High performance shape memory polymer networksbased on rigid nanoparticle coresJianwen Xu and Jie Song1

Department of Orthopedics and Physical Rehabilitation and Department of Cell Biology, University of Massachusetts Medical School, 55 Lake AvenueNorth, Worcester, MA 01655

Edited* by Carolyn R. Bertozzi, University of California Berkeley, Berkeley, CA, and approved March 11, 2010 (received for review October 28, 2009)

Smart materials that can respond to external stimuli are of wide-spread interest in biomedical science. Thermal-responsive shapememory polymers, a class of intelligent materials that can be fixedat a temporary shape below their transition temperature (Ttrans)and thermally triggered to resume their original shapes on de-mand, hold great potential as minimally invasive self-fitting tissuescaffoldsor implants. The intrinsicmechanismfor shapememorybe-havior of polymers is the freezing and activation of the long-rangemotion of polymer chain segments below and above Ttrans, respec-tively. Both Ttrans and the extent of polymer chain participation ineffective elastic deformation and recovery are determined by thenetwork composition and structure, which are also defining factorsfor their mechanical properties, degradability, and bioactivities.Such complexity has made it extremely challenging to achievethe ideal combination of a Ttrans slightly above physiologicaltemperature, rapid and complete recovery, and suitablemechanicaland biological properties for clinical applications. Here we report ashape memory polymer network constructed from a polyhedraloligomeric silsesquioxane nanoparticle core functionalized witheight polyester arms. The cross-linked networks comprising thismacromer possessed a gigapascal-storage modulus at bodytemperature and a Ttrans between 42 and 48 °C. The materials couldstablyhold their temporary shapes for>1 year at room temperatureand achieve full shape recovery ≤51 °C in a matter of seconds. Theirversatile structures allowed for tunable biodegradability andbiofunctionalizability. These materials have tremendous promisefor tissue engineering applications.

nanocompsite ∣ shape memory materials ∣ thermal responsive materials

A thermal-responsive shape memory polymer (SMP) can beimparted with a “permanent” shape above a critical transi-

tion temperature (Ttrans) when it is cast in a mold. For polymers,the Ttrans is either glass transition temperature Tg or meltingtemperature Tm. Such a permanent shape, formed at the elasticstate of the material without external stress, is retained (memor-ized) as the SMP cools to a temperature below its Ttrans. The SMPcan be deformed into a desired “temporary” shape by force atT > Ttrans, and this strained configuration can be fixed as the tem-perature cools below the Ttrans. When a thermal stimulus abovethe Ttrans is reapplied, however, the SMP recovers to its lessstrained permanent shape. This unique shape memory behaviorhas captured the imagination of the biomedical community asscientists strive to design smart implants and tissue scaffolds thatcan be delivered in a minimally invasive configuration and be sub-sequently reverted to a preprogrammed permanent shape in vivo.

Since the first demonstration of degradable SMPs for potentialtissue engineering applications (1, 2), many semicrystalline andamorphous polyester and polyurethane SMP networks have beenreported. Adjustment of thermomechanical properties anddegradability of these SMPs was accomplished by copolymerizingmultiple monomers (3), incorporating inorganic elements intopolymer backbones (4, 5), or applying different cross-linkingmethods (1, 5). The intertwined correlation between the compo-sition and structure of an SMP network and its derived physicalproperties, however, has made it difficult to achieve a combina-

tion of Ttrans, degradation profile, and mechanical strength suita-ble for biomedical applications in a single material. Existingbiocompatible SMP networks contain either untethered polymerchains resulting in plastic deformations and broad transitions orexcessive chain–chain interactions requiring extra energy toovercome. Consequently, they require harsh temperatures tofix temporary shape (<0 °C) (3, 6–8) or trigger shape recovery(>70 °C) that is often slow and incomplete (9–11). Moreover,few existing SMPs possess tunable biofunctionalizability andadequate mechanical strength at body temperature (12).

