Time Dependence of Dissipative and Recovery Processes in

May 15, 2013 - ABSTRACT: The strain rate effect on large strain dissipation and behavior recovery are presented to understand the toughening effect of...
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Time Dependence of Dissipative and Recovery Processes in Nanohybrid Hydrogels Séverine Rose, Alexandre Dizeux, Tetsuharu Narita, Dominique Hourdet, and Alba Marcellan* PPMD, Physico-Chimie des Polymères et des Milieux Dispersés (UMR 7615), UPMC-CNRS-ESPCI, 10 rue Vauquelin, 75005 Paris Cedex 05, France S Supporting Information *

ABSTRACT: The strain rate effect on large strain dissipation and behavior recovery are presented to understand the toughening effect of silica nanoparticles in nanohybrid hydrogels. Such nanohybrid gels combine a poly(N,N-dimethylacrylamide) (PDMA) covalent network and physical interactions by adsorption of polymer chains at the silica nanoparticle surface. A series of model nanohybrid gels has been designed to obtain a well-controlled architecture. First insights on the structure (SANS) demonstrated that silica nanoparticles were welldispersed in the gel, including after cyclic mechanical loading. The characteristic times involved in the nanoparticle/polymer association were investigated by large strain mechanical cycling varying the strain rate from 3 × 10−4 s−1 to 0.6 s−1. The mechanical behavior of the hybrid hydrogel varies tremendously over a relatively small range of strain rates, ranging from almost non dissipative (at slow strain rates) to highly dissipative at high strain rates. However, upon cycling over time-scales of tens of seconds, the strong physical interactions taking place between nanosilica particles and PDMA network chains enabled the hydrogel to recover its initial mechanical properties. The main feature of this work is the remarkable role played by silica nanoparticles in the network to promote transient and recoverable connectivity by reversible adsorption/desorption processes. The strong strain rate dependence suggests that toughening mechanisms operating at standard strain rates as often reported, maybe quite different at slower or larger strain rates.

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where Σ is the surface density of chains crossing the fracture plane, Nc is the number of monomer units between cross-links, and J is the energy required to rupture one covalent bond. Such argument is based on the assumption that the transmitted load throughout the network brings each bond in a polymer strand crossing the interface close to its maximal free energy (i.e., the dissociation energy of the weakest bond in the monomer unit). Of course in the case of highly swollen networks, the G0 decreases with polymer volume fraction and G0 values obtained for gels are typically ≅1 J·m−2 while they are 10 times higher for unfilled rubbers. Yet, experimentally, chemically cross-linked rubbers can have fracture energies, Gc that are several orders of magnitude higher (around 1−10 kJ·m−2) implying large viscoelastic dissipation10 processes in the crack tip vicinity, i.e., locally, inducing stress concentration and high strain rates. In contrast, classical chemically cross-linked hydrogels do not exhibit such dissipative processes and typical Gc values are quite in line with the Lake and Thomas’ elastic prediction. The fracture process is in that case far away from the monomer− monomer friction regime. Several groups have been devoted to

INTRODUCTION Gel toughening remains a crucial issue for developing new technological applications. Hydrogels are already involved in many bioapplications1−4 such as superabsorbants, in agriculture for holding moisture in arid areas, contact lenses, biosensors for encapsulation and drug delivery systems. Recently, emerging fields as microfluidics, actuation, tissue engineering demand gels with high mechanical performances. Besides the biocompatibility aspects, supplying tough gels could provide valuable rubbery materials with low cost and low ecological impact. Over the past ten years, since the works of Gong and coworkers,5 many groups have reported on the synthesis of hydrogels with significantly enhanced mechanical properties relative to standard chemically cross-linked gels. Authors have argued that the existence of dissipative mechanisms in the gels as they are deformed and fractured was responsible for gel toughening.6−8 Understanding how dissipation affects fracture toughness remains a key-aspect in mechanical toughening of soft materials. From the basic picture of a network formed by a collection of long and flexible chains, Lake and Thomas9 proposed an estimate of the threshold energy G0 to propagate crack in unfilled rubbers, i.e.

Received: March 1, 2013 Revised: May 3, 2013

G0 = ΣJNc © XXXX American Chemical Society

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adsorbed and unadsorbed polymer segments at the silica surface, the polymer/silica associations could play the role of exchangeable sacrificial bonds upon stretching. Adsorption could promote that way, “recoverable” dissipative processes depending on the probed time-scale relative to the particle/ polymer association characteristic time. These recoverable processes should be particularly important in controlling the properties of the material in large strain, a regime which is essential for fracture properties. The purpose of this paper is to probe the characteristic times of the bond exchange process and therefore how it impacts the dissipation and the recovery. More generally, understanding the interplay between time-dependent dynamic bonds and permanent bonds during deformation is also highly relevant for other materials, for example the theoretical works which have been recently reported on systems made from latex particles containing surface labile bonds.33,34 For that purpose, the hybrid gel composition was revisited to reach a better compromise between stretchability and to obtain well reproducible structures. The general architecture of this series of nanohybrid gels is schematically pictured in Figure 1

