Strain Rate-Dependent Viscoelasticity and Fracture Mechanics of

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Strain Rate Dependent Viscoelasticity and Fracture Mechanics of Cellulose Nanofibrils Composite Hydrogels Jun Yang, Changyou Shao, and Lei Meng Langmuir, Just Accepted Manuscript • DOI: 10.1021/acs.langmuir.9b01532 • Publication Date (Web): 12 Jul 2019 Downloaded from pubs.acs.org on July 18, 2019

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Strain Rate Dependent Viscoelasticity and Fracture Mechanics of Cellulose Nanofibrils Composite Hydrogels Jun Yang*, Changyou Shao, Lei Meng Beijing Key Laboratory of Lignocellulosic Chemistry, Beijing Forestry University, No. 35, Tsinghua East Road, Haidian District, Beijing, 100083, China *Corresponding

author: [email protected]

Tel: 86-10-62337223

Abstract In this work, the composite hydrogels toughening behaviors as manifested by strain rate dependent viscoelastic properties and enhanced fracture mechanics, i.e. suppressed catastrophic crack propagation with increased resistance, are systematically examined by using cellulose nanofibrils (CNFs) as fillers in polyacrylamide (PAAm) matrix. The uniaxial deformation tests show that the tearing energy increases with crack velocity and becomes dominated by the viscoelastic energy dissipation in front of the crack tip. The creep dynamics of the composite hydrogels under a constant stress is examined and the results indicate that the incorporation of CNF pronouncedly suppresses the creep deformation. In addition, the microdeformation and failure mechanisms are analyzed through the observation of morphology of arrested crack tips and the damage zone by TEM and SEM. By aligning the CNF along the crack direction, it is possible to focus on the study of interfacial slip mechanics and identify the role of interfacial slip during energy dissipation process. The results indicate that the CNFs are largely orientated parallel to the loading direction to maximize the energy dissipation, where the initiation of crack propagation is the primary fracture mechanism in composite hydrogels. The coarse feature on the composite fracture surface implies that the CNF initiates deflection of crack propagation fronts and thus increases the strain energy for continuation of fracture. It is envisioned that with incorporation of interdisciplinary strategies, one can rationally combine multiple approaches toward the creation of nanocomposite hydrogels with enhanced mechanical properties.

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Introduction Hydrogels are constructed through cross-linking of hydrophilic polymer chains within an aqueous environment, in which the cross-links can be achieved through various mechanisms, spanning from covalent cross-linking to physical entanglement of polymer chains.1 The water-rich property of hydrogels makes them attractive in many areas, including drug delivery, biomedical engineering, and soft actuators.2-4 However, due to the lack of appropriate dynamic features and structural complexity, most of the synthetic hydrogels are brittle and always exhibit poor mechanical performance.5 Thus, designing hydrogels with excellent mechanical properties is a critical task for soft material scientists. Recent pioneering reports of strong and tough polymeric hydrogels (such as nanocomposite hydrogels,6 double-network hydrogels,7 topological gels,8 and supramolecular gels9) have highlighted that their potential as structural materials can be realized by rational design at the molecular level and well-defined control over multiscale architectures. For example, hydrogels combining covalent polymer network with reversible cross-links for energy dissipation generally exhibit excellent toughness and stretchability. Thus, one aiming at obtaining a gel with high toughness needs to introduce efficient dissipative mechanisms and increase damping resistance.10 Due to the favorable interactions between amine groups and hydroxyl groups,11 adsorption of hydrophilic polymer chains, such as polyacrylamide (PAAm), to a lignocellulose surface has been demonstrated to be an effective way to create physically cross-linked cellulose networks. Cellulose nanofibrils (CNFs), extracted from plant cell wall with high elastic modulus and tensile strength in GPa range, have been widely applied in hydrogel materials, such as tissue engineering scaffolds and drug delivery systems,12,13 where the hydrogen bonding mediated fibril-fibril entanglement is favored for creating mechanically reinforced hydrogels. The large specific surface area of CNFs, rooted in their submicrometer length offers a great opportunity to dissipate energy in composites via interfacial sliding.14 In previous work, we have prepared a series of CNF reinforced PAAm composite hydrogels with viscoelastic network containing hydrogen bonding and covalent cross-links.15-18 The multiple hydrogen bonds, including both intramolecular and intermolecular hydrogen bonds on CNF surface, as well as the dense topological entanglements and lightly chemical cross-linking of the polymer chains synergistically stabilized the hydrogel networks. The weak bonds (hydrogen 2


