Reconfigurable Biodegradable Shape-Memory Elastomers via Diels

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Reconfigurable Biodegradable ShapeMemory Elastomers via Diels-Alder Coupling Chi Ninh, and Christopher John Bettinger Biomacromolecules, Just Accepted Manuscript • DOI: 10.1021/bm4002602 • Publication Date (Web): 16 May 2013 Downloaded from http://pubs.acs.org on May 21, 2013

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Reconfigurable Biodegradable Shape-Memory Elastomers via Diels-Alder Coupling Chi Ninh1 and Christopher J Bettinger1,2,3 * [1]

Department of Materials Science and Engineering

[2]

Department of Biomedical Engineering Carnegie Mellon University, Pittsburgh, PA 15213

[3]

McGowan Institute of Regenerative Medicine University of Pittsburgh, Pittsburgh, PA 15213

*To whom correspondence should be address: [email protected]

KEYWORDS: elastomer; polymer; biomaterial; shape memory; Diels-Alder; cycloaddition

ABSTRACT: Synthetic biodegradable elastomers are a class of polymers that have demonstrated far-reaching utility as biomaterials for use in many medical applications. Biodegradable elastomers can be broadly classified into networks prepared by either stepgrowth or chain-growth polymerization. Each processing strategy affords distinct advantages in terms of capabilities and resulting properties of the network. This work describes the synthesis, processing, and characterization of crosslinked polyester networks based on Diels-Alder coupling reactions. Hyperbranched furan-modified 1

ACS Paragon Plus Environment

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polyester precursors based on poly(glycerol-co-sebacate) are coupled with bifunctional maleimide crosslinking agents. The chemical and thermomechanical properties of the elastomers are characterized at various stages of network formation. Experimental observations of gel formation are compared to theoretical predictions derived from FloryStockmayer relationships. This crosslinking strategy confers unique advantages in processing and properties including the ability to fabricate biodegradable reconfigurable covalent networks without additional catalysts or reaction by-products. Reconfigurable biodegradable networks using Diels-Alder cycloaddition reactions permit the fabrication of shape-memory polymers with complex permanent geometries. Biodegradable elastomers based on polyester networks with molecular reconfigurability achieve vastly expanded properties and processing capabilities for potential applications in medicine and beyond.

INTRODUCTION Biodegradable elastomers have emerged as a unique class of polymers for potential use in biomedical applications such as tissue engineering scaffolds,1,2 drug delivery systems,3 and bioabsorbable medical devices.4 The unique combination of the following properties makes these materials ideally suited for many biomedical applications: low elastic moduli, high maximum strains, biodegradability, and relatively simple monomeric components. Other unique functionalities have been conferred upon biodegradable elastomers including luminescence5 and shape-memory capabilities.6 The desire to further expand the available materials properties has motivated the synthesis of new 2


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compositions. Complementary processing techniques have evolved to meet the requirements of application-specific biodegradable elastomers. Biodegradable elastomers can be generally categorized as one of the following general polymerization mechanisms: (1) step-growth or (2) chain-growth. Each network formation strategy affords distinct advantages that may be leveraged in designing application-specific materials. Preparing biodegradable elastomers formed from step-growth mechanism primarily involves thermal crosslinking through polycondensation. Precursor materials are commonly subjected to aggressive processing conditions (temperatures above 100 oC and pressures below 5 Pa).7 The temporary generation of a molten state and the evolution of reaction by-products limit the spectrum of possible three-dimensional geometries. Conversely, networks can also be formed through chain-growth mechanisms that employ free radical polymerization schemes. This strategy permits the formation of biodegradable elastomers with complex geometries at mild temperatures. However, elastomers prepared through chain-growth polymerization schemes afforded by free-radical crosslinking require additional components such as initiators. These crosslinking reactions generate extended aliphatic networks. The presence of secondary networks can negatively influence cytotoxicity profiles and preclude complete hydrolytic degradation of the networks.8 Potential limitations in processing biodegradable elastomers motivate the investigation of alternative network formation strategies that can obviate limitations associated with high temperature step-growth polymerization and free radical chain-growth schemes, both of which are widely used for this class of materials. Synthetic polymers composed of covalent adaptable networks (CANs) serve as an inspiration for next-generation biodegradable elastomers. CANs utilize reconfigurable 3


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covalent bonds to achieve capabilities such as self-healing and shape-memory.9 CANs may also be utilized as a strategy to synthesize biodegradable elastomers with enhanced processing capabilities. This work describes the design, synthesis, and characterization of a new class of biodegradable elastomeric polyesters based on reconfigurable Diels-Alder (DA) cycloaddition reactions.

