Thermoresponsive Supramolecular Hydrogels with High Fracture

Sep 14, 2018 - Supramolecular hydrogels formed from random copolymers of N-isopropylacrylamide (NIPAM) and 2-(N-ethylperfluorooctanesulfonamido)ethyl ...
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Thermoresponsive Supramolecular Hydrogels with High Fracture Toughness Fei Wang and R. A. Weiss* Department of Polymer Engineering, University of Akron, Akron, Ohio 44325, United States

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S Supporting Information *

ABSTRACT: Supramolecular hydrogels formed from random copolymers of N-isopropylacrylamide (NIPAM) and 2-(Nethylperfluorooctanesulfonamido)ethyl acrylate (FOSA) exhibit a volume phase transition due to a lower critical solution temperature (LCST). The LCST can be tuned between 32 and 5 °C by incorporating up to 14 mol % FOSA in the copolymer. The tensile modulus and strength of the hydrogels increased above the LCST as a consequence of an increase in the effective cross-link density due to the phase separation of water from the hydrogel. Below the LCST, the hydrogels exhibited extraordinary fracture toughness and crack blunting capability. Fracture energies of ∼8000 J/m2 were achieved, which is comparable or greater than that of cartilage and what has been previously achieved with synthetic supramolecular hydrogels. These gels exhibited large hysteresis behavior, though unlike tough double network hydrogels, the NIPAM/FOSA hydrogels can fully recover their dimensions (i.e., no permanent set) and properties after tensile deformations as high as 400% strain. Although the recovery of the macroscopic dimensions is relatively quick, the recovery of the microstructure requires times on the order of hours. The excellent energy dissipation behavior and recovery of the hydrogel is due to the reversible nature of the hydrophobic bonds and their aggregation into core−shell nanodomains, ∼6 nm in diameter. Under stress, the hydrophobic FOSA bonds break and the FOSA groups can rearrange either within the nanodomain or pull out of the nanodomain to dissipate energy. The bonds, however, reform when the stress is removed and the nanostructure heals. The hydrogels also exhibit stress-softening behavior (Mullins ef fect) as a consequence of the kinetics of the reversible structure and properties recovery following a deformation.

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INTRODUCTION

Supramolecular hydrogels are attractive alternatives to conventional covalently cross-linked hydrogels, because they can be melt- or solution-processed, recycled, and reversibly deformed. A caveat, however, is that by nature of the reversibility of the supramolecular bond, these materials are viscoelastic and can exhibit creep if under load for a sufficient amount of time for the viscous response to become appreciable. Creep, however, can be restrained by adding covalent cross-links to achieve a hybrid hydrogel of physical and covalent cross-links.20 Abdurrahmanoglu et al.15 were the first to demonstrate the ability of supramolecular bonds, in their case, hydrophobic bonds, to significantly improve the toughness of hydrogels. Since then, there have been many reports of tough hydrogels based on reversible bonds, such as hydrogen bonds, ionic bonds, hydrophobic bonds, and polymer−filler interactions.21 Our research group has been studying the structure and properties of strong, tough, supramolecular hydrogels and hybrid hydrogels prepared from amphiphilic random copolymers composed of an alkyl acrylamides and a fluorinated

Hydrogels are soft, water-swollen polymer networks that have a variety of uses, such as contact lenses,1 scaffolds for tissue engineering,2 drug delivery3 and artificial cartilage,4 agriculture,5 water remediation,6 and supercapacitors.7 In general, however, hydrogels have poor mechanical properties, especially fracture toughness, that limit the scope of applications. Conventional chemical hydrogels with covalent cross-links are brittle, with fracture energies less than 10 J/m2.8 In the past 2 decades, many advances have been achieved at improving the toughness of hydrogels,9 including double-networks,10 nanocomposites,11 slide-ring structures,12 and supramolecular networks.13 Supramolecular hydrogels rely on noncovalent, physical bonds, such as electrostatic interactions,14 hydrophobic interactions,15,16 and hydrogen bonds17 to develop a network in water. Those materials can have high tensile strength, ∼106 Pa, and high fracture toughness, ∼103−104 J/ m2, which is comparable to or even greater than that of human knee cartilage, ∼103 J/m2.18 Further improvements in toughness of hydrogels to values ∼3 × 104 J/m2 can be achieved by reinforcing a supramolecular hydrogel or a hybrid hydrogel, i.e., one that has both covalent and supramolecular cross-links, with macroscopic fibers.19 © XXXX American Chemical Society