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bonding) preferentially broke upon stress and acted as reversible sacrificial cross-links to dissipate energy, whereas the covalent cross-links served as permanent cross-links to maintain the integrity of the network over a long time scale.16 With progress of deformation, the coiled chains penetrated into hierarchical fibrils with relatively high dissociation energy began escape out of the fibrils and released the hidden length, playing another kind of sacrificial bond to substantially enhance the overall energy dissipation and facilitate the stress alleviation.18 For glassy polymers, crazing is a dominant deformation process that results in the occurrence of crack and final failure of the materials.19, 20 According to Balazs et al,21 the increased conformational entropy of chains in an enthalpically neutral system led to deformation through the emergence and propagation of a craze. Indeed, the propagation and breakdown of crazes, precursor and ahead of a crack tip, define the bulk fracture strength and toughness of nanocomposites.22 In this regard, understanding the crazing process in the presence of CNF is critical to interpret the mechanical performance. In this work, we continue to explore experimental results of alignment of CNFs during the onset and propagation of the craze and demystify the fundamental understanding of fracture mechanics and interfacial failure in cellulosic hydrogels. The results indicate that hydrogels with attractive interactions between CNF and matrix can lead to the slowdown of craze initiation and growth due to the balance of two competing mechanisms. Consider the breaking and re-forming of physical bonds are time-dependent,23,24 the mechanical properties of the gels are expected to demonstrate significant dependence on the relative relation between observation time and cross-link dynamics, i.e., deformation velocity dependent viscoelasticity. To validate this speculation, the cyclic tensile test for large strains and tearing test for fracture properties are conducted over a large range of deformation rate. Besides, consider the potential application of tough hydrogels as load-bearing materials,25,26 it is essential to examine the long-term mechanical stability of the materials, thus the creep behavior for the composite hydrogels under a constant load is also assessed in this work. It is expected that these efforts could significantly deepen the interpreting of failure mechanism of cellulosic nanocomposites and CNF spatial distribution under stress field.

Experimental Section Materials and Hydrogels Preparation The cellulose nanofibrils suspension with a solid content of 1.5 wt% were collected from bleached 3


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softwood Kraft pulp (DongHua Pulp Factory, China) through mechanical homogenization using a high pressure fluidizer, and the detail CNF extraction procedure and morphology were demonstrated in supporting information (Figure S1). All other chemicals that purchased from Beijing Chemical Reagent Company were analytical grade and used as received without further purification. The composite hydrogels were prepared by free radical polymerization of acrylamide (AAm) monomer in the presence of CNF according to our previous work.17 In brief, 20 mL of well dispersed aqueous suspension of AAm (2 g, 10 wt%), potassium persulfate (19 mg, 0.25 mol% against AAm), N, N’-methylenebisacrylamide (17 mg, 0.4 mol% against AAm), and CNF (volume fraction of 0.031-0.75%, corresponding to 0.05-1.2 wt% against aqueous solution) was degassed by centrifugation and then injected into a PTFE mold under N2 at 30℃ for 24 h. After polymerization, the as-prepared gels were immersed in a large amount of water for 48 h by replacing water every 12 h to remove soluble species. For the control of neat PAAm gels, they were synthesized by the same procedure in the absence of CNFs.

Characterization The mechanical properties were examined under pure shear profile using a Zwick Roell Testing Machine (Z005) with a load cell of 100 N. The rectangular specimens (length = 60 mm, width = 10 mm, and height = 4 mm) were fixed between two clamps and the uniaxial deformation was performed at a crosshead speed of 1.44-1140 mm/min (corresponding to stretch rate, έ, 0.0012-0.92 s-1). The stress (σ) was determined by the nominal value, which corresponded to the force per cross-sectional area of the specimen before deformation, and the strain (ε) was calculated by the deformed length relative to the initial length. The Young’s modulus (E) was calculated as the slope of the initial linear region (0-50% strain) of the stress-strain curves. For the cyclic measurement, the specimen was stretched to a preset strain and unloaded to the initial length at the same velocity. The hysteresis was determined from the area enclosed by the loading-unloading curves. Four measurements were conducted for each sample and results were the average of four trials. The tearing test was performed to obtain tearing energy of gels by the well-defined trousers test, where a precrack about 5 mm long with a sharp crack tip was made by pressing a fresh razor blade into the sample. The specimens with two arms were clamped tightly and set perpendicular to the surface of the clamps, which allowed the crack propagated along the central line of the 4


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specimen. The lower clamper was fixed, and the upper clamp was stretched at a constant velocity until the crack propagated through the whole specimen, where the corresponding stretched displacement (L) and tearing force (F) were recorded. The tearing energy (  ) was calculated as

