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Polymer 52 (2011) 4947e4954

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Polymer

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Impact of shape-memory programming on mechanically-driven recoveryin polymers

Christopher M. Yakackia,*, Thao D. Nguyenb, Roxanne Likosc, Robert Lamellc, Daniel Guigoud,Ken Gallc,d,e,f

aDepartment of Mechanical Engineering, The University of Colorado at Denver, Denver, CO 80217, USAbDepartment of Mechanical Engineering, Johns Hopkins University, Baltimore, MD 21218, USAcWallace H. Coulter School of Biomedical Engineering, The Georgia Institute of Technology, Atlanta, GA 30332, USAdGeorge W. Woodruff School of Mechanical Engineering, The Georgia Institute of Technology, Atlanta, GA 30332, USAe School Materials Science and Engineering, The Georgia Institute of Technology, Atlanta, GA 30332, USAfResearch and Development, MedShape Solutions Inc., Atlanta, GA 30318, USA

a r t i c l e i n f o

Article history:Received 28 June 2011Received in revised form19 August 2011Accepted 21 August 2011Available online 26 August 2011

Keywords:Shape memory polymersActivationRecovery

* Corresponding author. Department of Mechanical EColoradoatDenver,Denver,CO80217,USA. Tel.:þ13035

E-mail address: chris.yakacki@ucdenver.edu (C.M.

0032-3861/$ e see front matter � 2011 Elsevier Ltd.doi:10.1016/j.polymer.2011.08.027

a b s t r a c t

Shape-memory polymers (SMPs) are a class of mechanically functional “smart”materials defined by theirability to change shape upon exposure to an environmental stimulus. The shape-memory effect hastraditionally been activated by thermal mechanisms via heating the polymer above a transitiontemperature to increase chain mobility and initiate shape recovery. This study proposes a uniqueapproach to mechanically drive recovery in SMP networks using external forces to facilitate shapechange in a material with stored strain. SMP networks were synthesized from tert-butyl acrylate andpoly(ethylene glycol) dimethacrylate in three network compositions. Networks were tailored to maintaina constant glass transition temperature (w52 �C) with increasing crosslinking density, shown by rubberymodulus values of 1.2, 3.1, and 8.2 MPa. Hollow SMP cylinders were axially elongated (programmed) tostored strain levels of approximately 25%. A second set of samples was machined to match the pro-grammed dimensions of the SMP sample set. Compression testing revealed that the compressivestrength and energy required for deformation for the programmed SMP samples were on average 62%and 52% of the as-machined samples’ values, respectively. The ratios between programmed and as-machined samples’ compressive properties were independent of both crosslinking density andtemperature up to the onset of glass transition. Lastly, an interference-fit test model was used todemonstrate that mechanically-driven SMPs could immediately create and maintain a stronger fixationforce compared to as-machined samples and thermally-driven SMP samples. This study introduces anapproach to drive shape change that mitigates the time-temperature dependence and discusses thepotential of this mechanism for biomedical devices.

� 2011 Elsevier Ltd. All rights reserved.

1. Introduction

Shape-memory polymer (SMP) research has emerged as amulti-disciplinary focal point combining aspects of mechanics, polymerstructure, chemistry, bio-functionality, device design, composites,and predictive modeling. SMPs are defined as a class of “smart”materials that are able to adapt and respond to a stimulus. Ingeneral, SMPs are capable of storing large-straindeformations for anindefinite amount of time. When exposed to the proper stimulus,SMPs will recover their stored strains and return to their original

ngineering, The University of568516; fax:þ13035566371.Yakacki).

All rights reserved.

shape. If constrained, SMPs will exert forces upon the constraintsand can even be used to perform work. SMPs can be thermallyactivated through a variety of mechanisms [1] including heat [2,3],lasers [4e6], electricity [7,8], infrared absorption [7], and alternatingmagnetic fields [9e13]. Simple solvents, such as water, can also beused to athermally activate shape recovery by lowering the glasstransition temperature (Tg) below ambient conditions [14,15].