promoting dissipation processes within the hydrogel by modifying the architecture of the polymer network. Gong and co-workers pioneered the idea of introducing dissipation in chemically cross-linked gels. The concept relies on introducing sacrificial covalent bonds. Authors developed double network (DN) gels,5 which result from the synthesis of a primary highly cross-linked network with a secondary network, which is entangled and partially interconnected with the primary network.11 By sacrificing the integrity of the primary network, the second network ensures the stress transfer and enlarges the damage zone before catastrophic fracture occurs ; the results obtained demonstrated dramatic improvements of the fracture toughness.12 Although DN gels or more recently silica-grafted DN gels13 demonstrated large dissipative processes during fracture propagation, the macromolecular architecture was seen to be permanently damaged upon deformation. A large softening, analogous to the Mullins effect, was demonstrated after consecutive loadings without any complete recovery of the original properties over time.14 In order to overcome this disadvantage different ways of introducing physical and therefore reversible interactions in the network by coupling ionic cross-linking15 or hydrophobic interactions6,16−18 have been proposed. In contrast to these gels which imply a rather complex chemistry, Haraguchi and co-workers19−21 developed the first highly extensible nanocomposite gels (NC gels) and highlighted the strong reinforcement effect of clay nanofiller. Such NC hydrogels19−23 showed very interesting mechanical properties: very high deformability (up to 1000%) and variable gel stiffness (by varying the clay concentration) was achieved without sacrificing extensibility too much. Nevertheless for that system, a significant instantaneous residual strain was observed, mostly assigned to platelet residual alignement.19,24 Our group developed a similar approach combining the NC concept and the idea of double network.25−28 The strategy consisted of coupling a covalent network with reversible interactions using silica nanoparticles as a physical cross-linker. Indeed, based on pioneering studies in semidilute solutions, macromolecular assemblies demonstrated strong adsorption of polymer chains (i.e., poly(N-alkylacrylamide) or polyether) onto the surface of silica nanoparticles.29 Emphasizing on the strong particle-chain interaction by varying the silica volume fraction, previous papers25,26 reported on the very promising mechanical properties: initial modulus, dissipation, stretchability, and fracture energy were seen to be simultaneously enhanced by varying the silica nanoparticle volume fraction. Recovery effects varying the duration of rest30 or strain rate dependence on tensile behavior25,31 have already been reported on NC hydrogels or more recently on other viscoelastic gels.7 But in contrast, the strain rate effect on large strain cycling behavior was not addressed so far. Yet an important question remains: Hybrid gels showed large dissipation under cycling at standard strain rate,25 but are these dissipative processes still operating at shorter time-scales, i.e., higher strain rates, or at very long time-scales, i.e., low strain rates? The objective of the present study is to investigate the time-dependence of the mechanical response of nanohybrid gels, focusing on two aspects: strain rate dependent dissipation and conditions for modulus recovery. Evidently, adsorption can provide an efficient anchorage of polymer chains onto the silica particles, but conceptually this process is reversible.28,32 Thus, as described previously the silica nanoparticle acts as a multifunctional cross-link, but thanks to exchange processes between

Figure 1. Schematic representation of hybrid hydrogels combining covalent cross-links (orange •) and physical interactions sketched by polymer chains adsorbed at the surface of silica nanoparticles. Arrows sybolize the exchange processes between adsorbed and unadsorbed polymer segments at the silica surface. Note that structure is analogous to a double network: the chemical network controls the strain recovery thanks to the entropic restoring forces and silica nanoparticles promote transient and recoverable connectivity.

and three important points should be underlined: (1) A small amount of chemical cross-linker (0.1 mol %) was added to fix the network topology. This small value is a compromise to overcome the self-cross-linking process that takes place during DMA polymerization25 and to avoid a dramatic reduction of network extensibility. 26 (2) To be able to compare quantitatively samples with different silica content but equivalent polymer swelling ratios, the hybrid gel structure and mechanical behavior was systematically characterized at the gel preparation conditions. Thus, the gel matrix hydration is well-controlled (i.e., 88 wt %) and the silica volume fraction was varied independently. (3) Special attention was paid to the dispersion state of silica particles in the polymer network. After a short description of the synthesis procedure, we present first structural results of the silica dispersion state, then the mechanical characterization at small and large strain and at different strain rates. We will specifically focus on the identification of the characteristic times involved in the B

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Table 1. Nomenclature and Composition of Nano-Hybrid Hydrogelsa composition at gel preparation conditions nomenclature

silica volume fraction, ϕsi

mwater/g

msilica/g

mMBA/mg

mDMA/g

Q0

SP0_PW0.14-R0.1 SP2_PW0.14-R0.1 SP3.5_PW0.14-R0.1 SP5_PW0.14-R0.1

0 0.097 0.158 0.212

10.624 10.624 10.624 10.624

0.000 2.975 5.206 7.437

2.3 2.3 2.3 2.3

1.485 1.485 1.485 1.485

8.5 8.5 8.5 8.5

Qe 41 30 29 23

± ± ± ±

1 2 3 4

The matrix hydration at the preparation state was fixed at 87.7 wt %. For all compositions, the amounts of co-initiators were fixed at mKPS = 0.041 g and vTEMED = 22.5 μL. Qe and Q0 are respectively the measured equilibrium swelling ratio and the fixed initial swelling, corresponding to the preparation state. Silica volume fraction was defined as the volume ratio of silica particles to the total volume of the hybrid hydrogel (using ρSi = 2.3 × 106 g.m−3). a

above 98% and the amount of extracted polymer chains, when observable, was less than 1 wt % of the total organic content. Similarly, the silica content of the hybrid hydrogels was controlled by thermal gravimetric analysis (TGA). This analysis did not show any significant difference before and after swelling equilibrium in pure water showing that the silica particles cannot diffuse out of the gel. Small Angle Neutron Scattering (SANS). SANS experiments were performed at the Laboratoire Léon Brillouin (CEA-Saclay, France) with spectrometers PACE and PAXY dedicated to isotropic and 2D analyses, respectively. Contrast matching experiments were carried out with hybrid hydrogels initially prepared with a H2O/D2O volume ratio of 0.78/0.21 that matches the contribution of the organic matrix (ρs = ρPDMA = 0.94 × 10−6 Å−2). Absolute scattering intensities (I(q) in cm−1 or 10−8 Å−1), obtained from the standard procedure, were normalized (Icor(q)) using the invariant for minor corrections resulting from variations in the solvent composition:

reversible associations and on the structural recovery after unloading.