The shape memory capacity of polymers lies in the entropy-driven tendency for polymer chains to adopt a randomly coiledconfiguration. The intrinsic mechanism for shape memory beha-vior of polymers is the freezing and activation of the long-rangemotion of polymer chain segments below and above Ttrans, re-spectively. To achieve complete freezing of chain segment motionand thus prevent chain recoiling below Ttrans (temporary shapefixation) and full activation of chain recoiling above Ttrans (shaperecovery), a homogenous SMP network consisting of identicalchains with tunable chain–chain interactions would be ideal.We hypothesize that a network cross-linked from a well-definedstar-branched macromer containing a rigid nanoparticle corecould meet such requirements (Scheme 1). The rigid, symmetriccore defines the spatial distribution of polyester arms upon cross-linking and decreases excessive chain–chain interactions as oftenoccurred in linear polyester networks. The multiple reactive endsof the macromer are designed to achieve adequate mechanicalstrength via high-density cross-linking and desired bioactivityvia selective end-group functionalization. Here we report sucha nanoparticle-based homogeneous SMP network that exhibitsan extraordinary combination of stable temporary shape fixationand rapid and full shape recovery slightly above physiologicaltemperature with excellent mechanical properties.

Results and DiscussionMacromer Design: Polyhedral Oligomeric Silsesquioxane Core VersusOrganic Core. Previous studies on dendritic and hyperbranchedpolymers suggest that the core architecture (size and rigidity),molecular weight, and chain end composition of branchedpolymer systems could profoundly affect their physical properties(13, 14). Here we chose a polyhedral oligomeric silsesquioxane(POSS) nanoparticle as the core to prepare a star-shaped macro-mer building block for the SMP network (Fig. 1A). This design ismotivated by (i) the well-defined cubic geometry of POSS that en-ables the grafting of up to eight identical polymer arms, (ii) thecapability of the rigid POSS nanoparticle in controlling the graftedpolymer chain motions on a molecular scale (15), and (iii) thedemonstrated biocompatibility of POSS (16). Polylactides (PLAs)

Author contributions: J.X. and J.S. designed research, performed research, analyzed data,and wrote the paper.

The authors declare no conflict of interest.

*This Direct Submission article had a prearranged editor.1To whom correspondence should be addressed. E-mail: Jie.Song@umassmed.edu.

This article contains supporting information online at www.pnas.org/cgi/content/full/0912481107/DCSupplemental.

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of various lengths were grafted to the octahydroxylated POSS core(Fig. S1) via ring-opening polymerization of D,L-lactide (Fig. 1A)to give star-branched macromer POSS-ðPLAnÞ8 (n ¼ 10; 20; 40;the number of lactide repeating units) in near quantitative yieldsand low polydispersity (Table S1). End-group titration by proteinnuclear magnetic resonance (1H NMR) (17, 18) confirmed thesuccessful grafting of eight PLA arms to each core (Table S2and Fig. S2). The hydroxyl end groups of the macromer were thenreacted with hexamethylene diisocyanate (Table S3) to form cross-linked POSS-SMP via urethane linkages.

Mono- and difunctional POSS nanoparticles have been pre-viously utilized in SMP designs wherein interactions betweenthe particles themselves (crystallization tendency of POSS)were exploited to form percolating physical cross-links within achemically cross-linked system (4, 5, 19). Although the competi-tive crystallizations between POSS and polymer domains resultedin unconventional thermoplastic properties, neither the tempera-

tures nor the broadness of the thermal transitions were suitablefor biomedical applications. Our hope was that the octafunctionalPOSS-based macromers could be cross-linked to form an amor-phous network wherein the rigid POSS cores impart controlledinteractions between the tethered PLA arms. In addition tosynthesizing SMPs built from a rigid POSS core, we preparedSMPs comprising a flexible organic core (Fig. 1A). Functionalizedwith the same PLA arms (Table S2, Figs. S3 and S4) as the POSS-SMPs, this all-organic SMP network (Org-SMP) allowed forcomparative studies to determine the impact of core structureon physical properties.

Impact of Core Structures on Thermal-Mechanical and Shape MemoryProperties. We first compared the thermal-mechanical propertiesof the SMPs comprising the rigid POSS core with thosecontaining the flexible organic core. As representatively shownin Fig. 1B, POSS-SMP-20 and Org-SMP-20, cross-linked fromPOSS-ðPLA20nÞ8 and Org-ðPLA20Þ8, respectively, both pos-sessed gigapascal (GPa)-storage moduli at body temperature.Further, they exhibited similar temperature-dependent viscoelas-tic properties, with the storage modulus sharply descending fromthe GPa-glassy state to a megapascal-elastic plateau around theirrespective glass transitions, both within a narrow transition tem-perature range (width at the half-peak height of the transition, orWHPH, <10 °C). The narrow glass transitions observed in bothsystems support our hypothesis that a homogenous networkcross-linked from star-branched macromer building blockscontaining identical polymer chains could respond to the thermalstimuli more uniformly than structurally ill-defined networks.The glass transition temperature (Tg) of POSS-SMP-20, however,was >10 °C lower than that of Org-SMP-20, and the storage mod-ulus drop around the Tg was more pronounced for POSS-SMP-20. Given that POSS-ðPLA20Þ8 and Org-ðPLA20Þ8 had similarpolymer chain compositions, molecular weights, and polydisper-sity (Table S1) and identical numbers of urethane cross-linking

