Article pubs.acs.org/Macromolecules
Cellulose Nanocrystals Mechanical Reinforcement in Composite Hydrogels with Multiple Cross-Links: Correlations between Dissipation Properties and Deformation Mechanisms Jun Yang,* Chun-Rui Han, Xue-Ming Zhang, Feng Xu, and Run-Cang Sun Beijing Key Laboratory of Lignocellulosic Chemistry, College of Materials Science and Technology, Beijing Forestry University, Beijing 100083, China S Supporting Information *
ABSTRACT: Nanocomposites have drawn a great interest in materials science of elastomers in recent years, and tailoring interfacial interactions between fillers and polymer matrix plays a critical role in improving their mechanical properties. The synthetic platform of tough and stretchable cellulose nanocrystal−poly(acrylamide) (CNC−PAM) composite hydrogels was proposed and applied here to unravel the role of covalent network in PAM and physical interactions by CNC surface adsorption. The attractive physical interactions in the network were considered to increase the fracture strength of the hydrogels via reversible adsorption−desorption processes on the CNC surface. Stress-sensitive characteristic shifts of the Raman peak located at 1095 cm−1 indicated an efficient load transfer across the interface, where the tensile modulus was higher than the compression modulus. In situ transmission electron microscopy observation allowed to demystify the composites deformation process and interfacial bridging between CNC and polymer matrix. A detailed comparison of strain rate effect on large strain dissipation indicated that the viscoelastic behavior of the hydrogels varied remarkably over strain rates, ranging from little hysteresis at low strain rates to highly dissipative at high strain rates, suggesting a new, slow relaxation mode, most likely due to interfacial adsorption of polymer chains on the CNC surfaces. This study showed that polymer chains desorbed from the CNC surface under periodic strains would entangle with the free chains after the rest time via conformational rearrangements, consequently triggering a recovering mechanism during multiple crazing and shear relaxation processes.
1. INTRODUCTION Hydrogels with a cross-linked network swollen in water can be a good candidate for various applications ranging from biomedicine to daily chemicals, such as drug delivery systems, biosensors, and superabsorbent in diapers.1−4 While the major problem of this soft-and-wet material is its low fracture strength and poor flaw tolerance, many efforts have been contributed to enhance the mechanical properties in the past decade.5−7 Among them, one of the most promising strategies for reinforcement is the introducing of “sacrificial” bonds, where the dynamic and reversible noncovalent interactions complement the structural integrity ahead of a propagating crack.8−10 The physical interactions based on various temporary crosslinks, including hydrogen bonding,11 ionic interactions,12 or hydrophobic associations,13 have been widely introduced to the primary covalent bonds, which provided a unique method to allow the materials to dissipate a huge amount of energy under stretching and recovered the original shape when the stress was removed. Dispersing rigid fillers into an elastomer is a well-known strategy to reinforce polymeric composites because it significantly expands applications of the materials,14−16 and there are mainly three distinct theories that have been proposed to demystify the origin of reinforcement.17 At one extreme, © XXXX American Chemical Society
mechanical reinforcement is suggested to the agglomeration of particles, viz., stemming from a direct pathway for percolated agglomerates throughout the matrix.18,19 In contrast to this “particle-only” scenario, the “interaction zone” theory suggests that due to the favorable filler−polymer interactions, chain immobilization occurs around the fillers and decreases the segment dynamics.20 This immobilized layer further overlaps and leads to the formation of particle-mediated polymer networks with improved elasticity.20,21 Finally, a third proposal is that polymer chains bridge neighboring particles and form a particle network with origins in continuum mechanics.23 Indeed, the addition of microscopic fillers to polymer matrix would lead to a distortion in strain fields, resulting in a significant change in deformability of the composites. The description of interfacial interactions of composite hydrogels and in particular of their dissipation mechanisms is a perquisite to explore macroscopic properties with increased resistance to the applied deformation.8 Gong et al. pioneered the idea of introducing dissipation in chemically cross-linked networks and proposed a unique double-network (DN) gel, Received: April 8, 2014 Revised: May 24, 2014
A
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Figure 1. Schematic illustration and network microscopic structure of CNC−PAM hydrogels. (a) Well-dispersed aqueous suspension of CNC particles and gelation. (b) Proposed structural model for multiple cross-linked networks. (c) The networks combine the both covalent cross-links and noncovalent hydrogen bonds, which are analogous to a double network where the covalent bonds lead to elastic properties and noncovalent interactions promote viscoelastic and recovery properties. (d) Digital image of translucent CNC0.1 gels. (e) Cross-section morphology of CNC0.1 gels, showing interconnected porous structure. (f) Enlarged SEM image, showing homogeneous distribution of CNC fibrils (indicated by arrows).