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EXPERIMENTAL SECTION

Materials. N,N-Dimethylacrylamide (DMA, 99%, Aldrich), potassium persulfate (KPS, Acros Organics), N,N,N′,N′-tetramethylethylenediamine (TEMED, 99.5%, Sigma-Aldrich), and N,N′-methylenebis(acrylamide) (MBA, Fluka) were used as received. The silica nanoparticles (Ludox TM-50 from Dupont) were obtained from Aldrich. The spherical shape of the nanoparticles as well their radii (R ∼ 14 nm) were characterized by scanning electron microscopy and small angle neutron scattering. The crude silica suspension (52 wt % and pH = 9) was used as received. Gel Preparation and Composition. Hydrogels were prepared at 25 °C by free-radical polymerization (under nitrogen conditions) of DMA in an aqueous suspension of silica nanoparticles using KPS and TEMED as redox initiatiors. Compositions of hydrogels and their nomenclature are summarized in Table 1. Solid reagents were initially dissolved in water prior to synthesis: 4.5 wt % for KPS and 1.2 wt % for MBA. Then the MBA solution and DMA were mixed at room temperature in an aqueous suspension of silica particles. The homogeneous suspension was purged with nitrogen for 15 min under magnetic stirring. The aqueous solution of KPS was prepared and deoxygenated by nitrogen bubbling. KPS solution and TEMED were then added to the suspension under stirring and the mixture was transferred into appropriate molds. The samples were left for 24 h to allow the reaction to run to completion. For all syntheses, the (DMA)/(KPS) and (DMA)/(TEMED) molar ratios were fixed at 100. The cross-linking density was also held constant using a cross-linker/monomer molar ratio (MBA/DMA) of 0.1 mol %. The amount of silica particles was varied while maintaining the matrix composition. The nomenclature of hydrogels is SPX_PWYRZ; with S for silica, P for polymer, W for water and R for chemical cross-linker. X is the weight ratio between silica and polymer, Y is the weight ratio between polymer and water, and Z corresponds to the MBA/DMA percent molar ratio. The matrix composition is fixed throughout this study (i.e., Y = 0.14, Z = 0.1); samples in the next sections will be called SPX for an easier discussion. Swelling Measurements. Equilibrium swelling experiments were performed in 0.5 mol·L−1 LiNO3 solutions in order to screen electrostatic interactions. The samples, as prepared, were weighed and placed in a large excess of solvent at room temperature. The solvent was exchanged every day for 10 days and the swollen gels were reweighed. The swelling ratio Q, calculated assuming negligible extractible content and additivity of volumes, is:

Q=1+

1 mw p vspe mp

Icor(q) =

Q th Q exp

I(q) (2)

where Qexp and Qth are the experimental and theoretical values of the Invariant defined by:

Iexp =

∫0

∞

q2I(q) dq

(3)

and Q th = 2π 2(ρs − ρSi )2 ϕSi(1 − ϕSi)

(4)

where q is the scattering vector, ϕSi the volume fraction of silica and ρs and ρSi the scattering length densities of the solvent and silica particles, respectively. Details on the fitting procedure are given in Supporting Information, SM1. Specific sealed cells were used in order to prevent composition drift during experiment. Large Strain Mechanical Behavior. Tests were performed on an Instron tensile testing machine (model 5565) equipped with a 10N load cell (with a relative uncertainty of 0.16% in the range from 0 to 0.1 N) and a video extensometer which follows the displacement up to 120 mm (with a relative uncertainty of 0.11% at full scale). Sample dimensions were kept constant with a rectangular shape: gauge length (L0 = 25 mm), width (w = 5 mm), and thickness (t = 2 mm). The nominal strain rate was varied from 3 × 10−4 s−1 to 0.6 s−1, corresponding respectively to a displacement speed comprised between 0.4 and 900 mm·min−1. To avoid any drift in hydrogel composition during the experiment, mechanical tests at 3 × 10−4 s−1 were carried out under immersion in paraffin oil. For that purpose we designed a specific cell. A preload tensile force around 10 mN was applied before positioning and testing of the sample. Loading−unloading cycles were performed in order to investigate the characteristic times of the viscoelastic processes and the selfrecovery using three strain rates: 0.06, 0.6, and 3 × 10−4 s−1. Cycles were applied from 0 to a maximal nominal strain, ε0 = 0.25 and then the sample was unloaded. Strain was obtained from the optical extensometer and defined as the ratio (l − l0)/l0, where l and l0 are respectively the length during

(1)

is the specific volume of the dry polymer network (vpspe where −1 =0.95 mL.g ), mp and mw is the mass of the dry polymer network and the mass of absorbed water, respectively. The analysis of organic and inorganic extractibles was systematically carried out after swelling, by analyzing the external solution by size exclusion chromatography (SEC). For all the samples, the yield was vpspe

C

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Figure 2. Double logarithmic plots of the corrected scattered intensity obtained for (a) the silica suspension and the SP2 gel prepared at the same volume fraction, i.e., ϕSi = 0.099, and (b) PDMA/silica hydrogels with different amounts of silica: SP2, SP3.5, and SP4.8. The solid lines show the fit obtained from the hard sphere model. The curves are shifted from one to another by a factor of 10. The solvent used a H2O/D2O volume ratio of 0.78/0.21, that matches the organic matrix composition. Particle dispersion is not affected by the polymerization process. Details on the fitting procedure are given in the Supporting Information, SM1.

