Injectable Hydrogels from Segmented PEG-Bisurea Copolymers

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Injectable Hydrogels from Segmented PEG-Bisurea Copolymers Gajanan M. Pawar,† Marcel Koenigs,† Zahra Fahimi,‡ Martijn Cox,§ Ilja K. Voets,† Hans M. Wyss,‡ and Rint P. Sijbesma*,† †

Laboratory for Macromolecular and Organic Chemistry and Institute for Complex Molecular Systems, ‡Department of Mechanical Engineering and Institute for Complex Molecular Systems, Eindhoven University of Technology, P.O. Box 513, 5600 MB Eindhoven, The Netherlands § Xeltis BV, Eindhoven, The Netherlands S Supporting Information *

ABSTRACT: We describe the preparation of an injectable, biocompatible, and elastic segmented copolymer hydrogel for biomedical applications, with segmented hydrophobic bisurea hard segments and hydrophilic PEG segments. The segmented copolymers were obtained by the step growth polymerization of amino-terminated PEG and aliphatic diisocyanate. Due to their capacity for multiple hydrogen bonding within the hydrophobic segments, these copolymers can form highly stable gels in water at low concentrations. Moreover, the gels show shear thinning by a factor of 40 at large strain, which allows injection through narrow gauge needles. Hydrogel moduli are highly tunable via the physical cross-link density and the length of the hydrophilic segments. In particular, the mechanical properties can be optimized to match the properties of biological host tissues such as muscle tissue and the extracellular matrix.

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INTRODUCTION Hydrogels are widely investigated as biomimetic scaffolds, due to properties such as good biocompatibility and biodegradability, high water content, and excellent permeability for oxygen, nutrients, and metabolites.1−4 They are preferably designed in such a way that they are liquid-like during injection and form networks only in situ; such networks are termed injectable hydrogels.3−9 Injectable gels are known to flow under modest shear stresses and set promptly at the desired location in the absence of flow-induced stress. They can thus be used in minimally invasive surgical procedures and are also favorable for irregularly shaped defects.11−15 Most hydrogels are injected as a low viscosity liquid and subsequently cross-linked in vivo via chemical or physical interaction, which includes chemical linkers,16 enzymes,17 and photoinitiated polymerization,18 or by changes in pH,19 temperature,20 or ionic strength.21 Although there are several advantages of in vivo cross-linking related to the low viscosity during injection, there are also several limitations. These include toxicity of unreacted monomer, precross-linked polymers, and leakage during injection and dilution with body fluids.22 Alternative approaches that avoid these limitations make use of the shear thinning and yielding behavior of hydrogels.5−10 In this strategy, a solid-like hydrogel material is prepared ex vivo with the desired mechanical, morphological, and biological properties and then injected. The stress applied during injection leads to yielding and once flow has started, shear thinning may further facilitate injection. Once, the shear stress stops the material forms a solid-like hydrogel. As a result, the properties of the injected gel material © XXXX American Chemical Society

in vivo are similar to those of the original ex vivo hydrogel. Shear thinning has the additional advantage of giving rise to shear banding or even plug flow, which limits shear forces to a narrow region close to the needle wall, thus, preventing shearinduced damage to cells taken up in the gel.5 Recently, Guvendiren et al.6 reviewed shear-thinning hydrogels utilized for biomedical applications. They categorized these gels chemically into five main groups: (i) peptide-based hydrogels, (ii) protein-based hydrogels, (iii) hydrogels from blends, (iv) colloidal systems, and (v) hydrogels based on cyclodextrins and block copolymers. Here we introduce shear thinning hydrogels from segmented copolymers that are completely synthetic and have several advantages over hydrogels based on natural polymers. Various types of natural or seminatural polymers are known as biomimetic scaffolds for biomedical applications. These include cellulose,23 chitosan,7 xyloglucan,24 gelatin,25 and their derivatives, and synthetic polymers, including polyethers,30 block copolymers of polyethers and polyesters, and synthetic polypeptides.26−29,63 Naturally derived polymers have the potential advantage of biocompatibility and cell adhesion.34 However, they generally express limited mechanical properties. Moreover, large-scale synthesis of these materials is cumbersome and difficult. Therefore, simple and scalable syntheses are highly desirable for potential applications. Received: August 6, 2012 Revised: November 13, 2012

