Viscoelasticity and Structures in Chemically and Physically Dual

Article Cite This: Macromolecules XXXX, XXX, XXX-XXX

pubs.acs.org/Macromolecules

Viscoelasticity and Structures in Chemically and Physically DualCross-Linked Hydrogels: Insights from Rheology and Proton Multiple-Quantum NMR Spectroscopy Xueting Zou,† Xing Kui,‡ Rongchun Zhang,*,§ Yue Zhang,‡ Xiaoliang Wang,‡ Qiang Wu,† Tiehong Chen,∥,⊥ and Pingchuan Sun*,†,§,⊥ †

Key Laboratory of Functional Polymer Materials of Ministry of Education and College of Chemistry, Nankai University, Tianjin 300071, P. R. China ‡ Department of Polymer Science and Engineering, Nanjing University, Nanjing 210093, China § State Key Laboratory of Medicinal Chemical Biology, Nankai University, Tianjin 300071, P. R. China ∥ Institute of New Catalytic Materials Science, School of Materials Science and Engineering, Key Laboratory of Advanced Energy Materials Chemistry (MOE), Nankai University, Tianjin 300350, P. R. China ⊥ Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), Tianjin 300071, P. R. China S Supporting Information *

ABSTRACT: Hydrogels have received considerable attention as an innovative material due to their widespread applications in various fields. As a soft and wet material, its mechanical behavior is best understood in terms of the viscoelastic response to the periodic deformation, which is closely related to the microscopic chemically/physically cross-linked structures. Herein, a dualcross-linked (DC) hydrogel, where a physically cross-linked network by ionic coordination (Fe3+) is imposed on a chemically cross-linked poly(acrylamideco-acrylic acid) network, was studied in detail by rheology and proton multiplequantum (MQ) NMR spectroscopy. Rheology experiments revealed the diverse temperature- and strain-frequency-dependent viscoelastic behaviors for DC hydrogels induced by the dynamic Fe3+ coordination interactions, in contrast to the single chemically cross-linked (SC) hydrogels. During the shear experiment, the trivalent Fe3+ complex with moderate/weak binding strength might transform to those with strong binding strength and serve as permanent-like cross-linkages to resist the periodic deformation when a large strain frequency was applied. The viscoelastic behaviors of the DC hydrogels were strongly affected by the monomer ratio (CAAc/ CAAm) and Fe3+ concentrations; however, the chemically cross-linked density did not change with CAAc/CAAm, while the physically cross-linked density was greatly enhanced with increasing Fe3+ concentrations. Besides, the DC hydrogels have less contents of network defects in comparison to the SC hydrogels. The heterogeneous structural evolution with increasing the Fe3+ concentration and monomer ratio was also quantitatively determined and elucidated by proton MQ NMR spectroscopy. In addition, the moduli (G′, G″) of DC hydrogels were almost an order magnitude higher than that of the corresponding SC hydrogels, demonstrating the significant contribution of Fe3+ coordination to the mechanical properties, in consistent with the high activation energy of viscoelasticity for the physically cross-linked network as obtained from the variable-temperature shear rheology experiments. The experimental findings obtained from the rheology and proton MQ NMR experiments can be correlated with and complementary to each other. Herein, a combination of rheology and proton solid-state NMR is well demonstrated as an effective and unique way for establishing the relationship between microscopic structures and macroscopic viscoelastic properties. gels,15,16 macromolecular microsphere composite (MMC) hydrogels,17,18 polyampholyte hydrogels,19,20 hydrophobically modified hydrogels,21,22 and so on. Particularly, the concepts of sacrificial bonds, which were initially proposed to explain the molecular origin of the toughness of natural adhesives, fibers, and composites23 as well as bones,24 have been widely applied

1. INTRODUCTION Over the past decades, polymeric hydrogels have been the focus of intensive investigations due to their widespread applications in the biomedical, pharmaceutical, and industrial fields.1−8 However, since the hydrogels are intrinsically soft and wet, the poor mechanical properties often hinder their extensive use in many different applications. Therefore, multiple strategies have been proposed and developed in recent years for enhancing the mechanical performance, such as double-network (DN) hydrogels,9−11 nanocomposite hydrogels,12−14 slide-ring hydro© XXXX American Chemical Society

Received: August 28, 2017 Revised: November 13, 2017

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DOI: 10.1021/acs.macromol.7b01854 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Scheme 1. Procedures for the Preparation of Dual-Cross-Linked Hydrogels via Free-Radical Copolymerizationa

a

There are mono-, bi-, and trivalent Fe3+ coordination when there is excess Fe3+ loading, whereas soaking the hydrogel in excess deionized water will allow the reorganization of polymer chains to form a trivalent Fe3+ complex and to remove extra Fe3+ ions. The mono-, bi-, and trivalent Fe3+ complexes were indicated by transparent gold, blue, and purple rectangles, respectively.

with desirable structural properties such as fast self-healing, recovery, and dynamic exchange ability. Furthermore, the hierarchical hydrogel mechanics can be controlled by evenly mixing multiple kinetically distinct metal−ligand cross-links, where the strength of metal coordination complex can be measured by dynamic oscillatory rheology experiments adapting the Maxwell model assumption.35,49,50 Although substantial efforts have been devoted to improve the mechanical properties of hydrogels, such as mechanical strength, toughness, self-recoverability, stretchability, and so on, their viscoelastic properties and the corresponding structural origin were rarely investigated, which are the main focus of this study. For the mechanical enhanced hydrogel composed of both chemically (permanent) and physically (transient) cross-linked networks, the mechanical properties not only are determined by the network architectures but also depend on the deformation conditions, such as temperature, strain rate, and so on.51−54 Therefore, understanding the mechanical behaviors in terms of microscopic cross-linking structures and viscoelastic responses to periodic deformation could provide piercing insights into the fabrication of high performance hydrogels. For example, the toughness and/or strength of the DC hydrogels are codominated by the physical and chemical cross-linking density. Besides, if the strain rate is much smaller than the dissociation rate of the transient bonds, the hydrogel deformation will give rise to effective dissipation of the extern energy and thus the increase of toughness and modulus. In contrast, the hydrogel will behave as a loosely chemically crosslinked network if the strain rate is much higher than the dissociation rate.49 Yet, as far as we know, there are few studies on understanding the viscoelastic contributions of both chemically and physically cross-linked networks in the DC hydrogels.52,54 In addition, elucidating the viscoelastic response to the external deformation could shed a new light on the strength of transient Fe3+ coordination network as well as the organization of cross-linking structures. This study is aimed to

