The role of hydrogen bonding in alginate/poly(acrylamide-co-dimethylacrylamide) and alginate/poly(ethylene glycol) methyl ether methacrylate-based tough hybrid hydrogels

Abstract

The interpenetrating alginate-based hybrid hydrogel network is a tough yet recoverable material. This is believed to be caused by the combination of the strength of a covalent network, and the reversibility of an ionic network. However, hydrogen bonds are believed to also be responsible for the exceptional properties of these hydrogels. In this paper, the effect of varying the reactant concentrations on the mechanical properties of the hydrogels was first studied. By changing the monomer used (from polyacrylamide to polydimethylacrylamide) in the fabrication of the hydrogel, the effect of hydrogen bonding was studied. Compression test results showed that the presence of hydrogen bonds is critical for the high toughness of the hybrid hydrogel. Additionally, co-polymeric hybrid hydrogels were synthesized and shown to have improved mechanical properties over the original hybrid hydrogel, with an elastic modulus and compressive toughness of 350 kPa and 70 J m⁻³, respectively. The results of this experiment can be used to optimise the mechanical properties of future hybrid hydrogels.

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1. Introduction

Hydrogels are highly crosslinked, hydrophilic polymer networks that swell in water. Due to their ability to store water and good biocompatibility, hydrogels have been widely studied for many biomedical applications such as drug delivery,^1–4 cell scaffolds,^5–8 and artificial organs.⁹

Hydrogels can be classified into 2 main types depending on their crosslinks. Chemical hydrogels have permanent covalent bonds for crosslinks whereby the covalent network is formed by the addition of specific crosslinker molecules during the polymerization step (Scheme 1a). On the other hand, physical hydrogels have temporary physical interactions for crosslinks, and examples include ionic bonds¹⁰ and crystallite formation.¹¹ The types of crosslinks present in a network can affect the overall mechanical properties of a hydrogel.¹² However when compared to other classes of materials such as metals and even bulk polymers, hydrogels are still significantly weaker,¹³ and this can be explained by the lower volume fraction of polymer in the material due to the high water content in the hydrogels.^14,15 Hydrogels are also considered too weak for applications such as materials for artificial muscles and cartilages,^14,16,17 and hence this has generated a need for stronger hydrogels.

Scheme 1 (a) A covalent crosslink formed by N,N′-methylene(bisacrylamide) (BIS). The vinyl functional groups take part in the polymerization reaction, permanently joining 2 polymer chains together chemically. (b) Ionic bonds formed between divalent cations and alginate chains, resulting in an egg-box configuration. (c) A hybrid hydrogel network consisting of a covalently-linked network and an ionically-linked network. The networks are entangled with each other, forming an interpenetrating network.

In recent years, extensive work has been done to improve the mechanical properties of hydrogels. A notable method is the incorporation of a second, interpenetrating polymer network into the pre-existing hydrogel network.^18,19 This provides a new energy dissipation mechanism within the ‘double network’ hydrogels, giving rise to high toughness. However, this mechanism is based on the permanent rupturing of covalent bonds in the networks, and the inability of the hydrogels to recover after loading is a considerable weakness.

In a study published in 2012, Sun et al. addressed this issue by replacing one of the covalent networks with an alginate-based ionic network.²⁰ Alginate is a widely studied,^21,22 biocompatible^23,24 material, which is able to form ionic hydrogels^25–27 with multivalent cations such as Ca²⁺. It is made up of β-D-mannuronate (M) and α-L-guluronate (G) units which are stereoisomers. The G monomers are able to form supramolecular egg-box^28,29 complexes with Ca²⁺ ions (Scheme 1b), which are responsible for the toughness and good recovery of the hydrogels. Upon an applied stress, the egg-box configuration stretches and unzips into separate chains, dissipating energy²⁰ in the process. The chains are then able to re-zip together over time after the stress is released (Scheme 2).

Scheme 2 Energy dissipation mechanism of the hybrid hydrogel network. The entire network is stretched upon an incident stress (indicated by the arrows). The ionic crosslinks are preferentially broken, or unzipped, due to their weaker bond strengths. The ionic crosslinks, however, are temporary physical interactions which can re-form, or re-zip, over time, restoring the mechanical properties of the hydrogel.