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Fig. 1. Preparation and thermal-mechanical properties of SMPs containing POSS (POSS-SMP) versus organic (Org-SMP) cores: (A) synthesis and cross-linkingof macromers; (B) storage modulus (E′)-temperature and loss angle (Tan δ)-temperature (denoted by black arrows) curves of POSS-SMP-20 versus Org-SMP-20;(C) recovery rates of POSS-SMP-20 (red arrows) versus Org-SMP-20 (blue arrows) from an identical rolled-up temporary shape (Left) to fully extended rectangle(30.0 mm × 6.0 mm × 0.5 mm) in water at 51 °C.

cross-linking

Nanoparticle-strengthed homogenous network with tailored chain-chain interaction

Biodegradable polyester (e.g. PLA) arms

Nanoparticle core (e.g. POSS)

Bifunctional cross-link (e.g. urethane groups)

Hybrid star-branchedmacromer building block

Scheme 1. Depiction of a nanostructured SMP network.

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sites (Table S2), the observed difference in thermal-mechanicalproperties can be attributed to the different size and rigidityof POSS versus the organic cores. Indeed, grid search analysis(Fig. S5) revealed more torsional freedom at each polyol branch-ing point in the bulkier POSS core (1; 103-Å3 molecular volume)than in the organic core (539-Å3 molecular volume), suggestingthat the PLA arms could distribute more homogenously inthe cross-linked POSS-SMP network with less excessive chainentanglement.

We then examined the shape memory performance of theSMPs comprising different cores. Both POSS-SMP-20 andOrg-SMP-20 could be stably fixed at a temporary shape withinseconds upon cooling to room temperature, indicating completefreezing of chain segment motions below the Tg in both networks.At 51 °C, the rate of shape recovery of POSS-SMP-20 (<3 s) wasmuch faster than that of Org-SMP-20 (>20 s) (Fig. 1C andMovie S1), whereas at 73 °C, they recovered at a similar rate(<1 s; Movie S2). These observations are consistent with the no-tion that it is the increased polymeric segment motions (entropyelasticity) above the Ttrans (Tg in this case) that drive the shaperecovery process. The observed differential shape memoryperformances support our hypothesis that the nanostructuredmolecular network imposed by the POSS core could translate,on a macroscopic scale, into more rapid shape recovery at a lowertriggering temperature.

Tuning Thermal-Mechanical and Shape Memory Properties of POSS-SMPs. To explore the possibility of tuning thermal-mechanicalproperties and biodegradation rates of POSS-SMP, macromerscontaining various PLA arm lengths (n ¼ 10; 20; 40) were pre-pared and cross-linked. Differential scanning calorimetry(DSC) revealed a single narrow endothermic transition for eachcross-linked POSS-SMP with no crystalline phase transitionsdetected (Fig. 2A). This observation, as well as observed opticaltransparency, supports an amorphous network structure wherePOSS cores are well-dispersed rather than crystallizing as inthe case of the SMPs containing mono- or difunctional POSS(4, 5, 19). Unlike linear high molecular weight poly(D,L-lactide)(Tg ∼ 54 °C) (20), the Tgs exhibited by POSS-SMPs, 42.8–48.4 °C,are more suitable for biomedical applications. Conventionalcross-linked elastomers and the Org-SMPs (Fig. S6) have Tgs thatdecrease as chain lengths between cross-linking points increase.By contrast, the Tgs of POSS-SMPs (TDSC

g , Table 1) increased asthe PLA chain length of the macromer increased, suggesting amore profound impact of the rigid POSS core on the chain–chaininteractions within the cross-linked network than the flexibleorganic core. Dynamic mechanical analysis (DMA) revealed asimilar relationship of TDMA

g to chain length (Fig. 2B), suggestingthat the effect of POSS on chain–chain interaction becomes lesssignificant with longer PLA chains.