EXPERIMENTAL SECTION Synthesis and Characterization of Furan-modified Poly(glycerol-co-sebacate) Precursors. All chemicals were purchased from Sigma-Aldrich (Milwaukee, WI) and used as received unless stated otherwise. PGS pre-polymer was synthesized according to previous published methods.7 Briefly, glycerol and sebacic acid were combined in an equimolar ratio. The ratio was chosen to achieve a high degree of free hydroxyl groups for subsequent furan modification that can participate in DA cycloaddition reactions. The glycerol-sebacic acid mixture was combined at 130 oC under N2 gas for 3 hours followed by polycondensation under vacuum (100 mTorr) at 130 oC for an additional 48 hours. The resulting PGS pre-polymer was used for furan modification without any further purification. Pre-polymer PGS (20 g, 78 mmol of hydroxyl groups) was dried at 60 oC and 5 Pa for 2 hours, then dissolved in 200 ml of anhydrous chloroform with 4dimethylaminopyridine (DMAP). The reaction flask was cooled down to 0 oC in an ice bath under a positive pressure of N2. Furoyl chloride (20%-80% mol/mol of hydroxyl groups on PGS pre-polymer) was added dropwise and in parallel with an equimolar amount of anhydrous triethylamine (TEA). The reaction was allowed to reach room temperature and stirred for an additional 24 hours. Excess solvent was removed by rotary 4


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evaporation. The resulting product was re-suspended in ethyl acetate (100 ml), filtered, and dried at 40 oC and 5 Pa for 12 hours. The molecular weight distribution of the PGS pre-polymer was measured by the gel permeation chromatography (GPC) (Polymer Standards Services, Amherst, MA) using poly(methyl methacrylate) (PMMA) as the standard and dimethylformamide (DMF) as the eluent. Nuclear Magnetic Resonance (1H NMR) spectra were recorded in chloroform-d (Bruker 300MHz NMR spectrometer, Billerica, MA). The degree of substitution (DS) was calculated as the molar ratio of furan groups incorporated in the PGS pre-polymer per mole of sebacic acid. Formation and Characterization of Networks Using DA Cycloadditions. 1,1′(Methylenedi-4,1-phenylene)

bismaleimide

(DPBM)

was

recrystallized

in

2:1

dichloromethane/hexane before use. PGSF-DPBM networks were prepared on maltosecoated silicon oxide handling substrates by combining DPBM and PGSF in a 1:1 ratio of furan to maleimide functionalities. PGSF-DPBM melts were prepared and homogenized briefly at 120 oC. Crosslinked networks were prepared by accelerating forward DA cycloaddition reactions at 92 oC for 2-9 days. Films were delaminated from support substrates via dissolution of the release layer in ddH2O at 80 oC for 4 hours. Sol-free networks were generated by incubating films in ethanol and water in succession for 24 hours each with periodic sonication. Attenuated total reflectance – Fourier transform infrared spectroscopy (ATR-FTIR) spectra were recorded by placing PGSF pre-polymers or PGSF-DPBM mixtures (n = 4) directly on top of the ATR-FTIR crystal (Madison Instruments, Inc., Madison, WI). The time course of the absorbance intensity of the furan (1010 cm-1) and maleimide stretches (690 cm-1) were normalized to the peak at 1473 cm1

. The latter peak corresponds to the C-H bend and remains virtually constant during the 5


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forward DA cycloaddition crosslinking reaction.10 The viscoelastic properties of the networks were measured by a rheometer operated in parallel plate configuration (TA Instruments DHR-3, New Castle, DE). The material was briefly heated to 100 oC to ensure intimate contact between the plates. Amplitude sweeps were taken from 0.1% to 10% strain to identify the linear viscoelastic region. Frequency sweeps were conducted at 2% strain from 0.1 to 10 rad/s at 92 oC. The data was smoothed to discard minor noise fluctuations from the instrument by applying the local regression using weighted linear least squares and a 2nd degree polynomial model with a MATLAB function (LOESS).11 Physical Properties of PGSF-DPBM networks. The glass transition temperature (Tg) of the crosslinked networks was measured by differential scanning calorimetry (DSC, TA Instruments Q20-1853, New Castle, DE). Tensile tests were conducted on PGSF-DPBM dog-bone shaped films (L x W x T = 25 x 6 x 0.5 mm3). Tensile mechanical properties of hydrated samples were measured by straining samples to failure at a strain rate of 0.5 mm-min-1 (Instron 5943 equipped with BlueHill Testing version 3, Brockton, MA). The tensile Young’s modulus (E0) was calculated by linearizing the region between 0 and 3% strain. The toughness was determined by integration of the area under the stress-strain curve. The sol fraction and hydration (water uptake) in crosslinked networks were measured by graviometry. Briefly, the masses of dehydrated newly formed crosslinked networks (sol plus gel) were measured (W0). Samples were incubated sequentially in ethanol and water for 24 hours each to remove the sol fraction. Excess water was removed and the mass of the samples was recorded (Ws). Samples were dehydrated for 60 0