By incorporating this reversible network with a poly(acrylamide) (PAAm) covalent network (Scheme 1c), hybrid hydrogels capable of high toughness and recovery over time were obtained. Some other research has also suggested that there is an interaction between the metal ion and amide, which may also occur in the hybrid hydrogels in the present work.³⁰

However, while the toughness of the hybrid hydrogel has been attributed to the presence of an interpenetrating network of covalent and ionic components, a third form of crosslinking, that of hydrogen bonds, is believed to also be responsible for the exceptional toughness of these hydrogels.³¹ Acrylamide (AAm) monomers possess hydrogen bond donor groups (Scheme 3), and hydrogen bonds can be formed with water and other compatible functional groups within the alginate backbone.^31,32 Hydrogen bonds are weaker than ionic and covalent bonds and will thus preferentially break over the other bonds in the presence of an incident stress. This contributes to the energy dissipation mechanism of the material, and hence overall material toughness.

Scheme 3 The amine groups in acrylamide (AAm) monomers possess 2 hydrogen atoms, making them able to form hydrogen bonds with functional groups that possess available electron lone pairs.

Additionally, while this method has produced one of the toughest synthetic hydrogels, the mechanical properties are still significantly weaker than that of artificial cartilage and other human tissues.¹⁴ Nevertheless, the hybrid network has proved to be promising due to its recoverability, and this thus forms the motivation for this study. In this work, we will study the effect of hydrogen bonding on the toughness in hybrid hydrogels. The concentrations of different reagents, such as the alginate concentration and crosslinker concentrations are varied to study their effects on hydrogel properties. Consequently, PAAm-based copolymeric covalent networks are incorporated into the hybrid networks, in an attempt to improve on the abovementioned hybrid hydrogel network composition and to study the effect of hydrogen bonding on the mechanical properties of the hydrogels.

2. Experimental

Materials

Alginic acid sodium salt, N,N′-methylenebis(acrylamide) (BIS), ammonium persulfate (APS) and acrylamide (AAm) were purchased from Sigma-Aldrich and used without further purification. N,N,N′,N′-Tetramethylethylenediamine (TEMED) was purchased from Alfa Aesar and also used without further purification. Dimethyl acrylamide (DMAA) and poly (ethylene glycol) methyl ether methacrylate (PEGMA) (M_n 1100) were both purchased from Sigma-Aldrich and purified by dissolving the respective monomers in tetrahydrofuran (THF) (Tedia) and running the solution through a column of inhibitor remover (Sigma-Aldrich). THF solvent is then removed via rotary evaporation. Calcium chloride (CaCl₂) (Tractus) and sodium chloride (NaCl) solutions were prepared for soaking hydrogels. In all synthesis steps, deionised (DI) water (18.2 MΩ) was used. Polypropylene syringes (Braun) were used as gel moulds so as to obtain cylindrically-shaped samples.

Sample preparation

(i) Hydrogel synthesis. Photoinitiated radical polymerization was used to synthesize the covalent network within the hydrogels.^33,34 Sodium alginate and monomers were added in various weight ratios, and fully dissolved in a glass vial of DI water (86 wt%). APS and BIS were subsequently added as photoinitiator and crosslinker, based on a molar ratio to the amount of covalent monomers present. The solution was then immediately bubbled with nitrogen gas for 1 hour to remove trace oxygen and other impurities within the solution. 5 μL of TEMED, a crosslinking accelerator, was added, and the mixture poured into moulds. The gel solutions were then cured in a UV chamber with a mercury vapour light source (254 nm, 50 μW cm⁻²) for 2 hours. Following which, the fully polymerized hydrogels are left in their moulds and placed into 5 M calcium chloride (CaCl₂) solution for 2 days to form the calcium alginate (ALG) ionic network. After being able to retain their shapes, the hydrogels were removed from their moulds and soaked in CaCl₂ solution for an additional 3 days.

(ii) Equilibrium swelling. Hydrogels can be characterized based on their water content values, or swelling ratios, and this can affect their mechanical properties.³⁵ The water content can range between two extreme scenarios; a fully dehydrated state, or a fully solvated state, also termed the equilibrium swelling state.^36–38 However, as hydrogels are exposed to an aqueous environment for many of their applications,^1–3,5,6 they tend to approach the equilibrium swelling state over time, and it is thus important to characterize hydrogels based on this.