Received: March 7, 2018 Revised: August 15, 2018

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Macromolecules acrylate.20,22−26 As with the work of Okay and co-workers,15,27,28 the reversible cross-link in those materials are hydrophobic bonds, but instead of the hydrocarbon hydrophobes used by Okay and co-workers, our supramolecular hydrogels were based on fluorocarbon hydrophobes that provide a much stronger hydrophobic bond and produce much tougher hydrogels.16 In addition, the fluorocarbon-based physical hydrogels exhibit a microphase-separated microstructure consisting of ∼6 nm diameter core−shell nanodomains (a fluorocarbon core surrounded by a water-depleted alkyl acrylamide shell) connected to water-swollen alkyl acrylamide chains that comprise the continuous phase of the hydrogel,24 Figure 1.

Most of our prior work on supramolecular hydrogels focused on copolymers of DMA and 2-(Nethylperfluorooctanesulfonamido)ethyl acrylate (FOSA). This paper reports the mechanical properties and thermoresponsive nature of hydrogels prepared from NIPAM and FOSA. A major difference between the DMA and NIPAM-based hydrogels is that NIPAM hydrogels exhibit a volume phase transition (VPT) at ∼34 °C,35,36 which is a consequence of a lower critical solution temperature (LCST) where the polymer transforms from a swollen hydrated state to a collapsed partially dehydrated state upon heating. The only other tough thermoresponsive hydrogel that has been reported is that by Guo et al.,29,30 and as discussed above those materials had a completely different microstructure than the NIPAM-FOSA hydrogels discussed herein, i.e., bicontinuous vs discrete nanodomains and an order of magnitude larger size of the nanostructure. Another attribute of these hydrogels that is not directly addressed in this paper, but has been described previously,22 is that these gels are thermally processable and soluble in suitable solvents, so articles with complex shapes can be molded from a melt (e.g., injection molded) or cast from solution.

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EXPERIMENTAL SECTION

Materials. The NIPAM-FOSA random copolymers, NFx, used in this study were previously synthesized by a free-radical copolymerization in dioxane for 24 h at 60 °C.23 The chemical structure of the copolymers is shown in Figure 2, and the composition was

Figure 1. Schematic of water-swollen poly(alkyl acrylamide-cofluoroacrylate) copolymer hydrogels. The numbers denote (1) the water swollen poly(alkyl acrylamide) continuous phase and core− shell nanodomains with (2) fluoroacrylate core, and (3) waterdepleted poly(alkyl acrylamide) shell. The blue lines are the poly(alkyl acrylamide network chains.

The nanodomain microstructure of the fluorocarbonmodified poly(alkyl acrylamide) discussed in this paper and the use of nanodomains as a multifunctional cross-link in a hydrogel is unique, though we suspect that some of the other tough physical hydrogels reported in the literature may also be similarly microphase separated. With the notable exception of Guo et al.,29,30 who prepared tough, thermoresponsive hydrogels from graft copolymers of poly(N-isopropylacrylamide), poly(NIPAM) onto poly(dimethyl acrylamide), poly(DMA), no other research group appears to have characterized the microstructure of their tough supramolecular gels. The Guo et al. hydrogels had a bicontinuous microstructure with a periodicity of ∼88 nm, which differs from the thermoresponsive hydrogels in this paper, which have discrete nanodomains, Figure 1, that are an order of magnitude smaller in size. The lack of microstructure characterization for other tough supramolecular hydrogels, such as those with ionic bonds,14 is surprising in that nanodomain formation has also been previously reported in polyelectrolyte hydrogels,31,32 though those papers did not report toughness data. In addition to the excellent mechanical properties of hydrogels formed from copolymers of an alkyl acrylamide and a fluoroacrylate, the hydrogel can suppress or inhibit ice formation to temperatures as low as −70 °C33 and the copolymers can be electrospun into stretchable, tough hydrogel fiber mats.34

Figure 2. Chemical structure of NFx copolymers, where x indicates the mole fraction of FOSA.