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Figure 1. Architecture of the segmented copolymers used in this work. General Methods. All reactions were performed under a nitrogen atmosphere in a glovebox (MBraun LabMaster 130, MBraun, Garching, Germany) to achieve high molecular weight segmented copolymer. NMR spectra were acquired on a 400 MHz Varian Mercury Vx (400 MHz for 1H NMR, 100 MHz for 13C NMR). Proton and carbon chemical shifts are reported in ppm downfield of tetramethylsilane using the resonance of the deuterated solvent as internal standard. Splitting patterns are designated as singlet (s), doublet (d), triplet (t), multiplet (m), and broad (b). Infrared spectra were measured on a Perkin-Elmer 1600FT-IR. GPC and MALDI-TOF Analysis. GPC in dimethyl formamide (DMF) was performed on a PL-GPC 50 Plus of Polymer Laboratories with integrated refractive index detector, using a Polymer Standards Service Gram Linear M 8 × 300 mm, 10 μm particles column, and DMF with 10 mM LiBr as eluent at a flow rate of 1 mL/min (50 °C). Polyethylene glycol standards were used for calibration at 50 °C. Matrix-assisted laser desorption ionization time-of-flight mass spectroscopy (MALDI-TOF) was performed on a Perseptive DE PRO Voyager MALDI-TOF mass spectrometer using α-cyano-4-hydroxycinnamic acid as the calibration matrix. DSC Analysis. Differential scanning calorimetry measurements were performed on a Thermal Advantage Q2000 apparatus between −80 and 250 °C at a rate of 10 °C/min with a sample weight of 5−10 mg. Integration of the melting endotherm was performed with the TA Instruments Universal Analysis software. Rheology. Mechanical properties of these hydrogels were tested by rheology. Dynamic viscoelastic measurements were determined using a stress-controlled rheometer (Anton Paar, Physicia MCR501) equipped with a sand-blasted plate−plate geometry to prevent slippage. Measurement temperature was fixed at 15 °C unless stated otherwise and mineral oil was used to prevent evaporation. Small-Angle X-ray Scattering (SAXS). Small-angle X-ray scattering (SAXS) experiments were performed at the European Synchrotron Radiation Facility (ESRF) in Grenoble, France, at the high brilliance beamline ID02. An X-ray energy of 12.46 keV and one sample-todetector distance of 1.5 m was used to cover a q-range of 0.07 < q < 3.5 nm−1, with q being the magnitude of the scattering wave vector. The samples were contained in an aluminum plate with small holes sandwiched between Kapton tape at a temperature of 20.1 °C. The scattering data were corrected for background scattering, detector response and primary beam intensity fluctuations. The instrument scattering vector was calibrated using a silver behenate standard. Peak positions were determined by taking the midpoint between the local extremes in the first derivative of the scattering intensities. Atomic Force Micrographs. Atomic force micrographs were recorded under ambient conditions with silicon cantilever tips (PPPNCH, 300−330 kHz, 42 N/m from Nanosensors) using an Asylum Research MFP-3D-Bio in noncontact mode. For the atomic force microscopy, 2 μL of the solution at ambient temperature was drop-cast on freshly cleaved mica. Biocompatibility/Toxicity/Viability. The viability of human myofibroblast cells in the presence of the segmented copolymer was studied with XTT based In Vitro Toxicology Assay Kit, which was purchased from Sigma-Aldrich. Synthesis of PEG-Bisurea Segmented Copolymer. Bis(N-(tertbutyloxycarbonyl)-11-aminoundecanoyl)-poly(ethyleneglycol). In a 250 mL two-neck round-bottom flask 2.26 g (7.49 mmol) of N-(tertbutyloxycarbonyl)-11-aminoundecanoic acid, 1.44 g (7.49 mmol) of

The ideal hydrogel for tissue engineering combines biocompatibility, biodegradability, and allows tuning of its mechanical properties by control over the morphology.34−39 To have the most predictable structure−property relationship, structurally monodisperse segments are desirable as building blocks. Furthermore, it is useful to have “click on” functionality in order to facilitate specific recognition and binding of bioactive moieties and probes for tracing and analysis.31−33 Here we introduce bisurea-PEG segmented copolymers and show that they can provide many of the desired features of injectable hydrogels for tissue engineering. Recently, we described a powerful tool for the preparation of segmented copolymers possessing uniform (monodisperse) hard blocks based on urea groups and poly(tetrahydrofuran) (pTHF) containing soft blocks for developing thermoplastic elastomers.40,41 Following this concept, here we describe the preparation of injectable, biocompatible, and elastic hydrogels from a segmented copolymer with hydrophobic and hydrophilic segments. The structure of the present segmented copolymer combines a hydrophilic and swellable domain consisting of polyethylene glycol (PEG) with a hydrophobic, biodegradable, crystallizable, and nonswellable domain containing two urea groups (Figure 1). Urea groups are known to self-associate via hydrogen bonding, and this self-assembling motif has been used in the development of gelation agents.42−50 The role of the hydrogen bonding groups is 2-fold: First, they provide additional interactions between hydrophobic segments, increasing the strength and lifetime of the physical cross-links. Second, they are recognition sites for binding of functional molecules that are provided with two urea groups. In the segmented copolymers used in this study, the hydrogen bonding groups are separated from the PEG segments by 10 CH2 groups (methylene groups) to prevent the PEG from interfering with hydrogen bonding among the urea groups.44 They have the capability to gel in water via multiple hydrogen bonding. The mechanical behavior of our hydrogels was tested by oscillatory shear experiments and dynamic strain amplitude tests. Atomic force microscopy and small-angle X-ray scattering were used to characterize the morphology of the gels at small length scales. The viability of human myofibroblast cells in the presence of the segmented copolymer was studied with a XTT (2,3-bis-(2-methoxy-4nitro-5-sulfophenyl)-2H-tetrazolium-5-carboxanilide) assay.

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EXPERIMENTAL SECTION

Materials and Methods. Materials. Solvents used in synthesis were reagent grade. CH2Cl2, CHCl3, Et3N, and pyridine were distilled from CaH2. All PEG derivatives were dried in vacuum over P2O5 during at least 12 h. The reagents 11-aminoundecanoic acid, 1,4diisocyanatobutane, and 1,6-diisocyanatohexane were purchased from Aldrich, Fluka, or Acros and were used without additional purification. 11-tert-Butoxycarbonylamino-undecanoic acid was prepared according to literature procedures.51 B

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Scheme 1. Synthesis of PEG-bisurea Segmented Copolymer