To achieve equilibrium swelling, the synthesized hydrogels were soaked in DI water for 7 days. The time taken to achieve equilibrium swelling was determined by monitoring the daily weight of the soaked hydrogels, and identifying the point where the change in hydrogel weight from the previous day was less than 1%. This average time taken was found to be 5 days. To account for any potential changes in water uptake between different samples, 2 additional days of soaking were added, resulting in the abovementioned soaking time of 7 days.

(iii) ALG-P(AAm/DMAA) hybrid hydrogels. Poly(dimethyl acrylamide) (PDMAA) and PAAm have similar structures, except for the substitution of two –CH₃ methyl groups as seen in Fig. 1. This makes PDMAA units unable to form hydrogen bonds between one another.

Fig. 1 Chemical structures of (a) PAAm and (b) PDMAA. The substitutions of the 2 hydrogen atoms for –CH₃ methyl groups in DMAA make them unable to form hydrogen bonds.

To understand the effect of hydrogen bonding on mechanical properties, hybrid hydrogels with varying co-polymeric ratios of AAm and DMAA were synthesized using the abovementioned method, and subsequently swelled to the equilibrium state. The polymers constitute 14 wt% of the hydrogel, and the ratio of each monomer was determined by splitting the 14% into a covalent-to-ionic ratio (ALG/covalent), followed by a copolymer ratio (AAm/DMAA) within the covalent component. Samples were named based on the ratio of AAm and DMAA added.

(iv) ALG-P(AAm/PEGMA) hybrid hydrogels. ALG-P(AAm/PEGMA) hybrid hydrogels were synthesized by substituting DMAA with poly (ethylene glycol) methyl ether methacrylate (PEGMA) without further changes to experimental methods. The AAm/PEGMA copolymer ratio was varied up to 50 wt% PEGMA.

(v) Variation of reactant concentrations. The amount of other reactants were varied and studied as well. The ALG/covalent ratio is studied in Series B, while the crosslinker concentrations were represented by Series C.

(vi) ALG and PAAm control samples. ALG control samples were prepared by dissolving sodium alginate powder in DI water to form a 2.3 wt% solution which was poured into syringe moulds. CaCl₂ solution was added dropwise, and the samples were left to stand for 2 days, before being taken out of their moulds to soak in CaCl₂ solution. To investigate the effect of cationic concentration on mechanical properties, CaCl₂ was used in 2 different concentrations: 0.3 M and 5 M. The soaking time was also varied at 3 or 7 days.

PAAm control samples were prepared via photoinitiated radical polymerization, then left to soak in DI water for 7 days to reach equilibrium swelling. The effect of CaCl₂ on PAAm hydrogels was also studied by soaking PAAm samples in CaCl₂ followed by DI water for 7 days each.

Characterization

(i) Determination of elastic modulus and toughness. Due to the general unsuitability of tensile tests for hydrogels,¹⁴ compressions tests using cylindrically-shaped samples were done to obtain the elastic modulus and compressive toughness of the materials. Samples at equilibrium swelling were prepared into cylindrical pieces with a length-to-diameter ratio of 1 : 1. All samples were weighed at equilibrium swelling (W_eq) prior to mechanical testing.

Compressive tests were conducted with an Instron 5543 mechanical tester using a 100 N load cell, at a strain rate of 1.0 mm min⁻¹.

(ii) Determination of equilibrium swelling ratio. The crushed samples were dehydrated through a freeze-drying process, and their weights were measured to obtain the dry weight (W_d). The equilibrium swelling ratio in the hydrogel was calculated using the formula:
where a ratio of 2 would mean that the hydrogel absorbed water twice that of its own weight.

Results and discussion

Synthesis of hydrogels

Fully synthesized gels were nearly transparent and able to stand without support (Fig. S1†). In the initial stage, photopolymerization was determined to be complete upon a significant increase in viscosity in the sample, such that the gel solutions flowed minimally after the moulds were overturned for a few seconds.

Initial ionic crosslinking was verified by an obvious hardening of the hydrogel surface upon contact with CaCl₂ solution. Complete ionic crosslinking is determined by soaking ALG control samples in CaCl₂ solution for different amounts of time. This will be explained in greater detail in the following section.