Figure 3. (a) Tensile mechanical behavior at 0.06 s−1: effect of nanoparticle content for SP0 (without silica), SP2, SP3.5, and SP5 and (b) effect of silica volume fraction on initial modulus (for a range of strain rates comprises between ε̇ ∼ 3 × 10−4 s−1 and 0.6 s−1). The Guth and Gold’s prediction is represented as a dashed line with E0 = 10 kPa obtained from the SP0 sample. stretching and the initial length. Nominal stress, defined as engineering stress, was calculated from the tensile force and the initial cross section area (5 × 2 mm2). Rheology in the Linear Regime. Linear viscoelasticity was studied on a strain-controlled rheometer (Rheometrics RFSII, TAInstruments), in a parallel-plate geometry with a 25 mm diameter. Specific alumina plates were used to apply a stabilized normal force of 0.5 N (i.e., normal stress around 1 kPa). The gels were deformed at γ = 0.03 within the linear viscoelastic region, as determined using a frequency of 1 Hz. Spectra were recorded from 0.1 to 100 rad·s−1.

with a contrast matching of PDMA. In these conditions, only silica particles are seen. As shown in Figure 2a, the scattering plots obtained for the silica suspension and the nanohybrid gel prepared at the same particle concentration (SP2) are fully superimposed in the high-q regime. This behavior, as well as the q−4 dependence obtained in this regime (sharp interface), simply display that these particles are the same objects (similar size and morphology) which are observed in these two formulations. At low q, the slight excess of scattering intensity arising with the nanohybrid gel (SP2) points out weaker repulsions between silica beads possibly due to the bridging ability of the polymer network (see fitting parameters and procedure given in Supporting Information, Figure SM1 and Table SM1). Results demonstrated that all the formulations remain stable and that the polymerization process does not induce any particle aggregation effect (see Figure 2b). Silica as Multifunctional Strong Cross-Link. As shown in Table 1, the swelling ratio at equilibrium determined in these conditions clearly decreased with the amount of silica particles showing the existence of a constrained layer of polymer at the surface of the particles.25 As described previously in semidilute

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RESULTS AND DISCUSSION Hybrid Gel Structure: Uniform Dispersion of Nanoparticles in the Gel. The synthesis process of silica/PDMA hydrogels was shown to be simple, robust, and highly reproducible.25 This method allowed the preparation of hybrid hydrogels within a broad range of inorganic particle content (up to a silica volume fraction of 0.212 for SP5). As discussed above the gel composition after swelling shows that the polymerization is quantitative and particles do not leach out. To probe the structure, SANS experiments were carried out D

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Table 2. Strain Rate Effect on Initial Modulus (E), Dissipated Strain Energy (Ed), the Normalized Dissipated Strain Energy (Ed/ Ea). and Recovery in Initial Stiffness (R)a SP0 silica volume fraction, ϕsi ε̇ = 0.06 s−1

ε̇ = 0.6 s−1

ε̇ = 3 × 10−4 s−1

a

E (kPa) Ed (kJ·m−3) Ed/Ea R E (kPa) Ed (kJ·m−3) Ed/Ea R E (kPa) Ed (kJ·m−3) Ed/Ea R

SP2

0 10.2 ± 0.3 0 0 1 11.2 ± 0.3 0 0 1 8±1 0 0 1

0.097 14.0 ± 0.073 − 0.94 ± 28.1 ± 1.144 1.908 0.54 ± − − − −

0.5

0.03 0.6

0.06

SP3.5

SP5

0.158 35.0 ± 3.2 0.250 − − 55.0 ± 4.0 − − − − − − −

0.212 93.0 ± 8.1 1.430 2.821 0.98 ± 0.03 170 ± 9.3 5.387 8.984 0.31 ± 0.06 22 ± 5 0.216 0.384 0.98 ± 0.03

Ea is the applied strain energy for the fully relaxed gel, i.e., the area under the SP5 loading curve at 3 × 10−4 s−1.

solutions using adsorption isotherm29 (Γmax = 1 mg·m−2), the deswelling process observed is a qualitative indication that PDMA chains are strongly adsorbed at the nanoparticle surface, constraining that way the swelling equilibrium of the gel matrix. The characteristic time of the polymer/silica association can be probed with two different experiments: At small strains the spectrum of lifetimes of the bonds at iso structure can be probed, while at large strain, the role of these physical bonds on the hyperelastic behavior can be investigated. The large strain behavior was first investigated in monotonic uniaxial tensile tests to final failure. Typical stress−strain curves are given in Figure 3a for ε̇ = 0.06 s−1. The behavior is highly nonlinear and the pronounced “S-shaped” curve shows pronounced deviations from the classical model of rubber elasticity by both a pronounced softening at intermediate strains (around ε = 1) indicative of the breakup of a network structure and a strain hardening at large strain indicating finite extensibility of the polymer chains. The deformation process was homogeneous during the test: no necking was observed during extension. At low strains, as reported in Table 2, the initial modulus E is defined. The silica volume fraction in the gel (Figure 3b) leads to a significant increase of the initial stiffness. The impact of filler content on the small strain modulus of a rubbery matrix is a classical issue that has been addressed by Guth, Gold, and Simha.35,36 This model, well-known in rubber technology, implies an incompressible matrix, hard spherical particles and non adhering interfaces between the fillers and the matrix. The Guth and Gold equation, which is typically applied for filler volume fractions below ϕ < 0.35, is given as follows: E = E0(1 + 2.5ϕ + 14.1ϕ2)

However as shown by the tensile behavior in Figure 3a, and reported previously on more lightly25 or highly chemically cross-linked,26 interactions at the silica/polymer interface led not only to an increase in stiffness, but also to a large increase in the nominal strain at failure. This is a remarkable result within the fields of nanocomposite engineering as well as NC gels studies:19,37,38 here the strong enhancement of E by a factor 10 is combined with a significant increase in deformability at failure (by a factor 4 from SP0 to SP5). A simple estimate of the work of extension, given by the area under the tensile curve, yields Wext ≅ 8 kJ·m−3 for SP0 and Wext ≅ 210 kJ·m−3 for SP5, revealing that physical interactions efficiently delay crack initiation and propagation. To appreciate this increase in strain at failure it is important to note that the local average strain in the matrix domains is necessarily amplified over that of the macroscopic strain since the stiffer particles accommodate little of the applied macroscopic strain. A simple calculation leads to a strain amplification factor, A, as follows:39 A=

