Injectable Hydrogels from Segmented PEG-Bisurea CopolymersGajanan 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 MechanicalEngineering 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 forbiomedical applications, with segmented hydrophobic bisureahard segments and hydrophilic PEG segments. The segmentedcopolymers were obtained by the step growth polymerizationof amino-terminated PEG and aliphatic diisocyanate. Due totheir capacity for multiple hydrogen bonding within thehydrophobic segments, these copolymers can form highlystable gels in water at low concentrations. Moreover, the gelsshow shear thinning by a factor of 40 at large strain, whichallows injection through narrow gauge needles. Hydrogel moduli are highly tunable via the physical cross-link density and thelength of the hydrophilic segments. In particular, the mechanical properties can be optimized to match the properties ofbiological host tissues such as muscle tissue and the extracellular matrix.

■ INTRODUCTION

Hydrogels are widely investigated as biomimetic scaffolds, dueto properties such as good biocompatibility and biodegrad-ability, high water content, and excellent permeability foroxygen, nutrients, and metabolites.1−4 They are preferablydesigned in such a way that they are liquid-like during injectionand form networks only in situ; such networks are termedinjectable hydrogels.3−9 Injectable gels are known to flow undermodest shear stresses and set promptly at the desired locationin the absence of flow-induced stress. They can thus be used inminimally invasive surgical procedures and are also favorablefor irregularly shaped defects.11−15 Most hydrogels are injectedas a low viscosity liquid and subsequently cross-linked in vivovia chemical or physical interaction, which includes chemicallinkers,16 enzymes,17 and photoinitiated polymerization,18 or bychanges in pH,19 temperature,20 or ionic strength.21 Althoughthere are several advantages of in vivo cross-linking related tothe low viscosity during injection, there are also severallimitations. These include toxicity of unreacted monomer, pre-cross-linked polymers, and leakage during injection and dilutionwith body fluids.22 Alternative approaches that avoid theselimitations make use of the shear thinning and yieldingbehavior of hydrogels.5−10 In this strategy, a solid-like hydrogelmaterial is prepared ex vivo with the desired mechanical,morphological, and biological properties and then injected. Thestress applied during injection leads to yielding and once flowhas started, shear thinning may further facilitate injection.Once, the shear stress stops the material forms a solid-likehydrogel. As a result, the properties of the injected gel material

in vivo are similar to those of the original ex vivo hydrogel.Shear thinning has the additional advantage of giving rise toshear banding or even plug flow, which limits shear forces to anarrow region close to the needle wall, thus, preventing shear-induced damage to cells taken up in the gel.5 Recently,Guvendiren et al.6 reviewed shear-thinning hydrogels utilizedfor biomedical applications. They categorized these gelschemically 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 andblock copolymers. Here we introduce shear thinning hydrogelsfrom segmented copolymers that are completely synthetic andhave several advantages over hydrogels based on naturalpolymers.Various types of natural or seminatural polymers are known

as biomimetic scaffolds for biomedical applications. Theseinclude cellulose,23 chitosan,7 xyloglucan,24 gelatin,25 and theirderivatives, and synthetic polymers, including polyethers,30

block copolymers of polyethers and polyesters, and syntheticpolypeptides.26−29,63 Naturally derived polymers have thepotential advantage of biocompatibility and cell adhesion.34

However, they generally express limited mechanical properties.Moreover, large-scale synthesis of these materials is cumber-some and difficult. Therefore, simple and scalable syntheses arehighly desirable for potential applications.

Received: August 6, 2012Revised: November 13, 2012Published: November 14, 2012

Article

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© 2012 American Chemical Society 3966 dx.doi.org/10.1021/bm301242v | Biomacromolecules 2012, 13, 3966−3976

The ideal hydrogel for tissue engineering combinesbiocompatibility, biodegradability, and allows tuning of itsmechanical properties by control over the morphology.34−39 Tohave the most predictable structure−property relationship,structurally monodisperse segments are desirable as buildingblocks. Furthermore, it is useful to have “click on” functionalityin order to facilitate specific recognition and binding ofbioactive moieties and probes for tracing and analysis.31−33