We next sought to investigate how the strengths of these POSS-SMPs vary as a function of PLA chain length and temperature byusing DMA. We found that all three POSS-SMPs possessed simi-lar glassy state storage moduli, >2.0 GPa, at body temperature(E0

37 °C, Table 1). Interestingly, this value is ideal for cortical bonereplacement materials (21). It would allow POSS-SMPs to beused for weight-bearing in vivo applications. On the otherhand, the storage modulus of POSS-SMP in the elastic state(e.g., E0

85 °C), which is determined by the density of cross-links(22), decreased as the PLA arm length increased.

An exciting revelation from these experiments was the obser-vation of extremely narrow glass transitions (WHPH < 10 °C,Table 1) accompanied by sharp storage modulus changes of upto 3 orders of magnitude for POSS-SMPs. By contrast, previousSMP networks typically exhibited wide glass transitions(WHPH > 20 °C) with no more than 2 orders of magnitude mod-ulus changes around the Ttrans (9–11, 23). The steep and narrowthermomechanical transitions exhibited by POSS-SMPs make it

possible to achieve both temporary shape fixing and permanentshape reversion within a narrow physiologically relevant tempera-ture range. We realized this notion by stably fixing various tem-porary shapes at room and body temperatures for >1 year andinstantly (within a matter of seconds) recovering their permanentshapes≤51 °C (Fig. S7 andMovies S3–S5). Clinical thermal treat-ments employing a combination of such mild temperature andshort exposure time were shown to be well tolerated by humansclera (24, 25), epidermis (24, 25), and bony tissues (26–28).These properties are critical for implantation applications,wherein the implant material must be delivered in a stable mini-mally invasive shape and subsequently reverted to a permanentshape in vivo. Quantitative assessment of the shape memoryperformance through stress-controlled one-way shape memorycycles (4, 29) verified that all POSS-SMPs exhibited a highshape-fixing ratio (Rf > 91%) and shape recovery ratio

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(Rr≈100%) (Fig. 2C and Table 1), with POSS-SMP-40 achieving∼100%Rf and Rr after the second cycle.

Finally, we found that the temperature window for deforma-tion recovery (Fig. 2C) correlated well with the glass transitionprocess (Fig. 2B), supporting that entropy elasticity during glasstransition was the driving force for shape recovery. The number-averaged molecular weight between the cross-linking points (Mc),derived from Flory’s affine network model (22), also correlatedwell with the corresponding PLA arm lengths of the macromer(Table 1). This observation suggests that almost all PLA armswere tethered by urethane cross-links and could participate inthe elastic deformation and recoiling, thereby contributing tothe rapid and complete shape recovery.

Biofunctionalization of POSS-SMP. To demonstrate the possibility ofpresenting biological signals on POSS-SMP without compromis-ing its desired thermal-mechanical properties, we covalentlycoupled a model integrin-binding peptide Arg-Gly-Asp-Serknown to mediate biomaterial–cell interactions (30, 31) by usinga 2-step modification strategy. First, POSS-ðPLA20Þ8 wascross-linked in the presence of 3-azido-1-propanol (step 1,Fig. 3A) to introduce azido end groups in POSS-SMP-20-Az[see Fig. S8 for Fourier transform infrared (FTIR) spectroscopycharacterizations]. POSS-SMP-20-Az exhibited storage moduli of2.1 GPa at 37 °C and 1.2 MPa at 85 °C (Fig. 3B), with a slightlyreduced TDMA

g comparing to POSS-SMP-20, indicating that theincorporation of the azido end groups at the cost of 25% reduc-tion in urethane cross-links did not deteriorate the thermal-mechanical properties. No notable difference in temporary shapefixing but slightly slower shape recovery were observed forPOSS-SMP-20-Az at 51 °C, presumably because of the reductionin the number of tethered chains participating in the elastic de-formation and recovery. POSS-SMP-20-Az was then coupled witha fluorescently labeled alkyne-functionalized Arg-Gly-Asp-Ser, (4-pentynoic acid)-Gly-Arg-Gly-Asp-Ser-K(FTIC)-COOH,by using high-fidelity “click” chemistry (step 2, Fig. 3A) (32).The covalent attachment of the fluorescently labeled peptidewas confirmed by fluorescence microscopy (Fig. 3C). No changein thermal-mechanical properties except for a minor reduction inTDMAg (<1 °C, Fig. 3B) was detected upon the attachment of the

peptide to POSS-SMP-20-Az. This strategy can be extended tointroduce a wide range of bioactive molecules to POSS-SMPwhile maintaining its desired physical properties. Finally, wedemonstrated that POSS-SMPs can be engineered for variedhydrolytic degradation rates, e.g., 50% weight loss in 3–9 months,through the alteration of PLA chain length (Fig. S9). Together,these features make it possible to prepare POSS-SMP–based tis-sue scaffolds and implants tailored for patient- and defect-specificbiochemical environment and tissue repair/regeneration rate.

