation. Besides, decompress filtration, solvent evaporation, and rotary distillation could also be utilized for concentration of the CMCNF suspensions. Preparation of WPU/CMCNF Nanocomposites. The CMCNF dispersion was mixed with a WPU emulsion and ultrasonicated for 10 min to obtain WPU/CMCNF suspensions with different compositions, and the solutions were then stirred under vacuum for 1 h to remove air bubbles. Subsequently, the obtained mixtures were cast into Teflon molds and vacuum-dried at 60 °C for 3 h to achieve WPU/CMCNF nanocomposites. A series of WPU/CMCNF nanocomposites each with a thickness of approximately 0.1 mm were fabricated by changing the CMCNF content (1, 5, 10, 15, and 20 wt %), and the samples are denoted WPU/CMCNF-1, WPU/CMCNF2, WPU/CMCNF-3, WPU/CMCNF-4, and WPU/CMCNF-5, respectively. Characterization. Fourier transform infrared (FTIR) spectra were recorded on an FTIR spectrometer (Nicolet 5700 FTIR Spectrometer, MA) in the wavenumber range from 4000 to 400 cm−1 with a 2 cm−1 resolution and an accumulation of 32 scans. Solid-state 13C cross-polarization/magic angle spinning (CP/MAS) NMR spectra were collected using a Bruker Advance III (300 MHz) NMR spectrometer. The spinning speed and contact and delay times were 5 kHz, 20 ms, and 1 s, respectively. In addition, 2000 scans were performed. Wide-angle X-ray diffraction (WAXD) measurements were made on a WAXD diffractometer (D8-Advance, Bruker Corp.) using Nifiltered Cu Kα radiation with a wavelength of 1.542 Å. The voltage and current were set at 40 kV and 40 mA, respectively. The proportional counter detector was set to collect data at a rate of 3° min−1 over the 2θ range from 5° to 40°.

the cellulose chains with a controllable surface charge and degree of substitution. Subsequently, the bottom-up approach allows the formation of CMCNFs by the self-assembly of carboxymethylated cellulose chains via the reconstruction of the hydrogen bonding, hydrophobic interactions, and electrostatic repulsion. Furthermore, we demonstrated that the CMCNFs could be used as reinforcing nanofillers to enhance the properties of water-borne polyurethane (WPU) and be finely distributed in the WPU matrix. We performed dynamic thermomechanical analysis on the WPU/CMCNF nanocomposites to understand the reinforcement effect of CMCNFs with a percolation model. The sequential topdown and bottom-up strategies developed here provided the desired CMCNFs and will facilitate the reinforcement of polymer matrices, improving the practical exploitation of CMCNFs.

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EXPERIMENTAL SECTION

Materials. Cellulose (cotton linter pulp, CLP) with a viscosityaverage molecular weight of 9.2 × 104 g/mol was supplied by Sanyou Chemical Fiber Co., Ltd. (Tangshan, China) and was dried under vacuum at 60 °C for 24 h prior to use. The water-borne polyurethane dispersion (40% solid content) was supplied by Zhongshan Daoqum Chemical Co., Ltd. (Zhongshan, China). Lithium hydroxide monohydrate (LiOH·H 2 O), urea, sodium chloroacetate (ClCH2COONa), ethanol, tert-butanol, NaCl, and poly(ethylene glycol) (PEG2000) were purchased from Shanghai Chemical Reagent Co., Ltd., China. All chemicals were used as received without any further purification. Preparation of CMCNFs. The CLP was dissolved in an aqueous 4.6:15 LiOH/urea (w/w) solution that was precooled to −12 °C to form a transparent 1 wt % cellulose solution according to our previous work.41,42 Subsequently, sodium chloroacetate was added to the cellulose solution under continuous stirring. The molar ratio of sodium chloroacetate to the anhydroglucose units (AGUs) of cellulose ranged from 1:1 to 10:1, and the reaction time ranged 11851