surface area of nanofibers and hydrogen-bonded assembly of parallel chains in native cellulose−crystalline domains, it is envisioned that if load can be effectively transferred to the hard reinforcing phase, and then the modulus of the composites would be similar to that of randomly orientated rigid CNC with high strength and modulus.35 Although the dynamics of both chemical and physical gels are well characterized, new challenges arise when the two types of cross-links coexist, in particular, what is the dissipation mechanism and load transfer in composites, which are viewed as prerequisites to demystify their large deformation properties. Besides, the hydrogels with reversible cross-links normally exhibit rate-dependent viscoelasticity due to the multiple timescale energy dissipation mechanisms.8 Herein, we report on viscoelasticity and energy dissipation of the cellulose nanocrystals−poly(acrylamide) (CNC−PAM) hydrogels containing chemical cross-links and physical bonds (Figure 1), with the strain sensitivity by Raman peak shift. For this purpose, a small amount of chemical cross-links via covalent bonding is introduced to prevent the network from plastic flow and ensure the samples have a finite structure at rest, facilitating systematical mechanical evaluation thereby. This platform combines the concept of DN and NC gels by coupling a covalent network with reversible interactions using CNCs as physical cross-links. From a practical point of view, this study may lead to the development of reliable protocols for designing tough hydrogels from the perspective of multiple interactions, and the originality of this work resides in the discussion of dynamics of polymer chains combined to the attractive CNC interface. Fundamentally, this contribution rationally stimulates the development of next-generation tough hydrogels and understanding their mechanical performance.
where a primary densely cross-linked network entangled and interconnected with a second sparsely cross-linked network.24,25 The results demonstrated that the dramatically improved fracture strength was mainly attributed to the efficient stress transfer and enlarged damage zone upon stretching.25 However, the DN architecture was susceptible to fatigue once the bonds were broken and the gels soften permanently after large deformation, since all the sacrificial bonds of brittle networks were covalently cross-linked. As another example, Haraguchi et al.26 developed a series of highly extensible nanocomposite gels (NC gels) by in situ polymerization of hydrophilic polymer chains from clay surface, showing an impressive fracture elongation (up to 1400%) and tunable mechanical stiffness. Unfortunately, some notable residual strains were noted after continuous deformationresting process due to the “clay-brush” alignment. Native cellulose, one of main natural polysaccharides on earth, has been attracting great attention as a renewable biomass (extracted from cotton, wood, bacteria, and sea tunicates, etc.) due to its unique features, such as hydrophilicity, renewability, biodegradability, and high strength.27−29 Cellulose nanocrystal (CNC), with width of 5−30 nm and length ranging from 20 nm to several micrometers, is a highly rodlike crystalline component in cellulose that is isolated through strong acids hydrolysis, mechanical disintegration, or a combination of these.30 A large body of work has been pursued in elastomers for CNC as a reinforcing filler due to its amazing characteristics, such as nanoscale dimensions, high aspect ratio and Young’s modulus, low density, and surface functionalization (reactive surface of −OH side groups).31−33 Particularly, considering the good hydrophilicity and biocompatibility, CNCs have been broadly applied as fillers in reinforced composite hydrogels and explored the role of interfacial layer on the mechanical reinforcement.34,35 Because of the large B
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500 N load cell. Sample dimensions were kept constant with a rectangular shape: length = 40 mm, width = 4 mm, and height = 5 mm. The both ends of a sample were clamped along the deformation direction, and the stretching rate varied from 1.8 to 360 mm/min, corresponding to nominal strain rate (ξ) comprised between 1 × 10−3 and 0.2 s−1. To minimize dehydration, a thin layer of low viscosity of silicone oil was coated on the sample surface. Loading−unloading cycles were performed to examine the characteristic times of viscoelastic processes at three typical strain rates: 1 × 10−3, 0.0167, and 0.2 s−1. Cycles were performed from 0 to a maximal nominal strain (ε = 400%), and then the sample was unloaded at the same strain rate until zero stress was achieved. Strain was defined as the ratio (l − l0)/ l0, where l and l0 are the sample length during stretching and the initial length. Nominal stress, defined as applied stress, was calculated from the tensile force divided by initial cross section area (4 × 5 mm2). For each sample, at least four duplicates were tested to ensure good reproducibility (±5%). The fracture strength σ and fracture strain ε were the values at breaking point. The Young’s modulus was determined from the slope within the initial linear region of the stress−strain curves. The work of extension, a parameter to characterize the toughness of the sample, was determined by the area under the stress−strain curve until the fracture of