Conclusions and PerspectivesWe have prepared biodegradable SMP networks with an unpre-cedented combination of excellent mechanical properties(E0 > 2 Gpa) and stable temporary shape fixing at room and bodytemperatures, fast (<3 s) and complete shape recovery within a

Table 1. Summary of thermal-mechanical properties of POSS-SMPs

Sample name TDSCg (°C)* TDMAg (°C)† Tan δ‡ E037 °C (MPa)§ E085 °C (MPa)¶ WHPH (°C)∥ Mc (Dalton)** Rf (%)†† Rr (%)††

POSS-SMP-10 42.8 51.8 ± 0.4 2.34 ± 0.04 2,027.0 ± 38.3 4.2 ± 0.1 9.4 ± 0.3 876.9 96.0 100POSS-SMP-20 45.4 56.0 ± 0.8 2.68 ± 0.06 2,286.8 ± 62.7 2.3 ± 0.1 8.7 ± 0.2 1,563.3 91.6 100POSS-SMP-40 48.4 57.9 ± 0.5 2.69 ± 0.04 2,234.7 ± 17.1 1.2 ± 0.1 9.1 ± 0.3 2,996.3 100 95.2 (100)‡‡

*Glass transition temperatures as determined from the DSC scans.†Glass transition temperatures as determined from the Tanδ-temperature curves.‡Peak value of Tanδ-temperature curves.§Storage moduli of the glassy state determined at 37 °C.¶Storage moduli of the elastic state determined at 85 °C.∥Peak WHPH of the Tanδ-temperature curves.**The number-averaged molecular weight between the cross-linking points of the polymer network (Mc) calculated from E0 ¼ ρRT∕Mc, where ρ is the

density of the polymer network (ρ ¼ 1.27 g cm−3), R is the gas constant (R ¼ 8.314), T is an elastic state temperature (T ¼ 358 K), and E′ is storagemodulus measured at this temperature (E085 °C).

††Shape fixing ratio (Rf) and shape recovery ratio (Rr) calculated from the second one-way shape memory cycle shown in Fig. 2C.‡‡The Rf value shown in the parentheses was calculated from the third cycle.

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Fig. 3. Chemical modification of POSS-SMP with a bioactive peptide:(A) Synthetic scheme illustrating the introduction of azido groups duringthe covalent cross-linking of POSS-ðPLA20Þ8 and subsequent conjugation offluorescently labeled integrin-binding peptide to POSS-SMP via “click” chem-istry. (1) 100 ppm DBTDL, CH2Cl2, argon, r.t., 12 h; 75 °C, argon, 24 h; 75 °Cunder vacuum, 48 h. (2) Aqueous solution of CuSO4 (2.5 mM) and LðþÞ-ascor-bic acid sodium salt (7.5 mM), r.t., 24 h. (B) Storage modulus (E′;)-temperaturecurves and loss angle (Tan δ)-temperature curves (denoted by black arrows) ofPOSS-SMP-20, POSS-SMP-20-Az, and POSS-SMP-20-Peptide. (C) Differentialinterference contrast (DIC) and fluorescent (Fl) micrographs confirmingthe covalent conjugation of the fluorescently labeled peptide via clickchemistry. In the negative control (Left), POSS-SMP-20-Az was exposed tothe fluorescently labeled peptide in the absence of ascorbic acid underotherwise identical reaction conditions.

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narrow physiologically relevant temperature range (<51 °C), andtunable bioactivities for biomedical applications. The key struc-tural feature of the network is the well-defined star-branchedmacromer building blocks containing identical polymer chainsthat enabled high-density cross-linking, selective biofunctionali-zation, and more uniform response of the polymer chains tothermal stimuli. Compared to the all-organic polyhydroxyl core,the bulkier and more rigid POSS nanoparticle core is more effec-tive in minimizing excessive global entanglement of the tetherednetwork chains and in maximizing their participation in the shapememory process. The strategic use of well-defined nanoparticlesto mediate polymer chain–chain interactions and the bottom-upapproach towards the control over the structure, mechanicalproperties, and chemical functionalities may also be useful inother situations wherein the optimization of multiple propertiesis desired.