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ACS Sustainable Chemistry & Engineering Transmission electron microscopy (TEM) was performed on a JEM-2000 instrument (JEOL, Japan) at an acceleration voltage of 200 kV. The CMCNF suspension was diluted to a concentration of 0.01 wt %; then approximately 1 μL of solution was dropped onto copper mesh and dyed with sodium phosphotungstate for 30 s, and then the excess dyeing agent was removed. Scanning electron microscopy (SEM) observations were made on a field emission scanning electron microscope (SEM, Zeiss Sigma, Germany) at a voltage of 5 kV. The sample was coated with gold before observation. Nitrogen physisorption measurements were performed on an accelerated surface area and porosimetry analyzer (Micromeritics AsAp 2020, USA). The sample was degassed for 6 h at 13.3 Pa and 378 K. The BET analysis was conducted with relative vapor pressures of 0.05−0.3. The pore volume was obtained from the adsorption amount at P/P0 = 0.99, and the pore size distribution was obtained from the Barrett−Joyner−Halenda (BJH) desorption branch of the isotherm. Differential scanning calorimetry (DSC) analysis was performed on a Q20 thermal analyzer (TA Instruments, USA). The heating rate was 10 °C min−1, and the temperature range was from −80 to 80 °C. Dynamic mechanical analysis (DMA) temperature sweeps under oscillatory stress were measured on a DMA Q800 system (TA Instruments) in tensile mode at a heating rate of 5 °C/min and with a distance between the jaws of 10 mm over a temperature range of −100 to 150 °C with a frequency of 1 Hz. The light transmittance of the samples was tested with an ultraviolet−visible spectrometer (UV-6, Mapuda, China). The static water contact angles were measured using an OCA 20 contact angle meter (Contact Angle System, Germany). Deionized water (2 μL) was dropped onto the surface of the sample for 1 min, and then images of the water droplets were acquired. The angle was measured 3 times. The degree of substitution (DS) of the CMCNFs was determined by acidimetric titration (average of three measurements).44,45 The ζpotential was measured on a ZETA potentiometer (Nano-ZS ZEN3600, Malvern Instruments, U.K.). The NFC suspensions were diluted to 1 mg mL−1 before measurement and oscillated with a vortex mixer. The mechanical properties of the neat WPU and WPU/CMCNF nanocomposites were measured by a universal tensile machine (CMT6503, MTS/SANS, China) with a tensile rate of 2 mm min−1. The length and width of the samples were 40 mm and 8 mm, respectively. The tensile strength (σb), Young’s modulus (E′), and elongation at break (εb) of each sample were taken as the average of five tests. Thermogravimetric analysis (TGA) was performed on a Q500 thermogravimetric analyzer (TA Instruments) under nitrogen at a heating rate of 10 °C min−1 from 30 to 600 °C. The sample was dried at 105 °C for 3 h before analysis to remove moisture.

Table 1. Reaction Conditions and Physical Properties of the CMCNFsa sample

M

t, h

DS

ζ-potential, mV

S1 S2 S3 S4 S5 S6 S7 S8 S9

1:1 3:1 5:1 7:1 10:1 5:1 5:1 5:1 5:1

5 5 5 5 5 1 3 7 10

0.03 0.04 0.10 0.13 0.15 0.08 0.09 0.12 0.15

−3.5 −5.2 −15.7 −17.8 −18.1 −0.2 −7.3 −17.4 −18.2

a M is the molar ratio of ClCH2COONa to AGU, t is the reaction time, and DS is the degree of substitution.