determined by 1H NMR following the procedure described in ref 22. Four copolymers were used with FOSA concentrations of 5.4, 7.7, 10.2, and 14.2 mol %, denoted as NF5, NF8, NF10, and NF14. The copolymers were vacuum compression molded at ∼160 °C into 0.50 mm thick films, and transparent hydrogels were formed by immersing a film into deionized water and allowing the swollen film to equilibrate at the experimental temperature. The swelling ratio, SR, defined as the mass hydrogel/mass dry polymer was determined gravimetrically. The physical characteristics of the copolymers and the hydrogels are reproduced from ref 23 in Table 1. The properties of poly(NIPAM) homopolymer, poly(FOSA) homopolymer, and a covalently cross-linked poly(NIPAM) hydrogel, synthesized by a redox copolymerization of NIPAM and the cross-linker N,N′methylene bis(acrylamide), reported in ref 23 are also included in Table 1 for comparison. Note that the water absorption of the fluorocarbon core of the nanodomains is negligible,24 and although the composition of the core is essentially pure FOSA, it is not composed of poly(FOSA). Instead, the core is composed of FOSA groups from the different polymer chains that are associated due to hydrophobic interactions. Property Measurements. The effective cross-link density of the supramolecular network was calculated from the tensile modulus B

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Macromolecules Table 1. Properties of the Copolymers and Their Hydrogels dry copolymer23

hydrogel TLCST (°C)a

sample ID poly(NIPAM) poly(FOSA) Poly(NIPAM)gel NF5 NF8 NF10 NF14

FOSA (mol %) 0 100 5.4 7.7 10.2 14.2

Mn (104 Da)

Mw (104 Da)

Tg (°C)

4.8

7.2

145.1 44.5

3.4 4.2 4.3 4.8

6.4 7.2 7.2 7.5

130.6 128.2 126.0 117.0

SR5 °C

SR15 °C

SR22 °C

S

DSC

4.0 3.2 2.6 2.4

3.2 2.2 1.8

2.5 1.6 1.4

17.8 15.1 13.1 ∼11b

32.7 15.2 10.0 8.8 ∼7b

a

Values from swelling experiments (S) and differential scanning calorimetry (DSC)23 bEstimated by extrapolating values for the other three NFx samples and NF2 from ref 23.

Table 2. Tensile Properties of NFx Hydrogelsa

a Elongation rate = 50 mm/min. Unshaded boxes indicate T < TLCST (see Table 1). Lighter shaded boxes indicate T > TLCST. Darker shaded boxes indicate T ≈ TLCST.

using the theory of rubber elasticity37 and the assumption that the Poisson ratio was 0.5 (see the Supporting Information). Uniaxial tensile tests were carried out with an Instron universal testing machine, model 4204, using either a 5 N or a 100 N load cell, depending on the mechanical properties of the gels. An environmental chamber equipped with a liquid nitrogen gas purge system was used to control the test temperature at 5, 15, or 22 °C. For conventional tensile property experiments, dumbbell-shaped test specimens with a gauge length of 25.4 mm and width of 3.0 mm were cut from waterswollen NFx hydrogel films. The extension rate used was 50 mm/min, and five or more specimens were tested for each sample. Engineering values of stress (σ ≡ F/A0, where F is the instantaneous force required to stretch the sample) and strain (ε ≡ ΔL/L0) were used to calculate the tensile properties. Young’s modulus, E, was calculated from the slope of the initial, linear portion of the stress−strain curves. Tensile strength and ultimate strain were defined, as the stress and strain at break, respectively. Fracture energy, or fracture toughness, was measured using Suo et al.’s modification14 of a pure shear test developed for rubber by Thomas and Rivlin.38 Rectangular-shaped specimens with a width a0 ≈ 19 mm and thickness b0 = 0.680−1.14 mm were cut from the water swollen NFx hydrogel film. The initial sample length, L0, was ∼5 mm. Two types of hydrogel specimens were used, (1) un-notched and (2) notched with a razor edge cut of 9 mm (≈ 0.5a0) in the width direction in the middle of the gauge length, i.e., at L ≈ L0/2. The fracture energy was determined by stretching the specimen at 10 mm/ min. A video camera was used to record the critical strain, εc, at which the crack began to grow for the notched specimens. Five notched and unnotched sample measurements were made for each hydrogel. The area under the force-strain (F−ε) from ε = 0 to εc for the unnotched specimens is the work required to stretch the gel to the critical strain, U(εc), and the fracture energy was calculated from eq 1.14,38

ε

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U (εc) Γ= = a0b0

∫0 c FL0d(ε) a0b0

(1)