 =2F/w, where F was the constant tearing force at crack propagation and w was the thickness of the specimen. For uniaxial creep experiment, a nominal stress of 100 kPa was applied on the specimen and the creep modulus was calculated by the ratio of the initial stress to the creep strain. The creep test was performed in a closed chamber to prevent the hydrogel from drying out. Scanning electron microscopy (SEM) was conducted for the fracture surface observation of composites gel using a Hitachi S-3500 operating at an acceleration voltage of 6 kV. The fractured samples were freeze-dried for 24 h and sputter-coated with gold before observation. The transmission electron microscope (TEM, JEOL 2010) worked at 200 kV was applied to observe the CNF distribution in PAAm matrix, where the microtomed specimens were casted over TEM grid with holey carbon support film and stained with uranyl acetate solution (1.5 wt%) for 10 s before observation. The craze growth and propagation were examined by the copper grid technique which contained floating and transferring a sample film onto a copper gird substrate and subsequently subject to the uniaxial strain. Due to the low yield strain and plasticity of copper substrate, the grid could maintain the applied stress after the unloading, and the TEM observation was conducted to examine the craze distribution.

Results and Discussion Microscopic Structure and Swelling The structure of the composite hydrogels is firstly characterized by TEM (Figure 1a) and the image shows that the fibril-like CNFs are homogeneously dispersed within the hydrogels without apparent aggregation, inferring good compatibility of CNF and PAAm. Indeed, the PAAm chains undergo only weak amide-mediated hydrogen bonds,27 whereas the incorporating CNF at a small fraction in PAAm provides additional stronger interfacial interactions.15 The photos of attained composites gels are exhibited in Figure 1a and no macroscopic phase separation is observed. According to the time-dependent swelling ratio profiles (Figure 1b), one can note that the swelling ratio of composite hydrogels monotonically increases during the initial 10 h and then gradually attains the swelling equilibrium within the following 40 h. Although the CNFs present 5


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hydrophilicity and accommodate additional water, the addition of CNF restrains the swelling capacity. It is well known that water is a good solvent for PAAm while it is a nonsolvent for cellulose, thus the introduction of CNF into PAAm would increase water diffusion resistance in composite hydrogels. Indeed, due to the high aspect ratio of CNF (50-100), the inclusion of entangled CNF could lead to tortuous solvent path through fibril entanglement and increase the diffusion resistance thereby. Moreover, considering the attractive interfacial interaction between CNF and PAAm, the incorporation of CNF could impart constrained polymer layer near the CNF surface and leads to significant decrease in the free volume within the hydrogels, which in turn causes the restrained water transport permeability. This assumption of constrained polymer chain mobility near the CNF surface is corroborated by a slight increase of glass transition temperature (Tg) upon the incorporation of CNF (Figure S2). In fact, this result is consistent with other nanoparticle-filled polymers with significantly altered dynamics surrounding nanoparticles,28,29 where the surface immobilized chains are in favor of stress transfer and mechanical reinforcement.

Figure 1. (a) Optical images of swelling composite hydrogels and the corresponding TEM image exhibiting percolated CNF fibrillar network. (b) Time dependent swelling behavior for the gels with different CNF volume fractions.

Viscoelastic Properties The interconnected CNFs in PAAm matrix form a supramolecular network that is bridged by hydroxyl-amine hydrogen bonds. Contrast to the neat PAAm gels with sole covalent bonds, these 6


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hydrogen bonds are temporary and can dissociate at room temperature, consequently, enable the viscous properties that evidenced by the viscoelastic stress-strain curves.30-32 Thus, we conduct the strain rate dependent tensile tests to examine the viscoelasticity. The samples are initially stretched at a controlled strain rate of 0. 5 s-1 and the curves exhibit three distinct regimes (Figure 2a). At a small strain (ε＜50%, region  ), the stress rises fast in a nearly linear pattern; at the intermediate strain (50%＜ε＜400%, region  ), the stress increase rate slows down while still maintains the linear manner; with further increase in strain (ε＞400%, region Ⅲ), the rate of increase in stress almost levels off.

Figure 2. Mechanical properties of composite hydrogels with CNF volume fraction of 0.2%. (a) Uniaxial tensile curves with three distinct regimes depending on the extent of strain (έ = 0.5 s-1). (b) Stress-strain curves at different strain rates. (c) Correlation between Young’s modulus and strain rate.