the sample. The elastic recovery values of the hybrids were obtained from the cyclic tests. Rheology. The dynamic rheological properties of the hybrid hydrogels were measured using a TA AR2000 rheometer equipped with a Peltier device for temperature control. A 25 mm parallel-plate geometry and a gap of 1 mm were used to apply a stabilized normal force of 0.1 N (corresponding normal stress around 0.2 kPa) for all experiments at 25 °C. The large amplitude oscillatory shear measurements were conduct at a frequency of 1 Hz. The time dependency of the recovery was tested by fixing the strain amplitude at 300% and allowing the sample rested with increasing periods of time up to 900 s. Swelling Measurements. The cubic samples (1 × 1 × 1 cm3) were weighted and immersed into a large excess of water at 25 °C. After reaching swelling equilibrium for 72 h, the cubes were taken out and weighted. The swelling ratio (Qe) was defined as Qe = (Wt − Wd)/ Wd, where Wt and Wd represent the weight of swollen hydrogel and dry hydrogel, respectively. For all the samples, the amount of extracted polymer (e.g., linear PAM) analyzing by HPLC, if observable, was less than 1.5 wt % of the total organic content. Cryo-Scanning Electron Microscopy (Cryo-SEM). The hybrid gels were uniaxially stretched and “fixed” with tape, then cryo-fractured by immersing into liquid nitrogen, and then cut with a cold scalpel. The frozen and cut samples were sublimated at −90 °C for 5 min, followed by sputter coating with platinum at −130 °C for 20 s. Then coated samples were observed with a DSM 960 Zeiss SEM at −140 °C. Transmission Electron Microscopy (TEM). The small piece of dried gels were placed in a mold with cubic shape (∼1 cm3), and a hardening resin was added to the mold that the samples kept its position in the approximate middle of the resin. Then the mold with embedded sample was dried in an air-circulation oven for 48 h to obtain the dried samples in hard resin. The gel sample was fixed into the slice cutting machine, where the cross section of samples was sectioned at 90° relative to the surface (minimize the distortion of the structure) using a Leica Ultracut-E microtome that equipped with a diamond knife. The approximately 200 nm thick sections were collected, stained on a carbon support gird, and then observed by JEM-1010 (JEOL). Fibrous Network Deformation. The molecular orientation of individual fibrous in composite samples under compressive and tensile deformation was characterized using a Raman spectrometer (LabRAM XploRA, HORIBA Jobin Yvon) with a 785 nm near-infrared laser. The Olympus BH-1 microscope with 50× objective lens was applied to focus on the samples surface to a spot size of approximately 2 μm. The intensities of the Raman band at ∼1095 cm−1 were recorded using a least-squares fit of a mixed Gaussian/Lorentzian function.
2. EXPERIMENTAL SECTION Materials. A never-dried native bleached pulp used as original cellulose was obtained from Donghua Pulp Factory, China. The monomer acrylamide (AM) was recrystallized from a toluene/nhexane mixture and dried in a vacuum at 30 °C. The chemical crosslink N,N′-methylenebisacrylamide (MBA) and potassium persulfate (KPS) were obtained from Beijing Chemical Reagent Company, China. 2,2,6,6-Tetramethylpiperidine-1-oxyl (TEMPO, C9H18NO) was purchased from Sigma-Aldrich. All other reagents were analytical grade and used without further purification. Extraction of Cellulose Nanocrystals. Cellulose nanocrystals (CNCs) were prepared by the sulfuric acid hydrolysis procedure according to previous literature.35 In brief, 2 g of pulp was added into 100 mL of 55 wt % sulfuric acid at 50 °C for 1.5 h under mechanical stirring (300 rpm), where amorphous and paracrystalline regions were preferentially hydrolyzed. The suspension was diluted with ice cubes to stop reaction and washed by centrifugation and dialyzed to neutrality. The above CNC suspension was completed by ultrasonication treatment for 10 min, and finally the CNC powder was collected by freeze-drying (TEM image of CNCs in Supporting Information Figure S1). TEMPO-Mediated Oxidation of Cellulose Nanocrystals. TEMPO-mediated oxidation of CNCs was conducted according to a procedure described by Isogai et al.36,37 About 1 g of CNC was suspended in water (100 mL) containing TEMPO (0.016 g, 0.1 mmol) and sodium bromide (0.1 g, 1 mmol) and treated by ultrasonic dispersion at an ice−water bath for 10 min. Then desired amount of 10 wt % NaClO solution (5 mmol per gram of cellulose) was slowly dropped into the above suspension to start oxidizing reaction at room temperature by stirring at 300 rpm for 12 h. The system pH was maintained at 10 by adding 0.5 M NaOH using a pH stat until no NaOH consumption was observed. After oxidation, the reaction was quenched by adding ethanol (5 mL), and the oxidized CNC was washed thoroughly with water and then ethanol on a filter paper set in a Bϋ chner funnel. Finally, the TEMPO oxidized CNC was homogenized at 5000 rpm for 1 min and sonicated for 2 min to attain a CNC water dispersion (∼5 wt %). The mass recovery ratio of the TEMPO-oxidized cellulose was over 85%, and the carboxylate content of the oxidized cellulose was examined using an electric conductivity titration process and found to be 0.052 mmol/g. Preparation of CNC−PAM Hydrogels. Nanocomposite hydrogels were prepared at 35 °C by free radical polymerization of AM in an aqueous