ALG and PAAm control samples

ALG and PAAm control samples were treated in different conditions so as to study their effect on hydrogel properties. The compression test results were shown in Table S1.†

No significant changes in hydrogel properties were observed with a variation of soaking time. On the other hand, different soaking concentrations resulted in ALG hydrogels with noticeable differences in equilibrium swelling and mechanical properties (Fig. S2a and b†). This was believed to be due to the difference in gelation kinetics.²⁶ At lower cationic concentration, crosslinking occurs less rapidly, allowing the alginate chains to spread out to a larger extent, and thus resulting in larger water content in the alginate gels in 0.3 M solution. This reduces the volume fraction of polymer in the hydrogel, and thus weaker mechanical properties were observed.¹⁵ In the optimization of mechanical properties of hydrogels, introducing a higher cationic concentration to induce rapid crosslinking will be more effective.

PAAm samples on the other hand showed minimal response to CaCl₂ (Fig. S2a and b†), thus implying that any change caused by CaCl₂ occurred due to a dominant response by the ionic network. As compressive tests do not always have a definite point of failure,³⁹ mechanical failure was defined in this experiment to be at the point where (a) 70% strain was observed (Fig. 2a), or (b) visible fractures were observed in the sample, resulting in a significant dip in the stress–strain curve obtained (Fig. 2b).

Fig. 2 Stress–strain curves from the compression tests of different hydrogel samples. Elastic modulus was taken to be the linear region of the stress–strain curve, and was done in duplicate for test samples, and triplicate for control samples. The compressive toughness was taken to be the area under the stress–strain curve after (a) 70% strain, or (b) the strain at fracture, ε_f, of the sample, where fractures develop along the sample (see inset) and the stress–strain curve peaks. Compressive toughness values were taken from a single sample for each composition, and in triplicate for control samples.

The elastic modulus was obtained from the gradient of the most linear region of the stress–strain curve. Due to the ambiguity of the most linear region, compression tests for each sample were done in triplicate, and the average of the obtained elastic moduli values were used. Elastic moduli for control samples ALG and PAAm were done in triplicate. Compressive toughness was taken to be the area under the stress–strain curve, to a maximum strain of 70% (Fig. 2a). For samples that exhibit signs of fracture before 70% strain, the compressive toughness was taken to be the area under the stress–strain curve from 0% strain to the strain at fracture, ε_f, of the material, the point where stress is maximum (Fig. 2b). Compressive toughness for compositions were taken from single samples and averaged. It was also observed that ALG test samples became flattened after compression, and stayed in this physical state even after soaking in water (Fig. S3a†). Significant amounts of water were found around the ALG hydrogels on the sample stand of the mechanical tester. Weighing the crushed ALG samples immediately after showed that the samples had weight losses of up to 75%. In contrast to the response of the ALG hydrogels, PAAm hydrogels showed physical deformations but were still able to retain most of their original cylindrical shapes (Fig. S3b†). The samples were also found to have insignificant weight losses of less than 5% after compression.

Significance of hydrogen bonding in hybrid network

The list of samples and equilibrium swelling ratios of ALG-P(AAm/DMAA) hybrid hydrogels is shown in Table S2† and Fig. 3.

Fig. 3 The effect of DMAA concentration on hydrogel properties. (a) Elastic moduli and (b) toughness decreased steadily due to the absence of hydrogen bonding between DMAA monomers. *^Toughness values at DMAA concentrations of 50 and 90 wt% were measured until fracture strain at 65 and 59% strain respectively.

As the mechanical properties of hydrogels are dependent on both network composition and water content, samples should be compared at similar water content levels in order to effectively study the relationship between network composition and mechanical properties. As equilibrium swelling ratio remained within a small range of 3.07–3.64 (Fig. 3a), effects of water content were assumed to be insignificant, and a direct comparison of mechanical properties was done.

Modulus and toughness both decreased steadily with DMAA content (Fig. 3a & b). Additionally, both PAAm50-DMAA50 and PAAm10-DMAA90 had ε_f values below 70% (65 and 59% strain respectively). The significant reduction in mechanical properties supports the proposed effect of hydrogen bonds. As PAAm content goes down, the amount of hydrogen bond donor groups decreases, and this leads to a weakening of the material. From these results, it can be seen that hydrogen bonds play a significant role in contributing to material properties, and it is thus insufficient to only attribute them to the covalent and ionic networks. Furthermore, the pretreatment of the hydrogels by soaking them in urea solution led to poorer mechanical properties and also poorer toughness of the materials suggesting that hydrogen bonding plays an important role in the properties of the hydrogel.