εmatrix = (1 − ϕ1/3)−1 ε

(6)

Here, ϕ is the volume fraction of filler, εmatrix, the strain undergone by the matrix and ε, the nominal strain. In the case of SP5, eq 6 gives a strain amplification factor value of A = 2.5, meaning that the strain at failure within the gel matrix reaches approximately εmatrix ≅ 1700% (using ε = 700%). This is a strong increase in stretchability at break in comparison with the mean strain at failure obtained for the unfilled SP0 gel. Such toughening effect is relied on the ability of the hybrid network to relax stress concentration by dynamic exchanges between anchored polymer segments to the particles and unadsorbed polymer strands. Pointing out such exchanges and estimating the characteristic time of recovery and stress relaxation processes are crucial for understanding the synergistic impact of silica particle in the gel on both stiffening and on toughening. Indeed, the stiffening effect was observed to be highly time-dependent as shown in Table 2 and Figure 3b. Results reveal that the viscoelastic contribution is correlated to the content of physical cross-links. Increasing the nominal strain rate by a decade, the hydrogel becomes stiffer, reflecting that part of the relaxed chains at 0.06 s−1 are still unrelaxed at 0.6 s−1. Interestingly, if the nominal strain rate is strongly lowered (i.e., 3 × 10−4 s−1), the reinforcement effect of silica

(5)

Here E is the modulus of the filled material, E0, the unfilled matrix modulus, and ϕ, the volume fraction of filler. The model greatly underestimates the experimental modulus, indicating the presence of strong interactions between the gel matrix and the silica nanoparticles (Figure 3b). The reinforcement effect becomes more noticeable above 10 vol % of filler. As pictured in Figure 1, the silica nanoparticles mainly act as multifunctional cross-links where the anchoring of the polymer chains at their surface leads to an additional increase in the number of elastically active chains per unit of volume. These results are qualitatively in line with the swelling behavior previously described. E

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Figure 4. Storage (G′) and loss moduli (G″) obtained from dynamic measurements (γ = 0.03, 25 °C) for increasing amounts of nanoparticle: SP0, SP2, SP3.5, and SP5. G′: filled symbols. G′′: open symbols. Assuming the common relation G* = E*/3 with E* being the initial moduli obtained (at 0.06 s−1 and 0.6 s−1) from tensile tests, G* are figured as a guideline (⊗). Note that G* is given by G* = (G′2 + G″2)1/2.

Figure 5. (a) Effect of silica content on the first loading−unloading cycle (ε0 = 0.25 at ε̇ ∼ 0.06 s−1) ; general shape of 3 consecutive loading− unloading cycles is given for SP2. Note that the residual strain and mechanical behavior were fully recovered after 30 s of rest. (b) Strain rate effect of SP5 hybrid hydrogel: loading−unloading cycles for varied strain rates (from 3 × 10−4 s−1 to 0.6 s−1).

Reversible Associations: Dissipation and Recovery. Dissipative mechanisms were investigated by strain-controlled tensile loading−unloading cycles, from ε0 = 0 to ε0 = 0.25 to stay away from possible particle/particle compression in the transverse direction upon stretching. Indeed, above the critical strain of particle/particle interaction, recovery and dissipative processes are much more complex to analyze. As shown in Figure 5a for ε̇ = 0.06 s−1, the dissipated strain energy during the first cycle, corresponding to the area within the loop, is clearly induced by the particle/polymer associations, as was already reported on hybrid gels25 or NC gels.31,43 Without particles, the gel (SP0) behaves like a purely elastic system; dissipation of the matrix is therefore negligible. For hybrid hydrogels, the dissipated strain energy, Ed, given in Table 2, significantly increases with both silica volume fraction and strain rate. Figure 5b displayed the strain rate effect on dissipation for the SP5 hybrid gel. As Ed strongly depends on both strain rate and gel stiffness, it was normalized by the strain energy of the fully relaxed behavior (i.e., the area under the SP5 loading curve at 3 × 10−4 s−1). The SP5-normalized dissipated energies, plotted in Figure 6, show that the hybrid gel can dissipate more than 10 times the amount of energy than the pure matrix is able to store at the same strain level, which is quite remarkable. This dissipative capability is however highly strain rate dependent. Such a mechanism must be active in many gels toughened by labile sacrificial bonds and emphasizes

nanoparticles vanishes. The fully relaxed modulus obtained for SP5 at that low strain rate follows the Guth and Gold’s prediction. Within the experimental time-scale, the physical network appears to be fully relaxed and the mechanical response remains controlled by the chemical network. To complete these large strain experiments we carried out more classical dynamic mechanical experiments in a parallel plates rheometer to determine the viscoelastic properties within the linear viscoelastic regime. Results are given in Figure 4 and confirmed the strong strain rate effect induced by the particle/ polymer interaction. While large strain measurements are suited to characterize nonlinear strain dependent viscoelastic effects, the linear regime is better adapted to probe time dependent effects. Within the accessible strain rate regime (i.e., from 0.016 to 16 s−1), both storage and loss moduli were seen to increase with strain rate. The elastic modulus follows a similar power law as those reported in physical gels15,40 or hydrophobically modified gels.17,21,41 Values of the moduli are in quite good agreement with complex moduli obtained by tensile tests. Higher discrepancies were noticed for the stiffest gel, SP5, most likely due to slippage effects.42 A qualitative illustration of the difference in macroscopic bulk properties of the obtained gels is shown in Supporting Information, Figure SM2, comparing the falling of a bead on a pure chemically cross-linked gel (SP0) and on the hybrid gel (SP5). F

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strain after full unloading decreased as the strain rate was lowered. The integrity of the loading−unloading curves is plotted in Figure 7b. The initial modulus can be regarded as a global parameter to analyze the connectivity of the network, since the modulus is related to the density of elastically active chains. A decrease in stiffness can be considered as a reduction in the connectivity of the transient network and a recovery parameter, called R, is defined as follows: R=