Here we introduce bisurea-PEG segmented copolymers andshow that they can provide many of the desired features ofinjectable hydrogels for tissue engineering. Recently, wedescribed a powerful tool for the preparation of segmentedcopolymers possessing uniform (monodisperse) hard blocksbased on urea groups and poly(tetrahydrofuran) (pTHF)containing soft blocks for developing thermoplastic elasto-mers.40,41 Following this concept, here we describe thepreparation of injectable, biocompatible, and elastic hydrogelsfrom a segmented copolymer with hydrophobic and hydro-philic segments. The structure of the present segmentedcopolymer combines a hydrophilic and swellable domainconsisting of polyethylene glycol (PEG) with a hydrophobic,biodegradable, crystallizable, and nonswellable domain contain-ing two urea groups (Figure 1).Urea groups are known to self-associate via hydrogen

bonding, and this self-assembling motif has been used in thedevelopment of gelation agents.42−50 The role of the hydrogenbonding groups is 2-fold: First, they provide additionalinteractions between hydrophobic segments, increasing thestrength and lifetime of the physical cross-links. Second, theyare recognition sites for binding of functional molecules that areprovided with two urea groups. In the segmented copolymersused in this study, the hydrogen bonding groups are separatedfrom the PEG segments by 10 CH2 groups (methylene groups)to prevent the PEG from interfering with hydrogen bondingamong the urea groups.44 They have the capability to gel inwater via multiple hydrogen bonding. The mechanical behaviorof our hydrogels was tested by oscillatory shear experimentsand dynamic strain amplitude tests. Atomic force microscopyand small-angle X-ray scattering were used to characterize themorphology of the gels at small length scales. The viability ofhuman myofibroblast cells in the presence of the segmentedcopolymer was studied with a XTT (2,3-bis-(2-methoxy-4-nitro-5-sulfophenyl)-2H-tetrazolium-5-carboxanilide) assay.

■ EXPERIMENTAL SECTIONMaterials and Methods. Materials. Solvents used in synthesis

were reagent grade. CH2Cl2, CHCl3, Et3N, and pyridine were distilledfrom CaH2. All PEG derivatives were dried in vacuum over P2O5during at least 12 h. The reagents 11-aminoundecanoic acid, 1,4-diisocyanatobutane, and 1,6-diisocyanatohexane were purchased fromAldrich, Fluka, or Acros and were used without additional purification.11-tert-Butoxycarbonylamino-undecanoic acid was prepared accordingto literature procedures.51

General Methods. All reactions were performed under a nitrogenatmosphere in a glovebox (MBraun LabMaster 130, MBraun,Garching, Germany) to achieve high molecular weight segmentedcopolymer. NMR spectra were acquired on a 400 MHz VarianMercury Vx (400 MHz for 1H NMR, 100 MHz for 13C NMR). Protonand carbon chemical shifts are reported in ppm downfield oftetramethylsilane using the resonance of the deuterated solvent asinternal standard. Splitting patterns are designated as singlet (s),doublet (d), triplet (t), multiplet (m), and broad (b). Infrared spectrawere 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 Laboratorieswith integrated refractive index detector, using a Polymer StandardsService Gram Linear M 8 × 300 mm, 10 μm particles column, andDMF 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 spec-troscopy (MALDI-TOF) was performed on a Perseptive DE PROVoyager MALDI-TOF mass spectrometer using α-cyano-4-hydrox-ycinnamic acid as the calibration matrix.

DSC Analysis. Differential scanning calorimetry measurements wereperformed on a Thermal Advantage Q2000 apparatus between −80and 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 TAInstruments Universal Analysis software.

Rheology. Mechanical properties of these hydrogels were tested byrheology. Dynamic viscoelastic measurements were determined using astress-controlled rheometer (Anton Paar, Physicia MCR501) equippedwith a sand-blasted plate−plate geometry to prevent slippage.Measurement temperature was fixed at 15 °C unless stated otherwiseand 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 SynchrotronRadiation Facility (ESRF) in Grenoble, France, at the high brilliancebeamline ID02. An X-ray energy of 12.46 keV and one sample-to-detector 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 holessandwiched between Kapton tape at a temperature of 20.1 °C. Thescattering data were corrected for background scattering, detectorresponse and primary beam intensity fluctuations. The instrumentscattering vector was calibrated using a silver behenate standard. Peakpositions were determined by taking the midpoint between the localextremes in the first derivative of the scattering intensities.

