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Polymer 54 (2013) 2808e2820

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Polymer

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Evolution of microstructure during shape memory cycling of a main-chain liquidcrystalline elastomer

Kelly A. Burke a,b, Patrick T. Mather b,c,*aDepartment of Macromolecular Science and Engineering, Case Western Reserve University, 2100 Adelbert Road, Cleveland, OH 44106, USAb Syracuse Biomaterials Institute, Syracuse University, 121 Link Hall, Syracuse, NY 13244, USAcDepartment of Biomedical and Chemical Engineering, Syracuse University, 121 Link Hall, Syracuse, NY 13244, USA

a r t i c l e i n f o

Article history:Received 7 January 2013Received in revised form11 March 2013Accepted 22 March 2013Available online 28 March 2013

Keywords:Liquid crystalline elastomerShape memory polymerPolymer microstructure

* Corresponding author. Syracuse Biomaterials InstitUniversity, Syracuse, NY 13244-1200, USA. Tel.:þ1 3154

E-mail address: ptmather@syr.edu (P.T. Mather).

0032-3861/$ e see front matter � 2013 Elsevier Ltd.http://dx.doi.org/10.1016/j.polymer.2013.03.049

a b s t r a c t

The field of shape memory polymers (SMPs) has been dominated by polymeric systems whose fixingmechanism is based on crystallization or vitrification of the constituent chains, rendering such systemsstiff in comparison to elastomers, gels, and living tissues. Previously, we reported the synthesis andcharacterization of main-chain, segmented liquid crystalline elastomers (LCEs) that exhibit both bulk andsurface shape memory effects. These LCEs have excellent shape fixing and recovery characteristics withcompositionally-dependent transition temperatures that determine the fixing and recovery criticaltemperatures. Synthesis of the soft shape memory LCEs proceeded by hydrosilylation-linking of poly(dimethylsiloxane) oligomers with mesogenic dienes of two compositions and a tetravinyl crosslinker.The present report describes microstructural changes during ex situ shape memory deformation andrecovery of one such LCE. Wide-angle x-ray scattering showed that, once a critical deformation stresswas reached, the microstructure of the fixed, oriented LCEs was independent of the stress applied abovethe clearing (isotropization) transition. Stepped recovery of the fixed, oriented LCE showed additionalintermediate microstructures, however. Recovery was shown to proceed through changes in both thesmectic layer thickness and chevron architecture, while mesogen tilt angle remained unchanged. Themechanical and microstructural studies described herein give deeper insight to shape memory fixing andrecovery mechanisms of these unique materials, which offer potential for exploitation in areas such ascell and tissue culture, microcontact printing, and microfluidics.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Shape memory polymers [1e3] are a class of programmablematerials that may be fixed into a temporary shape and later trig-gered to recover a permanent, equilibrium shape established by acrosslinked configuration. Fixing of the temporary shape involvesdeforming the polymer above a critical temperature, Tcrit, and thenreducing temperature below Tcrit while under stress, kineticallytrapping the stressed state through network chain immobilization.The stress is then unloaded below Tcrit to yield the temporary shape.A polymer fixed in a temporary shape may recover its permanentshape by heating back through Tcrit under little or no applied stressand under the action of rubber elasticity.

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Shape memory has been reported for many different polymers,both thermoplastics and thermosets, but the materials may unifiedby categorizing them into one of four different classes that relateTcrit to the material’s transition temperatures [2]. Most shapememory materials fix a temporary shape by cooling through eitherthe glass transition (Tg) or crystallization (Tc) temperature of thepolymer. Because deformation into the temporary shape causesdeviation of the polymer chains away from their preferredconformation, material transitions like Tg and Tc are usuallyrequired to immobilize the polymer chains in the temporary shape.Without these transitions, the chains would recoil, a process drivenby a gain in conformational entropy. The need for a material tran-sition to immobilize the polymer chains in the temporary shapemeans that most shape memory polymers are either glassy orsemicrystalline. These materials are therefore quite stiff and are notsuited for some fields, such as microfluidics or soft lithography.Previously we have reported on the synthesis and characterization[4,5] of soft shape memory materials. Among these materialsare main-chain liquid crystalline elastomers (LCEs) in which the

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e2820 2809

mesogens are separated by oligomeric poly(dimethylsiloxane)(PDMS) spacers. These spacers give the crosslinked network a lowmodulus (typically less than 200 MPa for T as low as �70 �C), whilethe mesogens impart shape memory behavior to the network atsuperambient temperatures.