nowadays for the fabrication of a variety of high performance polymer materials25−30 including hydrogels.31−33 Bioinspired metal coordination interaction, which is often used to control and tune the mechanical properties of polymers as sacrificial bonds,34−38 is also widely incorporated into the physical crosslinkages for preparing high performance hydrogels. These sacrificial coordination bonds can dissociate and reassociate rapidly and reversibly and thus lead to efficient energy dissipation and recovery of the mechanical properties.39−41 As a matter of fact, a variety of cation ions have been used to enhance the mechanical performance of triblock polymer based hydrogels with a fracture stress up to 1 MPa.31 Suo and coworkers33 reported an extremely stretchable and tough DN hydrogel composed of a covalently cross-linked polyacrylamide network and a Ca2+ cross-linked alginate network. Such ultratough gels could be stretched beyond 20 times of the initial length.33 Moreover, they also demonstrated that hydrogels cross-linked by trivalent cations have much stronger mechanical strength than those cross-linked by monovalent/ divalent cations.42 Similarly, Zhou et al.43,44 prepared a dualcross-linked (DC) hydrogel that incorporated physically multivalent ion pairing of Fe3+−acrylic acid coordination into a covalently cross-linked poly(acrylamide-co-acrylic acid) (i.e., poly(AAm-co-AAc)) network. Thus, such hydrogel has ultrahigh mechanical strength, toughness, and good self-recovery property. The dynamic reversible Fe3+−acrylic acid coordination was also incorporated into a core−shell microgel covalently cross-linked network to significantly improve the toughness and recoverability of composite hydrogels45 and into a double-network hydrogel to improve the mechanical strength and fatigue resistance.46 Besides, completely physically crosslinked hydrogels with superior toughness and mechanical strength were also reported with the incorporation of Fe3+ coordination bonds.40,47,48 All these experimental findings well demonstrated that the sacrificial metal coordination bonds could not only impart desirable mechanical performance such as high strength and toughness but also provide the hydrogels B


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not specified. The above-obtained SC hydrogels were then immersed in the FeCl3 solution (with a weight ratio of 5:1, mferric solution/mgel) at room temperature for 16 h and thus to obtain dual-cross-linked hydrogels. Since the Fe3+ ions were excess, Fe3+ would form a mixture of mono-, bi-, or trivalent complex with carboxylic acid groups. Finally, the hydrogels were soaked in excess deionized water (with a weight ratio of 5:1, mwater/mgel) for 48 h, which allowed the reorganization of polymer chains to form trivalent complex between Fe3+ ions and carboxylic acid groups and to remove extra Fe3+ ions. The finally obtained DC hydrogels are referred as DX (e.g., D10, D20, ...), where D indicates dual-cross-linked network and X refers to the molar ratio of AAc to AAm in units of percentages. In order to understand the effect of Fe3+ concentration on the viscoelastic properties and structures of hydrogels, S20 hydrogel was immersed into a FeCl3 solution with variable Fe3+ concentration (0.03, 0.06, 0.12, and 0.18 mol/L) to prepare the corresponding D20 hydrogels. For the samples prepared for proton solid-state NMR experiments, heavy water (99.9% deuterium) was used in all the above procedures; otherwise, the polymer signals would be overwhelmed by the water signals in NMR experiments. Rheological Measurements. Dynamic shear rheology experiments were performed with a HAAKE Rheo-Stress 600 instrument, using a set of 20 mm diameter parallel plates. All the hydrogel samples were coated with vacuum grease on the edges to keep the sample hydrated and prevent desiccation. In all the dynamic shear experiments, the strain deformation was fixed at 0.5%, which was small enough to avoid the nonlinear response and large enough to have a reasonable signal intensity. The temperature dependence of the dynamic storage (G′) and loss (G″) modulus was measured at a fixed strain frequency of 6.28 rad/s and a fixed strain amplitude of 0.5% (all well within the linear viscoelastic range) in a temperature range of 25− 80 °C with a heating rate of 2 °C/min. Solid-State 1H NMR Experiments. All the solid-state NMR measurements were performed on a Bruker Minispec mq20 spectrometer at a 20 MHz proton resonance frequency, with a typical π/2 pulse length of about 3.2 μs and a receiver dead time of about 13 μs. All the experiments were measured at room temperature. T1 Measurements. The spin−lattice relaxation time (T1) was measured by a saturation recovery experiment66 with a variable relaxation delay. Since generally water and polymers have very different spin−lattice relaxation time, the obtained intensity buildup curve as a function of the variable delay could be fitted with the following function:

address the above challenges and to establish a detailed relationship between microscopic structures and macroscopic viscoelastic properties by utilizing a combination of rheology and solid-state NMR techniques. Solid-state 1H NMR spectroscopy is a powerful tool for elucidating structures and dynamics in polymer science,55−57 especially with the emerging rapid advances in high-resolution proton-detected solid-state NMR methodologies under ultrafast magic-angle-spinning conditions.58−60 However, in recent years, the implementation of advanced single-channel proton NMR experiments on a benchtop low-field NMR spectrometer has gained dramatic attention because it could provide quantitative information on the compositions, molecular dynamics, and network structures, albeit without chemical resolutions.61,62 In particular, proton multiple-quantum (MQ) NMR spectroscopy has been popular in gaining insights into the cross-linking structures and segmental constraints in gels.63−65 Herein, we focus on investigating a DC hydrogel, which is microscopically composed of a covalently cross-linked poly(AAm-co-AAc) network imposed with multivalent Fe3+− acrylic acid coordination, utilizing a combination of dynamic shear rheology experiments and advanced proton MQ NMR spectroscopy. A series of SC and DC hydrogels with different monomer ratios (CAAc/CAAm) and Fe3+ concentrations were prepared. The heterogeneous structural evolution with the monomer ratio and Fe3+ concentration and their influences on the linear viscoelastic behaviors were addressed in detail by proton MQ NMR spectroscopy and dynamic shear rheology experiments, respectively. The good agreement between the rheological results and proton NMR findings was well established and elucidated. To our knowledge, this is the first in-depth solid-state proton NMR and rheology study on a chemically and physically dual-cross-linked hydrogel.