Ei2nd Ei1st

(7)

with E2nd the initial modulus during the second loading cycle i and E1st i the initial modulus during the first loading cycle. The results presented in Figure 8 highlight that the recovery of the initial network connectivity was complete when studied

Figure 6. Time-dependence of the normalized dissipated strain energy, Ed/Ea, defined as the ratio of the dissipated strain energy at large strain during the first cycle (i.e., area under the loop with a maximal applied strain of ε0 = 0.25), normalized by the strain energy of the fully relaxed gel (i.e., the area under the SP5 loading curve at 3 × 10−4 s−1).

the need to carry out cyclic and fracture tests as different strain rates. In engineering materials single stress−strain curves are rarely the most useful since materials are subjected to repeated loadings. But more importantly, investigating at least several loading−unloading cycles and in particular the evolution of the modulus on successive loadings, is a relevant indicator of internal damage or of internal self-healing capabilities. Figure 7a shows that the hybrid hydrogel clearly displays a strain ratedependent recovery in initial modulus. Conversely for the unfilled system, as expected the stress−strain curve remained unaffected by the cycling, regardless of the strain rate. A large decrease in initial modulus is observed for the high strain rate (0.6 s−1) cycles as the number of cycles increases, whereas at low strain rates the modulus was unchanged. This trend was confirmed by loading−unloading cycles carried out at 3 × 10−4 s−1: the initial modulus was observed to remain constant with subsequent cycles. Furthermore, interestingly, the residual

Figure 8. Time-dependence of recovery in connectivity for SP5. the initial modulus Recovery is defined as the ratio between E2nd i during the second loading cycle and E1st i the initial modulus during the first loading cycle.

at low strain rates, while it strongly decreased at 0.6 s−1. Contrary to gels prepared with anisotropic clay fillers, no effect

Figure 7. Time-dependence of recovery in initial modulus for SP5. (a) Comparison of the initial modulus during three subsequent loadings. Note that the unloading portion of the cycle is not shown in order to simplify the reading. (b) Effect of 10 s of rest on the cycling behavior at 0.6 s−1. G

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Figure 9. 2D SANS patterns of SP2 gel prepared in H2O/D2O 21/79 (silica scattering) from left to right: initially at ε = 0, during an applied strain ε = 100% and at rest (ε = 0). Duration of recording is 1800 s for each applied strain steps. Stretching direction is horizontal.

interactions appears to be of the order of 10 s. Below this characteristic time t ≪ τ, typically for the case of consecutive cycles at 0.6 s−1, the stiffness decrease indicates a loss of around 70% of the elastically active chains. Only 10 s of rest enable the bulk material to self-reorganize and to recover an equivalent density of elastically active chains. The gel is mending in the bulk within this time-scale. Depending on the frequency window, the silica/polymer interactions can break and selfreform, enlarging therefore very efficiently the dissipative processes. Contrary to specific associative systems like telechelic polymers or more recently in hydrophobically modified gels,41 the dynamic mechanical measurements within the linear viscoelastic regime were not able to reveal a well-defined characteristic time but lead to a power law dependence of both the elastic and the loss moduli. Such a rather broad spectra of relaxation times was also reported for polyacrylamide−clay NC gels31,46 and was observed for many associating polymers with multiple stickers. Experimental estimates of the association lifetime in polymer solutions47 have shown to be complex. In our case, the chain detachment at the silica surface and exchange with other polymer chains that occur under strain probably proceed through multiple step mechanisms that are responsible for a such large distribution of relaxation times. Simultaneously to these strong dissipative processes, interestingly the residual strain is readily recovered. The chemical cross-links not only set the network topology at long time scales but also promote the strain recovery thanks to the entropic restoring forces. As depicted in Figure 1, the characteristic time of the exchanges between the adsorbed and unadsorbed state sharply controls the mechanical response since reinforcement, dissipation and recovery are highly timedependent. At high strain rates for t ≪ τ, the integrity of the double network (i.e., the covalent network and the frozen physical associations) is probed whereas for t ≈ τ, the stress relaxation is ensured by a release of strained polymer chains. Thus, at long time-scale for t ≪ τ, the behavior, in terms of initial modulus, reduction in both residual strain and dissipation, is mainly controlled by the chemical network. In such time window, the contribution of the nanoparticle/ polymer interactions is fully screened.

of residual orientation of the particles is expected in our system. Then the mechanical experiment is mainly probing the characteristic times involved in the polymer/particle dissociation/association. As shown in Figure 7b and Figure 8, the recovery in modulus is typically obtained in the range of 10 s. Within 30 s of rest, the residual strain fully disappeared and the cycling behavior was similar to that of the initial cycle. Finally, the characterization of the silica dispersion under cycling loading carried out by SANS demonstrated that the mechanical cycling did not induce any permanent change in structure, thus confirming the relevance of our recovery parameter, R as a an index of network connectivity. As shown in Figure 9, the isotropic scattering pattern of the initial unstrained material (ε = 0) deforms in an anisotropic manner for ε = 100%, with two lobes appearing in the stretching direction. This so-called “abnormal butterfly pattern”, already observed in many stretched materials like swollen networks, filled-reinforced elastomers,44 or NC gels,45 indicates that the silica particles are reassembled into anisotropic aggregates perpendicularly to the stretching direction. But unloading the sample at ε = 0, the scattering pattern returns to its original isotropic configuration. From Associative Systems to Double Networks. Nanohybrid gels can be depicted as a double network combining a permanent network ensured by chemical crosslinks and a transient network through associative junctions. The pure chemically cross-linked gel (SP0) behaves like a purely elastic network. Yet, SP5 strain cycling behavior results at very low strain rate demonstrated that the same chemical network which forms the covalent matrix of hybrid gels is also involved in associative interactions with silica nanoparticles. The relaxation processes observed in loading/unloading cycles can be interpreted as a release of constrained polymer chains under loading and we show that this release occurs in a characteristic time scale. Since the strength of the silica/polymer physical junctions is weaker than those of covalent bonds, the mechanical force accumulated under large strains within the polymer chains, above a critical strain energy, can be released through the yielding of these physical cross-links. The magnitude of macroscopic dissipation is then tuned by the amount of reversible cross-links and by the characteristic time involved in the polymer/silica association. More remarkably, such silica/polymer dissipative interactions also have the potential to be reformed within a given time-scale. Under large amplitude cycling, we demonstrated that the characteristic time needed to reform nanosilica/polymer