The existence of shape memory behavior at easily accessibletemperatures while maintaining a low modulus is unusual and hasbeen shown to permit localization of the shape memory cycle onthe surface of a LCE film [5]. Other examples of bulk shape memorynetworks with low modulus include ethylene propylene dienemonomer (EPDM) ionomers incorporating crystallizable fatty acidsalts [6], hydrogels with crystallizable alkyl side chains [7], andblends [8] and composites [9] prepared from semicrystallinepolymers and elastomers. While these materials display a modestlylow modulus and shape memory behavior, they rely on phaseseparation of the crystallizable component, which is typically onthe microscale at minimum, with the exception of nano-phaseseparation reported in Ref. (9). Phase separation is also hypothe-sized to be necessary for the shape memory behavior of the LCE,with the mesogen serving to fix the network chains and strainwhile the siloxane maintains network fluidity. An advantage to theLCE is that the phase separation occurs on the nanoscale and this, inprinciple, permits fixing of finer topographies on the surface of afilm or coating.

Previous studies have associated the Tcrit of the shape memoryLCEs to themesogen, with fixing and recovery occurring around thesuperimposed “mesogen glass transition” and isotropization tran-sition [5]. In these materials, mesogen glass transition refers to asuperambient transition that is not present in crosslinked PDMSand was detectable by a peak in the tan delta signal using dynamicmechanical analysis and a stepwise change in heat flow using dif-ferential scanning calorimetry. Mesogens with bulky pendantgroups are thought to be able to vitrify independently from thefluidlike siloxane chains in these LCEs, giving a nanoscale (smecticlayer) “glassy” phase within an elastomeric network. Whilemesogens form the minor phase in each LCE, they clearly have animportant role in engendering shape memory properties, ascrosslinked PDMS shows no shape memory behavior over thetemperature range tested. To maintain softness (as these LCEs do),the vitrifying mesogenic phase must spatially percolate and this isapparently possible in smectic elastomers. Much remains unknownconcerning the microstructure-shape memory relationships forthese materials, however. Thus, in this paper, we reveal the long-range microstructural transformations associated with strainfixing and recovery and relate them directly to macroscopic mea-surements, shedding light on the nanoscale mechanisms involved.

2. Experimental section

Main-chain liquid crystalline elastomers were prepared byreacting divinyl mesogens and a tetravinyl crosslinker with hy-dride end-capped poly(dimethylsiloxane) oligomers. The dienemesogens 5H and 5tB were synthesized according to methodsdescribed in literature [10]. The synthesis of E-5H805tB20 is givenbelow as an example of the main-chain LCE synthesis, noting thatother compositions have been prepared and characterized [5]. Thenomenclature used throughout this paper is as follows: E- denotesthat the material is an elastomer, 5H80 means that the LCE wassynthesized with a mesogen molar feed ratio of 80% 5H, and 5tB20

means that the LCE was synthesized with a mesogen molar feedratio of 20% 5tB. As described later, this particular LCE composi-tion, E-5H805tB20, was selected because both its mesogen glasstransition and isotropization temperatures are superambient,which allowed room temperature x-ray studies of the elastomer

to be conducted at a temperature well below these materialtransitions.