2. EXPERIMENTAL SECTION Materials. Acrylamide (AAm) and acrylic acid (AAc) were both purchased from Alfa Aesar. Co. Ltd. (Tianjin, China), while anhydrous ironic(III) chloride (FeCl3) powder was provided by Meryer Chemical Technology Co. Ltd. (Shanghai, China). Potassium persulfate (KPS) and N,N,N′,N′-tetramethylethylenediamine (TEMED) were purchased from Sigma-Aldrich Chemicals (Shanghai, China) and Heowns Biochemical Technology Co. Ltd. (Tianjin, China), respectively. N,N′Methylenebis(acrylamide) (MBAA) were purchased from Sinopharm Chemical Reagent Co. Ltd. (Shanghai, China). All the chemicals and solvents were used as received without any further purification or treatment. Preparation of Dual-Cross-Linked (DC) Hydrogels. The hydrogels were synthesized by following a three-step procedure similar to the one reported in the literature,43 which was shown in Scheme 1. First, 10 mL of 3 mol/L AAm solution was prepared in deionized water, mixing with different molar ratios of AAc (10%, 15%, 20%, and 25% molar ratio of AAc/AAm) and 0.0216% (molar ratio with respect to the total content of AAc and AAm) chemical crosslinker MBAA. The above mixture was stirred at 20 °C for 30 min to form a homogeneous solution and then deaired with dry nitrogen gas for 10 min. Subsequently, the initiator KPS (1 wt %, with respect to the total monomer weights) and the catalyst TEMED (8.6 μL) were added into the above solution. After complete mixing, the temperature was raised to 30 °C. Then the solution was transferred into a glass mold, which was placed at 30 °C for 16 h to form singly chemically cross-linked (SC) hydrogel (referred as SX, where the S indicates singly cross-linked network and the X refers to the molar ratio of AAc to AAm in units of percentages.). Second, a FeCl3 solution was prepared by dissolving anhydrous FeCl3 powder into deionized water, which was further treated with ultrasound for 5 min to obtain homogeneous FeCl3 solution with a concentration of 0.06 mol/L if

y = A p(1 − e−t / T1p) + A w (1 − e−t / T1w )

(1)

where Ap and Aw are the relative content of polymers and water, respectively. T1p and T1w indicate the spin−lattice time of polymer and water, respectively. Herein, the fraction of polymers in the hydrogel could be determined as f p = Ap/(Ap + Aw). However, such a method is only applicable for the SC hydrogels. In the DC hydrogels, the paramagnetic effect of Fe3+ has significantly accelerated T1 relaxation, rendering it difficult to distinguish the polymer and solvent signals. As a result, the fraction of polymer for the DC hydrogels could not be accurately determined by T1 experiments (Figure S1). Proton MQ NMR Experiment. 1H MQ NMR spectroscopy is one of the most versatile and robust techniques for providing quantitative information on the structures and dynamics of polymer network and melts.62 A complete detailed account of the principle of MQ NMR as applied to polymers and soft materials can be found in refs 67 and 68. Generally speaking, the isotropic segmental motions in the solution will average out the 1H−1H dipolar couplings, resulting in a highresolution site-specific 1H spectrum as typically obtained in the highfield solution NMR spectrometer. However, in some soft materials like hydrogels, due to the presence of constraints (e.g., cross-linking, entanglements), the segmental motions become anisotropic, leading to a residual dipolar coupling (Dres). Such unaveraged dipolar coupling interaction is typical for solid-state NMR spectra but often broadens the spectral peaks and leads to a complete loss of chemical resolutions in solution NMR spectra. As a matter of fact, proton MQ NMR mainly C


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Macromolecules utilizes Dres as a probe to extract quantitative information about the constraints (and thus structure and dynamics) of the hydrogel network. The Baum−Pines pulse sequence69 is adapted here to excite the MQ coherent signals due to its capability of producing pure double-quantum (DQ) Hamiltonian and the relatively high mobility of polymer segments in the hydrogels, while the three-pulse segment (90°x-τ-180°y-τ-90°x) has to be utilized for the molecular systems with rigid/immobile components.70,71 Typically, from the MQ NMR experiments, two data set could be obtained, IDQ and Iref, corresponding to the DQ and reference signal intensity, respectively. The sum of these two signals forms the sum MQ intensity (I∑MQ = Iref + IDQ), whose decay as a function of DQ evolution/excitation time (τDQ) is dominated by segmental fluctuations. Since the DQ intensity is subject to a severe relaxation effect at longer DQ evolution time, a point-by-point normalization protocol is required to obtain a normalized DQ intensity buildup curve:

InDQ (τDQ ) =

Figure 1. Storage modulus (G′, solid) and loss modulus (G″, open) of the S25 and D25 samples as a function of temperature with a heating rate of 2 °C/min. The measurements were performed with a strain rate (w) of 6.28 rad/s and a strain amplitude (γ) of 0.5%.

IDQ (τDQ ) I∑ MQ (τDQ )

(2)

Such a normalization procedure could eliminate the effect of temperature-dependent segmental dynamics (which generally induces significant changes of T1 and/or T2 relaxation), and thus the normalized DQ intensity buildup is exclusively dependent on Dres characterizing restrained structures of cross-linking networks. However, before the normalization, the signals of isotropic mobile components that do not contribute to the MQ coherent signals, such as dangling/sol chains and loops, have to be subtracted from the sum MQ intensity:

I∑ MQ = IDQ + Iref − fB e−2τDQ / T2B − fC e−2τDQ / T2C

the storage modulus G′ slightly decreases over the temperature range up to 80 °C, indicating a quite stable network structure. In addition, increasing the temperature is also expected to weaken the dynamical metal coordination bonds, which might slightly decrease the physical cross-linking density, leading to an increase in the swelling ratio. Indeed, the swelling capacity was enhanced at an elevated temperature as shown in Figure S2, where both S25 and D25 hydrogels have a higher swelling ratio at 50 °C than that at 25 °C. According to the affine network theory,73 G′ is proportional to the temperature and v2/3, where v is the volume fraction of polymers in the hydrogel. Therefore, the decrease of G′ can be ascribed to the increase of swelling ratio and thus the decrease of v. On the contrast, the G″ of two samples exhibits completely diverse temperature-dependent behaviors. For the S25 hydrogel, G″ slightly decreases with increasing the temperature like G′, whereas for the D25 hydrogel, G″ increases significantly by almost an order of magnitude with raising the temperature from 25 to 80 °C, indicating a substantial increase of viscous responses. In fact, the increase of G″ can be attributed to the fast dissociation and reassociation of the dynamic Fe3+ coordination bonds, leading to efficient energy dissipation without inducing significant changes of network structures, as indicated by the stable G′ with increasing the temperature. It is also worth noting that the introduction of Fe3+ coordination interactions has increased G′ by more than 5 times and G″ by more than 2 times in D25 hydrogel. Therefore, the introduction of the Fe3+ coordination network not only significantly increases the viscoelastic moduli due to the enhancement of segmental constraints but also dramatically improves the energy dissipation efficiency with increasing the temperature, leading to an enhanced toughness and thus excellent elongation.43 Frequency Dependence of the Linear Viscoelasticity. The frequency dependence of moduli (G′ and G″) for the DC hydrogels with variable CAAc/CAAm is shown in Figure 2A. For all the hydrogels, G′ is much larger than G″, indicating a solidlike and elastic nature of the hydrogels. In particular, with increasing the concentration of AAc monomers, both G′ and G″ are enhanced. On one hand, with increasing CAAc/CAAm, there are more carboxylic acid groups in the hydrogels, leading to an enhanced physical cross-linking density due to the Fe3+− acrylic acid coordination interactions. In addition, the capability of energy dissipation will be boosted up due to the increasing content of Fe3+ coordination complex, leading to the increase of G″. On the other hand, it is worth noting that the polymer mass

(3)

f B and f C are the fractions of potential two separable isotropic mobile components with an apparent spin−spin relaxation time of T2B and T2C, respectively. More examples of the application of proton MQ NMR spectroscopy could be found in a very recent review.62

3. RESULTS AND DISCUSSION This dual-cross-linked hydrogel contains both covalent crosslinking and trivalent Fe3+−acrylic acid coordination, leading to a superior mechanical performance with ultrahigh strength, toughness, and good self-recovery.43 Herein, rheological investigation on its viscoelastic behaviors in response to the periodic deformation can shed additional new light on the cross-linking structure as well as energy dissipation induced by the transient Fe3+ coordination network, which can also be quantitatively addressed in terms of cross-linking density and the distribution of segmental constraints as revealed by proton MQ NMR spectroscopy. Temperature Dependence of the Linear Viscoelasticity. For the associating polymers with dynamic metal coordination, the bulk mechanical properties are determined by the dissociation rate constant, regardless of the binding thermodynamics, yielding the widely applied principle “strong means slow”, which means that slow dissociation generally leads to stronger materials.50,72 Since the dissociation rate of Fe3+−acrylic acid coordination can be altered by the temperature, the temperature dependence of the linear viscoelastic behaviors of the DC hydrogel could provide qualitative information about the strength of the Fe3+ coordination as well as the stability of the hydrogel. Figure 1 shows the comparison of moduli (G′, G″) as a function of the temperature between S25 and D25 samples under a fixed strain frequency (ω) of 6.28 rad/s and a fixed strain amplitude (γ) of 0.5% (all well within the linear viscoelastic response range). The other samples show similar temperature-dependent behaviors, and thus data are not shown here. For both samples, D


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Figure 2. (A) Viscoelastic moduli (G′, solid; G″, open) and (B) complex viscosity (η*) as a function of the strain frequency for the DC hydrogels with a fixed CAAm (3 mol/L) but varying CAAc with a CAAc/CAAm of 10%, 15%, 20%, and 25%.

Figure 3. Tan δ as a function of the strain frequency for the DC (A) and SC (B) hydrogels with a fixed CAAm (3 mol/L) but varying CAAc with a CAAc/CAAm of 10%, 15%, 20%, and 25%.

fraction increases from ∼20% to ∼31% when CAAc/CAAm is increased from 10% to 25%, as shown in Figure S3. Actually, in the DC hydrogels, with an increasing concentration of AAc, the coordination interactions between Fe3+ ions and the carboxylic acid groups on polymer side chains may cause the contraction of polymer segments and thus squeeze the water out of the network. Therefore, the increase of both the physical crosslinking density and polymer content in the DC hydrogels leads to enhancement of the viscoelastic moduli. However, the influence of polymer content on the mechanical behaviors of hydrogels was rarely mentioned when the hydrogel composition was changed, e.g., in ref 43. The solidlike mechanical response of the DC hydrogels could also be well demonstrated by the frequency-dependent behaviors of the complex viscosity (η*, Pa·s) and loss factor (tan δ = G″/G′) as shown in Figures 2B and 3, respectively. For a perfectly covalent cross-linking network, a linear relationship between log η* and log w with a slope of −1 should be achieved.74 For all the DC hydrogels with different CAAc/CAAm, a slope of ∼−0.97 was obtained for all the samples, indicating a near-perfect elastic network over the probed frequency range as well as the presence of network defects. Similar results were also observed for the SC hydrogels (Figure S4). Since the DC hydrogels were fully swollen in the deionized water, there should not be any sol chains. Thus, the defects in the DC polymer network should be loops or dangling chains that do not involve forming cross-linking network. The loss factor (tan δ = G″/G′) as a function of the strain frequency for the DC and SC hydrogels can be obtained from Figure 2 and Figure S4, respectively. For the DC hydrogel, it is interesting to note that tan δ decreases with increasing the