through hydrogen bonding interactions, hydrophobic interactions, and electrostatic repulsion after dialysis. The resultant CMCNF suspension remained homogeneous upon standing for at least several months, and the concentrated CMCNF suspensions were viscous and jelly-like due to the hydrophilicity of CMCNFs (Figure 1c). The TEM image of the CMCNFs from a dilute suspension showed nanofibrils approximately 9 ± 1 nm in diameter and 385 ± 5 nm in length, giving an aspect ratio of approximately 43, which is obviously higher than that of cellulose nanoparticles regenerated from an ionic liquid (aspect ratio of approximately 10) (Figure 1d).46 Thus, the carboxymethylation of cellulose chains produced CMCNFs with a narrow width distribution that form a well-dispersed suspension due to electrostatic repulsion between the CMCNFs. Upon freeze-drying from the CMCNF suspension, a white and lightweight CMCNF aerogel was obtained (Figure 1e). The SEM image of the inner part of the CMCNF aerogel shows a hierarchical porous structure composed of interconnected cellulose nanofibrils and some thin films spanning the nanofibrils. The BET surface area of the CMCNF aerogel was 172 m2 g−1, translating to a cellulose nanofibril width of 14 nm (see Supporting Information Figure S1), which is consistent with the SEM image. Moreover, a transparent CMCNF film can be obtained by evaporating the water after filtering the CMCNF suspension under vacuum (Figure 1f). We optimized the reaction conditions for the carboxymethylation and self-assembly for the fabrication of CMCNFs. As expected, the molar ratio of ClCH2COONa to AGU and the reaction time were the main parameters influencing the homogeneous carboxymethylation and self-assembly of the carboxymethylated cellulose chains. We first set the reaction time to 5 h and evaluated the effects of the molar ratio of ClCH2COONa to the AGU of cellulose on the formation of the CMCNFs. As the molar ratio of ClCH2COONa to AGU was increased from 1:1 to 10:1, the degree of substitution gradually increased from 0.03 to 0.15 and the ζ-potential gradually decreased from −3.5 to −18.1 mV (Table 1, samples S1−S5; Figure 2a). Moreover, as the reaction time was increased from 1 to 10 h, the DS gradually increased from 0.08 to 0.15, and the ζ-potential gradually decreased from −0.2 to −18.2 mV (Table 1, samples S3 and S6−S9; Figure 2b). Notably, the presence of NaCl significantly suppressed microand nanogel formation and aggregation. At higher DS values, the CMCNF suspensions became transparent and had lower viscosities. According to the experimental results, the CMCNFs obtained with a ClCH2COONa to AGU molar

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RESULTS AND DISCUSSION Figure 1a shows a schematic of the preparation of CMCNFs by carboxymethylation and self-assembly of the cellulose chains from an aqueous LiOH/urea solution using sequential topdown and bottom-up strategies. Cellulose could be rapidly dissolved in an aqueous LiOH/urea solution to obtain a transparent 1 wt % cellulose solution (Figure 1b). This topdown approach of the dissolution of cellulose in an aqueous LiOH/urea solution destroys the intra- and intermolecular hydrogen bonding interactions between the cellulose microfibrils, resulting in a semistiff chain conformation.42 Due to the high chemical reactivity of the hydroxyl groups of cellulose chains, cellulose chains can be homogeneously carboxymethylated with sodium chloroacetate as the etherification reagent, and LiOH acted as the catalyst in the cellulose solution. Then, the bottom-up approach allows the carboxymethylated cellulose chains to readily self-assemble to produce CMCNFs 11852


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Figure 2. Degree of carboxymethylation and ζ-potential of the CMCNFs as a function of the ClCH2COONa/AGU molar ratio at 5 h (a) and the reaction time with a ClCH2COONa/AGU molar ratio of 5:1 (b). The insets show the corresponding CMCNF dispersions.

ratio of 5:1 formed more uniform and stable suspensions with a relatively moderate DS of 0.10 and ζ-potential of −15.7 mV; thus, this kind of CMCNF was used for further investigations. Compared to the conventional heterogeneous carboxymethylation of CMCNFs, which always involves multistep processes and organic solvents using high-pressure homogenizer, the homogeneous carboxymethylation of cellulose exhibited higher DS value at the same molar ratio of ClCH2COONa to AGU under mild stirring.47−49 To elucidate the changes in the chemical structure upon selfassembly of the cellulose chains from an aqueous LiOH/urea solution with the chosen carboxymethylation conditions, we analyzed the FTIR spectra of the CLP and CMCNFs (Figure 3a). The broad peak at 3409 cm−1 corresponds to the stretching vibration of the hydroxyl groups on the cellulose backbone. The absorption peaks at 1163, 1032, and 898 cm−1 were assigned to the asymmetric stretching vibration of the C− O−C glycosidic bonds, the stretching vibrations of C−O in primary C−O−H, and the C−O−C in-plane bending vibrations, respectively.50 The peaks at 1032 cm−1 decreased and broadened, indicating that the carboxymethylation occurred mainly at the hydroxyl groups of C-6 positions of cellulose backbone. Additionally, the most distinct difference between the spectra of CLP and the CMCNFs was observed at