2.1. Synthesis of E-5H805tB20

Into a 25 mL flame dried and inert gas cooled Airfree� reactiontube containing a magnetic stir bar, 0.205 mmol (99.8 mg) 5H and0.0514 mmol (27.8 mg) 5tB were weighed. The solids were purgedwith nitrogen gas for 30 min, after which 0.41 mL of dichloro-methane (purchased from Fisher Scientific and distilled over CaH2)was added to dissolve the solids at room temperature. The reactiontube was cooled to 0 �C before adding 12.5 ml (0.0257 mmol) of thecrosslinker, tetrakis(vinyldimethylsiloxy)silane (Gelest, Morrisville,PA), and 6 ml of Pt(0) catalyst (platinum(0)-1,3-divinyl-1,1,3,3-tetramethyldisiloxane complex in xylenes (Aldrich)). Next, 216 ml(0.308 mmol) of dihydride-terminated poly(dimethyl siloxane) (Mn652 g/mol as determined by 1H NMR, Aldrich) was added to thereaction, which was stirred to a homogeneous state. Expedienttransfer to a sealed beaker coated with a hexamethyldisilazanehydrophobic release layer for polymerization and crosslinking inthe desired film shape was required. For this, a 5 mL syringe andneedle chilled over dry ice were used to transfer the reaction to asealed 10 mL glass beaker that was equilibrated with dichloro-methane atmosphere. The use of a chilled syringe and needleprevented gelation prior to transfer. The reaction was then allowedto proceed for 24 h at room temperature. After 24 h, the beaker wasopened to reveal aMC-LCE film, whichwas easily removed from theglass floor of the beaker using tweezers. The elastomer was driedovernight in a fume hood before being dried further under vacuumat room temperature for 24 h. Once dried, the film was extractedfour times using amixed solvent system consisting of equal parts byvolume of dichloromethane and ethanol, which caused the elas-tomer to swell to approximately double its initial size, after whichthe film was once again dried at room temperature under vacuum.

2.2. Characterization

All characterization experiments were performed on the LCEafter extraction and relaxation, whichwas accomplished by heatingto 100 �C, a temperature well in excess of the isotropization tem-perature of E-5H805tB20, in a stress-free environment. Thermo-mechanical characterization was achieved using differentialscanning calorimetry (DSC) and dynamic mechanical analysis(DMA), and microstructure studies were conducted using wide-angle x-ray scattering. The onset of mass loss (decomposition)determined by thermogravimetric analysis was found to exceed themaximum characterization temperature by at least 250 �C. For DSC,samples with mass of 2e4 mg were encapsulated in a TA Tzerostandard aluminum pan and loaded into a TA Instruments Q200Differential Scanning Calorimeter. The thermal program was asfollows: cool rapidly to T ¼ �90 �C, hold isothermally for 1 min,heat at 10 �C/min to 100 �C (“first heat”), hold isothermally for1 min, cool to T ¼ �90 �C at 10 �C/min, hold isothermally for 1 min,heat at 10 �C/min to 100 �C (“second heat”), hold isothermally for1 min, and finally cool to �90 �C at 10 �C/min (“second cool”). Thesecond heating and second cooling traces of E-5H805tB20 are pre-sented, with heat flows normalized by sample mass. The broadendothermic transition observed during the DSC second heatingrun was resolved by running an annealing experiment. For thisexperiment, the LCE was heated rapidly to 100 �C and heldisothermally for 10 min before cooling at 10 �C/min to a tempera-ture, 50 �C (marked by an “x” on the second cooling trace of theunannealed sample), where the sample was held for 1 h. Afterthe isothermal annealing step, samples were cooled to �20 �C at10 �C/min, held for 10 min, and finally heated at 10 �C/min

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e28202810

to 100 �C. The heating trace following the cooling/annealing/cool-ing sequence is reported, with heat flow normalized by samplemass.

Temperature-dependent linear viscoelastic properties ofE-5H805tB20 were studied using a strain-controlled temperatureramp program in a TA Instruments Q800 Dynamic MechanicalAnalyzer (DMA). Prior to testing, a film of E-5H805tB20 was relaxedin a stress-free state by placing in 100 �C water. The sample wasloaded initially under slight tensile load (5.35 kPa) at room tem-perature to facilitate measurement of the sample length. Thisapplied static load was controlled at 108% of dynamic load tomaintain a state of tension in the sample during tensile strainoscillation. The temperature dependence of linear viscoelasticproperties were measured using a fixed oscillation frequency of1 Hz and strain amplitude <0.2% while heating at 3 �C/minfrom �70 �C to 150 �C. The tensile storage modulus, E’, and losstangent, tan(d), values are reported as functions of temperature.