Atomic Force Micrographs. Atomic force micrographs wererecorded under ambient conditions with silicon cantilever tips (PPP-NCH, 300−330 kHz, 42 N/m from Nanosensors) using an AsylumResearch MFP-3D-Bio in noncontact mode. For the atomic forcemicroscopy, 2 μL of the solution at ambient temperature was drop-caston freshly cleaved mica.

Biocompatibility/Toxicity/Viability. The viability of human myofi-broblast cells in the presence of the segmented copolymer was studiedwith XTT based In Vitro Toxicology Assay Kit, which was purchasedfrom Sigma-Aldrich.

Synthesis of PEG-Bisurea Segmented Copolymer. Bis(N-(tert-butyloxycarbonyl)-11-aminoundecanoyl)-poly(ethyleneglycol). In a250 mL two-neck round-bottom flask 2.26 g (7.49 mmol) of N-(tert-butyloxycarbonyl)-11-aminoundecanoic acid, 1.44 g (7.49 mmol) of

Figure 1. Architecture of the segmented copolymers used in this work.

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N-(3-dimethylaminopropyl)-N-ethylcarbodiimide hydrochloride and0.92 g (7.48 mmol) of dimethylaminopyridine were stirred in 100 mLof dry dichloromethane under argon for 20 min. To the resultingsolution was added 15 g (1.87 mmol) of poly(ethylene glycol) (Mn ca.8000) and the reaction mixture was stirred for 3 days. The solutionwas washed with 10% citric acid and precipitated by addition of diethylether, 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 of10 g (1.1 mmol) bis(N-(tert-butyloxycarbonyl)-11-aminoundecanoyl)-poly(ethyleneglycol) in 30 mL of dioxane and stirred at 0 °C for 1 hand subsequently at room temperature for 12 h. The solvent wasevaporated to yield 9.0 g (100%) of the product as its hydrochloricsalt.Bis(11-amonium chloride undecanoyl)-poly(ethylene glycol) (8k).

1H 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, CH2CO), 1.75−1.48 (m, 8H, CH2CH2N andCH2CH2CO), 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 1,4-

diisocyanatobutane in 1 mL of dichloromethane, solution of bis(11-aminoundecanoyl)-poly(ethylene glycol) (5.0 g, 0.55 mmol) andtriethylamine (0.1 mL, 0.96 mmol) in 20 mL of dichloromethane wasadded and stirred for 3 days. The solution was concentrated andprecipitated 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 (100MHz, 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 synthesizedaccording 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 = 295K) 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, CH2CH2Nand CH2CH2CO), 1.27 (bs, 19H, CH2). FT-IR (ATR mode, cm−1) ν2882, 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 = 295K) 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.

■ RESULTSSynthesis of the Segmented Copolymer. The prepara-

tion of segmented copolymers was achieved by the step-growthpolymerization of amino-terminated PEG and aliphaticdiisocyanate. The general synthetic protocol of segmentedcopolymers 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-

Scheme 1. Synthesis of PEG-bisurea Segmented Copolymer

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ethylene glycol I. Next, Boc groups were removed by treatmentwith 1 M HCl solution in 1,4-dioxane to obtain telechelicamino-terminated PEG material II. Finally, the correspondingPEG material II was copolymerized with tetramethylene orhexamethylene diisocyanate to develop segmented copolymerIII. To achieve high molecular weight and for stoichiometrycontrol, all polymerizations were performed under a nitrogenatmosphere in a glovebox. All segmented copolymers were fullycharacterized 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 couldbe 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 wasinvestigated using differential scanning calorimetry (DSC).The DSC measurements of the polymers showed twotransitions (Table 1). The first transition (around 50 °C)

corresponds to the melting point of the PEO segments,whereas the second transition state is assigned to melting of thehydrogen-bonded hydrophobic segments, in analogy with asimilar transition in bisurea pTHF segmented copolymers.35