N-(3-dimethylaminopropyl)-N-ethylcarbodiimide hydrochloride and 0.92 g (7.48 mmol) of dimethylaminopyridine were stirred in 100 mL of dry dichloromethane under argon for 20 min. To the resulting solution was added 15 g (1.87 mmol) of poly(ethylene glycol) (Mn ca. 8000) and the reaction mixture was stirred for 3 days. The solution was washed with 10% citric acid and precipitated by addition of diethyl ether, filtered off, and dried. Bis(N-(tert-butyloxycarbonyl)-11-aminoundecanoyl)-poly(ethyleneglycol) (8k). 1H NMR (400 MHz, CDCl3) δ 4.22 (t, 4H, CH2OCO), 3.8−3.3 (m, OCH2), 3.08 (br, 4H, CH2N), 2.31 (t, 4H, CH2CO), 1.60 (m, 4H, CH2CH2CO), 1.43 and 1.26 (s, 46H, C(CH3)3 and CH2). MALDI-TOF [M + Na+] = 9642.83 ± n × 44. GPC (CHCl3; PS standards): Mw = 12900 g/mol, PDI = 1.28. Bis(N-(tert-butyloxycarbonyl)-11-aminoundecanoyl)-poly(ethyleneglycol) (20k). 1H NMR (400 MHz, CDCl3) δ 4.21 (bs, 4H, CH2OCO), 3.8−3.3 (m, OCH2), 3.10 (br, 4H, CH2N), 2.31 (t, 4H, CH2CO), 1.61 (m, 4H, CH2CH2CO), 1.43 and 1.26 (s, 46H, C(CH3)3 and CH2). Bis(11-amonium chloride undecanoyl)-poly(ethylene glycol). To 30 mL of a 4 M HCl solution in dioxane was added to the solution of 10 g (1.1 mmol) bis(N-(tert-butyloxycarbonyl)-11-aminoundecanoyl)poly(ethyleneglycol) in 30 mL of dioxane and stirred at 0 °C for 1 h and subsequently at room temperature for 12 h. The solvent was evaporated to yield 9.0 g (100%) of the product as its hydrochloric salt. Bis(11-amonium chloride undecanoyl)-poly(ethylene glycol) (8k). 1 H NMR (400 MHz, CDCl3, T = 295K): δ = 7.78 (bs, 6H, NH3), 4.22 (bs, 4H, CH2OCO), 3.90−3.20 (m, OCH2), 2.89 (bs, 4H, CH2N), 2.32 (t, 4H, CH 2 CO), 1.75−1.48 (m, 8H, CH 2 CH2 N and CH2CH2CO), 1.28 (bs, 28H, CH2). MALDI-TOF [M + Na+] = 9350.27 ± n × 44. Bis(11-amonium chloride undecanoyl)-poly(ethylene glycol) (20k). 1H NMR (400 MHz, CDCl3, T = 295 K) δ 7.76 (bs, 6H, NH3), 4.24 (bs, 4H, CH2OCO), 3.90−3.20 (m, OCH2), 2.92 (bs, 4H, CH2N), 2.34 (t, 4H, CH2CO), 1.75−1.55 (bs, H2O + CH2CH2N + CH2CH2CO), 1.28 (bs, 26H, CH2). Poly(8kU4U). To a solution of 100.0 mg (0.58 mmol) of 1,4diisocyanatobutane in 1 mL of dichloromethane, solution of bis(11aminoundecanoyl)-poly(ethylene glycol) (5.0 g, 0.55 mmol) and triethylamine (0.1 mL, 0.96 mmol) in 20 mL of dichloromethane was added and stirred for 3 days. The solution was concentrated and precipitated by addition of diethyl ether, filtered off, and dried. Yield: 4.1 g (80%). 1H NMR (400 MHz, CDCl3, T = 295 K) δ 5.01 and 4.80

(bs, 2H, NH), 4.23 (bt, 4H, CH2OCO), 3.90−3.10 (m, OCH2), 3.15 (bm, 8H, CH2N), 2.32 (t, 2H, CH2CO), 1.65−1.40 (m, 12H, CH2CH2N and CH2CH2CO), 1.27 (bs, 24H, CH2). 13C NMR (100 MHz, CDCl3) δ 173.7, 158.7, 71.6, 70.5, 69.4, 69.1, 63.3, 40.4, 39.7, 34.1, 30.3, 29.4, 29.2, 29.1, 29.0, 27.5, 26.8, 24.8. FT-IR (ATR mode, cm−1) ν 2880, 2742, 1733, 1620, 1573, 1466, 1352, 1342, 1279, 1241, 1145, 1100, 1060, 962, 841. GPC (DMF; PEG standards): Mw = 160000 g/mol, PDI = 1. 90. Synthesis of Poly(8kU6U), Poly(20kU4U), and Poly(20kU6U). Poly(8kU6U), poly(20kU4U), and poly(20kU6U) were synthesized according to the procedure of poly(8kU4U). Poly(8kU6U). Yield: 85%. 1H NMR (400 MHz, CDCl3, T = 295 K) δ 4.76 (bd, 2H, NH), 4.22 (bt, 4H, CH2OCO), 3.90−3.33 (m, OCH2), 3.14 (bm, 8H, CH2N), 2.32 (t, 2H, CH2CO), 1.65−1.48 (m, 14H, CH2CH2N and CH2CH2CO), 1.27 (bs, 35H, CH2). 13C NMR (100 MHz, CDCl3) δ 173.7, 158.7, 71.6, 70.5, 69.4, 69.1, 67.2, 63.3, 40.3, 39.3, 34.1, 30.3, 29.8, 29.4, 29.2, 29.1, 29.0, 26.8, 24.8. FT-IR (ATR mode, cm−1) ν 3386, 2877, 1734, 1644, 1567, 1466, 1359, 1342, 1280, 1241, 1145, 1097, 1060, 961, 841. GPC (DMF; PEG standards): Mw = 134500 g/mol, PDI = 1.55. Poly(20kU4U). Yield: 85%. 1H NMR (400 MHz, CDCl3, T = 295 K) δ 4.19 (bt, 4H, CH2OCO), 4.04−3.24 (m, OCH2), 3.15 (bm, 8H, CH2N), 2.29 (b, H2O + CH2CO), 1.65−1.35 (m, 11H, CH2CH2N and CH2CH2CO), 1.27 (bs, 19H, CH2). FT-IR (ATR mode, cm−1) ν 2882, 1723, 1619, 1574, 1466, 1359, 1342, 1279, 1241, 1147, 1102, 1060, 962, 842. GPC (DMF; PEG standards): Mw = 86000 g/mol, PDI = 1.46. Poly(20kU6U). Yield: 75%. 1H NMR (400 MHz, CDCl3, T = 295 K) δ 4.22 (bt, CH2OCO), 3.98−3.20 (m, OCH2), 3.12 (bm, CH2N), 2.32 (t, CH2CO), 1.70−1.05 (m, CH2CH2N and CH2CH2CO, CH2). FT-IR (ATR mode, cm−1) ν 2879, 1650, 1549, 1466, 1359, 1342, 1279, 1241, 1146, 1099, 1060, 961, 841. GPC (DMF; PEG standards): Mw = 92000 g/mol, PDI = 2.05.