C under vacuum (5 Pa) for 24 hours and the mass of each sample was recorded (W1).

This dehydration time is sufficient to remove 100% of the water in the networks 6


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(Supplementary Information, Figure S1). The sol content (Solcont) was calculated using the following relationship:

Solcont (%) =

W 0 − W1 × 100 W0

(1)

The water content (H2Ocont, expressed as percent of Ws) was calculated by the following formula:

H 2Ocont (%) =

W s − W1 × 100 Ws

(2)

The mass density (ρ) was measured using pycnometer (25 ml nominal volume, Kimble Chase, Vineland, NJ). The volume density of active crosslinks (nc, moles of active network per unit volume) and the molecular weight between crosslinks (Mc) were calculated from the Young’s modulus and the mass density using the following equations:

n=

E0 ρ = 3RT M c

(3)

where R is the universal gas constant, and T is the temperature. Contact Angle and In Vitro Degradation of PGSF-DPBM Networks. Wettability of the surfaces was evaluated by water-in-air contact angle measurements. PGSF-DPBM thin films were prepared by spin-coating a 10% (w/v) solution of PGSF-DPBM in chloroform at 300 rpm on a glass cover slip. Polymer surfaces were cured as previously described. Droplet contours were recorded (Rame-Hart Instrument Co, Mountain Lake, NJ) and analyzed using NIH ImageJ. The masses of sol-free polymer slabs (L x W x T = 25 x 6 x 0.5 mm3) were recorded (W0) and incubated in 15 ml of 3M sodium acetate (pH

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= 5) at 37 oC (n = 4). The masses of both the dehydrated and hydrated samples were measured every 4 days throughout the time course of the degradation experiment. Molecular-Scale Network Reconfiguration. Complex geometries in PGSF-DPBM networks were fabricated by utilizing retro DA cycloaddition reactions. Planar sol-free PGSF-DPBM films (DS = 45%) were synthesized, hydrated, and configured into the final arbitrary complex geometry. The molecular scale network was scrambled by heating at 120 oC for 24 hours to induce the retro DA cycloaddition reactions. The newly formed permanent network was established by inducing forward DA cycloaddition at 92 oC for 24 hours. Samples selected for the negative control condition were prepared in an identical manner except for omission of the retro DA cycloaddition step at 120 oC. DSC analysis was performed on a pre-weighed PGSF-DPBM networks to provide insight into the progression of the retro and forward DA cycloaddition reaction. The Tg of PGSFDPBM network was measured after the second heating run (Tg1). The samples were held at 120 0C for 24 hours. The Tg of the network was measured again (Tg2). The samples were then heated at 92 0C for 24 hours and the Tg was measured for a third time (Tg3). All values of Tg were calculated by determining the temperatures at the midpoint of the phase transition.

RESULTS AND DISCUSSION Chemical Characterization of Furan-Modified Network Precursors. The weightaverage (Mw) and number-average molecular weight (Mn) of PGS pre-polymers are 3.79 kDa and 1.69 kDa, respectively, as assessed by gel permeation chromatography using a linear PMMA standard. The incorporation of pendant furan groups was verified by 1H 8