SMPs can typically be categorized into four different classifica-tionsbasedontheirpolymerstructure (chemically-crosslinkedglassythermosets, chemically-crosslinked semi-crystalline rubbers, physi-cally crosslinked thermoplastics, physically-crosslinked block copol-ymers) [16]; however, the shape-memory programming cycle isgenerally the same for every class. The polymer is first heated abovea transition temperature (Ttrans), such as a glass ormelting transition,

C.M. Yakacki et al. / Polymer 52 (2011) 4947e49544948

and deformed to a temporary configuration. The polymer is thencooled below Ttrans to reduce chainmobility and essentially freeze thepolymer in its temporarily deformed state. The polymer will remainin this configuration until heatedback to the vicinityof Ttrans, inwhichthe material is driven back to its original configuration via entropyelasticity [17]. It should be noted that process describes the basicsteps in the shape-memory cycle and that the specific details of eachstep can influence shape recovery. For example, SMPs have showna temperature-memory effect, in which the activation temperatureand rate of recovery can be influenced by the temperature of defor-mation [18e20]. Heating and cooling rates between steps have alsobeen shown to influence rate of recovery and magnitude of recov-erable stress [21]. Furthermore, researchers have intricatelyexploitedprogramming and activationmechanisms to enablemulti-way shaperecoveries [19,22,23]. As a result, SMPs can now possess more thanone recovery pathway from the temporary shape to the permanentshape. For example, He et al. demonstrated 5 recovery pathwayscould be achieved by controlling the spatial arrangement of iron-oxide (Fe3O4) nanoparticles and carbon nanotubes (CNTs), as selec-tive radiofrequencies anddirectheating could independentlyactivatedifferent regions of the SMP sample [23]. More in-depth discussionon the mechanisms of shape storage and recovery can be seen inprevious reviews [20,24,25].

Several different approaches have been taken to investigate therecovery characteristics of SMPs. Yakacki et al. investigated theinfluence of Tg and crosslinking density under both free- and fixed-strain recovery conditions [26,27]. Other researchers have takena modeling approach to predict recovery. These models range fromthe use of simple Voigt-Kelvin elements for amorphous lactidecopolymers [28] to in-depth thermo-viscoelastic models to incor-porate time-dependent effects in acrylate networks [29,30].Although, previous studies have primarily focused on the charac-teristics of the thermal transition in relation to the activationtemperature when focusing on shape recovery behavior.

For biomedical applications, the activation temperature must bein the vicinity of body temperature and shape recovery is typicallyrequired as an acute event [31]. The shape-memory effect is a timeand temperature dependent process, with previous studiesshowing the time for complete free recovery being affected by atleast 2 orders of magnitude by small changes in Tg and polymerstructure [26,27]. One of the major challenges in practically acti-vating a SMP biomedical device is to accurately deploy in a shorttime frame. Maitland et al. proposed laser activation of SMP wiresand foams for clot removal and embolism treatment to minimizetime dependence and premature activation associated with bodytemperature deployment [5,6,32]; however, this method may notbe practical for larger, bulkier samples.

This study proposes a method of mechanically driving recoveryin the shape-memory cycle to allow for temperature independenceand near-instantaneous recovery. Shape recovery is typically drivenby entropy elasticity and can be expressed by:

DG ¼ �TDS

where ΔG is the change in Gibbs free energy, T is temperature, andΔS is the change in entropy. Shape recovery occurs as it is ther-modynamically favorable for the network configuration to return toits original (equilibrium) shape, which increases entropy andlowers energy; however, this can only occur in the presence ofa thermal transition, as the polymer chains are kineticallyforbidden from free large-scale motions below the transition.Rather than solely relying on temperature and thermal transitionsto increase chain mobility and allow entropy elasticity, the authorspropose that mechanical energy (i.e. force) can be added into thesystem to overcome the intermolecular interactions and drive

shape change. The authors hypothesize that programmed SMPsamples will exhibit deformation at lower forces compared toidentically shaped non-programmed SMP samples due to shape-memory programming and polymer structure. Shape-memoryprogramming induces chain orientation and excess free volumedue to straining and quenching the polymer to lock in a temporaryshape. This has been shown to induce anisotropy and lower bothstrength and modulus [33e36]. Furthermore, the crosslink densityof a SMP network has been shown to generate higher values ofrecovery stress and initiate shape recovery at lower temperatures[27]. The purpose of this study was to investigate the fundamentalthermomechanics of mechanically-driven shape recovery asa function of temperature and network crosslinking density forprogrammed and non-programmed SMP samples.

2. Materials and methods

2.1. Materials

Tert-butyl acrylate (tBA), poly(ethylene glycol)n dimethacrylate(PEGDMA) with an molecular weight of Mn ¼ 550 for the PEGportion, diethylene glycol dimethacrylate (DEGDMA), and 2,2-dimethoxy-2-phenylacetophenone (DMPA) were ordered fromAldrich and used in their as-received condition without furtherpurification. The purity of TBA was 98%, while the other materials’purities were greater than or equal to 99%.