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RESULTS Synthesis of the Segmented Copolymer. The preparation of segmented copolymers was achieved by the step-growth polymerization of amino-terminated PEG and aliphatic diisocyanate. The general synthetic protocol of segmented copolymers is outlined in Scheme 1. In this multistep synthesis, Boc-protected amino acid was reacted with polyethylene glycol (8000 and 20000 g·mol−1), yielding bifunctionalized poly-

C

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concentration between 4 and 5 wt %. Above this characteristic concentration, the material behaves solid-like, with the storage modulus always larger than the loss modulus, G′ > G″, as shown in Figure 2b. Below 5 wt %, the magnitude of the moduli is dramatically lower, dropping by around 3 orders of magnitude compared to the 5 wt % sample. Nevertheless, for the low concentration samples, the storage modulus is still comparable or larger than the loss modulus, as accessed at an oscillation frequency of 1 rad/s. This indicates that even below 5 wt % the sample still has solid-like properties. However, it is difficult to discern the physical origin of this elastic-like response from the response at a single time scale, measured at just one single oscillation frequency. The frequency dependence of the storage (G′) and the loss (G″) modulus in the linear viscoelastic regime was measured by applying a 1% strain in the angular frequency range of 100 to 0.1 rad/s. In the linear regime, at a concentration of 10 wt %, gels formed from all our polymers poly(8kU6U), poly(8kU4U), and poly(20kU6U) exhibited solid-like rheological behavior, as shown in Figure 2c. At the same concentration, the elastic modulus of poly(8kU6U) was higher than poly(20kU6U). We hypothesize that this is because fewer physical cross-linking domains are present in poly(20kU6U). For all our gels, we find that the storage modulus was larger than their loss modulus and depended only weakly on frequency over the entire range of frequencies studied, a behavior typical for chemical or strongly interacting physical gels. However, the loss modulus shows a minimum at intermediate frequencies, indicating the presence of a structural relaxation process at very low frequencies, below the range accessible in our experiments. The significant changes in viscoelastic properties observed just below 5 wt % suggest that a qualitative change in the morphology of the gels is taking place at this characteristic concentration. We used atomic force microscopy (AFM) to investigate the surface morphology of 5 wt % poly(8kU4U) hydrogels (Figure 3). Samples were prepared in a wetted state to minimize the effect of drying on morphology. Height and phase mode images of the same sample showed that characteristic dimensions of the aggregates in the wet gel is in the order of 20−35 nm. In the height mode, no features with high aspect ratio are present. In the phase mode, elongated features are observed, but they are only present at the interface of the domains visible in height mode. Therefore, we conclude that, although the presence of aggregates of the hydrophobic domains can be deduced from the rheological properties of the gels, their shape (fibrous or globular) cannot unambiguously be determined from the AFM images. Further information on the structure of the hydrogels was obtained from small-angle X-ray scattering. The experiments were performed at the high brilliance beamline ID02 at the European Synchrotron Radiation Facility (ESRF) in Grenoble, France. To this end, samples were sandwiched between Kapton tapes glued onto two opposite sides of an aluminum sample holder with several small holes of a few mm’s in diameter that served as sample container. The scattering profiles of 10 wt % gels of poly(8kU4U), poly(8kU6U), and poly(20kU6U) are given in Figure 4a. A correlation peak was observed at intermediate q-values in all samples, signifying the existence of a characteristics distance, d = 2π/qmax in the material, where q is the magnitude of the scattering wave vector and qmax is the peak position in q-space. We attribute this characteristic distance to

ethylene glycol I. Next, Boc groups were removed by treatment with 1 M HCl solution in 1,4-dioxane to obtain telechelic amino-terminated PEG material II. Finally, the corresponding PEG material II was copolymerized with tetramethylene or hexamethylene diisocyanate to develop segmented copolymer III. To achieve high molecular weight and for stoichiometry control, all polymerizations were performed under a nitrogen atmosphere in a glovebox. All segmented copolymers were fully characterized by 1H and 13C NMR and IR spectroscopy. Weight average molecular weights (Mw) of these hydrogels were determined by using gel permeation chromatography (GPC) in DMF at 50 °C. The highest values of Mw that could be obtained for poly(8kU6U) were Mw = 134500 g/mol (PDI = 1.55), those for poly(8kU4U), poly(20kU6U), and poly(20kU4U) Mw = 160000 g/mol (PDI = 1.90), Mw = 92000 g/ mol (PDI = 2. 05), and Mw = 86000 g/mol (PDI = 1.46), respectively. Hydrogen bonding in the polymers was investigated using differential scanning calorimetry (DSC). The DSC measurements of the polymers showed two transitions (Table 1). The first transition (around 50 °C) Table 1. DSC Measurements of Polymers transition 1

transition 2

compounds

ΔT (°C/min)