The Q800 DMA was also used to characterize one-way shapememory behavior under a controlled force actuator mode, as pre-viously reported [5]. In these experiments, a film of LCE was loadedin tension above the combined mesogen Tg and isotropizationtransitions to a stress of 100 kPa, cooled under load through thecombined transition, unloaded to a slight tensile stress to reveal thetemporary shape, and finally reheated under slight tensile stress(5e6 kPa) through the combined transition to recover the LCE’spermanent shape. E-5H805tB20 was found to fix 99.7% of the tem-porary shape in both the first and second shapememory cycles, andthe LCE was found to recover 95.0% and 98.1% of the permanentshape in the first and second cycles, respectively.

One-way shape memory cycles were also used to prepareE-5H805tB20 for microstructural studies. To study the effect ofdeformation stress, controlled stress shape memory cycles wereused to prepare LCEs fixed at different amounts of strain. To preparethe deformation study samples, a film of E-5H805tB20 was loaded inthe DMA under a slight tensile load of 5.35 kPa. The sample washeated to 74 �C, which is 15 �C above the maximum of the clearingpoint peak observed in the tan(d) versus temperature trace ob-tained by thermomechanical characterization. The samplewas heldisothermally for 5 min before the sample length was measured foruse as the reference length for strain calculations. This preloaddeformed the film slightly (strain increased by 2.8%, as calculatedusing the modulus of the LCE at this temperature), but this strain issmall compared to the shape change from the deformation stepthat followed. The sample was deformed by increasing the stress ata rate of 10.7 kPa/min to a final stress that varied from 5.35 kPa to140 kPa. After a 5 min isothermal hold, the LCE was cooled at a rateof 2 �C/min to �40 �C, held isothermally for 5 min, unloaded to astress of 5.35 kPa at a rate of 10.7 kPa/min, and again heldisothermally for 5 min. The sample was then heated at 2 �C/min to25 �C, where it was held for 45 min under the preload stress. Thefinal isothermal hold allowed the sample to reach its “equilibrium”

strain at 25 �C, enabling room temperature x-ray scattering ana-lyses to be conducted without sample relaxation during the expo-sure. The details of the x-ray scattering measurements aredescribed in the paragraphs that follow. After such x-ray scatteringanalyses were completed, the LCE was relaxed in 100 �C water, andthis process was repeated using a different stress in the one-wayshape memory cycle. Previous calorimetry, thermomechanical,and x-ray scattering experiments have shown that a LCE may befully relaxed of orientation by heating to the isotropic state, whichdemonstrates that it was acceptable to reuse samples.

A LCE recovering from a fixed strained (temporary) state wasalso studied using x-ray scattering and stepped isothermal holds tounderstand how microstructure changes during shape memoryrecovery. In the recovery study, E-5H805tB20 was deformed using

the same one way shape memory protocol described in the aboveparagraph using a stress of 140 kPa. However, after the LCE wasanalyzed with room temperature x-ray scattering, it was notrelaxed in 100 �C water. Instead of heating through the clearingpoint, the filmwas placed back in the DMAwith only one end of thesample fixed to the instrument. The furnace was then closed, andthe sample was heated to a temperature Trec, at a rate of 1 �C/min,and held isothermally for 30 min. This allowed the sample to relaxwithout applied load at the given temperature for 30min. After thisisothermal hold, the samplewas removed from the DMA, measuredto enable strain calculation (described below) and analyzed withWAXS. This process was repeated for relaxation temperaturesranging from 27 �C to 100 �C, allowing the microstructure ofE-5H805tB20 to be studied ex situ during recovery using steppedisothermal experiments and room temperature WAXS.

It should be noted that, in the recovery study, the strain of theLCE at the end of the one-way shape memory cycle was 204% usinga reference length measured in the DMA at 74 �C and 5.35 kPastress. Once the sample was unloaded, it was marked with a per-manent marker where the tension clamps gripped the sample. Thedistance between these lines was measured after each recoverytemperature, and a strain was calculated using the distance be-tween the two lines after recovery at 100 �C as the reference length.This resulted in the initial strain of the experiment being reducedfrom 204% (as obtained from the DMA at 25 �C and 5.35 kPa stress)to 169% (as calculated from the markings on the sample). Whenstrain is referenced for the recovery study, it will be the straincalculated from the markings on the sample. However, when thedata for this sample is shown in the deformation study, the strainwill be reported at 204% to remain consistent between calculationmethods.