Mechanics and Structure as a Function of PolymerConcentration. The segmented copolymers have thecapability to form gels in water. Hydrogels from poly(8kU4U)were prepared by dissolving 3−7 wt % of the segmentedcopolymer in deionized water under sonication at 45 °Cfollowed by cooling to room temperature. The mechanicalproperties of the resulting material depend strongly on thepolymer concentration. The simplest test of the gel-likeproperties of a material is the tilting vial method, shown inFigure 2a. We observed that the poly(8kU4U) samples at 3 and4 wt % do not form gels strong enough to support their weightin the vials, whereas the samples above 5 wt % form strong gelsthat did not flow in the vials. We thus find a criticalconcentration between 4 and 5 wt % for the tilted vial method,above which the material no longer flows when the vial is tiltedover extended periods of time, thus indicating the formation ofa solid-like gel network.To investigate the viscoelastic properties of the formed

materials more closely, we also performed oscillatoryrheological measurements on gels that, after sonication, werekept overnight at room temperature to reach an equilibriumstate. All polymers were kept in the rheometer for at least 2 hbefore starting the measurements. Figure 2b shows the storageand loss modulus G′ and G″ of the system as a function ofpolymer concentration, as accessed at a oscillation frequency of1 rad/s and a strain amplitude of 1%. In agreement with thetilting vial method, these measurements show a dramaticchange in the magnitude of G′ and G″, which occurs at a

concentration between 4 and 5 wt %. Above this characteristicconcentration, the material behaves solid-like, with the storagemodulus always larger than the loss modulus, G′ > G″, asshown in Figure 2b. Below 5 wt %, the magnitude of the moduliis dramatically lower, dropping by around 3 orders ofmagnitude compared to the 5 wt % sample. Nevertheless, forthe low concentration samples, the storage modulus is stillcomparable or larger than the loss modulus, as accessed at anoscillation frequency of 1 rad/s. This indicates that even below5 wt % the sample still has solid-like properties. However, it isdifficult to discern the physical origin of this elastic-likeresponse from the response at a single time scale, measured atjust one single oscillation frequency.The frequency dependence of the storage (G′) and the loss

(G″) modulus in the linear viscoelastic regime was measured byapplying a 1% strain in the angular frequency range of 100 to0.1 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 rheologicalbehavior, as shown in Figure 2c. At the same concentration, theelastic modulus of poly(8kU6U) was higher than poly-(20kU6U). We hypothesize that this is because fewer physicalcross-linking domains are present in poly(20kU6U). For all ourgels, we find that the storage modulus was larger than their lossmodulus and depended only weakly on frequency over theentire range of frequencies studied, a behavior typical forchemical or strongly interacting physical gels. However, the lossmodulus shows a minimum at intermediate frequencies,indicating the presence of a structural relaxation process atvery low frequencies, below the range accessible in ourexperiments.The significant changes in viscoelastic properties observed

just below 5 wt % suggest that a qualitative change in themorphology of the gels is taking place at this characteristicconcentration. We used atomic force microscopy (AFM) toinvestigate the surface morphology of 5 wt % poly(8kU4U)hydrogels (Figure 3). Samples were prepared in a wetted stateto minimize the effect of drying on morphology. Height andphase mode images of the same sample showed thatcharacteristic dimensions of the aggregates in the wet gel isin the order of 20−35 nm. In the height mode, no features withhigh aspect ratio are present. In the phase mode, elongatedfeatures are observed, but they are only present at the interfaceof the domains visible in height mode. Therefore, we concludethat, although the presence of aggregates of the hydrophobicdomains can be deduced from the rheological properties of thegels, their shape (fibrous or globular) cannot unambiguously bedetermined from the AFM images.Further information on the structure of the hydrogels was

obtained from small-angle X-ray scattering. The experimentswere performed at the high brilliance beamline ID02 at theEuropean Synchrotron Radiation Facility (ESRF) in Grenoble,France. To this end, samples were sandwiched between Kaptontapes glued onto two opposite sides of an aluminum sampleholder with several small holes of a few mm’s in diameter thatserved as sample container. The scattering profiles of 10 wt %gels of poly(8kU4U), poly(8kU6U), and poly(20kU6U) aregiven in Figure 4a. A correlation peak was observed atintermediate q-values in all samples, signifying the existence of acharacteristics distance, d = 2π/qmax in the material, where q isthe magnitude of the scattering wave vector and qmax is the peakposition in q-space. We attribute this characteristic distance to