RESULTS AND DISCUSSION Swelling Properties. The swelling ratios, SR, for the NFx hydrogels at 5, 15, and 22 °C are listed in Table 1. The SR decreased as the FOSA content [FOSA] increased, which increases the cross-link density of the hydrogel due to the formation of more supramolecular bonds, and as temperature increased. The covalently cross-linked poly(NIPAM) hydrogel exhibited an LCST at ∼33 °C (measured by differential scanning calorimetry (DSC) from the onset of an endotherm associated with the LCST transition23), which is consistent with the reported value for poly(NIPAM) hydrogels.36 The supramolecular NFx hydrogels also showed LCST phase behavior, but the TLCST, was suppressed from that of the neat poly(NIPAM) hydrogels,23 see Table 1. Two values for TLCST for each NFx hydrogel are given in Table 1, one measured from equilibrium swelling experiments and one from the onset of the phase separation endotherm measured by DSC. Although it is not clear from the data in Table 1, a significant drop in the SR occurred at TLCST.23 Increasing the FOSA concentration lowered the TLCST and decreased SR at any temperature. Tensile Properties. The tensile properties of the NFx hydrogel at 5, 15, and 22 °C are summarized in Table 2. The shading of the data in Table 2 indicates whether that the test temperature was above, below, or near the TLCST. In general, Young’s modulus, E, and tensile strength, σu, increased with increasing [FOSA], i.e., as the cross-link density of the gel C

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Figure 3. Example tensile elongation experiments for NF10 hydrogel at 5 °C using an elongation rate of 10 mm/min: (a) stress vs strain curves for unnotched and notched geometries; (b) photos of tensile elongation of unnotched sample (in the first photo, the white dashed rectangle indicates the unstrained sample, L0 = 5 mm); (c) pure shear deformation of notched sample (white dashed rectangle in the first photo indicates the unstrained sample, L0 = 5 mm, and the solid white line signifies the initial 9 mm crack).

increased. E and σu also increased with increasing temperature, which was probably a consequence of the decrease in the water concentration, see Table 1. In particular, below TLCST the NFx hydrogels were strong and highly stretchable with mechanical properties competitive with the strongest physical hydrogels reported. The most significant changes in those properties occurred when T > TLCST, which is due to the collapse of the gel as the water phase separates. The trends for the ultimate strain, εu, of the hydrogels as a function of [FOSA] and temperature were not clear, except that εu decreased above T > TLCST, as the dehydrated gel became relatively brittle. Fracture Toughness. A typical tensile stress−strain curve for a pure shear deformation measured at 5 °C for a NF10 sample (19 mm width with a 9 mm edge-cut notch (crack) is shown in Figure 3a, and the photos in Figure 3b,c show how the crack progressed at various strains during stretching. Conventional covalently cross-linked hydrogels are brittle and cannot sustain much elongation without catastrophic failure in a pure shear fracture experiment. In contrast, the notched crack in the NF10 hydrogel shown in Figure 3 did not propagate prior to failure, which occurred at a 476% strain. Instead the crack was blunted as a result of the reversible nature of the supramolecular nanodomains, see the video in the Supporting Information. This mechanism of energy dissipation is similar to that observed with filled elastomers39 and other tough hydrogels, such as nanocomposite hydrogels40 and supramolecular hydrogels.14,41 The thermoresponsive, microphase separated hydrogels reported by Guo et al.29,30 also exhibited crack blunting behavior. The mechanism of toughening and the crack blunting observed for the NFx hydrogels involves the reversibility of the hydrophobic bonds between FOSA moieties, which provide the fundamental cross-link in these gels. The hydrophobic bond can dissociate (i.e., break) under stress and dissipate the strain energy, similar to the mechanism for other tough supramolecular hydrogels.6,8,14−16,21,27,28,33,42,47−51 As a result, the crack grew by tearing in the direction of extension but did not grow in the direction of the notch (normal to the