2.2. Polymer synthesis

A crosslinker solution (XLS) was made from the two meth-acrylated PEG crosslinking monomers by mixing 30 wt% DEGDMAand 70 wt% PEGDMA. The XLS solution was chosen and tailored tomaintain a constant glass transition temperature (Tg) when poly-merized with tBA, regardless of the amount of XLS, and allow forindependent tailoring of rubbery modulus. Three tBA-co-XLSsolutions were then mixed with 10, 20, and 40 wt% of XLS andremainder tBA. Polymer networks were synthesized via free-radical photo-polymerization by adding 0.1 wt% of DMPA to thetBA-co-XLS solution and curing the solution in an ultravioletcrosslinking oven (UVP, CL-1000). The crosslinking oven was pre-heated for 10 min prior to polymerization.

DMA samples were made by polymerizing the tBA-co-XLSsolutions between two glass slides separated by 1 mm spacers for10 min. Samples for mechanical compression testing were manu-factured from SMP rod stock. SMP rod stock was made by poly-merizing the tBA-co-XLS solution in polyethylene tubing. Duringpolymerization, the tubing was placed in an ice-water bath, whilethe UV crosslinker boxwas pulsed at intervals of 30 s on and 30 s offfor 10 cycles. After UV synthesis, thematerials were heated above Tgin an oven at 60 �C for 30 min to help ensure complete polymeri-zation and to relieve residual stresses.

2.3. Dynamic mechanical analysis (DMA)

DMA samples were cut to dimensions of 1�5� 20mm3 and theedges of the samples were wet sanded with 600 grit sandpaper.DMA was performed in tensile loading and run on a TA Q800. Theends of the samples werewrapped in aluminum foil to prevent gripfailure due to thermal expansion during testing. Samples wereequilibrated at 0 �C before testing with a preload of 0.001 N.Samples were then heated to 125 �C at a rate of 3 �C/min witha cyclic strain of 0.2% and force track setting of 150%. After testing,Tg and rubbery modulus were defined at the peak of the tan deltacurve and lowest point of the rubbery modulus plateau, respec-tively. Testing was performed with n ¼ 2 per polymer chemistry.

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2.4. Mechanical compression

Hollow cylinders were machined from the SMP rod stock to11.3 mm in outer diameter, 1.5 mm in thickness, and 13.5 mm inlength using a CNC lathe. Samples were then heated above Tgbetween 60 and 80 �C and elongated using three-piece sabots andring clamps (see Fig. 1). The samples were first elongated usinga 9.5 mm inner diameter sabot and then cooled back to roomtemperature. The samples were further elongated using an 8.5 mminner diameter sabot. During this second elongation-processingstep, the sample was placed firmly in the 8.5 mm sabot, reheatedabove Tg, further compressed diametrically, and then cooled. Thefinal dimensions of the elongated samples measured approxi-mately 8.5 mm in outer diameter, 5.0 mm in inner diameter, and17 mm in length. It should be noted that a two-step compressiontechnique was needed when compressing the samples with thesabots to avoid crimping and pinching of the sample, whichoccurred when attempting to compress the 11.3 mm diametersample directly with an 8.5 mm diameter sabot. Finally, sampleswithout shape-memory programmingweremachined tomatch thefinal dimensions of the elongated samples and to act as controls (i.e.“as-machined” samples).

Mechanically-driven recovery was tested using an Instron 5567equipped with a thermal chamber. The thermal chamber wasequilibrated at 20, 27, 34, and 41 �C before samples were placed inthe chamber. The test platens were coated with a small amount ofglycerin to reduce friction between the samples and the platens.Samples were held isothermally for 5 min with a preload of 10 Nbefore testing. Samples were compressed 3.5 mm at 0.1 mm/s,which returned the programmed samples to their original length.Compressive strength was defined at the yield point at the firstmaximum stress peak. The energy required tomechanically deformthe programmed and as-machined samples 1.77mm (i.e. 50% of theprogrammed samples’ stored strain) was also recorded. This metricwas chosen to capture and compare the yield-softening behaviorexperienced during compression of the samples. None of the

Fig. 1. Illustration of the shape-memory programming process for hollow SMP cylinders aprogramming, the SMP samples were loaded into a sabot, heated above Tg, and compressedcreate a temporary, programmed shape. Programmed samples can follow recovery pathway (this study, non-programmed samples were machined to the exact geometry of the prograhighlights the potential for programmed SMP samples to recover using a combination of therrecovered mechanically and fully recovered thermally.

programmed or machined samples experienced failure duringcompression.