Due to the relaxation of dynamic hydrogen bonds, the increase in deformation rate entails a faster dissociation process of reversible cross-links and requires a higher mechanical strength, leading to the increase in modulus and toughness. In this regard, the mechanical properties of composite hydrogels can be viewed as apparent parameters that are significantly dependent on the deformation conditions, such as temperature and stretch rate. Since we restrict the scope of current effort in discussing the deformation response of gels under isothermal condition and neglect possible water diffusion in the gels (relative low to the time scale of mechanical deformation), we attempt to model the following mechanical behavior at different strain rates (έ) with the rate ranging over near three orders of magnitude from 0.0012 to 0.92 s-1. One can note that for all the examined rates, the stress-strain curves are featured by the similar three-regime pattern that depending on the stretch ratio (Figure 2b). At the initial deformation stage, the Young’s modulus 7


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exhibits significant strain rate dependent behavior (Figure 2c). At small strain rates (έ ＜ 0.12 s-1), the modulus increases rapidly following a power law from 23 to 73 kPa (EⅠ∝έ0.43), whereas at larger strain rates (έ ＞ 0.12 s-1), the EⅠ gradually levels off at 85 kPa. This result can be rationalized by the hydrogen bond finite relax time where the bonds can break as the time scale is larger than the average lifetime.33 A smaller stretch rate provides adequate time for hydrogen bonds reconfiguration and leads to contribution to the network modulus since the fraction of unrelaxed reversible bonds increases with strain rate. However, the very large strain rates result in a short time scale where the hydrogen bonds become unrelaxed and the moduli under high strain rate gradually level off. Indeed, the crossover strain rate of 0.12 s-1 is close to the timescale order magnitude as the reciprocal lifetime of hydrogen bonds.34 To qualify the improvement of toughness driven by the presence of noncovalent junctions, cyclic loading-unloading tests are performed and the energy dissipation per cycle is calculated (Figure 3a). The neat PAAm gel with only chemical cross-link is chosen as a control, which is a widely used component in several recently well-defined tough hydrogels.35 For composite hydrogels, due to the presence of dense and reversible hydrogen bonds and rigid CNF itself, the dual-network of composite hydrogels composed of covalent cross-links and interfacial hydrogen bonds lead to a high Young’s modulus and toughness (78 kPa and 48 kJ/m3, respectively), which are close to the mechanical properties of some double-network (DN) gels.36 The loading-unloading curves of the composite hydrogels show a substantial hysteresis loop, which is related to the macroscopic manifestation of dissociation of hydrogen cross-links and leads to energy dissipation. Specifically, at the initial deformation stage, both chemical cross-links and hydrogen bonds are nearly intact, and their total amount of cross-links determines the Young’s modulus of the composite hydrogels (consider the constant elasticity of CNF during stretch). During the following stretch stage, the polymer chains near the CNF surface with interfacial hydrogen bonds start to dissociate and reassociate with neighboring relaxed chains, while polymer strands near the chemical cross-links remain strained. In this process, the chemical cross-links dominate the elasticity and maintain the integrity of network, while the dissociation and reassociation of hydrogen bonds lead to energy dissipation and determine the toughness by surviving the external stress at a high deformation level.37 During the unloading stage, the stress stored in the elastic chains of strained covalent network promotes the dynamic H-bonds to recover 8


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the initial stress-free state. Whereas for the neat PAAm hydrogels, their loading and unloading curves are almost overlapped. In fact, even the stretch rate is increased by one order of magnitude (from 0.05 to 0.5 s-1), the deviation between the loading and the unloading curves is less than 2.5 kPa (Figure S3), indicating the rate independence of PAAm gels.

Figure 3. Cyclic mechanical properties of hydrogels. (a) Tensile loading-unloading curves at a strain rate of 0.5 s-1. (b) Continuous loading-unloading curves without resting between each cycle (έ = 0.5 s-1). (c) Loading-unloading cycles at different strain rates and the corresponding hysteresis (d). The composite hydrogels with CNF volume fraction of 0.2% were examined.

Besides, due to the relaxation of transient cross-links, the continuous loading-unloading cycles lead to the decrease of hysteresis and the stress (Figure 3b). For example, the hysteresis and maximum stress decrease from 48 and 338 to 32 kJ/m3 and 286 kPa, respectively, after six cycles and leave some residual strains (~ 50%). This result shows good qualitative correlation with previous work of fully physically cross-linked hydrogels,36 which can be ascribed to the escape of hydrogen bonding domains and reattachment of polymer chains on CNF surface that slow down the recovery of elastic recovery of polymer chains. Contrast to conventional hydrogels with 9