Figure 3. FTIR spectra (a), XRD profiles (b), solid-state CP/MAS 13 C NMR spectra (c), and TGA curves under a nitrogen atmosphere (d) of CLP and the CMCNFs. The insert in panel d shows the DTG curves of CLP and the CMCNFs.

Figure 4. Photographs (a), light transmittance (b), and water contact angle (c) of the neat WPU and WPU/CMCNF nanocomposites with different CMCNF contents. Insets in panel c are photographs of a water droplet deposited on the surface of WPU (upper left) and WPU/CMCNF5 nanocomposites (bottom right). 11853


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Figure 5. FTIR spectra (a) and XRD patterns (b) of the neat WPU and WPU/CMCNF nanocomposites.

carbons, 74.4 ppm; and C6 carbon, 62.1 ppm).41 The chemical shifts of the C4 carbons of the CMCNFs shifted upfield and were less intense than those for CLP, suggesting a decrease in the degree of crystallinity upon the self-assembly of the cellulose chains from an aqueous LiOH/urea solution using sequential top-down and bottom-up strategies. Furthermore, a weak signal appeared at approximately 179.4 ppm corresponding to a low content of carboxymethyl groups on the CMCNFs, which is consistent with the results of the FTIR spectra and the degree of substitution by acidimetric titration.51,55 XRD measurements were carried out to evaluate the crystal structure and degree of crystallinity of the CLP and CMCNFs (Figure 3c). The XRD pattern of CLP showed diffraction peaks of 14.9°, 16.5°, and 22.8°, corresponding to the (11̅0), (110), and (200) reflections of cellulose I. However, the pattern of the CMCNFs showed several characteristic diffraction peaks with the major signals being at 11.9°, 19.9°, and 21.5°, which are very similar to those of the (11̅0), (110), and (200) reflections of cellulose II. These peaks indicate that the crystal form of the CMCNFs is cellulose II, which is obviously different from those of carboxymethylated cellulose nanofibrils with the cellulose I crystal structure.27,56,57 In addition, the crystallinities of the CLP and CMCNFs were 73% and 54%, respectively, suggesting that the self-assembly of cellulose nanofibrils from a cellulose solution reduced the order of the cellulose chains in the hierarchical structure of natural cellulose. Additionally, the maximum thermal decomposition temperature of the CMCNFs (352 °C) is higher than that of CLP (337 °C) (Figure 3d), demonstrating the higher thermal stability of the CMCNFs due to the stronger hydrogen bonding interactions in cellulose II.39 Therefore, these features strongly indicate that CMCNFs with the cellulose II crystal structure were successfully fabricated by homogeneous

Figure 6. DSC thermograms of the WPU and WPU/CMCNF nanocomposites.

1598 and 1420 cm−1, and this difference is attributed to the asymmetric stretching vibrations of the deprotonated carboxyl groups on the CMCNFs.28,51,52 Compared with CLP, the -OH stretching vibration absorption peak shifted from 3409 to 3429 cm−1, and the peak intensity decreased significantly, suggesting that the intramolecular and intermolecular hydrogen bonds were weaker in the CMCNFs. The solid-state CP/MAS 13C NMR spectra were then acquired to evaluate the structural differences between CLP and CMCNFs (Figure 3b). The spectra of the CLP show resonances characteristic of cellulose I, i.e., the C1 carbon in the anhydroglucose units (AGU) (105.5 ppm), the C4 carbon in the AGU (between 83.9 and 89.2 ppm); the C2, C3, and C5 carbons in the AGU (between 72.6 and 75.3 ppm); and the C6 carbon of the primary hydroxyl groups (between 63.2 and 65.8 ppm).53,54 However, the 13C NMR spectrum of the CMCNFs exhibits resonances characteristic of cellulose II (C1, 105.5 ppm; C4, between 83.0 and 87.1 ppm; C3, C5, and C2