2.5. Interference fixation testing

The cylindrical 40 wt% XLS samples were also tested for inter-ference fixation, in which cylindrical samples were tested to securea rope in a tunnel. This model represents a practical application ofSMPs, as recovery takes place in a partially constrained environ-ment. In this test, SMP samples were machined to 21.5 mm inlength and 6.8 mm in diameter, and were programmed using thesame technique illustrated in Fig. 1 to dimensions of 27.7 mm inlength and 5.7 mm in diameter. Again, as-machined controlsamples were created to match the processed dimensions of theprogrammed samples. Samples were placed in a 20-pcf syntheticsawbone block (Pacific Research Laboratories, Vashon, WA)measuring 40 � 40 � 40 mm3 with an 8.5 mm diameter tunneldrilled along the vertical z-axis. A flat nylon rope was placed withinthe tunnel along with the samples (Fig. 2). Programmed and as-machined samples were then mechanically compressed 3 mm(Fig. 1 e Pathways (b) and (c)), released, and held at 37 �C between0 and 30 min before pulling at the nylon rope at a rate of 1 mm/sec.A cylindrical stop was placed at the bottom of the samples to holdthe samples in the middle of the sawbone tunnel duringcompression. Temperature control was maintained in a thermalchamber built around the mechanical testing frame. The forcerequired to pull the nylon rope from the sawbone tunnel wasrecorded as pullout force. Lastly, to compare mechanically-drivenrecovery against traditional thermally activated SMPs, a set ofprogrammed samples were placed in the sawbone tunnel andallowed to thermally activate for a period between 0 and 30 minbefore testing the pullout strength of the rope (Fig. 1 e Pathway(a)). It should be noted that the sawbone blocks and nylon ropewere equilibrated at 37 �C before samples were placed in thesawbone assembly.

nd potential recovery routes using thermally- and mechanically-driven recovery. Fordiametrically; next, the samples were cooled below Tg and released from the sabot toa) or (b) by using thermal or mechanical energy to drive shape recovery, respectively. Inmmed samples and compressed following pathway (c) for comparison. Pathway (d)mal and mechanical mechanisms. For example, programmed samples could be partially

Fig. 2. Illustration of interference fixation test setup showing cross-sections of pro-grammed SMP sample being a). inserted and b). recovered into a sawbone block tosecure a nylon rope. Fixation is created as the sample radially pushes on the rope.

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3. Results

The thermomechanical behavior of the polymer networks wasfirst characterized to ensure shape-memory programming andtesting were performed at appropriate temperatures relative to theglass transition. Fig. 3 shows the storage modulus of the threenetworks through the glass transition. The networks were tailoredto maintain a constant Tg, while exhibiting increasing levels ofcrosslinking. The molecular weight between crosslinks (Mc), can becalculated by the equation:

Mc ¼ 3rRTEr

where r is the polymer density, R is the gas constant, T is theabsolute temperature, and Er is the rubbery modulus. The crosslinkdensity increases asMc decreases and thereby is proportional to therubbery modulus of the polymer network. It should be noted thatthis equation is only valid when a crosslinked polymer is in itsrubbery regime. The rubbery modulus values of the networks weremeasured as 1.2, 3.1, and 8.2 MPa, which corresponded to 8490,2748, and 1079 g/mol between crosslinks, respectively. Thenetworks showed a slight variation in Tg, which was measured at52, 53, and 49 �C for the 10, 20, and 40 wt% XLS networks,respectively; however, this variation was deemed acceptable fortesting. The onset of the glass transition (Tonset) is typically associ-ated with increased chain mobility and the initiation of shaperecovery, and was measured at 43, 42, and 39 �C for networks withincreasing crosslinker concentration.

Fig. 3. Storage modulus of the three SMP networks through the glass transition. Theglass transition temperature is marked by the peak in tan delta.