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permanent networks bridged by covalent cross-links, the transient hydrogen bonds reveal a finite relax time.34 Thus, the composite hydrogels are anticipated to illustrate strain rate dependent loading-unloading behavior. For this purpose, similar to the measurement in Figure 2b, another series of cyclic tensile test is conducted as the function of stretch rate with a maximum stretch range of 800% (Figure 3c). The results indicate that the hysteresis increases with strain rate following the distinct power law (with exponent of 0.36 and 0.24, respectively) at a crossover at the strain rate of 0.12 s-1 (Figure 3d). For the composite hydrogels with dual cross-links, the specific mechanical properties can be either soft-weak or rigid-tough depending on strain rate. Since the energy dissipation per unit volume is defined by the hysteresis area between the loading and unloading curves, the energy dissipation is related to the reversible dissociation of H-bond and exhibits the strain rate dependence that can be rationalized as follows. In detail, if the strain rate is smaller than the inherent dissociation rate of cross-links, the reversible dynamics of H-bond does not require outside mechanical stimulus, and the gels behave as soft elastic network and lead to a small amount of energy dissipation with a low toughness. In contrast, at strain rate higher than the dissociation rate, the cyclic loading-unloading process results in remarkable energy dissipation with increased apparent toughness.

Figure 4. Time dependent mechanical properties of composite hydrogels (CNF volume fraction of 0.2%). (a) Stress-strain curves with interrupted stress relaxation test. (b) Time evolution of normalized stress under different strain rates.

To further elucidate the strain rate dependent viscoelasticity, a series of sequential stress relaxation tests are conducted where the sample is stretched to a preset strain (100%, 190%, 305%, 10


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410%, 600%, and 800%) under a constant stretch rate of 0.2 s-1, then hold the strain for 2 min and repeat the holding period during next loading stage at the same stretch rate (stretch history scheme in Figure S4). Interestingly, the unique asymmetric strain rate dependent loading curves are noted with respect to Figure 4a and no rate sensitivity for unloading curves. Consider the time dependent relaxed stress (σ(t)) gradually relaxes to an asymptotic value σ∞ and define

 norm (t )  The normalized value of results indicate that

 (t )     (0)   

 norm (t ) as a function of stretch ratio is present in Figure 4b and the

 norm (t ) exhibits striking strain rate dependence of the normalized stress

relaxation. This strain rate dependent stress relaxation behavior is related to the hydrogen bond inherent dynamic of dissociation and reassociation.23,37 As the sample is stretched, the interfacial bonds in the composite hydrogels dissociate that leads to strain rate dependent loading curves. In contrast, for the unloading process where the strain is reduced, no additional dissociation of hydrogen bond occurs and results in the almost strain rate insensitivity.

Tearing Test The pure shear test is performed to investigate fracture process of composite hydrogels, and the typical curves of tearing force against deformation are shown in Figure 5a. Before the stable fracture process occurs, the tearing force attains a maximum value, which is corresponded to the highest resistance against fracture propagation. The composite hydrogels achieve a high tearing force and the maximum tear strength for the specimen with 10 mm width attains 0.71 N/mm (calculated by tearing force divided by the sample thickness), which is one order of magnitude greater than that of the neat PAAm gels (0.052 N/mm), indicating efficient energy dissipation role of CNF in composite hydrogels.17 In fact, according to the fracture cross-section of composite hydrogels by SEM images below (Figure 7), the pull-out fibril bundles bridge the neighboring matrix and lead to the great tear resistance against crack propagation. It has been proposed that the energy dissipation required to advance per volume of the material includes the intrinsic fracture energy of the cross-linked polymer chains ahead of the crack tip as well as the bulk viscoelastic energy dissipated around the tip,22 thus it is reasonable to find that for the composite hydrogels 11


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studied in this work, the tearing energy significantly depends on the tearing velocity and increases with increasing tearing velocity (Figure 5b). At relatively high tearing rates (＞0.15 s-1), however, this increase becomes saturated because most hydrogen bonds are unrelaxed and the modulus of the polymer chains under sufficient high shear attain nearly a constant level. Whereas for the neat PAAm gels, there also shows a slight increase with tearing velocity due to the weakly viscoelastic deformation that stems from chain entanglements or loops.27

Figure 5. Fracture properties of composite hydrogels with CNF volume fraction of 0.2%. (a) Tearing force-displacement curves at different tearing rates. (b) Tearing energy versus tearing rates. (c) Tearing rate dependence of crack propagation velocity (Vp) for the single edge notch test (the inset is the picture of tearing test).