Table 2. Physical Properties of WPU and the WPU/CMCNF Nanocompositesa CMCNF sample

wt %

% (v/v)

Tg, °C

σb, MPa

εb, %

E, MPa

Wf, MJ/m3

WPU WPU/CMCNF-1 WPU/CMCNF-2 WPU/CMCNF-3 WPU/CMCNF-4 WPU/CMCNF-5

0 1 5 10 15 20

0 0.7 3.3 6.7 10.3 13.9

−41.3 −42.0 −42.7 −42.9 −43.7 −42.7

4.8 8.5 12.0 13.1 15.6 16.8

700 722 648 526 454 429

4 10 17 24 109 265

14.3 23.9 34.7 32.6 34.2 35.8

The Tg values were determined by DSC. σb, εb, and E are the tensile strength, elongation at break, and Young’s modulus of WPU and the WPU/ CMCNF nanocomposites, respectively. Wf is the work of fracture of the nanocomposites.

a

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Figure 7. AFM topographical images (top) and phase-contrast images (bottom) of the surface of neat WPU (a, d), the WPU/CMCNF-3 nanocomposite (b, e), and WPU at the surface after being partially dissolved to expose the inner structure of the WPU/CMCNF-3 nanocomposite (c, f).

Figure 8. DMA temperature sweeps of the WPU and the WPU/CMCNF nanocomposites. Tensile storage modulus, E′ (a), and loss tangent, tan δ (b), of the nanocomposites as a function of the CMCNF content and temperature.

Figure 9. Storage moduli, E′ values, of the WPU/CMCNF nanocomposites at −75 (●) and 25 °C (■). The lines show the values predicted by the percolation model.

Figure 10. Stress−strain curves of the WPU/CMCNF nanocomposites with different CMCNF contents.

carboxymethylation and self-assembly of cellulose chains from an aqueous LiOH/urea solution using sequential top-down and bottom-up strategies. To further utilize the CMCNFs, we chose WPU as the polymer matrix to evaluate the potential applications of CMCNFs for reinforcing polymer nanocomposites. A CMCNF dispersion was mixed with a WPU emulsion and then cast into Teflon molds and vacuum-dried to achieve

WPU/CMCNF nanocomposites with different CMCNF contents. The incorporation of CMCNFs into the WPU matrix results in WPU/CMCNF nanocomposites with different light transmittance properties, as shown in Figure 4a. UV− vis spectroscopy was used to evaluate the light transmittance of the neat WPU film and WPU/CMCNF nanocomposites with different CMCNF contents (Figure 4b). The neat WPU film is 11855


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Figure 11. Thermogravimetric analysis (a) and derivative thermogravimetry (b) curves of neat WPU and the CMCNF and WPU/CMCNF nanocomposites under a nitrogen atmosphere.