suspension of CNCs using KPS as initiator. The solid reagents, including MBA, AM, and KPS, were initially dissolved in water prior to synthesis at room temperature. The homogeneous mixture of TEMPO-oxidized CNC, AM, MBA, and KPS was deoxygenated by nitrogen bubbling for 10 min under mechanical stirring and then transferred into molds under a nitrogen atmosphere. The mixture was left for 24 h to allow the reaction to run to completion, and the composition of hydrogels and their nomenclature are listed in Table 1. For all the synthesis, the amount of CNC was varied while maintaining AM and MBA compositions. The nomenclature of samples was named by CNCx, with x corresponding to the volume fraction of CNC. Large Strain Mechanical Measurements. The mechanical tests were conducted using a Zwick Z005 testing machine equipped with a
Table 1. Composition of CNC Gels code
CNC volume contenta (%)
mCNC (mg)
mAM (g)
mMBA (mg)
mwater (g)
mKPS (mg)
CNC0 CNC0.1 CNC0.2 CNC0.4 CNC0.6
0 0.1 0.2 0.4 0.6
0 40 80 160 240
4 4 4 4 4
15 15 15 15 15
25 25 25 25 25
12 12 12 12 12
a CNC volume fraction was defined as the volume ratio of CNC to the total volume of the composites using density of CNC = 1.6 g/cm3.28
C
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Figure 2. Mechanical properties of CNC gels with different CNC contents. (a) Monotonic stress−strain curves at ξ = 0.0167 s−1. (b) Young’s modulus versus CNC volume fraction; the dashed line represents the theoretical predictions for the Halpin−Tsai model.
3. RESULTS The general architecture of this series of CNC−PAM hydrogels (herein after abbreviated as CNC gels) is schematically shown in Figure 1, and three specific operating techniques are noted: (1) CNCs were prepared by TEMPO-mediated oxidation of pulp fiber and formed abundant sodium C6-carboxyl groups on the CNC surfaces, which facilitated the electrostatic colloidal stabilization and load transfer in composites.37 (2) A small amount of MBA as chemical cross-linkers was applied to fix topological network and excluded the PAM self-cross-linking process.38 (3) To be able to quantitatively compare samples with different CNC contents but equivalent swelling ratio, the mechanical properties were systematically evaluated at the gel preparation conditions; thereby the gel hydration parameter was excluded in discussion of network structure. The synthesis process of CNC gels was shown to be simple and highly reproducible, which allowed the preparation of hybrid hydrogels within a wide range of CNC content (up to a CNC volume fraction of 0.6% for CNC0.6). As forementioned, the gel composition after swelling equilibrium indicated that the polymerization was almost complete and polymer chains did not leach out. In comparison with the pristine PAM gels (CNC0 in Table 1), the hybrid hydrogels swelling ratio Qe decreased with increasing amount of CNC (Qe scaled with VCNC as Q ∼ VCNC−0.84 in Figure S2), suggesting that CNCs acted as multifunctional physical cross-links. To probe CNC dispersibility in PAM matrix, TEM images from fractured interior section of hydrogels were observed (Figure S3), where CNCs homogeneously dispersed in polymer matrix. Besides, the uniaxial deformation process was homogeneous during the tensile test, and no necking phenomenon was noted (Figure S4). Mechanical Properties of Hybrid Hydrogels. Given the existence of attractive interfacial bonding between CNC and matrix and homogeneous dispersion of CNC (Figure S5), one can expect some extent of strength reinforcement in the obtained CNC gels. The mechanical properties of the CNC gels were investigated by two types of tests: monotonic and cyclic tensile measurements. The typical monotonic stress− strain curves at different CNC contents are drawn in Figure 2a, and the specific mechanical properties are summarized in Table S1. The results showed that CNC gels possessed the both higher Young’s modulus and fracture strength than the pristine PAM and the composites elongated with increasing content of
CNCs. For example, the modulus of the CNC0.6 with 0.6% v/v of CNC increased by 6.1-fold from 5.5 to 33.8 kPa as compared to that of native PAM. On one hand, this increase in modulus is expected due to the addition of the rigid particles (the modulus of CNC estimated to be 110 GPa).28 On the other hand, this superior mechanical enhancement can be attributed to the strong interfacial adhesion between CNC and matrix, where polymer chains covalent cross-linking and interfacial noncovalent interactions lead to elastic properties as well as viscoelasticity.8 Furthermore, the increase in Young’s modulus of the hybrids was in agreement with theoretical values by Halpin−Tsai model that usually applied for the prediction of composites elastic properties based on the geometry and orientation of the filler (Figure 2b), implying homogeneously dispersed CNCs in the matrix with attractive interactions (see Supporting Information for model details). To further explore how the hybrid networks response to the applied stress, CNC0.1 gels were imaged with cryo-SEM at stretched state. Figure 3 shows the microstructure of CNC gels at a strain of 400%, where the crazes developed perpendicularly
Figure 3. Cryo-SEM images of CNC0.1 gels. The crazes with porous structures develop perpendicularly to the stretching direction (represent by arrows). The inserted image shows the cross-section pores where CNC acts as a bridge to prevent crack propagation: polymer chains around the active zone can be easily pulled into the fibrils of the crazes and withstand strain energy before failure via transient bonds. D