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NMR (300 MHz, CDCl3): δ1.2 (13H, q), δ1.65 (6H, q), δ2.35 (6H, t), δ3.2 (0.57H, m), δ3.7 (0.87H, m), δ4.2 (8H, m), δ5.1(0.23H, m), δ5.3 (0.39H, m), δ6.55 (1H, d), δ7.25 (1H, d), δ7.65 (1H, d) (Figure S1, Supporting Information). This synthesis procedure was used to produce PGSF precursors with four degrees of substitution (DS) for subsequent characterization (22%, 38%, 45%, and 59%). The DS was difficult to precisely define a priori because of the large polydispersity (PDI = 2.25) and complex topology of the globular PGS starting material. However, practical upper and lower limits for DS were established. PGSF precursors with DS of approximately 20% or smaller were not capable of robust network formation. Crosslinked networks formed from PGSF precursors with DS of approximately 60% or larger are highly hydrophobic and exhibit very slow degradation kinetics. Elucidation of Chemical Signatures During PGSF-DPBM Network Formation. Figure 1 displays ATR-FTIR spectra of PGS, PGSF, and PGSF-DPBM network. Furan conjugation of PGS produces an absorption peak at approximately 1010 cm-1, which corresponds to the furan C-O-C ether stretch. The =C-H bending of the maleimide group at 690 cm-1 was present in PGSF-DPBM melts during initial phases of crosslinking. Forward DA cycloaddition reactions in PGSF-DPBM melts were evident from the emergence of a new peak at 1500 cm-1, which corresponds to the C=C in the DA adduct.12 The reaction rates of forward DA cycloaddition reaction could be compared semi-quantitatively using ATR-FTIR. The absorbance of the peak centered about 1010 cm-1 was measured during network formation from PGSF with high DS (59%) and low DS (22%). Higher concentrations of furan groups induce relatively more rapid

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crosslinking kinetics as measured by the accelerated reduction in the peak absorbance at 1010 cm-1 (Figure 2). Gel Point Determination in Crosslinked Networks. Theoretical and experimental gel points were compared for PGSF-DPBM networks. The gel point was estimated using Flory−Stockmayer equation:13

pg =

1 r(1 − f M )(1 − f F )

(4)

where pg is the gel point conversion and r is the stoichiometric ratio. The value of r was fixed at unity for all network compositions. The degrees of functionality of furan- and maleimide-containing network precursors are represented by fF and fM, respectively. Bifunctional DPBM has a value of fM = 2. The effective functionality of PGSF was estimated by first modeling PGS as a linear polymer that consists of equal repeat units of glycerol and sebacic acid with a MWrepeat of 205 g/mol. The average number of repeat units was determined from GPC results (NGPC = 8.22). Kramers theorem for an ideal randomly branched polymer states that N ~ Rg4 , which exhibits a notable departure from the classic Debye result in which N ~ Rg2 .14 The adjusted degree of polymerization (N) for hyperbranched PGS networks was calculated to be NKramer = 67.5. The values pg were calculated for all PGSF compositions (Table 1). The gel point conversion was measured experimentally by small-amplitude oscillatory rheology. Winter and co-workers demonstrated that the real and imaginary moduli scale with frequency (w) identically at the gel point with n as the relaxation exponent, known as the Winter-Chambon criterion.15 This relationship is summarized as follows:

G'(w) ~ G''(w) ~ w n

(5) 10


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Eqn. 5 can be expressed in terms of the loss tangent tan δ as follows:

tan δ =

 nπ  G'' = tan   2  G'

(6)

The gel point is assigned to the time at which tan δ is temporarily frequency independent. The time course of tan δ across different frequencies for PGSF-DPBM networks (DS = 59%) indicates a crossover at 37 hours with tan δ = 1.53 (Figure 3). According to (Eqn. 6), n was calculated to be 0.63, which is in good agreement with the predicted value from the Rouse model (0.67).13,

16

The gel point determination of

networks formed from PGSF of various compositions is summarized in Table 1. The relaxation exponent ranges from 0.07 to 0.63 across PGSF compositions. The observed departure from Rouse model predictions likely occurs due to the high degree of chain overlap. Topologies with a high extent of overlap yield a relatively smaller relaxation exponent compared to the value predicted by the Rouse model.13, 17 The gelation time was verified by observing a local sharp increase in the loss (G’) and storage modulus (G”) of the network. Gradual increases in G’ and G” followed by a well-defined plateau were observed during the initial stages of network formation. The gel point was identified as the time when a clear increase in both G’ and G” was observed that dwarfed any temporal fluctuations. This increase corresponds to the precise time coordinate of network formation.18 This time point closely matches the gelation time extracted from tan

δ crossover measurements (Figure 3). Representative plots of the time course of the complex moduli over time was graphed for frequency of 10 rad/s. Similar plots for other frequencies are consistent with these observation (data not shown). Furthermore, the