ALG-P(AAm/PEGMA) hybrid hydrogels

In an attempt to synthesize tougher hydrogels, PEGMA was used to obtain ALG-P(AAm/PEGMA) hybrid hydrogels. PEGMA monomers possess a chain of 23 poly (ethylene glycol) (PEG) units (Fig. 4a). After polymerization, these chains are grafted onto the main polymer chain, resulting in a comb-like configuration within the network (Fig. 4b). Additionally, each PEG unit has a hydrogen bond acceptor group: an oxygen atom with 2 electron lone pairs. Thus, introducing PEGMA into the hybrid hydrogel network might lead to unique results. Table S3† lists the different ALG-P(AAm/PEGMA) samples synthesized.

Fig. 4 (a) Chemical structure of PEGMA. (b) A comb-like network structure was formed due to the introduction of PEGMA, which contains poly (ethylene glycol) side-chains. (c) Variation of equilibrium swelling ratio with differing AAm/PEGMA copolymer ratios, which is believed to be caused by the comb-like structure.

Weight measurements showed that the equilibrium swelling ratios increase with PEGMA content (Fig. 4c). This is believed to occur due to the combined effect of two features. Firstly, poly(ethylene glycol) (PEG) is very hydrophilic due to the presence of electron lone pairs for hydrogen bonding with water molecules.⁴⁰ Secondly, the grafted chains are inherently mobile.⁴¹ They thus adopt a fully extended conformation in water,⁴² allowing the hydrogel network to hold water more efficiently.

To compare the mechanical properties of ALG-P(AAm/PEGMA) compositions at the same water content, the hydrogels were first mechanically tested at various swelling ratios. This included the respective equilibrium swelling ratios for each composition, and the swelling ratios of hydrogels dried in an oven at 37.5 °C for 2 and 4 hours respectively. Results from the compression tests were then plotted against swelling ratio and are shown in Fig. 5a & c.

Fig. 5 (a) Scatter plot of modulus vs. swelling ratio for ALG-P(AAm/PEGMA) hydrogels. (b) At a swelling ratio of 3.0–3.15, modulus increased steadily with PEGMA weight concentration. (c) Scatter plot of toughness vs. swelling ratio for ALG-P(AAm/PEGMA) samples. (b) At a swelling ratio of 3.0–3.15, toughness increased steadily with PEGMA weight concentration. Legend: (PAAm100 (), PAAm90-PEGMA10 (), PAAm75-PEGMA25 (), PAAm50-PEGMA50 ()). The gray dashed lines in (a) and (c) represent a swelling ratio of 3.0.

Modulus and toughness values ranged between 200–600 kPa and 40–80 J m⁻³ respectively for compositions containing PEGMA concentration of 0–25 wt%. On the other hand, when the PEGMA concentration reached 50 wt%, PAAm50-PEGMA50, had a significantly lower modulus and toughness compared to the other compositions. At 50 wt% PEGMA content, the amount of PAAm hydrogen bond donor groups decreased so significantly that there is insufficient crosslinking to retain the hydrogel structure. This is similar to the discussion the preceeding section. In this case, the mechanical property of the hydrogel was much poorer than that of the DMAA or AAm containing hydrogels when the PEGMA content is very high. This shows that the brush like structure of the PEGMA as the interpenetrating network could disrupt the packing in the material thus leading to poorer mechanical properties. The incorporation of PEGMA is expected to improve the biocompatibility of the material as PEG is widely noted to be used to improve the biocompatibility of biomaterials. Yet, we need to be cautious about how much PEGMA we can add because the macro-monomeric unit such as this does not polymerize so well and does not form very effectively, an interpenetrating network.

Mechanical properties were then compared at a swelling ratio of 3.0 (dashed lines in Fig. 5a & c). Samples with the closest swelling ratio values were picked for comparison, resulting in a range of 3.00–3.15 and, their mechanical properties are presented in Fig. 5b & d. Due to the weak mechanical properties of PAAm50-PEGMA50, it was excluded from this comparison. Both elastic modulus and compressive toughness increased steadily to 350 kPa and 70 J m⁻³ respectively with PEGMA concentration. These trends were attributed to the amount of hydrogen bond crosslinking present. As PAAm content decreases, the number of hydrogen bond donors decreases as well. However for compositions with high PAAm content (PAAm90-PEGMA10, PAAm75-PEGMA25), the concentration of hydrogen bond crosslinking is expected to increase instead. With the addition of PEGMA, the PEG chains extend fully and become physically entangled with both the covalent and ionic networks. As this happens, it is speculated that hydrogen bonds can form between AAm and PEG units. While this replaces some of the pre-existing PAAm–alginate hydrogen bond crosslinks, a degree of PAAm–H₂O hydrogen bonds are also substituted by these interactions. The amount of hydrogen bond crosslinking thus increases in spite of the total number of hydrogen bonds in the network decreasing. However as the PAAm content becomes low, the effective number of hydrogen bond crosslinks is expected to peak at a particular PEGMA concentration, before decreasing due to the presence of hydrogen bonds becoming too sparse in the network.