As shown in the inset of Figure 5c, when the notched sample is stretched, crack propagation starts instantaneously from εc and the tearing force decreases gradually with time and tends to zero. Thus, consider the time interval that crack propagates throughout the sample, the velocity of crack propagation (Vp) can be estimated by Vp = (w-c)/t, where w and c is the sample width and initial notch width, respectively. By varying the tearing velocity over three orders of magnitudes from 0.0012 to 0.92 s-1, the result indicates that the crack propagation velocity of composite hydrogels exhibit significant tearing velocity rate dependent behavior with three distinct regimes (Figure 5c). At small tearing velocity (＜ 0.0085 s-1), the Vp is rate independent; at a relative large tearing velocity (0.0085-0.12 s-1), Vp increases rapidly with tearing rate that following a power law Vp~ε0.48 from 0.025 to 1.8 mm/s; with further increase of tearing velocity, the Vp becomes nearly constant at about 2.2 mm/s when the tearing velocity above 0.12 s-1. This result further highlights the importance of dynamic hydrogen bond at different time scales in achieving extremely high crack propagation resistance.24 In fact, this behavior can also be qualitatively rationalized by the 12


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relaxation

of

dynamic

hydrogen

bonds.

At

the

low

tearing

rates,

since

the

dissociation-reassociation rate of reversible bonds is much fast than the loading rate, it is difficult for polymer strands at CNF interface to accumulate large stretch before fracture and the required energy for crack propagation is low. Whereas at high tearing rates, i.e., the tearing experiment time scale is shorter than the average lifetime scale of the reversible bonds, the unrelaxed reversible bonds contribute to network strength, the required energy to propagate the crack tip increases and the critical stretch increases thereby. For an even higher tearing rate, however, this increase becomes saturated due to most hydrogen bonds become unrelaxed and the network rigidity against crack propagation becomes almost a constant level. On the other side, consider hydrogen bonds in composite gels can be suppressed in urea aqueous solution and leads to decreased toughness and strength of the gels, we further verify the role of hydrogen bonds in tuning crack dynamics in tearing test by equilibrating the gels in urea solution with different concentrations (0.02-0.12 M). As shown in Figure S5, the crack propagation resistance and tearing energy of composite hydrogels become strikingly decreased in urea solution at a given tearing velocity, implying a significant fraction of hydrogen bonds are destroyed. Based on the above result, one can conclude that interfacial hydrogen bond is critical for achieving the tough composite hydrogels.

Creep Properties Since the polymer chain mobility in composite hydrogels depends on the interfacial interactions and external loading condition, the elongational creep measurements under step loading stress are conducted at room temperature. From the profiles of creep strain versus time in Figure 6a at a constant step stress, the neat PAAm gels creep process includes three distinct stages: (1) initial slow elongation, (2) primary creep, and (3) secondary creep. The initial fast elongation is related to the elastic and plastic deformation of polymer chains once the constant load is applied. During the following primary creep stage, the creep stage occurs at a relatively low value and the creep strain increases with time, implying the slippage and reorientation of chains under persistent stress. Finally, the specimen falls into the secondary creep stage at a well-defined time (tc), at which the creep rate suddenly accelerates and leads to creep fracture. This tc corresponds to the occurrence of macroscopic crack propagation and the corresponding strain is defined as εc. 13


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Intriguingly, the addition of a small amount of CNF into PAAm gels significantly suppresses the growth of creep strain, where the composite hydrogels exhibit a much lower initial creep rate at the transition from the primary creep and secondary creep in comparison to the neat PAAm. In fact, the composite hydrogels with CNF fraction of 0.5% show negligible creep phenomenon in the observation window. This result can be ascribed to the fact that more conformation exchanges are required to alter chain topology in the network with a higher cross-linking density. The incorporation of CNF not only delays the creep flow but also stabilizes the hydrogen bonds to prevent the creep at a moderate stress level. For example, this dynamics difference leads to the creep strain boundary for the primary creep stage of composite hydrogels (CNF 0.2 % volume fraction) at around strains of 15 and 24%, respectively, which are much smaller than that of neat PAAm (amounting to 22 and 38%, respectively). Thus, the steady-state creep rate of composite hydrogels is about 8.7×10-5 min-1 in the primary stage, which is much lower than that of 20.3×10-5 min-1 for the neat PAAm hydrogels.

Figure 6. Creep deformation behavior of composite hydrogels. (a) Tensile creep strain and (b) creep modulus of the gels as a function of time under a constant creep stress of 100 kPa. The abrupt increase of the creep strain corresponds to the macroscopic crack growth as denoted in the dash line. (c) Creep deformation under step loading stress. (d) Step loading stress dependent 14


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failure time (tc) and failure strain (εc) within the observation of 24 h. The CNF volume fraction was the same for the corresponding curve colors in (a) and (b), and the composite hydrogels with CNF volume fraction of 0.5% was applied to creep test in (c) and (d).