chains. For comparison, Tg for the soft segments of the WPU in the WPU/CMCNF nanocomposites decreased slightly from −42.0 to −43.7 °C as the CMCNF content increased from 1 to 15 wt % (Table 2), suggesting that the presence of the CMCNFs strongly affected the microphase separation of the soft segments and hard segments of the WPU matrix. Further increasing the CMCNF content to 20 wt % increased the Tg of the WPU/CMCNF-5 nanocomposite to −42.7 °C, most likely due to the strong hydrogen bonding interactions between the CMCNF and the WPU chains, limiting the mobility of the WPU chains.59,60 Therefore, the results from FTIR, XRD, and DSC revealed that the formation of urethane linkages and hydrogen bonding interactions between the CMCNF and WPU chains affected the hydrogen bonding and microphase separation of the soft segments and hard segments of the WPU matrix in the WPU/CMCNF nanocomposites. The morphology and microstructure of neat WPU and the WPU/CMCNF nanocomposites were examined by AFM in tapping mode. The topographical and phase-contrast images of the surface of neat WPU revealed that the neat WPU film is relatively rough, most likely due to the microphase separation of the soft segments and hard segments of WPU (Figure 7a,d). In contrast, the WPU/CMCNF-3 nanocomposites revealed a uniform morphology with CMCNFs emerging from the WPU matrix (Figure 7b,e). Moreover, the images of the exposed surface of the WPU/CMCNF-3 nanocomposite show a uniform morphology and a homogeneous distribution of the nanofibrillar network structure, indicating that the CMCNFs did not undergo sedimentation or flocculation during the evaporation process. The CMCNFs were well-dispersed without any substantial agglomeration in the WPU matrix, resulting in a percolating CMCNF network structure within the WPU/CMCNF nanocomposites (Figure 7c,f). The dynamic thermomechanical properties of neat WPU and the WPU/CMCNF nanocomposites were evaluated by DMA (Figure 8). Neat WPU showed a glassy state in the range from −100 to −50 °C with a slight decrease in the tensile storage modulus (E′) due to restricted molecular motion (Figure 8a). With increasing temperature from −40 to 60 °C, E′ significantly decreased due to the glass−rubber transition of the soft segments of the WPU chains. Furthermore, E′ sharply declines above 65 °C, which is related to the dissociation of the hard segments of the WPU chains. Compared to neat WPU, the incorporation of CMCNFs has a great influence on E′ and the dynamic thermomechanical behavior of the WPU/ CMCNF nanocomposites. All of the WPU/CMCNF nanocomposites exhibit a modest increase in E′ below Tg of WPU;

transparent with a high light transmittance of 91% at 600 nm. As the content of CMCNFs increased from 1 to 20 wt %, the WPU/CMCNF nanocomposites became translucent with the light transmittance decreasing from 71 to 31%, most likely due to the difference in the refractive indices of the WPU and CMCNFs and the microstructure of the WPU/CMCNF nanocomposites.58 Additionally, the water contact angles of the WPU/CMCNF nanocomposites slightly decreased from 64° to 60° with increasing CMCNF content (Figure 4c), suggesting that the presence of CMCNFs within the WPU matrix results in more hydrophilic groups on WPU/CMCNF nanocomposites. FTIR spectra of the neat WPU and WPU/CMCNF nanocomposites are shown in Figure 5a. The FTIR spectrum of the WPU exhibited characteristic peaks at 3420 and 1743 cm−1 that correspond to N−H stretching and CO stretching vibrations, respectively.59,60 Compared to the neat WPU, the intensities of the N−H stretching and carbonyl absorption bands decreased with increasing CMCNF content, indicating the formation of urethane linkages (-NHCOO) between the hydroxyl groups of the CMCNFs and the isocyanate groups of the WPU chains and hindered hydrogen bonding interactions between the hard segments of the WPU in the WPU/CMCNF nanocomposites.61 In addition, the N−H stretching vibration of WPU shifted from 3420 to 3430 cm−1, suggesting the presence of hydrogen bonding interactions between the CMCNFs and the WPU matrix in the WPU/CMCNF nanocomposites. The XRD patterns of neat WPU and the WPU/CMCNF nanocomposites are presented in Figure 5b. Neat WPU exhibited a semicrystalline structure with a broad peak at 20.5°, resulting from the microphase separation of the soft segments and hard segments. Although the CMCNFs exhibited characteristic features of cellulose II crystals, as shown in Figure 3c, no obvious peaks of cellulose II were visible in the pattern of the WPU/CMCNF nanocomposites. Moreover, the peak shifted slightly (by approximately 0.3°) to a higher diffraction angle with increasing CMCNF content, most likely due to the covalent cross-linking at the interface of the CMCNFs and hydrogen bonding interactions between the CMCNFs and the hard segments of the WPU chains, improving the microphase separation compared to the neat WPU matrix. The DSC thermograms of the WPU and WPU/CMCNF nanocomposites are shown in Figure 6. Neat WPU exhibits a Tg characteristic of WPU at −41.3 °C, which corresponds to the glass−rubber transition of the soft segments of the WPU 11856