stretching direction). For the notched NF10 sample in Figure 3, the critical strain at which the crack began to propagate was εc = 476%. All of the samples exhibited crack blunting below the LCST but not above the LCST. An important difference between the mechanism of energy dissipation in the NFx hydrogels compared with all other tough supramolecular hydrogels, except for Guo et al.’s gels,29,30 is that the ∼6 nm diameter nanodomains represent a separate phase that acts as a multifunctional cross-link. The energy dissipation mechanism for most supramolecular hydrogels has been described as a “disengagement”15 or “unzipping”14 of the physical bond, the key point being that the cross-link disappears (though it can reform after the stress is removed). However, that is an oversimplification when the bond that breaks is one of many within a nanophase, such as with the NFx hydrogels. In that case, breaking a single supramolecular bond does not remove the nanodomain crosslink, though the hydrophobic groups involved with the broken bond may be removed from the nanodomain, i.e., they are pulled out by the retractive forces of the network chains. Since each FOSA group is attached to two network chains, pulling a FOSA group from a nanodomain lowers the cross-link density accordingly. Alternatively, it is not necessary that FOSA groups be removed from the nanodomains to dissipate the strain energy. The size of the nanodomains in the NFx hydrogels can accommodate ∼100 FOSA groups, so one can also envision energy dissipation by rearrangements of the FOSA groups participating in hydrophobic bonds within the nanodomains. Unfortunately, it is not possible based on the experiments reported herein or any other mechanical or scattering experiment to unambiguously determine whether or how many FOSA groups are pulled from the nanodomains, though SANS and stress relaxation experiments on similar hydrogels43 (discussed in the Hysteresis Behavior and Self-Recovery section below) suggest that some removal of FOSA groups from the nanodomains does occur. A similar energy dissipation mechanism has been previously proposed for ionomers, another microphase separated supraD

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Macromolecules molecular polymer with nanodomain sizes comparable to those in the NFx hydrogels. For ionomers, toughening is attributed to ion-hopping, that is supramolecular ionic or dipolar bonds break but are reformed by hopping44 of the ions or ion-pairs from one nanodomain to another. The preference for the ionic groups to be associated and not present as isolated ions or ionpairs is thermodynamic, i.e., the necessity of maintaining electronic neutrality for the ions and the large enthalpic penalty that is accrued by locating an ion-pair in the nonpolar polymer matrix (the continuous phase). By analogy, the mechanism of toughening in the NFx hydrogels may involve a hydrophobe-hopping mechanism, whereby the hydrophobes involved in a hydrophobic bond separate under stress, but reform, hydrophobic bonds when the stress is removed. It is not necessary that the reformed hydrophobic bond involves both of the original hydrophobes, so that a hydrophobe can hop from one position to another either within a nanodomain or from one nanodomain to another. The question of whether intra- or interdomain hopping occurs is challenging and cannot be easily resolved without a technique that can track the temporal location of specific hydrophobes. Unfortunately, we are not aware of any such experiment technique. Figure 4 and Table 3 show the temperature dependence of εc and the fracture energy, Γ, for the four NFx hydrogels,

Table 3. Temperature Dependence of the Fracture Energy for the NFx Hydrogelsa

a

Unshaded boxes indicate T < TLCST. Lighter shaded boxes indicate T > TLCST. Darker shaded boxes indicate T ≈ TLCST.

The general theory of rubber elasticity includes a swelling term, i.e., the volume fraction of polymer, and the effective cross-link density (including covalent and supramolecular cross-links, trapped chain entanglements and the concentration of physical intermolecular interactions that contribute to the modulus), νe, can be calculated from eq 2,37

E jij 2 zy = jj1 − zzzνeRT ϕ21/3 j 3 f z{ (2) k Equation 2 assumes a Poisson ratio, ν = 0.5, which is a reasonable assumption for flexible cross-linked networks and it applies for covalent and supramolecular cross-links. The only caveat is that for a supramolecular network, the stress needed to deform the material is dependent on the integrity of the network, i.e., the physical bonds can break during deformation, which lowers the effect cross-link density. However, the modulus used in eq 2 was measured from the initial slope of the stress−strain curve (i.e., the extrapolation to zero strain), where it is assumed that the physical bonds are intact. Thus, the cross-link density data in Figure 5 apply strictly to the cross-link density of the hydrogel at the beginning of the tensile experiment. The FOSA concentration and temperature dependences of the effective cross-link density, νe, for the NFx hydrogels are shown in Figure 5. At any temperature, νe increased with increasing FOSA concentration, which is a consequence of the greater concentration of supramolecular hydrophobic bonds that provide the physical cross-links. For NF5 below the LCST, the cross-link density decreased with increasing temperature. That was expected, since increasing temperature weakens the hydrophobic bonds, which decreases the concentration of supramolecular bonds. Note also that the SR decreased as temperature increased, Table 1, but if the number of supramolecular cross-links did not change with temperature, decreasing swelling ratio would increase the cross-link density. For each of the hydrogels, νe increased above the LCST. The arrows in Figure 5 indicate the LCST for each hydrogel as measured from swelling experiments.23 That result is a consequence of the large decrease in swelling of the gel due to phase separation of water as well as weakening of the supramolecular bonds. G=