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scenario is verified in Figure 5a. The CNC gels were first stretched to the strain of 400% and then recovered to initial state (ε = 0), where a prominent hysteresis was noted and the hysteresis loop at a high strain rate became larger than that of at the low rates (Figure 5b). Whereas for purely elastic CNC0 gels, they behaved like an elastic system and energy dissipation of the matrix was negligible. Generally, the addition of rigid fibrils to polymer matrix resulted in increases in both Young’s modulus and brittleness of the composites due to the poor interactions between hydrophilic cellulose and hydrophobic matrix.14,18 Interestingly, this study indicated that attractive interfacial interactions between CNC and PAM matrix contributed to highly stretchable and tough hydrogels. The appearance of the hysteresis suggested that a part of extension energy was dissipated during the deformation, and the residual strain at free became notable that was named as the Mullins effect.41,42 The polymer chains that were in closed coiled conformation started to be extensively elongated, and then the adsorbed chains were forced to dissociate and thereby increased the resistance against crack propagation. When the strain rate was higher, the strain hardening and hysteresis became more obvious (Figure 5b). This result further suggests that the CNC gels possess selfrecovery ability, which relates to the orientation of CNC and polymer chains (residual strain before yielding point and successive loading−unloading curves in Figure S8).31,32 The understanding evolution of modulus on successive loading−unloading cycles is a critical indicator for internal selfrecovery capabilities,8,9 and the strain-rate-dependent recovery curves are shown in Figure 6a. A large decrease in modulus was noted for the high strain rate cycles as the number of cycles increased, whereas at low strain rates the modulus was almost unchanged. Besides, the residual strain after full unloading decreased as the strain rate was lowered. Since the modulus is related to the density of elastic chains,10 the initial modulus can be viewed as a specific parameter to examine the connectivity of the elastomeric networks, Thus, the reduction in modulus is considered as a decrease in the connectivity of polymer chains across a unit area and the recovery parameter, R, is defined as
to the stretching direction in the sample, suggesting the fibrils around the active zone can be easily pulled out and dissipated large amount of energy without compromising the structural integrity (more images throughout the incremental loading process in Figure S6). Thus, the capacity of the crazes to undergo large lateral elongation was directly responsible for the tremendous elongation. It was held that the flexibility of the CNC loaded composites occurred when the efficient stress transfer was recognized at the polymer−CNC interface.31,33 In such a case, the CNCs acted as multifunctional physical crosslinks to transfer energy across the interface and prevented crack propagation. Besides, the anchored polymer chains at the CNC surfaces led to an additional increase in the number of elastically active chains per unit of volume, which was qualitatively in line with the swelling behavior described above. Therefore, when a crack occurs, the rearrangement of adsorbed chains can dissipate crack energy and suppress the macroscopic growth of the crack, which, in turn, leads to the reinforced gels.21,39 Strain Stiffening and Recovery. To rationally guide the design of viscoelastic hydrogels with high toughness, it is a pivotal task to study the behavior of CNC gels in dissipating stress energy under different strain rates. Herein, the large strain behavior was initially measured in monotonic uniaxial tensile test to final failure at different strain rates (ξ), and the typical stress−strain curves are illustrated in Figure 4. The
R=
E2nd E1st
where E1st and E2nd is the modulus during the first loading cycle and second loading cycle, respectively. One can note that the recovery of the network connectivity was complete when deformed at low strain rates, and it decreased remarkably at ξ = 0.2 s−1 (Figure 6b). Intriguingly, the recovery in modulus was typically attained within 15 min, where the residual strain decreased and the cycling behavior was similar to that of the initial cycle (Figure 6c). To further examine the time-dependent recovery, time evolution of small-amplitude oscillatory shear, where the samples were subjected to a strain amplitude from zero to γ (greater than the maximum limit of the linear viscoelasticity) and decreased to zero in a certain time range, was performed to study CNC gels dynamic behaviors in nonlinear region. The specific experimental protocol where samples subjected to varying resting timeline was depicted in Figure 7a with γ = 3 at frequency of 1 Hz. The results indicated that with rest time increasing from 10 to 900 s the storage moduli (G′) of CNC gels increased with waiting time (Figure 7b). This result can be explained that the continuous deformation−resting cycles
Figure 4. Tensile mechanical properties of CNC0.2 gels at different strain rates.