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much lower than the calculated critical volume fraction of CMCNFs, also demonstrated a significant reinforcement effect without the formation of a percolation network. Therefore, these features, together with the results from FTIR, XRD, AFM, and DMA, suggest that the formation of covalent crosslinking and strong hydrogen bonding interactions between the CMCNFs and WPU chains, as well as the rigidity of the percolated CMCNF network, allow the successful transfer of mechanical stresses, resulting in significant reinforcement of the mechanical properties of the WPU/CMCNF nanocomposites. The influence of the incorporation of CMCNFs on the mechanical properties of the WPU/CMCNF nanocomposites was further evaluated under a tension model at ambient temperature (Figure 10). The tensile strength, Young’s modulus, and elongation at break of neat WPU are 4.8 MPa, 4.2 MPa, and 710%, respectively, showing a typical rubbery state. As expected, when the content of CMCNFs was increased from 1 to 20 wt %, the tensile strength and Young’s modulus of the WPU/CMCNF nanocomposites gradually increased from 8.5 to 16.8 MPa and from 10 to 265 MPa, respectively; the highest tensile strength and Young’s moduli of the WPU/CMCNF nanocomposites were 3.5 times and 66 times greater than those of neat WPU. In addition, the elongation at break of the WPU/CMCNF nanocomposites decreased from 722% to 429% with increasing CMCNF content because of the higher cross-linking density between the CMCNFs and the WPU matrix. Moreover, the work of fracture of the WPU/CMCNF nanocomposites reached 35.8 MJ m−3, which is 2.5 times higher than that of WPU (14.3 MJ m−3), demonstrating a significant reinforcement in the toughness of the WPU/CMCNF nanocomposites. Therefore, the incorporation of CMCNFs into the WPU matrix can considerably enhance the mechanical properties of the WPU/ CMCNF nanocomposites. The thermal decomposition behaviors of neat WPU and the CMCNF and WPU/CMCNF nanocomposites were examined by thermogravimetry under a nitrogen atmosphere (Figure 11). The maximum thermal decomposition of the CMCNFs occurs at 352 °C, which is much higher than that of cellulose nanowhiskers (less than 150 °C) and similar to that of regenerated cellulose, demonstrating the relatively higher thermal stability of the CMCNFs. Neat WPU showed a three-step weight-loss process with maximum decomposition temperatures of 295, 355, and 460 °C. The maximum decomposition temperatures of the WPU/CMCNF nanocomposites are lower than those of CMCNFs and WPU, indicating that the formation of urethane linkages between CMCNFs and the hard segments of WPU accelerated the decomposition of the WPU/CMCNF nanocomposites.