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CONCLUSIONS Time-dependent dissipative mechanisms and recovery of highly stretchable nanohybrid hydrogels were investigated on a model nanohybrid hydrogels. The silica nanoparticles served as H

dx.doi.org/10.1021/ma400447j | Macromolecules XXXX, XXX, XXX−XXX

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and mechanical characterization of hydrogels and Pr. F. Boué, Pr. F. Cousin and , A. Hélary from Laboratoire Léon Brillouin (Saclay, France) for their help in SANS experiments. The authors thank Pr. C. Creton for his support and useful discussions.

multifunctional physical cross-linkers, allowing the transient anchoring of surrounding chains. Whereas the unfilled gel matrix behaved elastically, the nanosilica/polymer interactions were responsible for the viscoelastic character of the hybrid gels. We demonstrated that nanohybrid gels can be considered to some extent as double networks. The covalent network controls the elasticity recoverability and we verified that it controls the long time behavior. The physical network formed by adsorption of polymer chains on silica surface is responsible for the strong stiffening effect at small strain, but is highly time-dependent. The process involved in transient silica/polymer interactions clearly governs the dynamics of the gel: below the association characteristic time (for t ≪ τ), the hybrid gel response is governed by the double network whereas at long time-scale (for t ≪ τ), the physical network appears to be fully relaxed by a release in polymer chains constrained and the mechanical response remains purely controlled by the chemical network. In such time-scale, the contribution of the nanoparticle/polymer interactions vanishes. Interestingly, the silica nanoparticles, acting as transient cross-links, are able to interact reversibly with the polymer chains of the network. Thus, in the right frequency window, these hybrid gels are able to dissipate more than 10 times the amount of energy that the pure matrix can store while having a rapid recovery in strain (i.e., typically in the time scale of 1−10 s). Within this time scale, the network is able to self-reorganize in order to recover an equivalent density of elastically active chains. Thus, the “sacrificial” silica/polymer interactions may avoid irreversible damage of the covalent network by releasing stress, but more importantly can be reformed and promote very efficiently both dissipation and recovery processes. Such a combination of time-dependent bonds and permanent bonds may be relevant for other soft material systems33,48−51 where two populations of bonds exist and stress the need to understand the role played by these time scales on more complex processes such as fracture.

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(1) Buchholz, F. L.; Graham, A. T., Modern Superabsorbent Polymer Technology; Wiley-VCH: New York, 1997. (2) Hoare, T. R.; Kohane, D. S. Polymer 2008, 49 (8), 1993−2007. (3) Oh, J. K.; Drumright, R.; Siegwart, D. J.; Matyjaszewski, K. Prog. Polym. Sci. 2008, 33 (4), 448−477. (4) Peppas, N. A., Hydrogels in Medicine and Pharmacy; CRC Press: Boca Raton, FL, 1986. (5) Gong, J. P.; Katsuyama, Y.; Kurokawa, T.; Osada, Y. Adv. Mater. 2003, 15 (14), 1155−+. (6) Henderson, K. J.; Zhou, T. C.; Otim, K. J.; Shull, K. R. Macromolecules 2010, 43 (14), 6193−6201. (7) Sun, J. Y.; Zhao, X. H.; Illeperuma, W. R. K.; Chaudhuri, O.; Oh, K. H.; Mooney, D. J.; Vlassak, J. J.; Suo, Z. G. Nature 2012, 489 (7414), 133−136. (8) Yu, Q. M.; Tanaka, Y.; Furukawa, H.; Kurokawa, T.; Gong, J. P. Macromolecules 2009, 42 (12), 3852−3855. (9) Lake, G. J.; Thomas, A. G. Proc. R. Soc. London, Ser. A: Math. Phys. Sci. 1967, A300, 108−119. (10) Gent, A. N. Langmuir 1996, 12, 4492−4496. (11) Nakajima, T.; Furukawa, H.; Tanaka, Y.; Kurokawa, T.; Osada, Y.; Gong, J. P. Macromolecules 2009, 42 (6), 2184−2189. (12) Gong, J. P. Soft Matter 2010, 6 (12), 2583−2590. (13) Wang, Q.; Hou, R. X.; Cheng, Y. J.; Fu, J. Soft Matter 2012, 8 (22), 6048−6056. (14) Webber, R. E.; Creton, C.; Brown, H. R.; Gong, J. P. Macromolecules 2007, 40 (8), 2919−2927. (15) Kong, H. J.; Wong, E.; Mooney, D. J. Macromolecules 2003, 36 (12), 4582−4588. (16) Abdurrahmanoglu, S.; Can, V.; Okay, O. Polymer 2009, 50 (23), 5449−5455. (17) Hao, J. K.; Weiss, R. A. Macromolecules 2011, 44 (23), 9390− 9398. (18) Haque, M. A.; Kurokawa, T.; Kamita, G.; Gong, J. P. Macromolecules 2011, 44 (22), 8916−8924. (19) Haraguchi, K.; Li, H. J. Macromolecules 2006, 39 (5), 1898− 1905. (20) Haraguchi, K.; Li, H. J. J. Polym. Sci., Part B: Polym. Phys. 2009, 47 (23), 2328−2340. (21) Haraguchi, K.; Takehisa, T.; Fan, S. Macromolecules 2002, 35 (27), 10162−10171. (22) Messing, R.; Schmidt, A. M. Polym. Chem. 2011, 2 (1), 18−32. (23) Wu, C. J.; Gaharwar, A. K.; Chan, B. K.; Schmidt, G. Macromolecules 2011, 44 (20), 8215−8224. (24) Nishida, T.; Endo, H.; Osaka, N.; Li, H.; Haraguchi, K.; Shibayama, M. Phys. Rev. E 2009, 80 (3), 4. (25) Carlsson, L.; Rose, S.; Hourdet, D.; Marcellan, A. Soft Matter 2010, 6 (15), 3619−3631. (26) Lin, W. C.; Fan, W.; Marcellan, A.; Hourdet, D.; Creton, C. Macromolecules 2010, 43 (5), 2554−2563. (27) Lin, W. C.; Marcellan, A.; Hourdet, D.; Creton, C. Soft Matter 2011, 7 (14), 6578−6582. (28) Rose, S.; Marcellan, A.; Hourdet, D.; Creton, C.; Narita, T. Macromolecules 2013, DOI: http://dx.doi.org/10.1021/ma4004874. (29) Hourdet, D.; Petit, L. Hybrid Hydrogels: Macromolecular Assemblies through Inorganic Cross-Linkers. In Macromolecular Symposia; Patrickios, C. S., Ed. Wiley-VCH Verlag GmbH: Weinheim, Germany, 2010; Vol. 291−292, pp 144−158. (30) Lian, C. X.; Lin, Z. M.; Wang, T.; Sun, W. X.; Liu, X. X.; Tong, Z. Macromolecules 2012, 45 (17), 7220−7227. (31) Xiong, L. J.; Hu, X. B.; Liu, X. X.; Tong, Z. Polymer 2008, 49 (23), 5064−5071.