strain frequency and reaches a plateau at a critical strain frequency around 10 rad/s, as shown in Figure 3A. At a small strain frequency, the external energy could be dissipated by the dynamic dissociation and reassociation of Fe3+ coordination bonds, where the relaxation rate of the Fe3+ coordination network is comparable to or faster than that of the exerted strain frequency. When the strain frequency is large enough, the relaxation of Fe3+ coordination network is not fast enough to dissipate the stress energy instantly, and thus a plateau of tan δ is reached. Such a plateau was usually ascribed to the chemically covalent network. However, for the SC hydrogel where there is only a chemically covalent network, it is found that tan δ significantly increases with the strain frequency when the strain frequency is beyond a critical value at around 10 rad/s, indicating that the SC hydrogel network has been significantly deformed and does not recover instantly. That means, in the DC hydrogels, the plateau of tan δ cannot be simply attributed to the chemically cross-linking network. Therefore, it is speculated that in the DC hydrogels the Fe3+ coordination complex with moderate/weak binding strength may transform to those with strong binding strength during the shear experiments and may serve as permanent-like cross-linkages to resist the deformation under the large strain frequency. This speculation is supported by the repeated frequency-dependent shear experiments, where the storage modulus was enhanced with increasing the strain frequency sweep cycles, as shown in Figure S5. A similar conclusion was also obtained for a physical Fe3+-coordinated poly(acrylamide-co-acrylic acid) gels.40 As a result, a plateau of tan δ can be observed at the high strain frequency beyond 10 rad/s for the DC hydrogels. Besides, because of the conformation transition of the trivalent Fe3+ E


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Figure 4. Dynamic shear moduli (G′, solid; G″, open) of the D25 (A) and S25 (B) hydrogels as a function of frequency in a temperature range of 25−80 °C. (C) Master curves of both hydrogels were obtained by shifting the frequency sweep curves at different temperatures. (D) Arrhenius plot of the horizontal shift factor aT with a reference temperature of 25 °C.

complex, tan δ decreases with increasing the strain frequency when the strain frequency is below 10 rad/s. For the SC hydrogels, the covalent network could well resist the small deformations, and thus a plateau of tan δ is observed when the strain frequency is small, as is further shown by the almost constant moduli (G′ and G″) (Figure S4). However, at a large strain frequency, the covalent network cannot resist the periodic deformation, and thus the viscous response is taking a dominating role, since G″ increases faster than G′ with increasing the strain frequency (Figure S4). Therefore, tan δ increases with further increasing the strain frequency. Overall, the reversible Fe3+ coordination interaction is well demonstrated to be an effective energy dissipation route in the DC hydrogels, where during the deformation process the Fe3+ coordination complex with moderate/weak binding strength may transform to those with strong binding strength and serve as permanent-like cross-linkages to resist the deformation. In order to understand the Fe3+ coordination network strength and the capability of energy dissipation, frequencydependent rheological experiments were performed in a range of temperatures T = 25−80 °C, as shown in Figure 4. For the DC hydrogels, the temperature variation result in obvious changes in the viscoelastic moduli due to the temperaturedependent relaxation caused by the reversible Fe3+ coordination interactions. For the SC hydrogel, both G′ and G″ decrease slightly with increasing the temperature, which is ascribed to the differences in swelling ratio at different temperatures as explained above. It is worth noting that the viscoelastic moduli (G′, G″) follow the time−temperature superposition principle quite well and collapse onto a master curve by applying horizontal and vertical shift factors on the moduli curves at different temperatures, as shown in Figure 4C. Because temperature induces obvious changes on G′ and G″ of the DC hydrogel, the frequency range on the master curve is clearly

expanded compared to the individual curves at a specific temperature. However, there are much smaller changes on the frequency range for the master curve of S25, where the expanded small frequency range may be ascribed to the fast relaxation of network defects. The horizontal shift factors (aT) as a function of temperature are shown in Figure 4D for both D25 and S25 samples, which can be fitted using the Arrhenius equations ln(aT ) =

Ea ⎛ 1 1 ⎞ ⎜ − ⎟ R ⎝T Tref ⎠

(4)

where Ea is the activation energy and R is the gas constant. With Tref = 298.15 K, the activation energy of viscoelasticity was determined to be 15.3 and 60.8 kJ/mol for the S25 and D25 hydrogel, respectively. Assuming that the Fe3+ coordination does not have any obvious impacts on the covalent crosslinking structures, the contribution of the physical network of Fe3+ coordination to the activation energy of viscoelasticity will be around 60.8 − 15.3 = 45.5 kJ/mol. Since the probed temperature range is far away from the glass transition temperature of the polymer (Figure S6), the activation energy of viscoelasticity for the SC hydrogels should be attributed to the enhanced mobility of the network defects, such as the side chains (mainly carboxylic acid groups), terminal groups, etc. However, in the DC hydrogels, the content of side chains will be greatly reduced due to the Fe3+−acrylic acid coordination. Thus, the activation energy for the physical network is actually significantly underestimated. Still, this value is comparable to the free energy required to break the dimerization of 2-ureido4[1H]pyrimidinone (UPy)-based quadruple H-bonding groups (∼50 kJ/mol in toluene),75 which is often utilized to enhance the mechanical strength and toughness in polymer materials28,76 and hydrogels.77,78 Herein, the enhanced mechanical properties for the DC hydrogels are dominated by the strong F


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Macromolecules Fe3+ coordination network, since the activation energy of viscoelasticity for the physical Fe3+ coordination network is much larger than that of chemically covalent network. It is also expected that the activation energy for the physical network can be changed by varying the CAAc/CAAm or Fe3+ concentrations, since the strength of Fe3+ coordination network is closely related with the physical cross-linking density, as will be discussed below. Cross-Linking Structures Determined by Proton Multiple-Quantum NMR Spectroscopy. Proton MQ spectroscopy has been well established and proved to be a robust method for probing the heterogeneous structures and dynamics of polymeric materials.62 In particular, proton MQ NMR could provide quantitative insights into the anisotropic structures and related segmental mobility. Herein, proton MQ NMR is further utilized to investigate the network structures of SC and DC hydrogels as shown below. An illustrative example (S20 hydrogel) for obtaining normalized nDQ buildup curve is shown in Figure 5. The Iref −IDQ curve was used to fit the tail,

the tail in the MQ NMR experiments of DC hydrogels, it can be obtained DC DC Adefects + A solvent

DC ftail =

DC DC A polymer + A solvent

SC ≈ f defects =

SC Adefects SC SC A polymer + A solvent

(5)

where f indicates the polymer/solvent/tail fraction and A represents the corresponding polymer/solvent content. Therefore, it can be concluded that the fraction of polymer defects (with respect to all proton signals) in the DC hydrogels are much less compared to that of SC hydrogels, i.e. DC fdefects =