Table 1. DSC Measurements of Polymers

transition 1 transition 2

compoundsΔT

(°C/min)T1(°C)

ΔH1(J/g)

T2(°C)

ΔH2(J/g)

poly(8kU6U) 20 46.81 102.8 176.0 10.74poly(20kU6U) 20 42.85 111.6 160.0 5.23poly(8kU4U) 20 53.76 119.4 221.1 25.31poly(20kU4U) 20 63.76 136.0 233.1 0.89

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the spacing between hydrophobic domains within the hydrogel,which are dense in physical cross-links.49

The position of the correlation peak/shoulder is dependenton the molecular structure of the gel building blocks and theirconcentration; indeed, as shown in Figure 4b, with increasingconcentration of poly(8kU4U) we observe a systematic shift ofthe peak position toward higher q-values, corresponding to adecreasing characteristic length d. The characteristic length dsystematically increases as the length of the poly(ethyleneglycol) chain is increased, as shown in Figure 4a and Table 2.

The smaller value for d in the second entry of Table 2compared to the first entry is tentatively ascribed to the morehydrophobic U6U segments in the former compound. Asexpected, the correlation peak is more pronounced and is foundat higher q-values in samples prepared at the highestconcentrations; for the 15 wt % sample, a strong peak isobserved, indicating a well-defined characteristic length scale inthe structure of the material. The length scale indicates that thePEG spacers set the characteristic length d. However, atconcentrations as high as those of our gels, the available form

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).

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and structure factor models break down, thus, preventing usfrom obtaining more detailed information on the structure ofthe material.Using AFM and SAXS, it was not possible to determine

whether the structure consists of flowerlike spherical objects orof rodlike objects. However, the mechanical properties of thematerial ex vivo appear suitable for biomedical applications.Therefore, we investigated whether the material is indeedsuitable for use as an injectable hydrogel, as discussed in thesubsequent sections.Injectability and Flow Behavior. We started with the

simplest possible experiment, directly testing whether thematerial can be injected via a small diameter syringe needle.One key requirement for a suitable injectable material is thesize of the needle that could be used in the injection process.For patient comfort and maximum control, small diameterneedles are preferable. We thus tested the injectability of our

hydrogels using a syringe equipped with a thin 29 gauge needle(inner diameter dneedle = 0.184 mm). As shown in Figure 5, the10 wt % poly(8kU4U) hydrogel could readily be injectedthrough 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 needletip, we noted that the material appears solidlike almostimmediately after exiting the needle.This simple experiment shows directly that indeed the

material can be transferred through a thin syringe needle usingrealistic forces applied to the syringe piston. This requires twoseparate phenomena to occur in the syringe: First, the stressapplied to the material has to be large enough to lead to ayielding of the material, thereby transferring its behavior fromsolidlike to liquidlike. Second, the viscosity of the materialshould be small enough to enable transfer of the now fluidlikematerial at high enough flow rates and at reasonable levels offorce applied to the syringe.The yield stress was estimated by monitoring the lowest level

of applied force that leads to a observable flow of materialthough the syringe. This critical force was determined to be F ≈12 N, giving rise to a pressure of p ≈ (4F)/(πD2) in thesyringe; this is also a fair estimate for the typical shear stressexpected initially at the inlet of the syringe needle. From the

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 (○), andpoly(20kU6U)92k and (b) 4, 8, 10, and 15 wt % samples of poly(8kU4U)160k.