The creep modulus versus time in creep test also indicates that a significant modulus reduction occurs for PAAm hydrogels than that of composite hydrogels (Figure 6b), suggesting the incorporation of CNF leads to a much higher creep modulus in the whole range of creep. Thus, the dimensional stability and load bearing capacity is enhanced under creep load for the CNF reinforced gels. We further examine the effect of step of loading stress on creep dynamics (Figure 6c) and the results indicate that composite hydrogels exhibit a critical stress (50 kPa) below which no failure occurs in the observation time window of 24 h and this critical stress level increases with CNF fraction (Figure S6). For a better comparison, the loading stress dependences of failure time (tc) and failure strain (εc) are plotted in Figure 6d. It is interesting to find that the logarithm of failure time is linearly negatively correlated to the loading stress, implying the creep failure of composite hydrogels is governed by the thermal activated energy barrier of damaging of hydrogen bonds that strengthen the cross-linking network. For the correlation of failure strain with loading stress, one can note that the failure strain decreases with increasing loading stress, and there appears a small plateau of failure strain before steeply decreasing tendency with increase of loading stress from 50 to 200 kPa, implying the CNF is the origin of the delayed fracture in composite hydrogels.

Fractographic Analysis According to previous work, the formation of plenty of microcracks and the increased fracture area due to crack deflection are primary toughening mechanisms in the composite hydrogels.17,18 However, there are still some pending questions need to be answered: where does the micro-crack initiate, at the interface or within matrix, how is the evolution from the micro-crack towards macro-crack, and how crack tips develop around the interface. To clarify these mysteries, the arrested crack tip of composite hydrogels is examined by means of TEM and SEM. Initially, the crack tips reach the site where the fibril orientation is almost perpendicular to that of the crack propagation direction and arrest the crack propagation (Figure 7a). Although some microcracks 15


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nucleate in the front of the crack path, the adjoining fibril clusters do not break and give way to crack propagation and no significant plastic deformation of the matrix is observed near the crack. Thus, the crack preferentially deflects and penetrates outside the cluster, where the extended crack proceeds among neighboring fibrils and appears in the vicinity of large crack (Figure 7b). This phenomenon suggests the relatively poor interface interactions compared with interaction between cohesive fibrils.

Figure 7. Fractographic observation of composite hydrogels with CNF volume fraction of 0.2%. (a, b) Side-view SEM micrographs of the fracture surface, where the bridging zone with pull-out filamentous fibrils at the crack tip enhances the crack propagation resistance. (c, d) TEM images of long and narrow microcavities (marked by arrows) associated with fibrils in a region ahead of the arrested crack tip.

It has been recognized that the densely connected networks with homogeneous pores favor the stress to distribute evenly and result in the ductile behavior.38 The images from the vicinity in front of the arrested crack tip exhibit some incipient cracks that consist of continuous cavities and promoted stress distribution (Figure 7c). These cavities are closely associated with fibrils, in 16


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which long cavities associated with CNF are found in a region ahead of the arrested crack tip, as marked by the arrows in Figure 7c. Besides, most of microcracks are found to be orientated along the CNF-polymer interface or near the fibrils via delamination (Figure 7d). Taken together, this crack initiation process in the composite hydrogels can be described as follows. When the sample is subjected to a load, stress concentration occurs around the interface due to the difference of modulus and Poisson’s ratio of CNF and PAAm. Since the CNF-PAAm interfacial hydrogen bonding is weaker than the CNF cohesive strength,15 the microcrack tends to occur at the interface and some CNF strips bridging the crack tips are found due to the interface debonding. With further loading, the neighboring microcracks extend their length and grow along the stress direction that finally develops into a macroscopic crack.

Figure 8. Microdeformation evolution of crack initiation and propagation in CNF-PAAm composites dried film by TEM observation. At earlier stretching (a), the homogeneously dispersed CNFs occurred mutual slippage due to the breakage of interfacial hydrogen bonding, initiating crack propagation at the interface. With increasing of loading (b), the PAAm chains produced plastic deformation from initial coiled conformation to straight one and the pull-out fibrils bridged the crack, leading to the transition from cavitation to crazing (c). Meanwhile the bridged CNFs resulted in significant restriction of crack propagation and maintained the structural integrity (inset of c). With further increased loading stress, microscopic crack formed along the trajectory path that perpendicular to the stress direction (noted by the arrows) and the catastrophic fracture occurred (d).