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calculated gelation times for networks of each PGSF composition were consistent when determined by measuring both (1) the crossover point and (2) inflections in G’ and G”. The extent of the reaction p at gelation was calculated by combining rheology and spectroscopy data. There is strong congruence between the predicted and experimentally observed critical extent of reaction upon gelation for networks with PGSF compositions of DS = 38%, DS = 45% and DS = 59% (Table 1). However, a significant departure was observed in predicted and experimental values for extent of reaction upon gelation in networks with PGSF compositions of DS = 22% (Figure 4). The underestimation is likely to originate from the formation of a hybrid network that is formed from covalent DielsAlder linkages along with secondary bond formation based on hydrophobic interactions. The presence of secondary interactions can produce crosslinked networks as assessed by rheology despite ostensibly premature values of conversion. The sub-critical conversions in PGSF-DPBM networks (DS = 22%) are apparent by the rapid macroscopic disintegration of samples during in vitro degradation studies. These networks also exhibit a relatively low tensile Young’s modulus and reduced toughness (Table 2). PGSF with low degrees of substitution (DS = 22%) are likely capable of forming networks with subcritical values for conversion because of the high theoretical values required for conversion (Table 1). Physical Properties of PGSF-DPBM Networks. Hydrated PGSF-DPBM networks are elastomeric and exhibit composition-dependent values for virtually all of the relevant mechanical properties. Increasing the DS in PGSF-DPBM networks leads to monotonic increases in the elasticity, ultimate tensile strength, and toughness (Figure 5; Table 1). For example, the tensile Young’s modulus increases from 4 to 34 MPa and the maximum 12


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elongation at break increases from 5 to 58% as the DS rises from 22% to 59%. There is a notable departure in the trajectory of the stress-strain curves for PGSF-DPBM networks compared to classic elastomers. One source of this observed deviation is the absence of increased elasticity at large deformations due to strain-induced crystallization and chain alignment. These phenomena are not observed in amorphous hydrated PGSF-DPBM networks because of the presence of solvents that can disrupt secondary intramolecular interactions such as hydrogen bonding. The presence of water as a solvent will also reduce the strain-induced crystallization. Hydration may also be responsible for the subtle increase in the maximum elongation at break as the DS of PGSF-DPBM networks increases to 59% from 45%, a 1.33 fold increase in DS. The Tg of PGSF-DPBM networks ranges from -10 oC to 11.5 oC, which suggests that PGSF-DPBM networks exist in the rubbery state at room temperature (Table 2). Higher Tg values are observed for networks with higher DS. This observation can be attributed to the following trends. First, increasing the molar ratio of aromatic groups for both the furan-modified PGSF and DPBM crosslinker increases the Tg.19 Furthermore, increasing the volumetric crosslink density reduces chain mobility and increases the Tg as noted in previous reports.20 Finally, the increased mass density directly reduces the free volume of the network.19 The water content of hydrated networks ranges from 11 ± 2% to 44 ± 6% (w/w) and is inversely related to the DS. Networks prepared from PGSF with a relatively higher DS exhibit an increase in mass density, a reduced effective Mc, and thus a reduction in the extent of hydration. Increasing the DS of PGSF from 22% to 59% also reduces the sol content of crosslinked PGSF-DPBM networks (Table 2). PGSF-DPBM networks are hydrolytically labile: the normalized mass of networks drops to 73 ± 10%, 70 ± 8%, and 13


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95 ± 5%, respectively for DS = 38%, 45%, and 59% after 32 days in hydrolytic buffer. PGSF-DPBM networks prepared from PGSF with DS = 22% lose macroscopic mechanical integrity after 20 days of in vitro hydrolytic degradation with roughly 20% mass loss (Figure 6). Networks with DS = 45% lose approximately 36% of their mass after 64 days (data not shown). The rates of mass loss and network disintegration of this class of materials are expected to be more rapid when used as an implantable material. The presence of enzymes and the formation of reactive oxygen species will likely accelerate in vivo degradation as previously described.7 Fabrication of complex geometries. Samples with complex helical geometry were fabricated by utilizing molecular scale reconfigurability of PGSF-DPBM networks (Figure 7). The values of Tg1, Tg2, and Tg3 are 6, 0, and 6 oC, respectively (Figure S6, Supporting Information). The observation of Tg1 > Tg2 after heating at 120 oC for 24 hours suggests that the retro DA cycloaddition reaction proceeds as expected. High temperature environments employed to induce the retro DA cycloaddition reactions can produce several side reactions in addition to the cleavage of the DA adduct. Such reactions include homopolymerization of the bifunctional maleimides or irreversible side reactions related to the formation of imino heterocycles.21 However, recent work suggests that crosslinked networks formed from monosubstituted furan heterocycles exhibit a high degree of reversibility without deterioration.22 Furthermore, reversibility in PGSF-DPBM adducts can likely be attributed to the bias towards an endergonic system due to the presence of a strong electron-withdrawing groups (esters in the case of PGSF) as substituents on furan heterocycles.23 The substituents utilized in PGSF-DPBM networks therefore can actively promote the retro DA cycloaddition. These observations contrast 14