T1 (°C)

ΔH1 (J/g)

T2 (°C)

ΔH2 (J/g)

poly(8kU6U) poly(20kU6U) poly(8kU4U) poly(20kU4U)

20 20 20 20

46.81 42.85 53.76 63.76

102.8 111.6 119.4 136.0

176.0 160.0 221.1 233.1

10.74 5.23 25.31 0.89

corresponds to the melting point of the PEO segments, whereas the second transition state is assigned to melting of the hydrogen-bonded hydrophobic segments, in analogy with a similar transition in bisurea pTHF segmented copolymers.35 Mechanics and Structure as a Function of Polymer Concentration. The segmented copolymers have the capability to form gels in water. Hydrogels from poly(8kU4U) were prepared by dissolving 3−7 wt % of the segmented copolymer in deionized water under sonication at 45 °C followed by cooling to room temperature. The mechanical properties of the resulting material depend strongly on the polymer concentration. The simplest test of the gel-like properties of a material is the tilting vial method, shown in Figure 2a. We observed that the poly(8kU4U) samples at 3 and 4 wt % do not form gels strong enough to support their weight in the vials, whereas the samples above 5 wt % form strong gels that did not flow in the vials. We thus find a critical concentration between 4 and 5 wt % for the tilted vial method, above which the material no longer flows when the vial is tilted over extended periods of time, thus indicating the formation of a solid-like gel network. To investigate the viscoelastic properties of the formed materials more closely, we also performed oscillatory rheological measurements on gels that, after sonication, were kept overnight at room temperature to reach an equilibrium state. All polymers were kept in the rheometer for at least 2 h before starting the measurements. Figure 2b shows the storage and loss modulus G′ and G″ of the system as a function of polymer concentration, as accessed at a oscillation frequency of 1 rad/s and a strain amplitude of 1%. In agreement with the tilting vial method, these measurements show a dramatic change in the magnitude of G′ and G″, which occurs at a D

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Figure 2. Mechanical properties of hydrogels of poly(8kU4U) as probed by (a) tilting vial method and (b) oscillatory rheology at 293 K (1 rad/s, strain amplitude of 1%). (c) Frequency dependent oscillatory rheology at a concentration of 10 wt % for poly(8kU4U) (G′ (■), G″ (□)), poly(8kU6U)(G′ (●), G′′ (○)), and poly(20kU6U) (G′ (⧫), G′′ (◊)).

Figure 3. Height image (a) and phase image (b) of the wet hydrogels of 5 wt % poly(8kU4U).

The smaller value for d in the second entry of Table 2 compared to the first entry is tentatively ascribed to the more hydrophobic U6U segments in the former compound. As expected, the correlation peak is more pronounced and is found at higher q-values in samples prepared at the highest concentrations; for the 15 wt % sample, a strong peak is observed, indicating a well-defined characteristic length scale in the structure of the material. The length scale indicates that the PEG spacers set the characteristic length d. However, at concentrations as high as those of our gels, the available form

the spacing between hydrophobic domains within the hydrogel, which are dense in physical cross-links.49 The position of the correlation peak/shoulder is dependent on the molecular structure of the gel building blocks and their concentration; indeed, as shown in Figure 4b, with increasing concentration of poly(8kU4U) we observe a systematic shift of the peak position toward higher q-values, corresponding to a decreasing characteristic length d. The characteristic length d systematically increases as the length of the poly(ethylene glycol) chain is increased, as shown in Figure 4a and Table 2. E

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Figure 4. Small angle X-ray scattering profiles (I(q) vs q) of (a) 10 wt % hydrogels of samples of poly(8kU4U)160k (■), poly(8kU6U)134k (○), and poly(20kU6U)92k and (b) 4, 8, 10, and 15 wt % samples of poly(8kU4U)160k.

hydrogels using a syringe equipped with a thin 29 gauge needle (inner diameter dneedle = 0.184 mm). As shown in Figure 5, the 10 wt % poly(8kU4U) hydrogel could readily be injected through this narrow needle by applying a gentle force of (∼12− 20 N) to the piston of a syringe with D = 4 mm inner diameter. By qualitatively probing the material manually with the needle tip, we noted that the material appears solidlike almost immediately after exiting the needle. This simple experiment shows directly that indeed the material can be transferred through a thin syringe needle using realistic forces applied to the syringe piston. This requires two separate phenomena to occur in the syringe: First, the stress applied to the material has to be large enough to lead to a yielding of the material, thereby transferring its behavior from solidlike to liquidlike. Second, the viscosity of the material should be small enough to enable transfer of the now fluidlike material at high enough flow rates and at reasonable levels of force applied to the syringe. The yield stress was estimated by monitoring the lowest level of applied force that leads to a observable flow of material though the syringe. This critical force was determined to be F ≈ 12 N, giving rise to a pressure of p ≈ (4F)/(πD2) in the syringe; this is also a fair estimate for the typical shear stress expected initially at the inlet of the syringe needle. From the