The shape-memory programming of the SMP cylinders resultedin an average axial elongation and strain of 3.6 mm and 26.6%,respectively. To measure the effect of mechanically-drivenrecovery, a second set of samples were machined to match thedimensions of the programmed SMP samples. These samples weremade from the samematerials as the programmed samples but willbe designated ‘as-machined’ samples to denote they did notundergo shape-memory programming. Fig. 4 is an example of thecompressive behavior of the 20 wt% XLS programmed and as-machined samples tested at 27 �C. The programmed sampleshave a compressive yield of 25.2 MPa, which is 60% of thecompressive yield of the as-machined samples measuring42.2 MPa. The force required maintaining flow past yield remainslower for the programmed sample throughout its mechanicalrecovery. The energy to mechanically recover 50% of the pro-grammed sample’s stored strain was measured at 2.35 J/mm3,while the energy to deform the as-machined sample the sameamount of distance measured 4.33 J/mm3. Additionally, these datashow a good degree of repeatability with the test method.

Representative curves of programmed and as-machinedsamples tested at four different temperatures for the threedifferent polymer networks can be seen in Fig. 5. The as-machinedsamples for the three networks tested showa transition fromvisco-plasticity to elastomeric behavior with increasing temperature. Thisis indicated by a decrease in modulus and compressive yield. InFig. 5c, the Tg of the 40wt% XLS networkwas slightly lower than theremaining networks; therefore, the highest testing temperature of41 �C was above the Tonset of the network by 2 �C. This resulted ina full linear elastomeric response, which is not seen by the othernetworks as the highest testing temperature was just below theirTonset. The behavior of the programmed samples showed a similartrend; however, exhibited their behavior at markedly lower stresslevels. The programmed samples showed a noticeable amount offree-strain recovery at the highest test temperature, as the samplesrecovered 2.6, 1.0, and 16.5% for the networks with increasingcrosslinking density, respectively. The increased amount of free-strain recovery of the 40 wt% XLS network could also be attrib-uted to its Tg being slightly lower compared to the other networks.

The changes in compressive mechanical properties wereanalyzed for relationships between crosslinking density andtemperature. A summary of the decreased compressive strengthand strain energy requirements for the mechanically-recovered

Fig. 4. Example of compressive stress-strain behavior and decreased-force benefits ofmechanically-driven recovery of programmed samples versus compression of as-machined samples.

Fig. 5. Representative compressive stress-strain behavior for SMP and as-machined samples made from a). 10 wt% XLS, b). 20 wt% XLS, and c). 40 wt% XLS.

C.M. Yakacki et al. / Polymer 52 (2011) 4947e4954 4951

programmed samples can be seen in Table 1. The ratio ofcompressive strength in programmed to as-machined samplesranged from 51.9 to 73.5%, while the ratio of compressive energyranged from 35.6 to 62.5% for all samples except for 40 wt% XLSsamples tested at 41 �C. The 40 wt% XLS sample could not berepeatedly measured at this elevated temperature due to excessivefree-strain recovery. Fig. 6 compares the compressive properties ofprogrammed and as-machined samples with temperature relativeto Tonset of each network. Both compressive strength and energydecreased linearly with increasing temperature, while there did notappear to be any trends in the ratio of properties between

Table 1Summary of average compressive strength and energy for programmed and as-machined1.77 mm, which corresponds to 50% axial-strain recovery of the programmed samples.

XLS Temperature (�C) Programmedcompressivestrength (N)

As-machinedcompressivestrength (N)

10 wt% 20 32.1 � 3.6 49.5 � 0.927 26.4 � 4.3 42.7 � 1.434 20.5 � 2.6 31.9 � 0.941 8.7 � 1.4 15.8 � 1.6

20 wt% 20 27.5 � 0.6 52.5 � 1.227 25.2 � 0.4 42.2 � 2.134 17.7 � 1.0 34.2 � 4.641 8.1 � 1.1 10.6 � 1.2

40 wt% 20 23.9 � 1.5 35.9 � 7.627 17.5 � 0.8 31.2 � 2.934 10.6 � 3.0 14.4 � 1.641 n/a n/a

programmed and as-machined samples as a function of crosslinkerconcentration. Graphical summaries of these data can be seen inthe online supplement (Fig. 1S).