The recovery study was designed under the assumption that,after a sample is cooled back to room temperature after recovery atTrec, it will follow the same strain versus temperature path when itis heated to the next recovery temperature and that no additionalrecovery occurs in the sample until the temperature exceeds theprevious Trec. This assumption was tested by conducting a “ratchet”shape memory cycle where E-5H805tB20 was loaded in the DMAunder a slight tensile load of 5.62 kPa. The sample was heated to74 �C and held isothermally for 5 min before the stress wasincreased to 100 kPa at a rate of 11.2 kPa/min. The sample was thenheld isothermally for 5 min before it was cooled at a rate of 2 �C/min to �40 �C and again held isothermally for 5 min. The samplewas then unloaded to a stress of 5.62 kPa at a rate of 11.2 kPa/min,whichwas followed by an isothermal, constant force hold for 5min.The sample was then heated at 2 �C/min to 25 �C, where it was heldfor 30 min. The cycle to this point is very similar to the one-wayshape memory cycles used to prepare samples for the deforma-tion and recovery x-ray scattering studies; however, the steps thatfollow will describe the “ratchet” behavior of the cycle. The samplewas heated to a recovery temperature, Trec, at 2 �C/min, heldisothermally for 30 min, cooled to 25 �C at 2 �C/min, and held for5 min. At the end of this isothermal, it was heated to the next re-covery temperature, which was 2 �C above the previous recoverytemperature, at 2 �C/min, where it was held for 30min before it wascooled to 25 �C at 2 �C/min and held isothermally for 5 min. Thisratchet process thus proceeded, with recovery temperaturesspaced at 2 �C intervals from 25 �C to 53 �C. After the 53 �C ratchet,the sample was heated at 2 �C/min to 74 �C to completely recoverthe sample. The strain of the sample was calculated using thelength of the sample at 74 �C and 5.62 kPa and was plotted versustemperature and time to demonstrate the ratchet shape memorybehavior.

To analyze microstructure using wide-angle x-ray scattering,film samples weremounted in the instrument (Rigaku S-MAX 3000

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e2820 2811

Pinhole SAXS systemwith adjustable sample-image plate distance)and exposed with Cu Ka radiation (1.54 Å) for 1 h. Two differentsample-to-detector distances were required to study the reflectionsof the smectic layers, siloxane, and mesogens. The 236.4 mmsample to detector distance captures reflections from the smecticlayers and the siloxane backbones, and the 65.7 mm sample todetector distance captures the siloxane and mesogen in the net-work’s backbone. Intensities were averaged azimuthally to produceplots of intensity versus two theta position for d-spacing calcula-tions according to Bragg’s law (Eq. (1)).

d ¼ l=ð2sinðqÞÞ (1)

The intensity over a selected two theta range was averaged foreach azimuth and plotted to give intensity versus azimuth plots.The positions of the peaks in the azimuthal scans indicate orien-tation, and these measurements were used to construct schematicsof the LCE’s microstructure. The first step in constructing theschematics was to determine the position of the smectic andmesogen reflections. To this end, two angles were defined: F wasdefined as the angle between the smectic layer normal and thestrain axis, and 4 was defined as the angle between the mesogenmajor axis and the strain axis. These angles are illustrated inScheme 2. To determine these two angles, the positions of thepeaks were determined by fitting Gaussian peaks to the azimuthalscan data using PeakFit software (Systat Software, Inc, San Jose, CA).The maximum of the fitted peaks was taken to be the location,measured in degrees, of the peak. Once the peak locations werenoted, F and 4 were determined (see SI for details). Finally, the tiltangle of the mesogen within the smectic layers is defined as theangle between the mesogen major axis and the smectic layernormal and was determined using Eq. (2), which stems from thegeometry of the problem.

b ¼ Fþ f (2)

The mesogen tilt angle and the angles the strain axis forms with1) the smectic layer normal (F) and 2) the mesogen major axis (4)were thus determined for each pattern from the deformation andrecovery studies. This enabled a detailed understanding of micro-structural changes in the LCE during these processes.