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

hydrogel qmax (nm−1) d (nm)

poly(8kU4U)160k, 10 wt % 0.26 24.2poly(8kU6U)134k, 10 wt % 0.35 17.9poly(20kU6U)92k, 10 wt % 0.17 37.0

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

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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 theapproximate flow rate of material through the needle atconstant applied force. During application of a constant force ofF ≈ 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 canthen be obtained directly from the applied pressure p ≈ (4F)/(πD2), the flow rate, the length L = 20 mm, and the diameterdneedle of the syringe needle as67

ηπ

= ≈d p

LQ646 Pa sneedle

For injectability, a material that exhibits yielding and shear-thinning behavior is highly beneficial, as during the injectionphase the material is transformed into a fluidlike state. Studyingthe nonlinear response to large applied strains and strain ratesis, thus, of key importance. Therefore, the viscoelastic responseof the hydrogels was measured as a function of the appliedstrain. Strain-dependent measurements were performed at fixedangular frequency (1 rad/s) by varying the strain amplitudefrom 1 to 100%, shown in Figure 6. The linear viscoelastic

regime, where the storage and loss modulus is independent ofthe applied strain amplitude of the material, is limited to strainsbelow a critical yield strain. At larger strains, the material showsa yielding behavior with both the storage and the loss modulusdecreasing with increasing strain amplitude. In this regime also,the response becomes nonharmonic, so strictly speaking theviscoelastic moduli are no longer properly defined. The valuesrepresented here are based solely on the first harmonic of thestress response; nevertheless, we can clearly identify a yieldstrain for the material. The decrease in these moduli above acritical level of strain indicates a transition from a solidlike to amore fluidlike material response for all hydrogel materialsstudied. For poly(8kU4U) and poly(20kU6U), this yield strainis around 30%, where the onset toward a purely liquidlikebehavior at larger strains is observed. However, the mostpronounced yielding behavior is observed for poly(8kU6U),where the yield strain is only around 10% and at the largestaccessed strain the material already shows a clearly fluidlike

behavior. As a result, while both poly(8kU4U) and poly-(8kU6U) show similar elasticlike properties at low strains at astrain amplitude of 100%, the viscoelastic moduli of poly-(8kU6U) are significantly lower than those of poly(8kU4U). Insummary, for all materials studied, we observe the onset ofyielding from a solidlike to a liquidlike behavior occurring at acritical strain of less than 100%. However, the strain level wherethe material truly becomes liquidlike with G″ < G′ should occuronly at larger strains. We thus estimate the critical yield stressneeded 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 thestress 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 lowenough to enable material to flow with a sufficient flow rate atreasonable levels of applied pressure. From the syringeexperiment, we had seen that this is indeed possible and hadestimated a viscosity of around 6 Pa s. Remarkably, from themagnitude of the loss modulus in the oscillatory rheologicalmeasurements a much higher level of viscosity would beexpected; the Cox−Merz rule suggests that steady shearmeasurements are directly linked to oscillatory measurementsand that the viscosity is well approximated by η(γ = 1/ωγ) ≈G″(ω)/ω. Given the oscillatory measurements of the poly-(8kU4U) hydrogel, for a value of 2000 Pa for the storagemodulus, a strain of 100% and a frequency of 1 rad/s, aviscosity of around 2000 Pa s is expected. This value is morethan 2 orders of magnitude larger than the 6 Pa s estimatedfrom the syringe experiment.This dramatic difference indicates that the viscosity of the

material is significantly reduced as the material starts to flowthrough the needle.To investigate this behavior in more detail, we perform

viscosity measurements on the poly(8kU4U) hydrogel as afunction of the applied shear rate. Indeed, as shown in Figure 7,

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.

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

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we observe a significant degree of shear thinning of thematerial. The viscosity decreases by almost 2 orders ofmagnitude upon increasing the shear rate from 1 to 1000 s−1.A significant degree of syneresis is observed betweenconsecutive 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 shearthinning behavior indicate that the structure of the material ischanging significantly under shear.To make a direct comparison with the simple syringe

experiment, the typical shear rate in the syringe needle wasestimated to be γ = 8Q/πdneedle

3 ≈ 32 s−1. The estimatedviscosity of 6 Pa s is shown as a red star in the same figure withrate-dependent viscosity measurements. A remarkably goodagreement is found, indicating that the rate-dependent viscositymeasurements are suitable for predicting the injectability of thematerial. It is, thus, the shear thinning behavior of our hydrogelmaterials that enables their use as an injectable hydrogelmaterial; in the absence of such dramatic shear thinning, a forceon the order of kilo-Newtons would have to be applied toobtain a comparable flow rate.