The above discussion can explain the peak in mechanical properties, since they are directly related to the amount hydrogen bond crosslinking: the increase in hydrogen bond crosslinking causes a stiffer network to form, thus increasing the elastic modulus; toughness increases as well, as there were more crosslinks to dissipate energy. In this case, the peak is believed to occur before 50 wt% PEGMA, explaining the extremely poor mechanical properties of PAAm50-PEGMA50.

From these results, significantly tougher hybrid hydrogels were obtained from the addition of PEGMA to form a copolymeric network. PEGMA is biocompatible,⁴³ and is thus an ideal addition to the network, since it does not affect the potential utility of the obtained hydrogels.

Variation of ionic crosslinking

In order to study the effects of ionic crosslinking concentration on mechanical properties, the ALG/covalent ratio was varied and represented by Series B. Copolymeric compositions were based on PAAm90-PEGMA10, and results shown in Table S4† and Fig. 6.

Fig. 6 Swelling ratio increased with alginate content due to ionic inter-chain repulsion. (b) Moduli and (c) toughness followed an increasing trend with ionic content as more ionic bonds are now required to be broken per unit strain. *^Toughness values at alginate concentrations of 20 and 33.3 wt% were measured until fracture strain at 69 and 48% strain respectively, which account for the similar toughness values.

Equilibrium swelling ratio increased steadily with ionic crosslink concentration. The increase in swelling ratio was believed to be due to inter-chain ionic repulsion. Alginate comprises negatively charged carboxylate groups which repel one another, thus forming a loose hydrogel network capable of holding larger volumes of water.

While mechanical properties of Series B samples were not compared at identical swelling ratios, the derived trends still show the effect of increasing ionic crosslinking. Contrary to the understood inverse relation between swelling ratio and mechanical properties,¹⁵ B2 and B3 still showed higher moduli and toughness over B1 (Fig. 6b & c). In this case, the higher amount of ionic crosslinks increase network stiffness and improve the energy dissipation mechanism as well. Thus, increasing the ionic crosslinking concentration might result in a tougher hydrogel.

Variation of covalent crosslinker

The effect of covalent crosslinking concentration was studied in Series C, where the molar fraction of covalent crosslinker, BIS, was varied with respect to the number of moles of covalent monomer. Copolymeric compositions were based on PAAm75-PEGMA25. Results are presented in Table S5† and Fig. 7.

Fig. 7 (a) An increase in covalent crosslinker concentration resulted in an overall tighter network. This caused a lower equilibrium swell ratio since the network can only stretch by a certain extent before polymer chains become too taut to hold addition water. Since gel composition was identical across the series, all changes in mechanical properties were believed to be solely caused by this effect. (b) The decreasing swell ratio resulted in an increase in material stiffness. (c) The tighter network also resulted in a reduction in the energy dissipation of the hydrogel, since the ionic network was unable to unzip and re-zip as much as before. *Toughness values for C1 at crosslink concentration of 1.4 mol% was measured until fracture strain at 44% strain.

Equilibrium swelling ratios for Series C samples were observed to decrease with BIS concentration (Fig. 7a). This is in agreement with prior literature.^35,37,44 As crosslinking increased, the polymer network became stiffer, and this reduced its ability to expand and absorb water.

As network composition was constant throughout the series, mechanical properties are expected to only be affected by the equilibrium swelling ratio of the samples. In agreement with the decreasing swelling ratio, the elastic modulus increased steadily with covalent crosslinking concentration due to the lower swelling ratio (Fig. 7b).