3. Results and discussion

E-5H805tB20 was synthesized by feeding 80% 5H and 20% 5tBinto the reaction outlined in Scheme 1, as described in the experi-mental section, resulting in a liquid crystalline elastomer with a gelfraction of 93%. We previously reported on the synthesis andcharacterization of 5H-5tB shape memory elastomers with fourdifferent compositions, so the thermal and thermomechanicalbehavior of E-5H805tB20 is briefly described as relevant to thiswork. The reader is directed to our previous manuscript [5] formore detailed characterization of this and additional 5H-5tB LCEcompositions. The temperature-dependent linear viscoelasticproperties of E-5H805tB20 are shown in Fig. 1(a), where the LCEshowed a step-down in tensile storagemodulus as thematerial washeated through its broad isotropization transition. Above this steptransition, the modulus of the LCE is similar to an isotropic rubberand is dominated by the crosslink density of the material. Using theisotropic modulus and a model of rubber elasticity, the molecular

1 The model of rubber elasticity used to calculate molecular weight betweencrosslinking points for the elastomer studied was G ¼ 4=5rRT=Mc, whereT ¼ 343 K, r was 1030 kg/m3, and G was calculated from E measured using DMA at343 K using the relationship E ¼ 3G.

weight between crosslinks was calculated to be 39 kg/mol.1 Attemperatures below this step, the modulus is two orders ofmagnitude higher than the isotropic modulus, but is a full order ofmagnitude less than what is expected for a glassy polymer(w1 GPa). Closer inspection of the trace shows that the dramaticdrop in storage modulus did not occur in one continuous step:there is a low temperature sectionwith a smaller slope and a highertemperature section with a larger slope. The loss tangent tracesupports this observation: a broad shoulder exists that is not fullyseparated from the more intense higher temperature peak. Wepreviously reported that the higher temperature transition, markedby a peak in the loss tangent trace and correlated with an endo-therm observed in the DSC heating trace (Fig. 1(b), trace (i) and(iii)), marks the isotropization of E-5H805tB20. Below the iso-tropization transition, the mesogens are organized into a smectic-Cliquid crystalline phase, as will be discussed in detail later. Uponheating through this transition, the mesogens mix with thesiloxane spacers, resulting in a detectable endothermic transitionand the formation of an isotropic phase. The lower temperaturetransition, marked by the low temperature shoulder of the losstangent and a stepwise change in heat flow in the DSC heating trace(Fig. 1(b), trace (iii)), is attributed to the glass transition of themesogens, which are thought to be able to vitrify within thesmectic layers. The smectic layers may be thought of as a phase-separated block copolymer, where the siloxane chains and meso-gens form two different blocks. In this case, the siloxane chainsallow the fluidity of the network to be maintained at all tempera-tures shown, but the phase-separated mesogens could indepen-dently vitrify within their layers if cooled sufficiently.

E-5H805tB20 has excellent one-way shape memory fixing andrecovery properties. Our previous work [5] has quantified the shapememory fixing and recovery properties of this LCE. In these ex-periments, a tensile stress of 100 kPa was applied to the LCE filmabove the isotropization transition, which caused the elastomer todeform. Cooling through the combined transition under loadcaused the elastomer to elongate, part of a reversible thermoelasticphenomenon common to LCEs [11e16] and referred to by someresearchers as “two-way shapememory”. Unloading the film belowthe combined transition resulted in fixing of the LCE into a tem-porary shape, with a fixing ratio of 99.7%, indicating almost perfectfixing of the temporary shape. Heating of the elastomer underminimal tensile load through the combined transition resulted inrecovery of greater than 95% of the permanent shape. Such shapememory results were repeatable for subsequent cycles, a charac-teristic common to covalently crosslinked shapememory polymers.For this material, the fixing and recovery ratios for the second cyclewere 99.7% and 98.1%, respectively [5]. In these studies, we alsonoted that E-5H805tB20 is highly extensible, with a failure strainexceeding the instrumental limit (380%).