characteristic time of CNC/polymer chains dynamics can be probed at this large strain, where the toughening properties were affected by ability of the network to relax stress by dynamic interactions between adsorbed chains and desorbed segments on CNC surface. Increasing the strain rate led to the more pronounced reinforcement, suggesting the part of chains relaxed at 1 × 10−3 s−1 became unrelaxed at 0.2 s−1. Besides, if the strain rate was further decreased (e.g., ξ = 4 × 10−4 s−1), the reinforcement properties of CNC on matrix become negligible (Figure S7). In fact, the reinforcement mechanism for DN gels with bicontinuous phases has been suggested that the rigid and brittle polyelectrolyte network acted as sacrificial bonds to increase the resistance against crack propagation.25 In this regard, we assume that the toughness of CNC gels roots in the sacrificial bonds between CNC and polymer matrix. This E
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Figure 5. Stress−strain curves of CNC gels. (a) Loading−unloading cycles with different CNC contents at strain rate 0.0167 s−1. The inset shows the general shape of continuous loading−unloading cycles for CNC0.6 gels. (b) Loading−unloading cycles at different strain rates for CNC0.4 gels.
Figure 6. (a) Strain rate dependence of continuous tensile curves and (b) modulus recovery using the same sample (CNC0.4) and (c) cycling behaviors versus resting time.
after 12 min resting, the fracture strength of CNC0.4 gels increased from 62.9 to 148 kPa, suggesting the continuous deformation−resting process led to the denser adsorbed polymer chains and solid-like recovery. Load Transfer and Deformation. It is well-known that a shift in Raman spectroscopy at 1095 cm−1 at deformation, corresponding to C−O stretching along the chain axis, indicates deformation along the backbone of cellulose chains.43,44 The Raman study of interfacial characteristics in CNC gels was performed, and the strain-induced vibrational frequencies change in Figure 8 mirrored the load transfer to the CNCs. The shift in the Raman peak with strain in compression is positive and in tension is negative and about 0.6−0.8 wavenumbers/percentage applied strain. Interestingly, it is noted that gradient of the Raman peak (1095 cm−1) in the
strength the both entanglement of desorbed chains with free dangling and rearrangement of PAM chains during resting. When the polymer chains are stressed at large strain amplitudes, the physical interactions between PAM and CNC desorbed from the interface, whereas the shear stops for some resting time; the polymer chains relax and readsorb to the interface. Accordingly, this continuous cycle generates a denser adsorbed region and leads to a stiffening at the interface. The loss factor tan δ = G″/G′, representing the ratio of dissipated energy to stored energy during deformation,42 was also calculated for each cycle, and the results indicated deviation from the viscous-like behavior to solid-like property with resting time (Figure 7c). Indeed, this notable increase in the stress−strain curves was another support of this interfacial recovery via attractive interactions (Figure 6c). For example, F
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Figure 7. (a) Scheme of strain−rest time measurement at 25 °C. (b) Elastic moduli and (c) loss tangent of CNC0.4 gels in nonlinear deformation experiments at 300% strain amplitude and frequency 1 Hz. Inset in (b) indicates the normalized elastic modulus versus rest time.