A major roadblock in translating scaffold-based tissue engi-neering into clinical practice is the lack of materials combiningtissue-like mechanical and biochemical properties with clinicallyrelevant deployability to enable their safe delivery and integrationwith target tissue (33–36). The POSS-SMPs reported here havegreat potential as self-fitting tissue scaffolds and implants wheretheir unique properties could address unmet medical challenges,for instance, in the reconstruction of skeletal and craniofacialdefects that are characterized with complex and irregular geome-tries, particularly at weight-bearing locations. Conventionalprefabricated weight-bearing scaffolds (e.g., stiff polymers andceramics) do not readily conform to such defects. On the otherhand, injectable formulations that could penetrate into suchdefects and solidify in situ are known for health concerns becauseof the exothermic solidification process (e.g., leading to tissuenecrosis) and potential leaks. A POSS-SMP scaffold that is castin vitro with the desired size/shape on the basis of the MRI orradiographic scans of the defect can potentially overcome suchlimitations by enabling its delivery in a less invasive compressedconfiguration and subsequently conforming to the defect uponbrief and safe thermal triggering (e.g., via catheter heating).Coupled with its excellent mechanical properties at body tem-perature, the more securely anchored POSS-SMP implant mayalso reduce the need for auxiliary metallic fixators, which oftenrequire a second surgery to remove and obscure postoperativeradiographic monitoring of the osteointegration of the implant.Finally, the tunable biofunctionality and degradability of POSS-SMPs opens the possibility of locally delivering therapeuticsexpediting the healing of the defect while enabling the implantedscaffold to “vanish” after fulfilling its function.

Materials and MethodsDetailed methods and results on molecular modelling of POSS and organiccores, gel permeation chromatography (GPC), NMR, FTIR, and high resolutionmass spectrometry (HRMS) characterizations, end-group titration ofPOSS-ðPLAnÞ8 and Org-ðPLAnÞ8 macromers, formulations for preparingPOSS-SMP and Org-SMP, and thermal-mechanical properties of Org-SMPcan be found in SI Text. Detailed protocol for stress-controlled cyclic ther-mal-mechanical testing to quantify shape memory properties is also includedin SI Text, and the movies demonstrating the shape memory processes areavailable online.

Synthesis of Macromer Building Blocks POSS-ðPLAnÞ8. Macromer buildingblocks POSS-ðPLAnÞ8 (n ¼ 10; 20; 40) were synthesized from an octahydroxy-lated POSS core as summarized in Fig. S1 and Table S1.

Octa(3-hydroxypropyldimethylsiloxy)octasilsesquioxane (POSS core).The POSS core was synthesized according to the literature with slight mod-ifications. PSS-Octakis(dimethylsilyloxy) substitute (≥97%, Aldrich, 5.012 g,4.923 mmol) was dissolved in 25 mL dry toluene and bubbled with argonfor 20 min upon the addition of ally alcohol (98.5%, Aldrich, 3.430 g,59.07 mmol). Platinum (0)-1,3-divinyl-1,1,3,3-tetramethyldisioxane complex[Pt(dvs), 3 wt % solution in xylene, Aldrich], diluted by dry toluene into2.0-mM solution, was then introduced to the reaction mixture (1.5 mL,

0.003 mmol) by a syringe. The reaction proceeded for 1 h at room tempera-ture (r.t.) under argon atmosphere and another 1.5 h under reflux at 90 °C.Upon cooling to r.t., the bottom phase of the mixture was separated and con-centrated under vacuum. The crude product was washed with toluene(50 mL, 3 times) to yield white solid octahydroxylated POSS core in >90%yield. 1H NMR (400 MHz, CDCl3): δ3.57 (16H, t, J ¼ 7.2 Hz), 2.85 (b), 1.63(16H, p, J ¼ 8.0 Hz), 0.60 (16H, t, J ¼ 8.5 Hz), 0.15 (48H, s) ppm. HRMS forC40H104O28Si16Na½Mþ Naþ�: calculated 1,505.2921; observed 1,505.2928.