DC Adefects DC DC A polymer + A solvent

SC < f defects =

SC Adefects SC SC A polymer + A solvent

(6)

In addition, water content in the DC hydrogels was dramatically reduced, since the Fe3+−acrylic acid coordination interactions would cause the contractions of polymer segments and thus squeeze the water out of the polymer network. Thus SC DC A solvent < A solvent

(7)

Besides, there are a large number of carboxylic acid groups on the side chains (i.e., dangling chains) in the SC hydrogels, which do not involve in forming chemically cross-linking network. However, those groups will form coordination complex with Fe3+ ions and thus contribute to the physical cross-linking network in the DC hydrogels. Furthermore, the unreacted monomers, which are part of defects in SC hydrogels, would be removed when the SC hydrogels were soaked in FeCl3 solution for the preparation of DC hydrogels. Thus Figure 5. An illustrative example (S20 hydrogel) of obtaining normalized DQ buildup curves from the obtained IDQ and Iref data set. The Iref−IDQ curve was also plotted, which amplified the isotropic relaxation behavior and enabled easier identification of the contributions from network defects. The Itail curve was obtained by fitting the tail of the Iref−IDQ curve with eq 3. The blue solid line indicates the fitting result on nDQ buildup curve obtained from eq 14, while other solid lines are just used as the guidelines for the eyes.

DC SC SC SC Adefects = Adefects − A sol chains − Acarboxylic acid groups

(8)

DC SC SC A polymer = A polymer − A sol chains

(9)

As a result, the fraction of polymer defects with respect to the total polymer content in the DC hydrogel can be obtained as DC pdefects =

which enabled an easier identification and subtraction of the network defect contributions.79 In the MQ experiment for the SC hydrogel, a recycle delay of 0.2 s was used between consecutive scans. Since there is a significant contrast for the T1 of polymer and solvent (HDO) in S20 sample (Figure S1A), a short recycle delay of 0.2 s could actually well filter out the proton signals of solvents. Therefore, the tail signals in the MQ experiment actually completely indicate the contributions of polymer defects (loops and dangling/sol chains) in the network. In fact, for all the SC hydrogels, network defect signals take up around 80% of the polymer (Figure S7), which is around 48% with respect to the total proton signals since the polymer fraction is around 60% as determined by proton T1 experiments (Figures S1A and S3C). In contrast, the T1 of solvent in the DC hydrogels is quite short (Figure S1B). As a result, a recycle delay of 0.2 s is not able to filter out the solvent signals in the proton MQ NMR experiments, and thus the tail fraction directly indicates the fraction of polymer defects with respect to the total proton signals in the DC hydrogels, around 49% in average as shown in Figure S7. Interestingly, the network fraction in DC gels (∼49%) is nearly the same as that in SC hydrogels (∼48%). Since all water signals contribute to

DC Adefects DC A polymer

SC < pdefects =

=

SC SC SC Adefects − A sol chains − Acarboxylic acid groups SC SC A polymer − A sol chains

SC Adefects SC A polymer

(10)

The derivation of eq 10 can be found in the Supporting Information. Herein, the fraction of the polymer defects with respect to the polymer content is less in the DC hydrogels in comparison to the SC hydrogels. Besides, it is also worth noting that for the SC hydrogels around 80% of the polymer signals does not contribute the MQ coherent signals, indicating a very heterogeneous and low cross-linking density in the chemically covalent network. Both facts could also explain why the SC hydrogel has very low mechanical strength and viscoelastic moduli.43 In the cross-linked polymer network, Dres is correlated with the order parameter of constrained segments according to61 Sb = k

Dres 3r 2 = Dstat 5N

(11)

Here, Sb is the segmental order parameter, and Dstat is the proton dipolar coupling constant in the static limit (∼30 kHz). G


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Figure 6. Normalized DQ intensity as a function of DQ excitation time for the SC (A) and DC (C) hydrogels. The corresponding Dres distribution curves are shown on the right for SC (B) and DC (D) hydrogels. The solid lines indicate the fitting results obtained from eq 14 in panels A and C, whereas they are used as guides for the eyes in panels B and D. All measurements were performed at room temperature.

distribution, log-normal distribution,82,83 is adapted here due to its capability to characterize the Dres distribution width (σ) as well as the averaged Dres values in terms of the median value (Dm) as discussed below. The probability function for the lognormal distribution is

k is a scaling factor that takes into account the fast segmental dynamics within a single statistical (Kuhn) segment. r2 = r2/ ⟨r2⟩0, where r is the end-to-end vector of segments separating the constraints, while ⟨r2⟩0 is the unperturbed melt-state value of the vectors. N is the number of Kuhn segments between constraints. Herein, Dres is a direct measure of the length of network chains in terms of segmental end-to-end distances and number of segments N; thus, it is proportional to the density (inverse molecular weight) of network chains in units of mol/ kg. However, the proportionality factor is quite dependent on the properties of the investigated polymers, including localscale segmental dynamics, details of proton dipolar couplings in the monomer unit, and so on.80 For a synthetic polymer, it is impossible to determine such proportionality factor especially when the network is quite heterogeneous. But still the Dres distribution in the cross-linked hydrogel can well reveal the structural heterogeneity in terms of segmental lengths between cross-linking nodes in the sample. In order to obtain such quantitative information, numeric integration needs to be implemented to fit the normalized DQ (nDQ) buildup curves as a function of DQ excitation/evolution time (τDQ), as below. For a rather homogeneous network, the nDQ intensity as a function of DQ excitation time (τDQ) can be fitted with a generic empirical function:81

P(Dres) =

(13)

Here, σ is the standard deviation and dimensionless, reflecting the inhomogeneity of the distribution and roughly corresponding to the fwhm (full width at half-maximum) of the Dres distribution on a logarithmic scale (eq 15). Dm is the median value for the Dres distribution, indicating the residual dipolar couplings with the largest probability. The finalized nDQ intensity buildup can thus be calculated by a finite-step integral according to InDQ (τDQ ) =

InDQ (τDQ , Dres) = 0.5[1 − exp{− (0.378DresτDQ )1.5 } × cos(0.583DresτDQ )]