however, these nanocomposites display significant mechanical reinforcement over that of neat WPU when the temperature is above Tg of the soft and hard segments of WPU. For instance, at 60 °C, the E′ value of the WPU/CMCNF-5 nanocomposite with 20 wt % CMCNFs is 1.39 GPa, which is more than 4600 times greater than that of neat WPU (0.3 MPa). Moreover, the DMA temperature sweeps also revealed that the incorporation of the CMCNFs significantly extends the rubbery plateau of the soft segments of WPU to 100 °C for the WPU/CMCNF nanocomposites with CMCNF contents higher than 5 wt %, demonstrating a significant mechanical reinforcement of WPU, most likely due to the strong hydrogen bonding interactions between CMCNFs and the hard segments of WPU, limiting the mobility of the soft segments of the WPU chains. In addition, the maximum loss tangent (tan δ) of the WPU/ CMCNF nanocomposites exhibited a lower relaxation temperature with increasing CMCNF content (approximately −35.5 °C for WPU/CMCNF-5 nanocomposite compared to approximately −14.2 °C for the neat WPU). Moreover, the intensity of tan δ of the WPU/CMCNF nanocomposites decreased by a factor greater than the change in the CMCNF content, most likely because of the confinement effects from covalent cross-linking and strong hydrogen bonding interaction between the CMCNFs and the WPU matrix.11,62−66 We further evaluated the CMCNF content dependence of the storage modulus of the WPU/CMCNF nanocomposites at −75 °C (below Tg of the soft segments of WPU) and 25 °C (above Tg of the soft segments of WPU) by the percolation model (Figure 9). The volume fraction of CMCNFs in the WPU/CMCNF nanocomposites was determined gravimetrically using the densities of CMCNFs and WPU (1.62 and 1.05 g/cm3), respectively. For the percolation model, E′ of the nanocomposites can be expressed as follows.11,62,67−69 E′ =

(1 − 2ψ + ψX r )Es′Er′ + (1 − X r )ψ (Er′)2 (1 − X r )Er′ + (X r − )ψEs′

(1)

with

ij X − Xc yz zz ψ − X r jjj r j 1 − Xc zz k { Xc =

0.7 f

0.4

(2)

(3)

where Es′ = E′ of the neat polymer, Er′ = E′ of the percolating cellulosic network, Xr = volume fraction of the nanofibrils, Xc = volume fraction of the nanofibrils at the percolation threshold, Ψ = volume fraction of the nanofibrils contributing to percolation (“effective skeleton”), and f = aspect ratio of the nanofibrils. As a consequence, above the critical volume fraction of CMCNFs (approximately 1.6% (v/v) with an aspect ratio of 43) at the percolation threshold, the E′ values for the WPU/ CMCNF nanocomposites agree well with the percolation model when using experimental values for Xr and E′s (1.8 GPa at −75 °C (below Tg of the soft segments)) and 21 MPa at 50 °C (above Tg of the soft segments) for neat WPU) and fitting values from the experimental data for Er′ (150 GPa at −75 °C and 28 GPa at 25 °C for CMCNFs), further confirming that CMCNFs indeed form a percolating network in the WPU/ CMCNF nanocomposites. Notably, the WPU/CMCNF-1 nanocomposite with only 0.65% (v/v) CMCNF, which is

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CONCLUSION In summary, CMCNFs with the cellulose II crystal structure were successfully fabricated by homogeneous carboxymethylation and self-assembly of cellulose chains from an aqueous LiOH/urea solution using sequential “top-down and bottomup strategies. The CMCNFs could be used as reinforcing nanofillers to enhance water-borne polyurethane (WPU). CMCNFs were evenly distributed in the WPU matrix, leading to a significant enhancement of the tensile storage modulus above the glass transition temperature of the soft segments of WPU. The reinforcement effects can be explained by the percolation model. Additionally, the WPU/CMCNF nano11857


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composites also demonstrated high mechanical properties and moderate thermal stability. The strategy developed here produce CMCNFs and could be utilized in the development of commercial packaging materials, flexible devices, and biomaterials and combined with other polymers to obtain superior functional CMCNF-based composite nanomaterials.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acssuschemeng.9b02424. Nitrogen adsorption−desorption of CMCNF aerogels (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +86-27-6878-9321. ORCID

Lina Zhang: 0000-0003-3890-8690 Jie Cai: 0000-0002-0660-4740 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research was funded by the National Natural Science Foundation of China (Grants 21875170, 51373125, and 21422405). We thank the Wuhan Morning Light Plan of Youth Science and Technology (Grant 2017050304010312), the Special Fund for the Development of Strategic Emerging Industries of Shenzhen City of China (Grants JCYJ20180507184711069 and JCYJ20170818112409808), and Fundamental Research Funds for the Central Universities (Grants 2042017kf0195 and 2042015kf0259) for the facility support.

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New Approach for the Fabrication of Carboxymethyl Cellulose

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