ASSOCIATED CONTENT

S Supporting Information *

Figure SM1, double logarithmic SANS plots for nanohybrid gels SP2, SP3.5, and SP4.8; Table SM1, fitting parameters related to silica scattering obtained from hybrid hydrogels prepared in solvent in solvent with a H2O/D2O volume ratio of 0.78/0.21; and Figure SM2, chronophotographs demonstrating a 25 mm steel bead falling on hydrogels, which illustrates the impact of nanoparticles on the mechanical behavior of the nanohybrid hydrogels. This material is available free of charge via the Internet at http://pubs.acs.org/

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: (A.M.) [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Ph.D. scholarship of S.R. was provided by the Ph.D. school “Chemistry and Physics of Materials” of University Pierre and Marie Curie (ED 397, UPMC Paris, France). The authors would like to acknowledge Dr. G. Ducouret, F. Martin, Dr. D. Martina, and L. Olanier for their help and advice in synthesis I

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(32) Pefferkorn, E.; Haouam, A.; Varoqui, R. Macromolecules 1989, 22 (6), 2677−2682. (33) Salib, I. G.; Kolmakov, G. V.; Gnegy, C. N.; Matyjaszewski, K.; Balazs, A. C. Langmuir 2011, 27 (7), 3991−4003. (34) Iyer, B. V. S.; Salib, I. G.; Yashin, V. V.; Kowalewski, T.; Matyjaszewski, K.; Balazs, A. C. Soft Matter 2012, 9 (1), 109−121. (35) Guth, E.; Simha, R. Kolloid-Z. 1936, 74 (3), 266−275. (36) Guth, E. J. Appl. Phys. 1945, 16, 20−25. (37) Liu, Y.; Zhu, M. F.; Liu, X. L.; Zhang, W.; Sun, B.; Chen, Y. M.; Adler, H. J. P. Polymer 2006, 47 (1), 1−5. (38) Wu, Y. T.; Xia, M. G.; Fan, Q. Q.; Zhang, Y.; Yu, H.; Zhu, M. F. J. Polym. Sci., Part B: Polym. Phys. 2011, 49 (4), 263−266. (39) Bueche, F. J. Appl. Polym. Sci. 1961, 5 (15), 271−281. (40) Ng, T. S. K.; McKinley, G. H.; Ewoldt, R. H. J. Rheol. 2011, 55 (3), 627−654. (41) Abdurrahmanoglu, S.; Cilingir, M.; Okay, O. Polymer 2011, 52 (3), 694−699. (42) Meyvis, T. K. L.; De Smedt, S. C.; Demeester, J.; Hennink, W. E. J. Rheol. 1999, 43 (4), 933−950. (43) Gaharwar, A. K.; Dammu, S. A.; Canter, J. M.; Wu, C. J.; Schmidt, G. Biomacromolecules 2011, 12 (5), 1641−1650. (44) Rharbi, Y.; Cabane, B.; Vacher, A.; Joanicot, M.; Boue, F. Europhys. Lett. 1999, 46 (4), 472−478. (45) Shibayama, M. Polym. J. 2010, 43 (1), 18−34. (46) Abdurrahmanoglu, S.; Okay, O. J. Appl. Polym. Sci. 2010, 116 (4), 2328−2335. (47) Regalado, E. J.; Selb, J.; Candau, F. Macromolecules 1999, 32 (25), 8580−8588. (48) Haraguchi, K.; Uyama, K.; Tanimoto, H. Macromol. Rapid Commun. 2011, 32 (16), 1253−8. (49) Wang, Q.; Mynar, J. L.; Yoshida, M.; Lee, E.; Lee, M.; Okuro, K.; Kinbara, K.; Aida, T. Nature 2010, 463 (7279), 339−343. (50) Cordier, P.; Tournilhac, F.; Soulie-Ziakovic, C.; Leibler, L. Nature 2008, 451 (7181), 977−980. (51) Hui, C. Y.; Long, R. Soft Matter 2012, 8 (31), 8209−8216.

J

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