On the other hand, compression toughness was observed to be significantly low at 23.7 J m⁻³ for C1, before peaking at 64.7 J m⁻³ at C2, then decreasing steadily (Fig. 7c). As the covalent network became more densely crosslinked, a stiffer network resulted, and this caused less allowance for the network to undergo deformation. Ionic crosslinks were unable to unzip and re-zip to the same extent as networks with covalent crosslinks of lower density. This reduced the ability of the network to dissipate energy, and thus toughness decreased. The significantly low toughness for C1 occurred due to insufficient covalent crosslinks to hold the network, thus resulting in the failure of the material at an strain of 44%.

Conclusion

The design of tough hydrogels is still highly anticipated due to the many potential applications that can be developed from them, and the hybrid hydrogel network holds great promise due to its unique characteristics of high toughness and good recovery. In this paper, alginate and PAAm-based hybrid hydrogels were successfully synthesized by photoinitiation. Compression test results on control samples revealed that a soaking time of 3 days in 5 M CaCl₂ solution was the most ideal, and this was carried out for all subsequent samples. The replacement of PAAm for its non-hydrogen bonding analogue, DMAA, showed the significance of the presence of hydrogen bonds towards the high toughness obtained in the hybrid hydrogels. PAAm/PEGMA copolymerized hybrid hydrogels were then synthesized and tested, and it was found that the addition of PEGMA up to a concentration of 25 wt% improved both modulus and toughness. An increase in ionic crosslinking was also found to improve mechanical properties significantly, while covalent crosslinking concentration was concluded to only have an effect on equilibrium swelling ratio, leading to a resultant changes in mechanical properties due to swelling ratio of the hydrogels. In view of the importance of hydrogen-bonding in causing the high toughness of the hybrid hydrogels, the incorporation of additional forms of physical interactions might also improve toughness, and future experiments can explore the introduction of different polymeric materials capable of other physical interactions such as π–π stacking and host–guest interactions.^45,46 Lastly, the results in this experiment can be used in the optimization of tough hybrid hydrogels in the future.