Star-Shaped Macromer POSS-ðPLAnÞ8. PLAs of varying lengths(n ¼ 10; 20; 40) were grafted to the POSS core via ring-opening polymeriza-tion of D,L-lactide. For the synthesis of POSS-ðPLA20Þ8, POSS core (0.604 g,0.408 mmol) and D,L-Lactide (≥97%, 4.707 g, 3.266 mmol) were degassedat r.t. for 1 h in a 25-mL Kjeldahl reaction flask and then melted at 125 °C. Tin(II) 2-ethylhexanoate [SnðOctÞ2, 95%, 0.265 mg, 0.654 μmol] wasintroduced by a syringe and the melt was reacted at 130 °C for 20 h. Uponcooling, the resulting solid was dissolved in chloroform (60 mL) and precipi-tated in hexane (200 mL). The precipitation purification was repeated 3 timesbefore the product was dried in vacuum oven at r.t. for 24 h and then at 90 °Cfor 48 h. Transparent solid POSS-ðPLA20Þ8 was obtained in near quantitativeyield (95–99%) with a polydispersity index of 1.19. 1H NMR (400 MHz, CDCl3):δ5.24–5.12 (152H, m), 4.42–4.25 (8H, b), 4.15–4.00 (16H, b), 1.68–1.49 (496H,m), 0.65–0.44(16H, b), 0.16–0.05 (48H, s) ppm. 13C NMR (100 MHz, CDCl3):δ169.78–169.28, 69.31–69.13, 67.79, 66.83, 22.27, 20.65, 16.97–16.77, 13.47,−0.32 ppm. Macromers POSS-ðPLA10Þ8 and POSS-ðPLA40Þ8 were synthesizedand purified in a similar fashion with near quantitative yield. Representative1HNMR spectra and summary of GPC characterizations of the purifiedmacro-mers are shown in Fig. S1 and Table S1.

Synthesis of Macromer Building Block Org-ðPLAnÞ8. Macromer building blockOrg-ðPLAnÞ8 was synthesized from a multihydroxylated organic core asshown in Fig. 1A.

Acetal-protected organic core. 4-Dimethylamino-pyridinium-p-toluene-sulfate (DPTS) and 2,2,5-trimethyl-1,3-dioxane-5-carboxylic acid (TMDC)were prepared according to the literature. TMDC (10.79 g, 61.97 mmol),di(trimethylolpropane) (3.30 g, 13.18 mmol), and DPTS (3.88 g, 13.18 mmol)were dissolved and stirred in 200 mL anhydrous pyridine under N2. Dicyclo-hexylcarbodiimide (14.06 g, 68.17 mmol) was added after the solutionbecame clear and stirred at r.t. for 20 h under N2. The mixture was filtered,and the filtrate was concentrated under vacuum. The resulting light-yellowsolids were redissolved in hexane and filtered, and the filtrate was concen-trated in vacuum for further purification by flash chromatography (silica gel,Merck grade 9384, 230–400 mesh; ethyl acetate/hexane 2∶3) to yield 4.5 gcolorless oil (Rf ¼ 0.3, 40%). 1H NMR (400 MHz, CDCl3): δ4.15–4.12 (8H, d,J ¼ 12.0 Hz), 4.07 (8H, s), 3.62–3.59 (8H, d, J ¼ 12.0 Hz), 3.29 (4H, s), 1.49–1.43 (4H, q, J ¼ 7.4 Hz), 1.39 (12H, s), 1.33 (12H, s), 1.12 (12H, s), 0.87–0.84(6H, t, J ¼ 7.4 Hz) ppm. 13C NMR (100 MHz, CDCl3): δ 173.97, 98.27, 71.11,66.25, 64.31, 42.56, 42.30, 25.37, 23.15, 22.32, 18.80, 7.72 ppm. HRMS forC44H78NO17½Mþ NHþ

4 �: calculated 892.5270; observed 892.5273.

Organic Core. Acetal-protected organic core (3.2 g, 3.66 mmol) was depro-tected with 15.0 g resin (Amberlite IR-120, Hþ form, 16–45 mesh) in 75 mLmethanol for 15 h. After removing the resin by filtration, the filtrate wasconcentrated in vacuo to give colorless oil, which was redissolved in 16 mLmethanol and precipitated in 100 mL anhydrous ethyl ether. The precipitatewas filtered and dried under vacuum over P2O5 to give white powder (2.2 g,84%). 1H NMR (400 MHz, CD3OD): δ 4.06 (8H, s), 3.70–3.67 (8H, d,J ¼ 10.9 Hz), 3.62–3.59 (8H, d, J ¼ 10.9 Hz), 3.37 (4H, s), 3.31 (s) 1.54–1.49(4H, q, J ¼ 7.4 Hz), 1.16 (12H, s), 0.94–0.90 (6H, t, J ¼ 7.4 Hz). 13C NMR(100 MHz, CD3OD), δ 175.11, 70.76, 64.71, 63.99, 50.66, 42.12, 23.01,16.20, 6.75 ppm. HRMS for C32H59O17½Mþ Hþ�: calculated 715.3749; observed715.3752.