2 2 1 e−(ln(Dres / Dm)) /2σ 2π σDres

∫ P(Dres)IDQ (τDQ , Dres) dDres

=

∫

2 2 1 e−(ln(Dres / Dm)) /2σ IDQ (τDQ , Dres) dDres 2π σDres

=

∫

2 2 1 e−(ln(Dres / Dm)) /2σ IDQ (τDQ , Dres) d ln Dres 2π σ

=

∫ P(ln Dres)IDQ (τDQ , Dres) d ln Dres (14)

(12)

where

where Dres is the apparent residual dipolar coupling due to the anisotropic segmental dynamics imposed by the cross-linking/ entanglement structures. However, in most cases, heterogeneity is present in the system, either structural/conformational difference among polymer chains or dynamics gradient/ variation in different components. Herein, an asymmetric

P(ln Dres) =

1 −(ln(Dres / Dm))2 /2σ 2 e 2π σ

(15)

As is shown in Figures 6A and 6C, the nDQ buildup curves can all be fitted with eq 14 quite well. The obtained Dres distribution curves for the SC and DC hydrogels are shown H


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Macromolecules in Figures 6B and 6D, respectively, where the obtained Dm and σ values for all the hydrogels are summarized in Table 1. It is

concentration, leading to enhanced constraints of polymer segments. Interestingly, it is found that the standard deviation of the Dres distribution (σ) decreases with increasing CAAc/CAAm (Table 1), suggesting that the hydrogel network is becoming more homogeneous. In fact, since there are excess Fe3+ for the preparation of DC hydrogels, the polymer chains with abundant carboxylic acid groups could adjust the conformations accordingly to form tight trivalent Fe3+ complex. As a result, the polymer segments are more restricted and tend to form a homogeneous Fe3+ complex structure, leading to enhanced viscoelastic moduli (Figure 2). Such a process could be clearly demonstrated from the MQ experiments on D20 hydrogels with variable Fe3+concentrations, as shown in Figure 7. By changing the Fe3+ concentration of the ferric solution during sample preparation, it is possible to control the Fe3+ content in the DC hydrogels, where the corresponding rheological and DQ results are shown in Figure 7. As is clearly shown, with increasing the Fe3+ concentration, the nDQ intensity builds up faster. When the concentration is beyond 0.06 mol/L, further increasing the concentration does not significantly enhance the nDQ intensity buildup rate, indicating that an equilibrated physical network is almost achieved. A quantitative comparison is presented in Figure 7B,C. With increasing the Fe3+ concentration, Dm also increases obviously, whereas the increase of D m is very small when the concentration is beyond 0.06 mol/L, as shown in Figure 7C. Interestingly, the Dres distribution curve becomes broader on the logarithmic scale when the Fe3+ ions were incorporated into the hydrogels (Figure 7B), indicating an enhanced structural heterogeneity of the DC hydrogels compared to the SC hydrogels. To a large degree, this can be attributed to the partial formation of Fe3+ coordination complex, which imposes additional constraints on the polymer segments. However, the distribution width on the logarithmic scale decreases with further increasing the Fe3+ concentrations from 0.03 to 0.06

Table 1. Dm and σ Obtained from the Fitting on the Corresponding nDQ Buildup Curves of Hydrogelsa samples

Dm/2π (kHz)

σ

SC D10 D15 D20 D25

0.14 0.58 0.80 1.32 1.59

0.58 1.21 1.16 0.95 0.89

Dm and σ represent the median value and standard deviation of the Dres distribution, respectively.

a

interesting to note that the nDQ curves are basically the same for the SC hydrogels with different CAAc/CAAm (Figure 6A), indicating the same chemically cross-linking density. Such results are expected since AAc and AAm monomers have quite similar chemical structures, and the cross-linker (MBAA) has a constant content with respect to the total content of AAc and AAm monomers. The corresponding Dres distribution for the SC hydrogels is shown in Figure 6B, where a median value of 0.14 kHz and standard deviation of 0.58 are obtained, indicating a very heterogeneous structure (since σ/Dm ∼ 4) . Such heterogeneity can be ascribed to the presence of a large fraction of network defects, such as unreacted monomers, dangling chains, or loops. For the DC hydrogels, the nDQ intensity builds up faster with higher CAAc/CAAm (Figure 6C), and thus Dm increases with increasing CAAc/CAAm. It can also be seen from the Dres distribution curves in Figure 6D, where the distribution curves shifts to higher Dres values with increasing CAAc/CAAm. This is expected since there are more carboxylic acid groups to form coordination complex with Fe3+ ions with a higher AAc

Figure 7. Normalized DQ intensity (A) as a function of DQ excitation time for D20 hydrogels prepared with variable Fe3+ concentration. (B) The Dres distribution curves as obtained from the fitting on the normalized DQ curves for D20 hydrogels prepared with a variable Fe3+ concentration. (C) The median value (Dm) and standard deviation (σ) as a function of Fe3+ concentration for the D20 hydrogels. (D) Dynamic shear moduli (G′, solid; G″, open) as a function of the strain frequency for the D20 hydrogel with variable Fe3+ concentrations. The solid lines in (A) indicate the fitting results obtained from eq 14, while others are just guides for the eyes. All experiments were performed at room temperature. I