To study the ex situ recovery of E-5H805tB20 withWAXS, the LCEwas deformed, fixed, and recovered in 2 �C increments. It was hy-pothesized that holding the LCE isothermally at Trec for 30 minwould allow the film to reach an intermediate “equilibrium” shapeso that the LCE will not continue to relax during the x-ray exposureat room temperature. Further, it was hypothesized that, once anintermediate strainwas reached at a particular Trec, the LCE will notrecover more until the temperature is increased to a new Trec. A“ratchet” shape memory test was designed to test these hypothe-ses, where the LCE was deformed and fixed into a temporary shapeand heating for recovery was interrupted. Here, the recovery pro-ceeded by heating to a Trec, holding isothermally for 30min, coolingback to 25 �C, holding isothermally for 5 min, and then heating tothe next Trec that was 2 �C higher than the previous Trec. The strainversus temperature curve from the ratchet shape memory experi-ment is shown in Fig. 2(a). When the sample was cooled from the

Scheme 1. Hydrosilylation synthesis of a polydomainmain-chain liquid crystalline elastomer using the 5H and 5tBmesogens. Themesogen and crosslinkerwere dissolved and cooled to0 �C where the Pt(0) catalyst and hydride-terminated poly(dimethylsiloxane) (H-PDMS) were added. This mixture was transferred to a crosslinking cell, where the reaction ran for 24 h.

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e28202812

recovery temperature to 25 �C, held isothermally, and then heatedto the next recovery temperature, the strain versus temperatureplot has a negligible slope, signifying that the sample did notchange in length and the LCE did not recover during the cooling and

reheating parts of the cycle. The vertical separation between thesetwo horizontal lines is the amount of strain the LCE recoveredduring the isothermal hold at the higher recovery temperature.These horizontal lines are spaced irregularly on the plot: they

Fig. 2. Strain versus temperature (a) and stress, strain, and temperature versus time(b) of the E-5H80-5tB20 ratchet shape memory cycle. The LCE was deformed at 74 �C toa stress of 100 kPa before cooling to �40 �C at 2 �C/min and unloading to a slighttensile stress. The sample was then heated to 25 �C and held isothermally before thetemperature was increased. The sample was held isothermally at this new recoverytemperature before cooling to 25 �C, holding, and then heating to the next recoverytemperature. (a) Strain versus temperature shows that strain recovers during theisothermals, not during the cooling and reheating steps. (b) Strain versus experimentaltime shows that strain recovery occurs exponentially during the isothermals and thatexpansion/contraction of the sample occurs around the isotropization transition due tothe two-way shape memory effect.

Fig. 1. (a) Tensile storage modulus (E’) and loss tangent (tan(d)) trace of E-5H805tB20.The LCE was relaxed in a stress-free state at 100 �C prior to testing (3 �C/min heatingrate, 10 mm amplitude, 1 Hz frequency, 108% force track). (b) DSC traces of E-5H805tB20

at 10 �C/min. Trace (i) is the second heat, and trace (ii) is the second cool. Trace (iii) isthe first heating trace after cooling from the isotropic state to 50 �C (denoted by the “x”on the second cooling trace) and annealing for 1 h.

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e2820 2813

are closer together at low recovery temperatures than they are atthe higher recovery temperatures, indicating that more strainrecovered during the isothermal holds at higher temperatures.

Strain, temperature, and stress are plotted against experimentaltime from the ratchet shape memory experiment in Fig. 2(b). Thestress versus time trace shows that stress was ramped during thedeformation step, held constant at 100 kPa during the cooling step,and reduced to the preload stress of 5.62 kPa during the unloadingand recovery steps. The temperature trace shows that temperaturewas held constant during the deformation step, reduced linearlyduring the cooling step, held constant during the unloading step,and finally cycles between heating, holding, cooling, holding, andheating for the ratchet recovery steps. The strain trace shows thatthe strain reached ca. 75% from the deformation step and increasedto ca. 175% from cooling under 100 kPa stress due to the two-wayshape memory effect. It is noted here that the sample did creepduring the isothermal hold under high stress, but the increase instrain due to creep was small (2.7%), and thus its contribution to theoverall shape change is only 1.5% of the total and is considerednegligible. The strain decreased to ca. 165% when it was heated to25 �C and held isothermally for 30 min. The strain response to thetemperature cycling shows that, upon reaching each new recoverytemperature, the strain decreased with the form of an exponentialdecay. After the isothermal recovery, the strain did increase upon

cooling, if the recovery temperature was around the clearing point,due to the two-way shape memory effect; i.e., order-inducingelongation. Strain did not change once the sample reached 25 �C.Upon heating to some recovery temperatures, the sample con-tracted due to the two-way shape memory effect, but no additionalstrain was recovered until the temperature exceeded the previousrecovery temperature, where an exponential dependence withtime was again assumed once the new recovery temperature wasreached.