Figure 4. Temperature dependence of the critical strain for a pure shear fracture test for the NFx hydrogels. The open symbols and filled symbols denote measurements below TLCST and above TLCST, respectively. The two colored red/black symbol for the NF8 sample signifies that T ≈ TLCST.

respectively. In general εc and Γ increased with increasing [FOSA] and decreased with increasing temperature, but as with the SR data described above the number of temperatures used in this study was insufficient for clearly resolving any abrupt change in either property at TLCST. At 5 °C, which was below TLCST for each of the hydrogels, the fracture toughness of the NFx hydrogels was remarkably high, varying from ∼2000 to 8000 J/m2 as the [FOSA] increased from 5.4 to 14.2 mol %. These Γ-values exceed those reported for doublenetwork hydrogels (0.1−1 × 103 J/m2)45 and are comparable with the toughest hydrogels that have been reported in the literature, e.g., hybrid hydrogels with covalent and ionic crosslinks (1−16 J/m2),14,46−48 polyampholyte hydrogels (4−8 × 103 J/m2),49,50 ionically cross-linked Fe(III)-neutralized poly(acrylamide-co-Fe(III)-acrylate) hydrogels (∼1 × 103 J/m2).51 E

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Figure 7. Effect of recovery time between successive loading and unloading cycles (extension rate = 50 mm/min) on the mechanical hysteresis behavior of NF8 at 5 °C. The recovery time is the interval between the start of the cycle and the end of the previous cycle.

Figure 5. Temperature dependence of the effective cross-link density for the NFx hydrogels. The dashed lines are only intended to improve the clarity of the figure. Where no error bars are visible, they were smaller than the symbols. The arrows denote the LCST (from swelling measurements21) for each hydrogel.

εmax = 1200%)49 and hybrid hydrogels (∼1.0 × 106 J/m3 at εmax = 300%).14 Figure 6 shows that a large residual deformation remained immediately after a tensile loading and unloading cycle, but the strain completely recovered to zero soon after the load was removed. The reversible nature of the supramolecular crosslinks is responsible for the residual deformation of the sample after a loading and unloading cycle as well as for the recovery behavior. The large hysteresis is a consequence of dissociation (i.e., breaking) of the physical cross-links and the finite time necessary for the microstructure to recover. The kinetics of the microstructure recovery is demonstrated by Figure 7, which shows that although the strain recovered quickly after a loading and unloading cycle, the stress−strain behavior required a much longer time period to recover. Figure 7 indicates that if a rest period was used between successive loading and unloading cycles, the mechanical response partially recovered toward the original response of the hydrogel and the extent of the recovery increased with increasing rest time. For a rest period of 1 h, the recovery was 96%. Figure 7 shows that after a loading−unloading cycle, the modulus (the slope of the stress−strain curve as ε → 0) of the gel measured on the subsequent loading−unloading cycle

Hysteresis Behavior and Self-Recovery. As expected for a physically cross-linked networks, the NFx hydrogels exhibited pronounced hysteresis, as shown by the cyclic tensile data shown in Figures 6 and 7. Figure 6a shows the loading and unloading curves for a single NF10 sample at 5 °C for tensile experiments where the sample was deformed to four different maximum strains (100, 200, 300, and 400%). The area of the hysteresis loop between the loading and unloading curves is the energy dissipated per unit volume for the cycle. The hysteresis was defined as the ratio of the hysteresis energy loss and the energy required to stretch the sample, i.e., the area under the loading curve. The hysteresis was remarkably high, varying from ∼50% to ∼80% as the maximum strain changed from 100 to 400%, Figure 6b. Those values corresponded to energy losses of 0.201 to 2.13 MJ/m3. Also included in Figure 6b are the hysteresis and the energy loss data for NF8, which exhibit a similar trend. The hysteresis and energy loss increased with increasing [FOSA] or cross-link density of the physical network. The energy losses associated with the hysteresis of these hydrogels are comparable to the values reported for other supramolecular polyampholyte hydrogels (∼1.5 × 106 J/m3 for

Figure 6. Hysteresis of NF10 hydrogel at 5 °C: (a) tensile loading and unloading cycles (extension rate = 50 mm/min) for four different maximum strains indicated in legend (a fresh sample was used for each tensile loading and unloading cycle) and (b) percentage hysteresis and energy loss for the loading and unloading cycles shown in part a. Circles denote NF10 and triangles represent NF8. F