The practical utility of mechanically-driven recovery was testedin an interference fixation model (Fig. 7). As-machined sampleswith no programmed strain were unable to fixate the nylon ropewith pullout strength greater than 50 N when compressed 3 mmin the interference model. Programmed samples that werecompressed (i.e. mechanically-recovered) experienced initialpullout strength around 200 N, which was maintained over 30 minat 37 �C. Thermally activated programmed samples initially had

samples. Compressive energy represents the energy required to deform the samples

Ratio (%) Programmedcompressiveenergy (J/mm3)

As-machinedcompressiveenergy (J/mm3)

Ratio (%)

64.8 2.91 � 0.4 4.66 � 0.02 62.561.7 2.27 � 0.6 4.08 � 0.02 55.764.4 1.58 � 0.6 3.02 � 0.03 52.455.2 0.62 � 0.4 1.59 � 0.08 39.352.4 2.48 � 0.3 5.03 � 0.06 49.359.8 2.53 � 0.1 4.33 � 0.29 58.451.9 1.74 � 0.1 3.03 � 0.19 57.676.3 0.43 � 0.0 1.20 � 0.12 35.666.7 2.24 � 0.3 3.58 � 0.66 62.556.1 1.57 � 0.1 3.06 � 0.22 51.273.5 0.81 � 0.5 1.57 � 0.17 51.6n/a 0.002 � 0.001 0.23 � 0.17 0.7

Fig. 6. a). Compressive strength and b). compressive energy as a function of temperature relative to Tonset of all the samples tested. Data points represent mean values.

C.M. Yakacki et al. / Polymer 52 (2011) 4947e49544952

pullout strength of 100 N, which steadily increased to 158 N after30 min. On average, mechanically-recovered programmed samplesincreased in diameter 0.40 mm over the course of the experiment,while mechanically-compressed as-machined samples andthermally-activated programmed samples only increased 0.12 and0.22 mm, respectively. In terms of axial-strain recovery, pro-grammed samples were mechanically compressed 3 mm torecover 48% of their programmed strain, however, on averagemaintained 38% recovery when the load was removed due toelastic spring-back. In comparison, thermally activated samplesexhibited only 27% recovery by 30 min. As-machined sampleswere compressed 3 mm and experienced 2.44 mm of spring-back.An additional increase of 0.26 mm of viscoelastic spring-back wasexperienced over the course of the test, which explains why theas-machined samples demonstrated a slight decrease in pulloutforce over time. The data demonstrates that mechanical recoveryof stored strain has advantage over both thermal activation ofstored strain and simple deformation of a sample with no storedstrain.

Fig. 7. Fixation force as a function of time and recovery mechanism (thermal versusmechanical) for an interference model. Programmed samples were mechanicallyrecovered by being compressed 3 mm longitudinally. As-machined samples were alsocompressed 3 mm longitudinally, but they contained no stored strain. A second set ofprogrammed SMP samples were allowed to thermally activate at 37 �C.

4. Discussion

The purpose of this study was to investigate a new method ofshape recovery in polymers that utilizes mechanical energy to helpovercome the intermolecular frictional forces that typically restraina SMP from recovering to its original shape at temperatures belowTtrans. In amorphous SMP networks, these molecular frictionalforces are alleviated at the onset of the glass transition, in whichthere is an increase in free volume and micro-Brownian motion;however, relying on thermal activations of the shape-memoryeffect at temperatures near Tg is a highly time and temperaturedependent process. The proposed method to mechanically driverecovery in SMPs is a novel approach to rapidly induce shapechange in SMPs by eliminating the time-temperature dependenceof the shape-memory effect.

Compression testing showed that hollow cylindrical samplesprogrammed for shape-memory demonstrated a substantialdifference in compressive behavior compared to identically shapedsamples that were not programmed for shape memory. As a result,shape-memory programmed samples exhibited a drop incompressive strength and energy required for deformation onaverage by 62 and 52%, respectively. One explanation for thisdifference is that the programmed SMP samples were not in theirequilibrium configurations, as shape-memory programminginduces stored strains in the material. The SMP materials wouldtherefore be at a higher energy (lower entropy) state, in whichentropy elasticity would aid the external mechanical forces inreturning the programmed SMP samples to their original shape andresult in lower force requirements for deformation compared tosamples not programmed for shape memory. Unfortunately,entropy elasticity cannot account for the reduction in compressiveproperties because the networks are kinetically forbidden to returnto an equilibrium state below Tg. Fig. 6 showed there was a consis-tent difference between SMP and as-machined samples indepen-dent of crosslinking density. Previous studies have shown that theamount of recoverable force of SMPs under constraint (i.e. fixed-strain recovery) is proportional to the crosslink density of thenetworks, which is indicated by the rubbery modulus [27]. In thisstudy, rubber modulus was varied between w1e8 MPa. Therefore,the contribution of entropy elasticity for the highest crosslinkednetwork should be approximately 8 times greater than the networkwith the lowest crosslink density; however, the results suggest thatthe crosslink density of the networks does not affect differences incompressive behavior below Tg. This is also supported by the factthat all three networks have equivalent glassy modulus values.