Table 2. Characteristic Distance d = 2π/qmax, as Determined from the Scattering Profiles Depicted in Figure 4a hydrogel

qmax (nm−1)

d (nm)

poly(8kU4U)160k, 10 wt % poly(8kU6U)134k, 10 wt % poly(20kU6U)92k, 10 wt %

0.26 0.35 0.17

24.2 17.9 37.0

and structure factor models break down, thus, preventing us from obtaining more detailed information on the structure of the material. Using AFM and SAXS, it was not possible to determine whether the structure consists of flowerlike spherical objects or of rodlike objects. However, the mechanical properties of the material ex vivo appear suitable for biomedical applications. Therefore, we investigated whether the material is indeed suitable for use as an injectable hydrogel, as discussed in the subsequent sections. Injectability and Flow Behavior. We started with the simplest possible experiment, directly testing whether the material can be injected via a small diameter syringe needle. One key requirement for a suitable injectable material is the size of the needle that could be used in the injection process. For patient comfort and maximum control, small diameter needles are preferable. We thus tested the injectability of our

Figure 5. Transfer of a 10 wt % poly(8kU4U) hydrogel though a narrow 29 G syringe needle. F

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critical force, a critical yield stress of around σyield ≈ (4F)/(πD2) = (4 × 12 N)/(π(4 mm)2) ≈ 235 kPa is estimated. A rough estimate of the viscosity of the material during transfer through the needle was obtained by measuring the approximate flow rate of material through the needle at constant applied force. During application of a constant force of F ≈ 17 N to the syringe piston, a fluid flow rate of around Q ≈ 0.184 μL/s was obtained. A typical value for the viscosity can then be obtained directly from the applied pressure p ≈ (4F)/ (πD2), the flow rate, the length L = 20 mm, and the diameter dneedle of the syringe needle as67 η=

behavior. As a result, while both poly(8kU4U) and poly(8kU6U) show similar elasticlike properties at low strains at a strain amplitude of 100%, the viscoelastic moduli of poly(8kU6U) are significantly lower than those of poly(8kU4U). In summary, for all materials studied, we observe the onset of yielding from a solidlike to a liquidlike behavior occurring at a critical strain of less than 100%. However, the strain level where the material truly becomes liquidlike with G″ < G′ should occur only at larger strains. We thus estimate the critical yield stress needed to fluidize the material to be larger than σyield > G′γc ≈ 20 kPa. Indeed, from the simple syringe experiment, a yield stress of around 200 kPa was estimated, significantly larger than the stress at the edge of the linear viscoelastic regime, where still G′ > G″. After the sample has yielded and is transformed into a liquidlike state, the viscosity of the material should also be low enough to enable material to flow with a sufficient flow rate at reasonable levels of applied pressure. From the syringe experiment, we had seen that this is indeed possible and had estimated a viscosity of around 6 Pa s. Remarkably, from the magnitude of the loss modulus in the oscillatory rheological measurements a much higher level of viscosity would be expected; the Cox−Merz rule suggests that steady shear measurements are directly linked to oscillatory measurements and that the viscosity is well approximated by η(γ̇ = 1/ωγ) ≈ G″(ω)/ω. Given the oscillatory measurements of the poly(8kU4U) hydrogel, for a value of 2000 Pa for the storage modulus, a strain of 100% and a frequency of 1 rad/s, a viscosity of around 2000 Pa s is expected. This value is more than 2 orders of magnitude larger than the 6 Pa s estimated from the syringe experiment. This dramatic difference indicates that the viscosity of the material is significantly reduced as the material starts to flow through the needle. To investigate this behavior in more detail, we perform viscosity measurements on the poly(8kU4U) hydrogel as a function of the applied shear rate. Indeed, as shown in Figure 7,

πdneedlep ≈ 6 Pa s 64LQ

For injectability, a material that exhibits yielding and shearthinning behavior is highly beneficial, as during the injection phase the material is transformed into a fluidlike state. Studying the nonlinear response to large applied strains and strain rates is, thus, of key importance. Therefore, the viscoelastic response of the hydrogels was measured as a function of the applied strain. Strain-dependent measurements were performed at fixed angular frequency (1 rad/s) by varying the strain amplitude from 1 to 100%, shown in Figure 6. The linear viscoelastic

Figure 6. Yielding behavior in oscillatory rheological measurements. Strain sweeps at a frequency of 1 rad/s for poly(8kU4U) (G′ (■), G″ (□)), poly(8kU6U) (G′ (●), G′′ (○)), and poly(20kU6U) (G′ (⧫), G′′ (◊)), performed at 10 wt % concentration.

regime, where the storage and loss modulus is independent of the applied strain amplitude of the material, is limited to strains below a critical yield strain. At larger strains, the material shows a yielding behavior with both the storage and the loss modulus decreasing with increasing strain amplitude. In this regime also, the response becomes nonharmonic, so strictly speaking the viscoelastic moduli are no longer properly defined. The values represented here are based solely on the first harmonic of the stress response; nevertheless, we can clearly identify a yield strain for the material. The decrease in these moduli above a critical level of strain indicates a transition from a solidlike to a more fluidlike material response for all hydrogel materials studied. For poly(8kU4U) and poly(20kU6U), this yield strain is around 30%, where the onset toward a purely liquidlike behavior at larger strains is observed. However, the most pronounced yielding behavior is observed for poly(8kU6U), where the yield strain is only around 10% and at the largest accessed strain the material already shows a clearly fluidlike