Article

Macromolecules mol/L and almost remains constant when the Fe 3+ concentration is beyond 0.06 mol/L. This is also obviously shown from the variation of the standard deviation (σ) of the Dres distribution when the Fe3+ concentration changes from 0 to 0.18 mol/L (Figure 7C). It clearly indicates that a high concentration of Fe3+ ions will benefit the formation of a homogeneous coordination network with large physical crosslinking density. The enhanced strength of the Fe3+ coordination network also corresponds well with the viscoelastic moduli variation with Fe3+ concentrations, as shown in Figure 7D. The viscoelastic moduli increase with increasing Fe3+ concentrations, where G″ changes more obvious than G′ because the reversible Fe3+ coordination is efficient in energy dissipation and thus the viscous responses are enhanced. It is worth noting that when the Fe3+ concentration is 0.18 mol/L, G″ changes little with increasing the strain frequency. That clearly indicates that all the trivalent Fe3+ complexes have strong binding strength and may serve as permanent-like cross-linkages to resist the deformation at the probed strain frequency range. In contrast, when the Fe3+ concentration is not high enough, the binding strength of trivalent Fe3+ complex may be weak or medium, even though the Fe3+ content is quite enough for all the carboxylic acid groups to form trivalent coordination with Fe3+. As a result, as shown (Figures 2A and 3) and explained above, during the shear experiment, the trivalent Fe 3+ coordination complex with weak/medium binding strength may transform to tight ones with strong binding strength. As a result, G″ will decrease and reach a plateau with increasing the strain frequency. When the Fe3+ concentration is beyond 0.06 mol/L, the hydrogel has a nearly equilibrated physical Fe3+ coordination network due to the slight increase of Dm and almost constant Dres distribution width shown in Figure 7. It is noteworthy here that inductively coupled plasma optical emission spectroscopy (ICP-OES) can be used to accurately determine the Fe3+ content in the dried hydrogels, as shown in Figure S8. When the Fe3+ concentration is increased from 0.06 to 0.18 mol/L, the Fe3+ content in the dried D20 hydrogels actually only increased around 0.8 wt %, and thus there is only slight enhancement of Dm as shown in Figure 7C. In addition, in principle, each Fe3+ ion will coordinate to three carboxylic acid groups, which means that the molar ratio of CFe3+/CAAc should be around 1/3 (∼33.3%) in the saturation state. Here, CFe3+/CAAc is defined as the molar ratio between the final Fe3+ content and the initial AAc concentration. However, it is interesting to note that even at a high Fe3+ concentration of 0.18 mol/L, where the Fe3+ content (corresponding to ∼1.4 g of FeCl3) in the ferric solution during the sample preparation is much higher than the theoretical need (corresponding to ∼0.26 g of FeCl3) for forming the trivalent Fe3+ complex, the CFe3+/ CAAc is still around 31% (Figure S8). This may be ascribed to the difficulty of the diffusion of Fe3+ ions into the central part of the bulk hydrogel samples. On the other hand, the CFe3+/CAAc value may also be underestimated, since the possibly unreacted AAc/AAm monomers and the cross-linker/initiators/catalyst weights were not taken into consideration. Furthermore, for the D20 hydrogels, Dm and the final Fe3+ content in the hydrogels can also be correlated, as shown in Figure S9. In fact, Dm increased ca. 2.5 times when CFe3+/CAAc increased from 17% to 32%. However, an attempt to establish a quantitative correlation between Dm and actual Fe3+ content (or the corresponding number of physical coordination bonds) is rather difficult and complex, since the final content of carboxylic groups and the actual coordination state (mono-

valent, bivalent, or trivalent) are difficult to determine. In this regard, a correlation between Dm and the concentration of the FeCl3 solution shown in Figure 7C will provide more straightforward hints and insights for the sample preparation.

4. SUMMARY AND CONCLUSION In this study, dynamic shear rheology and solid-state proton NMR experiments were performed to investigate the linear viscoelasticity and network cross-linking structures, respectively. For the SC hydrogels, the cross-linking density does not change with increasing CAAc/CAAm as revealed by proton MQ NMR spectroscopy. Because of high contents of polymer defects (∼80%) in the SC hydrogels, the viscoelastic moduli and mechanical performance were much poorer than those of DC hydrogels. For the DC hydrogels, the viscoelastic moduli were enhanced with increasing CAAc/CAAm, which were ascribed to the increased content of polymers and carboxylic acid groups for forming Fe3+ coordination complex as well as the reduction of network defects in the hydrogels. Variable temperature frequency-dependent shear experiments were performed to determine the activation energy of viscoelasticity. The activation energy of viscoelasticity for the physical network was even higher than that of the chemically covalent network in the D25 hydrogel, which was also comparable to the free energy required to break the UPy-based quadruple H-bonding groups. Hence, it clearly demonstrated the significant contribution of Fe3+ coordination to the enhancement of viscoelastic moduli and mechanical performance of DC hydrogels. Furthermore, the diverse temperature-dependent behaviors of loss modulus were also observed for the SC and DC hydrogels, which can be attributed to the dynamic dissociation and reassociation of Fe3+ coordination bonds for energy dissipation in the DC hydrogels. In addition, during the shear experiment, the Fe3+ complex with moderate/weak binding strength might transform to those with strong binding strength and serve as permanent-like cross-linkages to resist the periodic deformation when a large strain frequency was applied. Thus, the loss modulus (G″) and loss factor (tan δ) will decrease first and then reach a plateau with increasing the strain frequency. However, when the Fe3+ concentration is large enough during the preparation of DC hydrogels, the binding strength between Fe3+ ions and carboxylic acid groups will be strong enough to resist the deformation, and thus G″ nearly changes in the probed strain frequency range. Proton MQ NMR spectroscopy also quantitatively determines the heterogeneous structural evolution of the DC hydrogels with increasing the Fe3+ concentration, where a nearly equilibrated Fe 3+ coordination network was formed when the Fe 3+ concentration was beyond a critical concentration of 0.06 mol/L. To conclude, the microscopic origin of the viscoelasticity of the DC hydrogels can be well elucidated by the experimental findings obtained from the proton MQ NMR experiments. The correspondence between the rheological results and NMR findings is well addressed in this study and thus will be beneficial for understanding the relationship between microscopic structures and macroscopic viscoelastic properties of hydrogels. To a large degree, such a wellestablished structure−property relationship could further provide guidance for fabrication of high performance hydrogels with precisely controllable microstructures and mechanical behaviors. J


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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.7b01854. Figures S1−S9 (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*(R.Z.) E-mail: [email protected]. *(P.S.) E-mail: [email protected]. ORCID

Rongchun Zhang: 0000-0002-2480-2652 Xiaoliang Wang: 0000-0001-7820-4706 Pingchuan Sun: 0000-0002-5603-6462 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Prof. Kay Saalwächter for pointing us to the lognormal distribution. We are indebted to the reviewers for the detailed and insightful comments, which greatly help improve our data analysis and interpretation. The authors are grateful for the financial support by the National Natural Science Foundation of China (NSFC) (Nos. 21534005, 21374051, and 21704046), China postdoctoral Science Foundation (No. 2016M601249), PCSIRT (IRT1257), and the 111 Project (B12015).

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