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Figure 8. Effect of five successive loading and unloading cycles to 200% strain on the hysteresis of NF10 at 5 °C. Stretching rate = 50 mm/min. Each cycle was started immediately following the end of unloading for the previous cycle: (a) hysteresis curves and (b) energy loss in each cycle normalized by the energy loss for the first cycle.

calcium alginate) hydrogels exhibit a small but finite permanent deformation due to the irreversible rupture of some of the chemical cross-links.14 The time-dependent recovery of the strain and stress in the NFx hydrogels, and perhaps for the polyampholyte hydrogels49, is a consequence of the finite relaxation times for the network chains to re-equilibrate and the microstructure to reform after the load is removed.43 The stress-softening phenomenon shown in Figures 7 and 8 is commonly referred to as the Mullins effect,54 where the mechanical response of a material depends on the prior loading history. The Mullins effect for the NF8 hydrogel is more clearly seen in Figure 9,

decreasd. According to eq 1, that is due to a reduction of νe. The water content was not remeasured prior to every loading− unloading cycle, but according to eq 1 a loss of water during the tensile experiment would produce an increase of νe. The lowering of νe is directly a consequence of breaking the hydrophobic bonds during loading and the finite recovery time required for the microstructure to heal. The microstructure healing process is demonstrated by the increase of modulus and νe in a subsequent loading−unloading cycle when the rest time between cycles increased. The healing of the microstructure of the NFx gels in not complete for any of the data in Figure 7. The best that was achieved was 96% for a rest period of 60 min. A time for full recovery of the NFx hydrogel following a loading−unloading cycle was estimated using the data from a previous study of small angle angle neutron scattering (SANS) and stress relaxation of a dimethyl acrylamide−FOSA hydrogels (DF10) with a similar FOSA composition and similar nanodomain microstructure24,43 In that case, following a step strain, full recovery of the stress in the network chains and the nanodomain spacing required ∼10 h. That prediction appears to be consistent with the hysteresis data in Figure 7. Figure 8a shows the stress−strain results for successive loading and unloading cycles for the NF10 hydrogel at 5 °C. The loading step of each cycle began immediately following the unloading step of the previous cycle. The hysteresis and energy loss of each cycle decreased from that of the previous cycle, though the most significant drop in those values occurred between the first and second loading and unloading cycle, Figure 8b. Each successive cycle also exhibited a larger residual strain, Figure 8a. The residual deformation of the hydrogel, however, could be recovered after any deformation if the hydrogel was allowed to re-equilibrate for about an hour, though, as discussed above, the recovery of the properties took much longer. The effect of loading/unloading cycles on the hysteresis of the NF8 hydrogel was similar, Figure 8b. Similar hysteresis behavior was also reported for polyampholyte hydrogels.49 Double network (DN) hydrogels show similar hysteresis behavior upon successive loading/unloading cycling if the maximum strain is kept below the yield point, but for higher strains a permanent deformation occurs (i.e., plastic flow) due to the irreversible breaking of the first, more brittle network.52,53 Hybrid hydrogels that contain covalent and physical cross-links, such as cross-linked poly(acrylamide-co-

Figure 9. Demonstration of the Mullins softening effect for NF8 hydrogels at 5 °C. The black dashed curve represents a single loading−unloading cycle where the hydrogel was stretched to 400% strain. The other curves represent two different experiments: (1) successive loading−unloading cycles to maximum strains of 100% (purple curves) and 300% (red curves) and (2) successive loading− unloading cycles to maximum strains of 200% (green curves) and 400% (blue curves). A fresh specimen was used for each of the three experiments. The extension rate was 50 mm/min.

which shows two separate successive loading/unloading cycle experiments: (1) one where the hydrogel was first stretched to 100% strain and unloaded (purple curves) and then immediately restretched to 300% strain and unloaded (red curves) and (2) one where the hydrogel was stretched to 200% strain and unloaded (green curves) and then immediately restretched to 400% strain and unloaded (blue curves). The G