Figure 7. Viscosity as a function of shear rate for 10 wt % poly(8kU4U). Significant shear thinning is observed, with the viscosity decreasing by almost 2 orders of magnitude upon increasing the shear rate from 1 to 1000 s−1. A significant degree of syneresis is observed between consecutive runs with increasing (full squares) and decreasing shear rates (open circles). A simple estimate of the shear rate and viscosity in the syringe injection experiment (marked as a red star) agrees remarkably well with the viscosity measurement. G

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Figure 8. (a) Storage and loss modulus of 50k (G′ (●), G′′ (○)) and 160k (G′ (⧫), G′′ (◊)) molecular weight bearing 10 wt % poly(8kU4U). (b) Storage and loss modulus of 5 wt % poly(8kU4U) (G′ (●), G′′ (○)), 6 wt % poly(8kU4U) (G′ (⧫), G′′ (◊)), 10 wt % poly(8kU4U) (G′ (■), G′′ (□)), and 15 wt % poly(8kU4U) (G′ (★), G′′ (☆)).

Figure 9. Dynamic strain amplitude test of hydrogels of poly(20kU6U) (10 wt %) for four cycles.

we observe a significant degree of shear thinning of the material. The viscosity decreases by almost 2 orders of magnitude upon increasing the shear rate from 1 to 1000 s−1. A significant degree of syneresis is observed between consecutive runs with increasing and decreasing shear rates. The values of viscosities are indeed much smaller than those estimated from the oscillatory measurements using the Cox− Merz rule. This observation as well as the dramatic shear thinning behavior indicate that the structure of the material is changing significantly under shear. To make a direct comparison with the simple syringe experiment, the typical shear rate in the syringe needle was estimated to be γ̇ = 8Q/πd3needle ≈ 32 s−1. The estimated viscosity of 6 Pa s is shown as a red star in the same figure with rate-dependent viscosity measurements. A remarkably good agreement is found, indicating that the rate-dependent viscosity measurements are suitable for predicting the injectability of the material. It is, thus, the shear thinning behavior of our hydrogel materials that enables their use as an injectable hydrogel material; in the absence of such dramatic shear thinning, a force on the order of kilo-Newtons would have to be applied to obtain a comparable flow rate.

To further investigate the mechanical properties of bisureabased hydrogels, we proceeded to study the influence of molecular weight on hydrogel stiffness. To do so, we prepared 50k and 160k molecular weight bearing poly(8kU4U) segmented copolymer hydrogels and studied their mechanical properties by using rheology. As shown in Figure 8a, the effect of the overall molecular weight on the viscoelastic moduli of the hydrogel is small. The effect of concentration was also investigated using oscillatory rheological measurements, as shown in Figure 8b, for poly(8kU4U) at concentrations of 5 wt %, 6 wt %, 10 wt %, and 15 wt %. As expected, the viscoelastic moduli of the hydrogels clearly increase with increasing concentration. This is due to both the increase in chain density and cross-link density with increasing concentration. Further, the yield strain of the material moves to lower strains as the concentration of poly(8kU4U) is increased, indicating that the polymer gel becomes more brittle. Self-Healing of Gels. The strain-dependent measurements show that the material exhibits a yield strain and transforms to a more liquidlike behavior at larger deformations, a behavior essential for injectable gels. However, it is also important that H

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Figure 10. Dynamic strain amplitude test of poly(8kU4U) (10 wt %) for four cycles.

these materials can recover or heal from a fluidlike state back to a solidlike gel state once the deformation is removed. To characterize this behavior we performed rheological tests targeted toward quantifying the recovery from a yielded, weakly viscoelastic state to a recovered, strongly solidlike state. Oscillatory rheological measurements were performed where periods of low applied strain amplitude (γ = 0.1%) were alternated with periods where the applied strain amplitude is well outside the linear viscoelastic regime (γ = 100%), shown in Figure 9. Recovery of the viscoelastic properties was studied for poly(20kU6U) hydrogels at 10 wt %; the initial G′ in the yielded state was around 1 kPa. As the low strain cycle (γ = 0.1%) was initiated, G′ rapidly increased to around 4 kPa, a 4fold increase. In turn, upon repeated initiation of the high strain, cycle, G′ rapidly decreased within 20 s to its initial value of 1 kPa. The yielding and recovery behavior remains unchanged after many cycles, four cycles are shown in Figure 9, indicating that yielding and recovery in these materials is a completely reversible process. However, studying yielding behavior of some of the hydrogel materials by applying a large strain was difficult, because they form brittle gels that break into pieces and are ejected from the measuring geometry when a large deformation is applied. The yielding and recovery behavior of these samples was therefore investigated by applying a series of cycles where both the temperature and the strain amplitudes were varied, as shown in Figure 10 for the poly(8kU4U) 10 wt % sample. At elevated temperatures the material becomes softer and less brittle; therefore it can be deformed at large strain amplitudes without breaking up inside the rheometer. Thus, in our experiments, the hydrogel was first heated to 50 °C for 400 s and then 100% strain at 1 rad/s was applied at the same temperature. The recovery behavior of the hydrogels was then monitored in a cycle at 1% strain and 1 rad/s frequency at a temperature of 15 °C. Again, results indicate that the storage modulus recovers very rapidly from a yielded state, here achieved by a combination of large strain deformation and elevated temperature. Such recovery after mechanical yielding is often referred to as “self healing”. In the recovery measurements shown in Figure 10, the initial storage modulus was observed to be around 20 kPa at low strain amplitude (γ = 1%) and low temperature (15 °C). Upon increasing the temperature from 15 to 45 °C the storage

modulus rapidly drops to around 1 kPa. Applying a large strain (100%) at 50 °C causes another decrease by half. By returning the strain to small values (1%) and returning to a temperature of 15 °C, the storage modulus recovers quickly within 80s to its initial value, a 40-fold increase. The rapid recovery of hydrogel strength was repeatedly confirmed for four cycles and in each cycle it recovers to almost 97% of its initial value, indicating that the gel structure after healing is not significantly affected by the breakup and recovery process.