Macromer Org-ðPLAnÞ8. Organic macromers were prepared in the sameway as POSS-ðPLAnÞ8. Representatively, organic core (1.073 g, 1.50 mmol)and D,L-lactide (17.35 g, 120.38 mmol) were reacted at 130 °C withSnðOctÞ2 (25.02 mg, 61.75 μmol) to prepare Org-ðPLA20Þ8 in near quantitativeyield. 1H NMR (400 MHz, CDCl3): δ 5.25–5.12 (159H, m), 4.40–4.28 (16H, m),4.24–4.20 (8H, b), 4.03 (8H, b), 3.25 (4H, s), 2.84–2.60 (OH, b), 1.75–1.71 (4H,m), 1.69–1.45 (m), 1.23 (12H, b), 0.90–0.81 (6H, m) ppm. 13C NMR (100 MHz,CDCl3): δ175.31–175.19, 169.90–169.36, 72.66, 69.38–69.19, 66.87, 66.82,46.65, 42.33, 20.70, 20.25, 16.94–16.87, 15.95, 7.68 ppm. Representative

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1H NMR spectra and summary of GPC characterizations of the purifiedmacromers are shown in Fig. S3 and Table S1.

Preparation and Characterization of POSS-SMPs, Org-SMPs, and Azido-Functio-nalized POSS-SMP-20-Az. In a typical procedure, POSS-ðPLAnÞ8 or Org-ðPLAnÞ8,hexamethylene diisocyanate (HDI, ≥98.0%, Fluka), and 3-azidopropan-1-olwere mixed (molar ratios shown in Table S3) in 2.5 times (wt∕wt) dichloro-methane. A catalytic amount (100 ppm) of dibutyltin dilaurate (DBTDL,≥95%, Aldrich) was added. The solution was stirred for 2 h at r.t. before beingpoured into Teflon molds. The solvent was evaporated at r.t. overnight underAr, and the material was further cross-linked at 75 °C under Ar for 24 h. Thefinal product was heated at 75 °C under vacuum for 48 h to remove residuevolatiles. The complete conversion of HDI to urethane cross-links wasconfirmed by the disappearance of the FTIR absorption at 2; 280 cm−1 (forisocyanate) upon cross-linking (Fig. S8). To determine the efficiency ofcross-linking, POSS-SMPs were extracted in chloroform (100 mL∕g) for 12 hand then dried under vacuum for 24 h. Gel content, defined by the ratioof dry weight before and after the solvent extraction, was calculated(Table S3).

Preparation of POSS-SMP-20-Peptide. A specimen of POSS-SMP-20-Az(30.0 mm × 6.0 mm × 0.5 mm) was immersed into a 50-mL aqueous solution

of (4-pentynoic acid)-Gly-Arg-Gly-Asp-Ser-Lys(FTIC)-COOH (BiomerTechnol-ogy; 1.0 mg∕mL), to which 0.8 mL CuSO4 aqueous solution (2.5 mM) wasadded. The mixture was degassed under argon for 1 h before 0.8 mL de-gassed solution of LðþÞ-ascorbic acid sodium salt (7.5 mM) was injected.The reaction was carried out at r.t. under argon for 24 h. The peptide-modified specimen was washed with water and ethanol for 1 h, respectively,and dried under vacuum.

DMA. The dynamic mechanical properties of the POSS-SMPs and Org-SMPwere determined on a Q800 DMA (TA Instruments) equipped with tensilefilm clamps. Specimens with dimensions of 30.0 mm × 6.0 mm × 0.5 mmwere used for testing. The temperature was ramped from r.t. to 110 °C ata heating rate of 2.0 °C∕min. A 0.02% strain amplitude and 1.0-Hz frequencywere applied. Three specimens were tested for each sample.

ACKNOWLEDGMENTS. This work was supported by National Institutes ofHealth Grants R01AR055615 (to J.S.) and R01GM088678 (to J.S.) and Amer-ican Cancer Society Grant IRG 93-033 (to J.S.). Core resources supported byDiabetes Endocrinology Research Center Grant DK32520 and National Centerfor Research Resources Grant S10 RR021043 were also used. J.S. is a memberof the University of Massachusetts Medical School Diabetes and Endocrinol-ogy Research Center (DK32520).

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