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Macromolecules

toughness is due to physical cross-links formed by nanodomain aggregates of hydrophobically associated FOSA moieties, which can dissipate substantial energy by either deforming or reversibly releasing FOSA groups and their associated network chain from the cross-link. The extraordinary toughness is demonstrated by the ability of the hydrogel to blunt cracks, preventing crack propagation such that the failure mechanism is tearing of the hydrogel perpendicular to the crack orientation. The lack, however, of any permanent set from tensile deformation as high as 400% strain indicates that the microstructure is reversible, which would require any pulledout chains to return to their parent nanodomain. Although the deformed gels recover their shape and mechanical properties, the recovery requires a holding time of at least an hour after removal of the stress for the strain to recover and much longer for the mechanical properties to recover. The two separate time periods are believed to be related to the characteristic relaxation times for re-equilibration of the conformation of the network chains and of the nanodomain structure, respectively. In addition, the NFx hydrogels exhibit a Mullins stresssoftening effect, though in contrast with other materials that show a Mullins effect, the NFx hydrogels fully recover from a loading−unloading cycle without any permanent deformation.

dashed black curves in Figure 9 are the loading and unloading data for a fresh NF8 hydrogel sample stretched to 400% strain. For each two successive cycle experiments, the stress required to restretch the specimen to either 300% or 400% strain after a prior loading and unloading cycle to 100% or 200% strain, respectively, was considerably lower than for the initial loading and unloading cycle. In both experiments, the loading traces of each original cycle and the section of the second cycle where the strain was greater than the maximum in first cycle followed the path of the single 400% loading cycle, i.e., the black dashed curve. However, for the part of the second cycle for which the strains were the same as in the first cycle, the stresses were considerably lower than the dashed loading curve. That is, subsequent loading and unloading cycles are softer over the strain range corresponding to the previous stress−strain history, but for strains greater than the maximum strain previously achieved, the sample deforms exactly like a fresh sample and is unaffected by its past mechanical history. The Mullins effect is most commonly observed for filled elastomers and neat elastomers that can strain-crystallize,55 but it has also been reported for unfilled polyampholyte hydrogels,49 microphase-separated thermoplastic elastomers,56 DN hydrogels,52 and clay,57 graphene,58 and silica59 nanocomposite hydrogels. Various explanations have been proposed for the origin of the Mullins softening effect, including detachment of chains from filler particles, slippage of chains bound to the reinforcing filler, rupture of filler clusters, chain disentanglement (which requires either large scale movement of the particles to which the chains are bound or breakage and reformation of covalent or physical bonds) and changes of a “double layer” microstructure composed of carbon black aggregates surrounded by polymer layers.55 With the notable exceptions of the supramolecular NFx and polyampholyte hydrogels, one thing common with all the other materials that exhibit a Mullins effect is that the Mullins effect coincides with a permanent deformation of the material. All of the mechanisms of stress softening delineated above have in common either the breaking or jumping (i.e., breaking and reforming) of a reversible bond or a microstructure change. For the supramolecular hydrogels that mechanism could be the reversible breaking of the supramolecular bonds or orientation of nanodomains. Ascertaining the mechanism of the Mullins effect for supramolecular hydrogels was beyond the scope of this paper, but the ability of the supramolecular hydrogels to fully recover the deformation, though kinetically dependent, indicates that either the reformation of the surpramolecular bonds occurs spatially at the same position in the gel where it was originally or the deformation of the microstructure is elastic. In either case, this reversibility represents a notable advantage of these materials.

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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/acs.macromol.8b00490.

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Calculation of cross-link densities shown in Figure 6, based on a phantom network model using the dynamic modulus data from linear viscoelastic dynamic mechanical shear measurements (PDF) Video of tensile test of notched NF10 hydrogel (AVI)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Fei Wang: 0000-0001-5344-8931 R. A. Weiss: 0000-0002-5700-6871 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was funded by a grant from the Civil, Mechanical and Manufacturing Innovation (CMMI) Division within the Engineering Directorate of the National Science Foundation, Grant CMMI-1300212.

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CONCLUSIONS Supramolecular hydrogels (NFx) made from linear random copolymers of a fluoroacrylate (FOSA) and NIPAM exhibit thermoresponsive behavior and extraordinary fracture toughness. The thermos-responsiveness occurs due to a lower critical solution temperature (TLCST) at which the gel goes through a volume phase transition, and TLCST can be tuned to between ∼5 and 32 °C by varying the FOSA concentration in the copolymer from 0 to 14 mol %. The NFx hydrogel are mechanically strong and stretchable and have fracture toughness as high as 8000 J/m2, which is comparable to the toughest hydrogels that have been reported. The substantial

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DOI: 10.1021/acs.macromol.8b00490 Macromolecules XXXX, XXX, XXX−XXX