References

1. T. Inoue, G. Chen, K. Nakamae and A. S. Hoffman, J. Controlled Release, 1997, 49, 167–176.
2. Y. Qiu and K. Park, Adv. Drug Delivery Rev., 2001, 53, 321–339.
3. N. A. Peppas, P. Bures, W. Leobandung and H. Ichikawa, Eur. J. Pharm. Biopharm., 2000, 50, 27–46.
4. Z. Li and X. J. Loh, Chem. Soc. Rev., 2015, 44, 2865–2879.
5. J. Patterson and J. A. Hubbell, Biomaterials, 2011, 32, 1301–1310.
6. J. L. Drury and D. J. Mooney, Biomaterials, 2003, 24, 4337–4351.
7. N. A. Peppas, J. Z. Hilt, A. Khademhosseini and R. Langer, Adv. Mater., 2006, 18, 1345–1360.
8. X. J. Loh, W. C. D. Cheong, J. Li and Y. Ito, Soft Matter, 2009, 5, 2937–2946.
9. A. Chortos and Z. Bao, Mater. Today, 2014, 17, 321–331.
10. A. Chenite, C. Chaput, D. Wang, C. Combes, M. D. Buschmann, C. D. Hoemann, J. C. Leroux, B. L. Atkinson, F. Binette and A. Selmani, Biomaterials, 2000, 21, 2155–2161.
11. F. Yokoyama, I. Masada, K. Shimamura, T. Ikawa and K. Monobe, Colloid Polym. Sci., 1986, 264, 595–601.
12. K. R. Kamath and K. Park, Adv. Drug Delivery Rev., 1993, 11, 59–84.
13. S. Naficy, H. R. Brown, J. M. Razal, G. M. Spinks and P. G. Whitten, Aust. J. Chem., 2011, 64, 1007–1025.
14. P. Calvert, Adv. Mater., 2009, 21, 743–756.
15. X. Zhao, Soft Matter, 2014, 10, 672–687.
16. X. J. Loh, H. X. Gan, H. Wang, S. J. E. Tan, K. Y. Neoh, S. S. J. Tan, H. F. Diong, J. J. Kim, W. L. S. Lee, X. T. Fang, O. Cally, S. S. Yap, K. P. Liong and K. H. Chan, J. Appl. Polym. Sci., 2014, 131 DOI:10.1002/app.39924.
17. P. Thoniyot, M. J. Tan, A. A. Karim, D. J. Young and X. J. Loh, Adv. Sci., 2015, 2, 1400010.
18. J. P. Gong, Y. Katsuyama, T. Kurokawa and Y. Osada, Adv. Mater., 2003, 15, 1155–1158.
19. S. Naficy, J. M. Razal, G. M. Spinks, G. G. Wallace and P. G. Whitten, Chem. Mater., 2012, 24, 3425–3433.
20. J.-Y. Sun, X. Zhao, W. R. K. Illeperuma, O. Chaudhuri, K. H. Oh, D. J. Mooney, J. J. Vlassak and Z. Suo, Nature, 2012, 489, 133–136.
21. K. Y. Lee and D. J. Mooney, Prog. Polym. Sci., 2012, 37, 106–126.
22. M. George and T. E. Abraham, J. Controlled Release, 2006, 114, 1–14.
23. J. Lee and K. Lee, Pharm. Res., 2009, 26, 1739–1744.
24. G. Orive, S. Ponce, R. M. Hernández, A. R. Gascón, M. Igartua and J. L. Pedraz, Biomaterials, 2002, 23, 3825–3831.
25. B. Dupuy, A. Arien and A. P. Minnot, Artif. Cells, Blood Substitutes, Biotechnol., 1994, 22, 71–82.
26. D. F. Matyash Marina, I. Chrysanthy and G. Michael, Tissue Eng., Part C, 2014, 20, 401–411.
27. C. H. Yang, M. X. Wang, H. Haider, J. H. Yang, J.-Y. Sun, Y. M. Chen, J. Zhou and Z. Suo, ACS Appl. Mater. Interfaces, 2013, 5, 10418–10422.
28. O. Smidsrød and G. Skja°k-Br˦k, Trends Biotechnol., 1990, 8, 71–78.
29. I. Donati, S. Holtan, Y. A. Mørch, M. Borgogna and M. Dentini, Biomacromolecules, 2005, 6, 1031–1040.
30. L. H. Zhang, N. R. Brostowitz, K. A. Cavicchi and R. A. Weiss, Macromol. React. Eng., 2014, 8, 81–99.
31. D. Solpan, M. Torun and G. Gueven, J. Appl. Polym. Sci., 2008, 108, 3787–3795.
32. C. B. Xiao, Y. S. Lu, H. J. Liu and L. N. Zhang, J. Macromol. Sci., Part A: Pure Appl.Chem., 2000, 37, 1663–1675.
33. L. H. Gan, G. R. Deen, X. J. Loh and Y. Y. Gan, Polymer, 2001, 42, 65–69.
34. X. J. Loh, G. R. Deen, Y. Y. Gan and L. H. Gan, J. Appl. Polym. Sci., 2001, 80, 268–273.
35. K. S. Anseth, C. N. Bowman and L. Brannon-Peppas, Biomaterials, 1996, 17, 1647–1657.
36. S.-E. Park, Y.-C. Nho and H.-I. Kim, Radiat. Phys. Chem., 2004, 69, 221–227.
37. K. Pal, A. K. Banthia and D. K. Majumdar, Des. Monomers Polym., 2009, 12, 197–220.
38. X. Pang, J. Wu, C.-C. Chu and X. Chen, Acta Biomater., 2014, 10, 3098–3107.
39. ASTM, Standard Test Method for Compressive Properties of Rigid Plastics, D695-10, ASTM International, West Conshohocken, PA, 2010, http://www.astm.org,  DOI:10.1520/D0695-10.
40. D. Eagland, N. J. Crowther and C. J. Butler, Eur. Polym. J., 1994, 30, 767–773.
41. R. Yoshida, K. Uchida, Y. Kaneko, K. Sakai, A. Kikuchi, Y. Sakurai and T. Okano, Nature, 1995, 374, 240–242.
42. K. Tasaki, J. Am. Chem. Soc., 1996, 118, 8459–8469.
43. J.-F. Lutz, J. Polym. Sci., Part A: Polym. Chem., 2008, 46, 3459–3470.
44. B. D. Johnson, D. J. Beebe and W. C. Crone, Mater. Sci. Eng., C, 2004, 24, 575–581.
45. E. A. Appel, J. del Barrio, X. J. Loh and O. A. Scherman, Chem. Soc. Rev., 2012, 41, 6195–6214.
46. X. J. Loh, Mater. Horiz., 2014, 1, 185–195.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra09926a

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