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the polymer networks described by Laita et al.21 In the aforementioned work, furan groups are directly incorporated into polymer backbone. Aliphatic substituents on furan heterocycles increase the stability of the DA adduct: retro DA cycloaddition reactions are not observed in the polymer networks. Taken together, PGSF-DPBM networks heated at 120 oC will preferentially undergo retro DA cycloaddition reactions compared to other side reactions because of the composition of the substituents. The crosslinked network is recovered by inducing the forward DA cycloaddition reaction in PGSF-DPBM networks at 92 oC. The equilibrium of forward and retro DA cycloaddition reactions is highly sensitive to temperature. Retro DA cycloaddition reaction could proceed at 92 oC. However, the kinetics of retro DA cycloaddition reaction is significantly slower compared to the forward reaction. It is assumed that the retro DA cycloaddition reactions are negligible compared to the forward DA cycloaddition reactions at 92

o

C.13 The extent of crosslinking in PGSF-DPBM polymers was

comparable before and after network reprogramming as inferred by Tg1 = Tg3. Molecularscale reprogramming of PGSF-DPBM networks was demonstrated through the preparation of shape-memory networks. A sample with helical geometry is fixed into a temporary planar geometry below its Tg. When immersed in water at room temperature, the sample was heated above its Tg and returned to its permanent helical shape (Supporting Video). PGSF-DPBM networks prepared without the reprogramming step are not able to recover the complex geometry upon glassy-rubbery transitions at elevated temperature. Molecular scale configurability via DA cycloadditions may also enable novel functionalities in biodegradable elastomers. For example, cycloaddition reactions may serve as convenient routes to bioconjugation with small molecules or peptides. 15


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Furthermore, forward DA cycloaddition reaction may also be used as an orthogonal chemistry for the efficient preparation of dual interpenetrating networks.24 Temperatures used to process PGSF-DPBM networks greatly exceed physiological temperatures. Elevated processing temperatures present some potential limitations regarding the incorporation of bioactive components and broad in vivo applications of this new class of biodegradable elastomers. However, it is worth noting that mild temperatures (T = 37 oC) may also be used to achieve forward Diels-Alder cycloaddition. Although crosslinking is physically observed at lower temperatures, the kinetics is much slower. The elevated temperatures employed in the processes described in this work were selected to (1) accelerate the forward crosslinking and (2) achieve reconfigurability and subsequent shape-memory capabilities with complex geometries in a timely manner.

CONCLUSIONS Reconfigurable covalent polyester networks based on DA cycloaddition reactions represent a new strategy for the synthesis and processing of biodegradable elastomers. Crosslinkable biodegradable elastomers that utilize DA cycloaddition reactions can overcome many of the traditional limitations in previously reported systems based on step-growth or chain-growth polymerization mechanisms. The materials described herein can be processed into hydrolytically labile networks with complex geometries and shapememory properties. These capabilities are made possible by incorporating molecular reconfigurability via covalent adaptable networks based on hyperbranched polyesters. This novel class of biodegradable elastomers will have broad utility as mechanically

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robust biomaterials for potential applications in scaffolds for tissue engineering, matrices for drug delivery systems, and temporary bioabsorbable surgical materials.

ASSOCIATED CONTENT Supporting Information Sorption and desorption kinetics of water in polymer networks, proton NMR spectra, additional rheological data, differential scanning calorimetry thermograms. This information is available free of charge via the Internet at http://pubs.acs.org/

Supporting Video Video S1. The shape-memory capabilities of PGSF-DPBM networks (DS = 45%) are demonstrated. The linear temporary shape is fixed in the glassy state by reducing the temperature. The complex permanent helical geometry is recovered immediately upon incubating the sample in a water bath at room temperature. The resulting hydrated network is elastomeric, robust, and retains its permanent shape after temporary physical deformation. This information is available free of charge via the Internet at http://pubs.acs.org/