Fibril pull-out has been identified as one of the key mechanisms responsible for the high toughness of fiber-based composites.38 Theoretical analysis based on shear stress transfer model 17


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has indicated that the aspect ratio of the fiber acts a critical role in improving strength of the composite,39 since the larger aspect ratios provide efficient load transfer to CNF and thus increase the strength of composite hydrogels. Indeed, compare to the relatively short and rigid rod-shaped CNC, more flexible and longer CNF (Figure 8a) with unique local wrinkled surface is credited as a favorable medium for creating a robust interface that able to induce mechanical interlocking with the matrix due to its surface roughness.40 The microcracks tend to penetrate through the matrix and merge together with the progress of loading, where the crack path is quite tortuous (Figure 8b). The narrow zones with highly deformed lamellae configuration appear that crack propagation direction is perpendicular to the applied stress, and the CNF repulsion from craze fibrils in length results in CNF entanglement between craze fibrils (Figure 8c). During this merging process, the repelled CNFs are entrapped by propagating craze fibrils and are noted as large clusters in the center of the craze. With massive microcracks coalesce into the macrocrack and the increase of fracture surface area due to crack deflection, the catastrophic fracture occurs (Figure 8d). According to the interfacial slip micromechanics model by Glaz et al.,41 once the load transfer from the matrix to the fibrils is obtained, the resulting interfacial shear force can be determined via interfacial slip mechanics. While it is important to recognize that the interfacial slip responsible for energy dissipation is also detrimental to composite strength and stiffness. This result highlights the importance of interface dynamics in engineering composite properties. Based on the above results, the microdeformation and toughening mechanisms of composite hydrogels can be summarized into the following two interface-dominated aspects. One is the microcrack deflection at the interface, where the staggering randomness to a certain extent alters the shear stress at the interface and leads to a progressive interface failure via a highly meandering crack path. In fact, due to the large aspect ratio of CNF in this study (50-100), such parameter promotes a large amount of interface failure region to dissipate elastic energy. The observed multiple crack bifurcation in Figure 8c can be viewed as a proof of highly meandering microcrack stemmed from interface failure. The other primary toughening mechanism is crack bridging at a larger scale (inserted image in c). With increase of the applied stress, the interface failure is dominated by interface sliding with friction, in which the crack tips bridge the crack and form a matrix bridged zone. One should note that even though CNFs have the potential to improve the mechanics of 18


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composites, this ability is critically dependent on the interfacial interactions and requires efficient load transfer between the filler and matrix. Indeed, according to the shear lag model by Cox,42 the long straight fibrils increase the viscoelastic energy dissipation associated with stress concentration at the ends of fibril due to the elastic mismatch between the filler and matrix. This model further highlights the importance of interfacial factors, such as hydrogen bonds, van der Waals interactions, mechanical interlocking, and contraction at the interface: once the shear stress exceeds a critical value, the filler would debond from the matrix and energy is dissipated by frictional sliding at the interface.

Conclusions The CNF hydrogels toughening mechanisms are explored by examining viscoelastic properties as well as its alignment in morphology distribution under stress. The mechanical properties of composite hydrogels with dual cross-linked network and asymmetric strain rate dependent loading and unloading profiles are manifestation of dynamics of reversible hydrogen bonds. In the small deformation regime, the modulus and strength of gels increase with strain rate and show the remarkable Mullins effect during unloading-unloading cycles with large hysteresis loops: at lower rate, the chemical cross-links dominate the mechanical properties and result in small hysteresis loops; at higher strain rate the gels show significant energy dissipation that stems from strain-induced dissociation of H-bond. The creep fracture time decreases exponentially with loading stress, implying the unique thermally activated hydrogen bonding dissociation mechanism. The evidence of toughening mechanisms from morphological insights, including fibril pullout, crack bridging, and nucleation of voids surrounding debonded fibrils are observed on CNF reinforced composite hydrogels fracture surface. The large aspect ratio of these fibrils, combined with their attractive interfacial bonding with matrix and remarkable tensile properties, could act to arrest crack growth. This work promotes further efforts toward tough hydrogels through the integration of covalent and reversible bonds at molecular scale and the applications of cellulosic hydrogels in high-loading condition required fields.

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Supporting Information Extraction of CNF and its size distribution, DSC test of composite hydrogels, cyclic tensile test of PAAm gels under different strain rates, stretch history of sequential stress-relaxation experiment, tearing properties in urea aqueous solution, and CNF volume fraction dependent crack propagation critical stress.

ACKNOWLEDGMENTS This work was financially supported by National Natural Science Foundation of China (21674013).

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