To further investigate the mechanical properties of bisurea-based hydrogels, we proceeded to study the influence ofmolecular weight on hydrogel stiffness. To do so, we prepared50k and 160k molecular weight bearing poly(8kU4U)segmented copolymer hydrogels and studied their mechanicalproperties by using rheology. As shown in Figure 8a, the effectof the overall molecular weight on the viscoelastic moduli of thehydrogel 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 thehydrogels clearly increase with increasing concentration. This isdue to both the increase in chain density and cross-link densitywith increasing concentration. Further, the yield strain of thematerial moves to lower strains as the concentration ofpoly(8kU4U) is increased, indicating that the polymer gelbecomes more brittle.

Self-Healing of Gels. The strain-dependent measurementsshow that the material exhibits a yield strain and transforms to amore liquidlike behavior at larger deformations, a behavioressential for injectable gels. However, it is also important that

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.

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these materials can recover or heal from a fluidlike state back toa solidlike gel state once the deformation is removed.To characterize this behavior we performed rheological tests

targeted toward quantifying the recovery from a yielded, weaklyviscoelastic state to a recovered, strongly solidlike state.Oscillatory rheological measurements were performed whereperiods of low applied strain amplitude (γ = 0.1%) werealternated with periods where the applied strain amplitude iswell outside the linear viscoelastic regime (γ = 100%), shown inFigure 9. Recovery of the viscoelastic properties was studied forpoly(20kU6U) hydrogels at 10 wt %; the initial G′ in theyielded state was around 1 kPa. As the low strain cycle (γ =0.1%) was initiated, G′ rapidly increased to around 4 kPa, a 4-fold increase. In turn, upon repeated initiation of the highstrain, cycle, G′ rapidly decreased within 20 s to its initial valueof 1 kPa. The yielding and recovery behavior remainsunchanged after many cycles, four cycles are shown in Figure9, indicating that yielding and recovery in these materials is acompletely reversible process.However, studying yielding behavior of some of the hydrogel

materials by applying a large strain was difficult, because theyform brittle gels that break into pieces and are ejected from themeasuring geometry when a large deformation is applied. Theyielding and recovery behavior of these samples was thereforeinvestigated by applying a series of cycles where both thetemperature and the strain amplitudes were varied, as shown inFigure 10 for the poly(8kU4U) 10 wt % sample. At elevatedtemperatures the material becomes softer and less brittle;therefore it can be deformed at large strain amplitudes withoutbreaking up inside the rheometer. Thus, in our experiments, thehydrogel was first heated to 50 °C for 400 s and then 100%strain at 1 rad/s was applied at the same temperature. Therecovery behavior of the hydrogels was then monitored in acycle 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 oflarge strain deformation and elevated temperature. Suchrecovery after mechanical yielding is often referred to as “selfhealing”. In the recovery measurements shown in Figure 10, theinitial storage modulus was observed to be around 20 kPa atlow 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 returningthe strain to small values (1%) and returning to a temperatureof 15 °C, the storage modulus recovers quickly within 80s to itsinitial value, a 40-fold increase. The rapid recovery of hydrogelstrength was repeatedly confirmed for four cycles and in eachcycle it recovers to almost 97% of its initial value, indicatingthat the gel structure after healing is not significantly affected bythe breakup and recovery process.

■ BIOCOMPATIBILITY/TOXICITY/VIABILITYAn XTT assay on the poly(8kU6U) has been performed toinvestigate the toxicity of the material. According to the ISO-norm 10993−5:2009, the viability of human myofibroblast cellstreated with extract of the material must be above 70%. As canbe seen in Figure 11, the segmented copolymer showedexcellent biocompatibility with cells.

For the XTT measurements, the blank was a XTT solutionincubated without cells. The positive control indicates healthycells treated with normal medium (treated similar to the extractof the material). The negative control was cells that wereexposed to the latex (which is highly toxic to the cells). A pieceof 25 × 25 mm (23.29 mg) was used for the high concentrationstudy and a strip of 5 × 25 mm (4.95 mg) was used for the low

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

Figure 11. XTT of poly(8kU6U).