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Table 1. Theoretical and experimental gel point determination in crosslinked PGSFDPBM networks from Flory-Stockmayer theory and experimental results from rheology and FT-IR kinetics study. DSa

Predicted gel point, pg, theory b

Loss tangent, tan δ

Relaxation exponent, n

22%

0.27

0.11

38%

0.20

45% 59%

Experimental gel point, pg, exp c

0.07

Gelation time, tgel (hours) 166

3.47

0.82

96

0.18

0.16

1.88

0.68

61

0.17

0.18

1.53

0.63

37

0.19

a

Degree of substitution of PGSF on a per sebacic acid basis. From Flory-Stockmayer equation (Eqn 4).13 c Measured by correlating rheology and FT-IR spectroscopy data. b

18


0.16

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Table 2. Thermomechanical properties of covalently crosslinked PGSF-DPBM networks. DS

H2Ocont

Solcont

E0

Toughness

Tg

ρ

nc

Mc

θ

(%)

(%)

(%)

(MPa)

(MPa)

(oC)

(g/cm3)

(mol/m3)

(g/mol)

(o)

22

44 ± 6

44 ± 3

4.0 ± 1.4

0.003 ± 0.001

-10.0

1.01

547

1858

78.3 ± 2.4

38

26 ± 2

31 ± 3

10.5 ± 1.8

0.087 ± 0.007

-8.5

1.04

1441

725

73.1 ± 2.9

45

24 ± 4

18 ± 1

11.5 ± 1.8

0.532 ± 0.151

6.0

1.05

1578

665

74.3 ± 1.0

59

11 ± 2

11 ± 1

38.5 ± 1.8

0.694 ± 0.350

11.5

1.07

5263

204

73.4 ± 0.6

19


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Scheme 1. Preparation of covalent adaptable networks based on PGSF-DPBM. a) Poly(glycerol-co-sebacate) pre-polymer (PGS) is prepared from polycondensation reactions. b) PGS is modified with pendant furan groups using furoyl chloride to yield the hyperbranched multivalent PGSF network precursors. c) PGSF is coupled with bifunctional DPBM crosslinker through forward DA reactions to form the final crosslinked network.

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Figure 1. Representative ATR-FTIR spectra of PGS pre-polymer, furan-modified PGS (PGSF), and crosslinked PGSF-DPBM networks (DS = 45%). The absorption peak centered about 1010 cm-1 corresponds to the furan C-O-C ether stretch. The C=C in the DA adduct is present at the 1500 cm-1 is visible after crosslinking thereby confirming the formation of PGSF-DPBM networks based on forward DA cycloadditions.


Biomacromolecules

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Figure 2. a) Evolution of the furan peak at 1010 cm-1 over time for PGSF-DPBM (DS = 59%). b) Evolution of the furan peak at 1010 cm-1 over time for PGSF-DPBM (DS = 22%). Spectra were recorded every 24 hours for both compositions. The spectra at the top of each graph represents the initial recording (t = 0 hr). The subsequent plots represent the spectra for samples cured for 24 and 48 hours, respectively.


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Figure 3. a) Time course of the loss tangent (tan δ) during crosslinking reaction between PGSF (DS = 59%) and DPBM at 92 OC, obtained at different frequencies. The crossover point at 37 hours indicates a temporary frequency independence of the loss tangent and is taken as the critical gel point of the network. b) Time evolution of the storage and loss modulus. The time at which there was a sharp increase in the moduli coincides with the time a temporary frequency independence of the loss tangent occurred. c) The gel times of PGSF-DPBM networks with different DS are compared across two types of rheological measurements. 23 ACS Paragon Plus Environment

Biomacromolecules

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Figure 4. Comparison between pg theoretical calculated from Flory-Stockmayer relationship (ptheory, Eqn. 4) and pg measured from rheology and spectroscopy measurements (pexperimental). The normalized gel point is defined as ⁄ . A value of unity represents congruence between the experimental and predicted values of the extent of reaction at the gel point. Good agreement between experimental and theoretical gels points is observed for PGSF-DPBM networks prepared with DS of both 45% and 59%. Values of 1 are observed for PGSF-DPBM networks prepared using smaller DS. These data suggest that lower DS leads to an overestimation of the theoretical gel point (See Text).