between PAM and rod-like CNCs is sufficient for effective load transfer. Many CNC composites present high ultimate strength and Young’s modulus and the primary reason is strong interfacial interaction.31−33 The intriguing feature in Figure S3 is that in contrast to native cellulose fibrils (Figure S1), CNCs remain curved and interwoven in the polymer matrix that can act as effective local crack arresters, leading to flexibility and efficient load transfer. Thus, the strength and modulus show positive correlations with the CNC content, suggesting the percolated CNC network is the major load-carrying phase with strong interfacial adhesion. Confluence of Toughening Mechanisms. Mechanistically, composites break under filler fracture mode are strong but brittle and thus flaw-intolerant, whereas composites that fail under the filler pull-out mode are usually weak but ductile and thus flaw-tolerant.46 In this regards, it is necessary to distinguish a mutual competition between intrinsic rupture that proceeds ahead of the crack tip to facilitate propagation and extrinsic crack tip shielding mechanisms that normally act behind the crack. 8 For this purpose, a series of complementary deformation measurements were performed by in-situ TEM observation. The condensed electron beam onto the carbon support film resulted in local thermal stress and initiated cracks in the composites. The growth of crack could be controlled by varying the electron beam flux onto the samples. Thus, the images were recorded at intervals along the crack to observe deformation in the composites as a function of crack opening displacement (Figure 9). One can note that the cracks initiated at low CNC density regions and propagated along the weak CNC−PAM interface. It should be pointed that the CNC aligned almost perpendicular to the crack extension (indicated by arrows) and bound the neighboring cracks, providing efficient energy dissipation via interfacial stick sliding at interface. With increasing stress, the CNCs began to break and pulled out of the matrix when the crack opening displacement exceeded approximately 350 nm. Thus, the pullout of the CNCs in a bridging zone in front of the crack could largely dissipate mechanical energy in CNC gels upon crack extension, which is different from other mechanical dissipation mechanisms, such as transformation of chains or cross-linkers,
Figure 8. Cellulose Raman peak shift located at 1095 cm−1 as a function of applied uniaxial strain in tensile loading compared to compression.
tensile model is slightly larger than that in compressive model, and this difference in Raman response in tension and compression may be due to the efficient load transfer of CNC in tension compared to that in compression.43 One explanation for this result is that while it is difficult to load the nanofibrils in compression, those orientated at the angle to loading direction are easily transferred in tension.44 In tension, the load is efficiently transferred into the inner layers of the CNCs through easy buckling and the bent section of the fibrils, where the slippage of the polymer layers is promoted due to the dynamic interactions between CNC and polymer.45 In fact, the straight and stiff cellulose molecules have a considerable tendency to rearrangement to their parallel conformations at stress, which are the basic elements of reversible network dynamics.44,45 Our results indicate that the interfacial interactions between CNCs and the polymer matrix play a crucial role in the load transfer efficiency from the ductile polymer phase to the strong filler phase. Since the CNC-bridged network builds a continuous fibril network with attractive interfacial features added in terms of ductile property, the interfacial interaction G
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Figure 9. TEM observation of crack nucleation and propagation in matrix (a), where the cracks nucleate along weak CNC−PAM interfaces (b, c) and the CNCs tend to pull out of the matrix (d), indicating external load can be effectively transferred to the CNCs. Schematic showing fracture behavior in term of CNCs pull-out upon crack propagation in the fiber-reinforced composites.
friction of filler with matrix, and rapture of polymer chains or cross-linkers.8 Several groups have shown that the efficiency of load transfer and toughening mechanisms in composites is critically dependent on the quality of interfacial interactions on differing length scales.9−11 Indeed, in this pursuit of high damage tolerant materials, mechanisms for alleviating local stress regions through shielding of fillers within the matrix to provide for extrinsic toughening (e.g., crack bridging and deflection, interface sliding) can generate a very potent source of toughening on multiple mechanisms.46,47 This so-called “pullout” of toughening by inter-CNC shearing in the wake of the crack tip through the soft phase associated with limited sliding is particularly potent in extensive crack-tip shielding and hence the prime origin of the fracture resistance. Consequently, CNCs that are aligned more parallel to the crack propagation direction tend to bridge between the crack faces, which may contribute to unprecedented flaw tolerance.