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

■ DISCUSSION

Most of the synthetic hydrogels studied for injectable materialsin tissue engineering are based on peptide amphi-philes.52,53,55−61,63 However, the synthesis of the peptideamphiphiles limits the scalability. Synthesis of the PEG-bisureasegmented copolymers is straightforward and can be performedon a large scale. Moreover, the moduli of the hydrogels can becontrolled to match the properties of biological host tissues bythe length of the hydrophilic segments which influences densityof physical cross-links, as is seen from comparison of moduli ofpoly(8kU6U) and poly(20kU6U). These features bear asimilarity to the PEG-based supramolecular networks recentlyreported by Dankers et al.3

The PEG-bisurea-based hydrogels show rheological charac-teristics of physically cross-linked networks, mainly indicated bythe thixotropic recovery of the mechanical properties. With thebisurea PEG-derived segmented copolymer storage moduli upto 2 × 104 Pa were obtained, significantly higher than otherinjectable systems obtained with peptide, alginate, or protein-based materials (from 50 to 5000 Pa),52 although higherconcentrations are needed (10 wt %). The poly(8kU4U)hydrogel (G′ = 2 × 104 Pa) shows shear thinning by close to afactor of 40 at large strains (γ = 100%), higher than thetelechelic protein based hydrogel (G′ = 3.8 × 103 Pa) reportedby Olsen et al.,5 which shows shear thinning by a factor of 30 atidentical strain, while both materials show rapid recovery of themodulus after removal of deformation stress. In the poly-(8kU4U) hydrogel, complete recovery to the original modulusis observed within 80 s, much faster than in many other shearthinning systems.56,61,62

Pronounced shear thinning in combination with rapidrecovery of the poly(8kU4U) hydrogel allows injection byhand through a 29 gauge diameter needle by applying gentlepressure, with rapid self-healing after extrusion from the tip.Thus, the material is injectable through a narrow bore needle ata higher storage modulus (G′ = 20 kPa) than the peptidenanofiber hydrogels studied by Bakota et al. (a hydrogel withG′ = 480 Pa was injected through a 21-gauge needle53) oralginate-based hydrogel studied by Aduado et al. (a hydrogelwith G′ = 58.1 Pa was injected through 28-gauge needle54). Inthe initial tests, the poly(8kU6U) gels display good propertiesin terms of cell viability.

■ CONCLUSIONS

The combination of scalable, straightforward synthesis, yielding,and significant shear thinning make the PEG-bisurea-basedhydrogels competitive materials for biomedical applicationswhere injectability is required. The modulus of the hydrogel istunable thorough the length of the PEG segments. In someapplications, interactions with cells (e.g., adhesion) are notrequired, and the high PEG content of the material ensuresweak cell interaction. In other applications, cell adhesion isessential. The bisurea motif of the PEG-bisurea-based hydro-gels, embedded in the hydrophobic segments provides thepossibility to “click in” functionality in a noncovalentmanner31,50,64−66 to provide binding of probes and bioactivemoieties. Current work is aimed at optimizing cell adhesion tothe materials in this manner.

■ ASSOCIATED CONTENT

*S Supporting Information1H and 13C NMR and DSC data of the materials. This materialis available free of charge via the Internet at http://pubs.acs.org.

■ AUTHOR INFORMATION

Corresponding Author*Fax: 31-40-2451036. Tel.: 31-40-2473111. E-mail: r.p.sijbesma@tue.nl.

NotesThe authors declare no competing financial interest.

■ ACKNOWLEDGMENTSThe small-angle X-ray experiments were performed at theEuropean Synchrotron Radiation Facility (ESRF) in Grenobleat the high brilliance beamline ID02. We are grateful for thebeamtime and we kindly acknowledge T. Narayanan and I. deFeijter for technical support and assistance with the SAXSmeasurements. We are also thankful to Leonie Grootzwagersfor the XTT measurements. I.K.V. acknowledges financialsupport from The Netherlands Organisation for ScientificResearch (NWO−VENI Grant: 700.10.406) and the EuropeanUnion through the Marie Curie Fellowship Program FP7-PEOPLE-2011-CIG, Contract No. 293788). This researchforms part of the Project P1.04 SMARTCARE of the researchprogram of the BioMedical Materials Institute, cofunded by theDutch Ministry of Economic Affairs, Agriculture andInnovation. The financial contribution of the NederlandseHartstichting is gratefully acknowledged.

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