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BIOCOMPATIBILITY/TOXICITY/VIABILITY An XTT assay on the poly(8kU6U) has been performed to investigate the toxicity of the material. According to the ISOnorm 10993−5:2009, the viability of human myofibroblast cells treated with extract of the material must be above 70%. As can be seen in Figure 11, the segmented copolymer showed excellent biocompatibility with cells.

Figure 11. XTT of poly(8kU6U).

For the XTT measurements, the blank was a XTT solution incubated without cells. The positive control indicates healthy cells treated with normal medium (treated similar to the extract of the material). The negative control was cells that were exposed to the latex (which is highly toxic to the cells). A piece of 25 × 25 mm (23.29 mg) was used for the high concentration study and a strip of 5 × 25 mm (4.95 mg) was used for the low I

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concentration. For both concentrations we found a viability of around 100%, thus, indicating that the material is not cytotoxic and should indeed be suitable as an injectable material.

Article

ASSOCIATED CONTENT

S Supporting Information *

1

H and 13C NMR and DSC data of the materials. This material is available free of charge via the Internet at http://pubs.acs.org.

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DISCUSSION Most of the synthetic hydrogels studied for injectable materials in tissue engineering are based on peptide amphiphiles.52,53,55−61,63 However, the synthesis of the peptide amphiphiles limits the scalability. Synthesis of the PEG-bisurea segmented copolymers is straightforward and can be performed on a large scale. Moreover, the moduli of the hydrogels can be controlled to match the properties of biological host tissues by the length of the hydrophilic segments which influences density of physical cross-links, as is seen from comparison of moduli of poly(8kU6U) and poly(20kU6U). These features bear a similarity to the PEG-based supramolecular networks recently reported by Dankers et al.3 The PEG-bisurea-based hydrogels show rheological characteristics of physically cross-linked networks, mainly indicated by the thixotropic recovery of the mechanical properties. With the bisurea PEG-derived segmented copolymer storage moduli up to 2 × 104 Pa were obtained, significantly higher than other injectable systems obtained with peptide, alginate, or proteinbased materials (from 50 to 5000 Pa),52 although higher concentrations are needed (10 wt %). The poly(8kU4U) hydrogel (G′ = 2 × 104 Pa) shows shear thinning by close to a factor of 40 at large strains (γ = 100%), higher than the telechelic protein based hydrogel (G′ = 3.8 × 103 Pa) reported by Olsen et al.,5 which shows shear thinning by a factor of 30 at identical strain, while both materials show rapid recovery of the modulus after removal of deformation stress. In the poly(8kU4U) hydrogel, complete recovery to the original modulus is observed within 80 s, much faster than in many other shear thinning systems.56,61,62 Pronounced shear thinning in combination with rapid recovery of the poly(8kU4U) hydrogel allows injection by hand through a 29 gauge diameter needle by applying gentle pressure, with rapid self-healing after extrusion from the tip. Thus, the material is injectable through a narrow bore needle at a higher storage modulus (G′ = 20 kPa) than the peptide nanofiber hydrogels studied by Bakota et al. (a hydrogel with G′ = 480 Pa was injected through a 21-gauge needle53) or alginate-based hydrogel studied by Aduado et al. (a hydrogel with G′ = 58.1 Pa was injected through 28-gauge needle54). In the initial tests, the poly(8kU6U) gels display good properties in terms of cell viability.

AUTHOR INFORMATION

Corresponding Author

*Fax: 31-40-2451036. Tel.: 31-40-2473111. E-mail: r.p. [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The small-angle X-ray experiments were performed at the European Synchrotron Radiation Facility (ESRF) in Grenoble at the high brilliance beamline ID02. We are grateful for the beamtime and we kindly acknowledge T. Narayanan and I. de Feijter for technical support and assistance with the SAXS measurements. We are also thankful to Leonie Grootzwagers for the XTT measurements. I.K.V. acknowledges financial support from The Netherlands Organisation for Scientific Research (NWO−VENI Grant: 700.10.406) and the European Union through the Marie Curie Fellowship Program FP7PEOPLE-2011-CIG, Contract No. 293788). This research forms part of the Project P1.04 SMARTCARE of the research program of the BioMedical Materials Institute, cofunded by the Dutch Ministry of Economic Affairs, Agriculture and Innovation. The financial contribution of the Nederlandse Hartstichting is gratefully acknowledged.

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CONCLUSIONS The combination of scalable, straightforward synthesis, yielding, and significant shear thinning make the PEG-bisurea-based hydrogels competitive materials for biomedical applications where injectability is required. The modulus of the hydrogel is tunable thorough the length of the PEG segments. In some applications, interactions with cells (e.g., adhesion) are not required, and the high PEG content of the material ensures weak cell interaction. In other applications, cell adhesion is essential. The bisurea motif of the PEG-bisurea-based hydrogels, embedded in the hydrophobic segments provides the possibility to “click in” functionality in a noncovalent manner31,50,64−66 to provide binding of probes and bioactive moieties. Current work is aimed at optimizing cell adhesion to the materials in this manner. J

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