CNC/PAM interactions can break and recover through efficient energy dissipative processes. It is proposed that the resistance of polymer adsorbed particles to crack propagation and reversible interactions are probably important factors that affect the interfacial stress transfer.8 Since the mechanical properties of CNC gels are dominated by the fraction of CNC and MBA, it is necessary to clarify what are the roles that MBA and CNC act in the elastomers at different stretching stages. Before stretching, CNCs are homogeneously dispersed in PAM matrix. When the uniaxial elongation test starts at a nonlinear regime, the deformation is attributed to the randomly cross-linked chains at relative small strains. As the strain becomes higher, the physical interactions between CNC and PAM begin to dominate the viscoelasticity through hydrogen bondings. With further stretching, the stress is transferred to the CNCs while resisting fracture through the large and strong interfacial plane, where physical cross-links break and hence relax the high stress before crack tip, leading to the notable strain recovery phenomenon. Notably, due to the soft phase could connect neighboring rigid CNC strands, the extensibility of the polymer segments also attribute to the strain recovery before the final fracture of the elastomers. It has known that polymeric composites usually fail in three modes: brittle fracture, crazing, and shear yielding.15 For pristine PAM, the covalent bonds dominate the brittleness and the damage principally takes the form of microcracks in the absence of cracked-ligament bridging (Figure S9). In contrast, CNC−PAM hybrids possess sufficient chain segments mobility to allow local plastic flow and result in notable viscoelastic behaviors, where CNC bridges effectively span the crack and transfer load that would otherwise be used to promote macroscopic crack. In particular, there are two main approaches of load transfer from a filler to a matrix, including micromechanical interlocking, and chemical bonding or noncovalent bonding (e.g., van der Waals, hydrogen bonds) between fillers and the matrix.48 The first one is not guaranteed due to the
4. DISCUSSION The CNC gels here can be depicted as a double network combining a permanent network through chemical cross-links and a transient network ensured by surface adsorption. Since the strength of physical junctions between CNC and polymer matrix is weaker than those of covalent bonds, the stressing energy accumulated under large strains would be gradually released through the yielding of these physical interactions, and the magnitude of dissipation is determined by the amount of reversible cross-links as well as the characteristic time involved in the CNC/polymer association. Furthermore, the CNC− PAM dissipative interactions have the ability to recover within a given time scale. For example, under the large amplitude loading−unloading cycles, it was demonstrated that the characteristic time required for recovery appeared to be the order of 103 s, which enabled the network to self-heal and reformed an equivalent density of effectively elastic chains (Figure 7b). Therefore, depending on the time window, the H
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CNC relative smooth surface.28 The latter may be more reasonable in this contribution due to the high surface area of nanofibers creates a large interfacial region via attractive interactions between CNCs and PAM matrix. In fact, according to molecular dynamic simulation by Gersappe,49 the nanoscale fillers acted as temporary cross-links, providing localized regions of enhanced properties and retarding growth of cracks. Thus, the formation of plastic zone in front of cracks prevents crack propagation that is absent in the pristine PAM, and the tortuous crack paths make a major origin in toughness.
Program of Higher Education of China (20120014120006), Chinese Ministry of Education (113014A), National Science Fund for Distinguished Young Scholars (31225005), and Program for New Century Excellent Talents in University (NCET-12-0782).
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5. CONCLUSIONS Time-dependent dissipative mechanisms and recovery of highly stretchable hybrid hydrogels were investigated on a model of the CNC−PAM system. The CNCs acted as multifunctional physical cross-linkers, facilitated the transient anchoring of surrounding polymer chains, and provided viscoelastic character of the composites. The elastomers can be considered to some extent as double networks, where covalent network dominated the elasticity and physical network formed by adsorption of polymer chains on CNC surface attributed to the time-dependent stiffening effect. Given the right time window, the composites were able to dissipate more amount of energy than the chemical matrix can store and the network was able to recovery. Thus, the sacrificial interfacial interactions avoided permanent damage of the covalent network by releasing stress and promoting efficient energy dissipation processes. Microscopy observation revealed that CNC inside the craze can be extends to large deformations before final failure. It is assumed that such a high reinforcement is due to the effective interfacial interaction between the CNC and PAM as well as homogeneous dispersion of CNC, which maximizes the area of the load transfer between CNC and PAM. This elastomer had a strain-to-failure as high as 450% due to efficient loadcarrying feature of the continuous CNC network. The general principle here for designing tough hydrogels should contain substantial amount of mechanical energy dissipation under large deformation and maintain the network original configurations after deformation. Extending this work to functionalized CNC assemblies and polymer interfaces would pave the way for creating tough hydrogels at multiscales.
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ASSOCIATED CONTENT
S Supporting Information *
TEM images of acid hydrolysis obtained CNC, tensile properties, TEM observation of CNC gels, swelling ratio as a function of CNC contents, optical images of CNC gels deformation, description of Halpin−Tsai equation, stress− strain curves of CNC gels at a superlow strain rate, cryo-SEM images of CNC gels at stretching, and native PAM crack propagation. This material is available free of charge via the Internet at http://pubs.acs.org.
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REFERENCES
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected] (J.Y.). Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was financially supported by Fundamental Research Funds for the Central Universities (TD2011-10), Doctoral I
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