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Figure 5. Stress-strain curves for PGSF-DPBM elastomers prepared from using PGSF with different DS. a) The tensile Young’s modulus range from 4 to 34 MPa while (b) The maximum strain at failure increases from 5 to 58% as the DS increases. (c) The values of tensile Young’s moduli are plotted for elastomeric PGSF-DPBM networks prepared from different values of DS. 25 ACS Paragon Plus Environment

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Figure 6. In vitro degradation of PGSF-DPBM networks prepared from PGSF with (a) DS = 22% and DS = 38%; (b) DS = 45% and DS = 59%. After 64 days, the networks with DS = 45% lost approximately 36% of their masses (data not shown in graph). Networks with DS = 22% lost macroscopic mechanical integrity after 20 days with approximately 20% mass loss (See Text). 26 ACS Paragon Plus Environment

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Figure 7. a) Schematic representation of fabrication process of elastomer with complex geometries. (i) PGSF-DPBM is prepared into planar films containing stress-free networks. (ii) Elastomeric networks are shaped into arbitrarily complex geometries at room temperature thereby inducing residual stresses within the polymeric network. (iii) The retro DA cycloaddition reaction is inducted by heating networks to 120 oC for 24 hours. (iv) Scrambling and subsequent forward DA cycloaddition reaction by cooling the network to 92 oC for 24 hours relieve the residual stress in the networks. This process in total reprograms the networks into the complex permanent geometries. b) Photographic images of PGSF-DPBM networks (DS = 45%) demonstrate the macroscopic configuration of the (i) planar temporary shape and (iv) the complex helical permanent geometry.


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ACKNOWLEDGMENTS Funding

provided

by

the

following

organizations:

American

Heart

Association

(12SDG12050297); American Chemical Society (PRF51980DN17); and the Carnegie Mellon University School of Engineering. The authors would also like to thank the CMU Thermomechanical Characterization Facility. NMR instrumentation at CMU was partially supported by National Science Foundation (CHE-0130903 and CHE-1039870).

REFERENCES 1. (a) Kibbe, M. R.; Martinez, J.; Popowich, D. A.; Kapadia, M. R.; Ahanchi, S. S.; Aalami, O. O.; Jiang, Q.; Webb, A. R.; Yang, J.; Carroll, T.; Ameer, G. A., Citric acid-based elastomers provide a biocompatible interface for vascular grafts. Journal of Biomedical Materials Research Part A 2010, 93A (1), 314-324; (b) Nieponice, A.; Soletti, L.; Guan, J.; Hong, Y.; Gharaibeh, B.; Maul, T. M.; Huard, J.; Wagner, W. R.; Vorp, D. A., In vivo assessment of a tissue-engineered vascular graft combining a biodegradable elastomeric scaffold and muscle-derived stem cells in a rat model. Tissue engineering. Part A 2010, 16 (4), 1215-23. 2. (a) Sundback, C. A.; Shyu, J. Y.; Wang, Y.; Faquin, W. C.; Langer, R. S.; Vacanti, J. P.; Hadlock, T. A., Biocompatibility analysis of poly(glycerol sebacate) as a nerve guide material. Biomaterials 2005, 26 (27), 5454-5464; (b) Engelmayr, G. C.; Cheng, M.; Bettinger, C. J.; Borenstein, J. T.; Langer, R.; Freed, L. E., Accordion-like honeycombs for tissue engineering of cardiac anisotropy. Nature Materials 2008, 7 (12), 1003-1010. 3. Amsden, B., Curable, biodegradable elastomers: emerging biomaterials for drug delivery and tissue engineering. Soft Matter 2007, 3 (11), 1335-1348. 4. Webb, A. R.; Yang, J.; Ameer, G. A., Biodegradable polyester elastomers in tissue engineering. Expert Opinion on Biological Therapy 2004, 4 (6), 801-812. 5. Yang, J.; Zhang, Y.; Gautam, S.; Liu, L.; Dey, J.; Chen, W.; Mason, R. P.; Serrano, C. A.; Schug, K. A.; Tang, L., Development of aliphatic biodegradable photoluminescent polymers. Proceedings of the National Academy of Sciences 2009, 106 (25), 10086-10091. 6. Serrano, M. C.; Carbajal, L.; Ameer, G. A., Novel Biodegradable Shape-Memory Elastomers with Drug-Releasing Capabilities. Advanced Materials 2011, 23 (19), 2211-2215. 7. Wang, Y.; Ameer, G. A.; Sheppard, B. J.; Langer, R., A tough biodegradable elastomer. Nature Biotechnology 2002, 20 (6), 602-606.


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TABLE OF CONTENTS FIGURE


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Reconfigurable Biodegradable Shape-Memory Elastomers via Diels

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