C.M. Yakacki et al. / Polymer 52 (2011) 4947e4954 4953

The reduction in compressive properties may better beexplained by examining the consequences of shape-memoryprogramming with respect to non-equilibrium volume and chainorientation. G’Sell and McKenna previously studied the effects ofaging on quenched polymer networks [33]. Amorphous networksquenched from the rubbery state are not in thermodynamic equi-librium and have excess volume, as the network structure does nothave sufficient time to rearrange when cooling through the glasstransition and experiencing a collapse in free volume. This processresults in the network experiencing a reduction in both modulusand strength. These properties predictably increase as the polymerstructure slowly relaxes and densifies over time through “physicalaging.” The programming of the shape-memory effect involvesa thermal cycle of heating and quenching analogous to the physicalaging process. Therefore, the thermal cycle in shape-memoryprogramming is likely to induce excess volume after cooling tothe glassy state after the SMP has been deformed to its temporarystate, which would contribute to the reduction in properties duringmechanically-driven recovery. In ongoing studies by the authors,physical aging of programmed SMP samples has been shown toinfluence the free-strain recovery response of the networks duringheating, by delaying the onset of shape recovery and increasing therate of recovery, indicating that the amorphous SMP networksundergo structural reconfiguration/relaxation over time; however,to the best of the authors knowledge, no studies to date havequantified the relative amount of non-equilibrium relaxationcaused by quenching alone compared to deformation andquenching, as experienced in shape-memory programming.

The reduction in compressive properties may also be contrib-uted to anisotropy caused by chain orientation during shape-memory programming, in which the polymer chains are alignedin the direction of deformation. As a consequence, tensile proper-ties are known to increase in the direction of deformation and chainalignment. Lee et al. investigated the influence of cold drawing andcold extrusion on polymeric rods with respect to both tensile andcompressive properties [35]. Compressive modulus and strengthwere shown to decrease, while tensile properties increased, alongthe axis of deformation. Arruda et al. demonstrated reducedcompressive properties in polymethylmethacrylate when a samplewas recompressed perpendicular to the direction of originalcompression; this was attributed to the oriented chain configura-tion offering less resistance to deformation as they returned toa recoiled state [34]. Their study also showed that strain hardeningwas delayed until the polymer chains recoiled to their originalconfiguration, and it is important to note that strain hardening wasnot observed during flow for mechanically-recovered SMP samplesin the present study. The absence of strain hardening in thedeformation of programmed samples helps contribute to loweroverall energy required for deformation.

Though polymers typically exhibit anisotropy with deformationprocessing techniques, such drawing or extrusion, these techniquesare limited to the amount of strain applied to the material. Highlyamorphous thermoplastics show limited anisotropy due to defor-mation processing, as it is difficulty to obtain high chain alignmentcompared to semi-crystalline polymers [37]. Shape-memoryprogramming is designed to optimize the strain capabilities ofa material and achieve large strains, often over 100% [38,39].Though the process of mechanically-driven recovery is similar tothe results of deformation processing techniques, as shape-memory programming can be considered a deformation process,this study is the first to propose taking advantage mechanicalanisotropy of deformed polymers coupled with the shape-memoryeffect to achieve shape recovery with minimal time-temperaturedependence. An interference-fit test model was used to demon-strate the utility of this approach (Fig. 7). A thermally-activated

SMP plug showed a gradual increase in fixation force over time,while the mechanically-recovered SMP plugs demonstratedincreased initial fixation force that was maintained over the courseof the experiment. It is interesting to note that the thermally acti-vated sample initially exhibited a response closest to a mechan-ically compressed, as-machined material at the beginning of thetest. The mechanical performance of thermally and mechanically-recovered samples then began to converge as time continued, asboth materials should eventually reach the same steady-statepullout-force value.

Many researchers have proposed exploiting the shape-memoryeffect in polymers for biomedical applications [31,40e42]. SMPsnaturally lend themselves to minimally invasive surgery, as bulkydevices can be compacted in size, inserted through a small incision,and recover to a larger structure in vivo. Examples of SMP-inspireddevices include stents for sclerotic arteries, interference devices fortendon and ligament fixation, coils for aneurysm and embolismtreatment, gastro-intestinal fillers for weight management, andeven trans-cervical devices for birth control. One of the reasonsSMPs have yet to significantly break through the medical devicemarketplace is the inherent difficulty in transferring SMPs from thelaboratory to the operating room (OR). Conditions that seemcommonplace in the research environment can become enormouspractical challenges for the clinician. For example, controlling thetemperature of a SMP sample and monitoring strain recoveryhardly seem like a technical feat in the laboratory; however, doingso in an arthroscopic procedure where OR-temperature saline isbeing pumped through a joint while under significant timerestraints poses a much more difficult problem. As a benchmark,the average time it takes for an orthopedic surgeon to implant andtie a suture anchor in a rotator cuff repair is between 225 and 380 s[43]. SMP must be implanted, activated, and functioning withinstandard surgical time limits in order to gain clinical acceptance.

Mechanically-driven recovery of programmed strains mayprovide a means to enable more practical application of SMPs intothe medical device marketplace. Medical procedures, such asangioplasty-balloon expanded stents, currently utilize mechanicaldeformation of materials in vivo [44]. In these procedures, high-pressure inflatable balloons are used to plastically deform andexpand a compacted stent to support an arterial wall and preventrestenosis [45,46]. Elastic recoil and localized areas of high plasticstrains are two issues arising from this approach [44].Mechanically-driven recovery of SMPs may provide a novel meansof maintaining the current approach while eliminating thesematerial problems. The process of mechanically deforming a pro-grammed SMP back to its original shape returns the polymerstructure to a recoiled and lower energy configuration, thus shouldlimit elastic recoil and serve to lower strain energy. One of the keypoints of mechanically-driven recovery in potential biomedicalSMP devices is that the material will be relieved of internal strain,rather than inducing strain, with deformation in vivo. This couldlead to better long-term performance by reducing stress-inducedenvironmental interactions and increase fatigue resistance.Furthermore, a SMP device could offer the benefit of initialmechanical recovery followed by long-term thermal activation toaccount for growth or remodeling (Fig. 1 e Pathway (d)).Mechanically-driven recovery also would allow for near-instantaneous shape recovery of SMPs with Tg’s that are typicallytoo high for thermal activation at body temperature, whichconversely would help the SMP maintain a higher modulus in vivo.All of these factors warrant further investigation for use inbiomedical devices.

This study focused on the compressive properties of pro-grammed SMP cylinders for a tBA-co-PEGDMA polymer system.Though a change in polymer structure (i.e. crosslinking density)

C.M. Yakacki et al. / Polymer 52 (2011) 4947e49544954

was not shown to influence this affect, the influence of polymerchemistry was not investigated. In a recent study systematicallycomparing (meth)acrylate polymers demonstrated that polymerside group chemistry influenced the thermomechanical propertiesof the networks [47]. It is likely that polymer chemistry wouldinfluence the characteristics of mechanically-driven recovery suchas the magnitude of reduction in compressive properties betweenprogrammed and machined samples; however, the general processshould be exhibited by all SMPs, as the mechanisms of thisphenomenon have been attributed to the effects of non-equilibrium and deformation processing. Mechanically drivingthe shape-memory effect may also find potential applicationsoutside of biomedical devices. High-strength polymers, such aspoly aryls, are defined by their resistance to deformation; however,mechanical recoverymay offer a newapproach tomanipulate high-strength materials with low-force requirements.

5. Conclusions

Shape-memory polymers (SMPs) can bemechanically recoveredby using external forces to return the sample from its programmedshape to its original shape. The compressive strength and energyrequired to recover 50% of the stored strain for amorphous SMPnetworks were on average 62 and 52% lower when compared to as-machined samples, respectively. The crosslink density of thenetworks did not influence the forces required to mechanicalrecover the programmed SMP samples. The reduction in mechan-ical force requirements for the programmed SMP samples wereattributed to excess volume due to quenching and chain orientationdue to deformation, both of which cause non-equilibrium condi-tions and are artifacts of programming the shape-memory effect.Mechanically-driven recovery can be used to quickly recover storedstrains of SMPs, while also being able tomaintain applied force overtime under partially constrained conditions.

Acknowledgments

The authors would like to thank Kurt Jacobus, PhD, Jack C. Griffis,and Carl Frick, PhD for their contributions to this study.

Appendix. Supplementary material

Supplementary data associated with this article can be found, inthe online version, at doi:10.1016/j.polymer.2011.08.027.

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