The most important finding from this experiment was that thestrain at the end of the isothermal recovery segment was reachedagain after cooling to 25 �C, holding, and reheating. This means thatthe cooling and reheating steps do not affect the strain the samplereaches, and that the once the sample has recovered at a temper-ature, it will not recover further at lower temperatures, aside fromthe reversible changes in strain due to the two-way effect. Thesefindings confirm that cooling to room temperature and reheating tothe next recovery temperature will not affect the strain recovery ofthe LCE, and they validate ex situ design of the recovery experiment.

Wide-angle x-ray scattering of E-5H805tB20 fixed at differentstrains using one-way shape memory cycles conducted at differentstresses are shown in Fig. 3(a). Note that the stretch axis is hori-zontal and that, like all otherWAXS data presented, two sample-to-detector distances (236.4 mm and 65.7 mm) were used for thedeformation and recovery studies in order to resolve differentfeatures over a wide range of d-spacings. The unstretchedE-5H805tB20 patterns (Fig. 3 (a(i, ii))) show isotropic rings for the

Fig. 3. Wide angle x-ray scattering of E-5H805tB20 (a) deformed and fixed at different strains using controlled force shape memory cycles and (b) recovering from an initial strain of169% using interrupted isothermal heating. Each sample was studied at two different sample to detector distances, 236.4 mm (images i, iii, v, vii, and ix) and 65.7 mm (images ii, iv,vi, viii, and x), to capture the reflections from the smectic layers, siloxane chains, and mesogens. (a) In the deformation study, each column was stretched to a different strain usingthe stress given in parentheses: 0% (0 kPa) (i, ii), 10.1% (5 kPa) (iii, iv), 27.1% (10 kPa) (v, vi), 40.2% (15 kPa) (vii, viii), 179% (100 kPa) (ix, �). (b) In the recovery study, each column wasrecovered at a different temperature, given in parentheses, to produce the strain noted: 164% (25 �C, initial state) (i, ii), 95.1% (51 �C) (iii, iv), 57.3% (52 �C) (v, vi), 15.2% (53 �C) (vii,viii), 0% (100 �C) (ix, �). Strain axis is horizontal.

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e28202814

smectic layers (smaller bright ring in Fig.3(a(i))), siloxane backbone(larger, diffuse, bright ring in Fig. 3(a(i)) and the smaller bright ringin Fig. 3(a(ii))), and mesogen reflections (larger, diffuse, bright ringin Fig. 3(a(ii))). Isotropic rings are expected for these polydomainelastomers because, though there is liquid crystalline order withina domain, there is no overall orientation to the domains within theelastomer. As the stress of the shape memory cycle was increased,the smectic layer reflection split into meridian reflections(Fig. 3(a(iii))), which then separated into a four point pattern that istypical of a chevron pattern commonly observed in smectic-Csystems [10] in samples deformed to higher strains. The mesogenreflections also suggested orientation changes with increasingsample strain. The 65.7 mm sample-to-detector pattern from the10 kPa one-way shape memory cycle (Fig. 3(a(vi))) shows that themesogen reflection splits into two broad equatorial reflections.These broad reflections become more intense for the 15 kPa cycle(Fig. 3(a(viii))) and finally each splits into three spots for the100 kPa cycle (Fig. 3(a(x))). The three mesogen spots in the 100 kPapattern differ in their two theta position: two of the spots (thesewill be called “M2” reflections and are labeled by the horizontalarrows in Fig. 3(a(x))) are located at a smaller two theta than theother spot (this will be called the “M3” reflection and is labeled bythe upright arrow in Fig. 3(a(x))). The splitting of the M2 reflectionsuggests that the mesogens were tilted with respect to the strainaxis, which was horizontal in this case, while the lack of splitting in

the M3 reflection indicates that the mesogen was oriented alongthe strain axis. It should be noted that there was another reflectiondue to the mesogens that appeared very close to the siloxanereflection. This reflection, called the “M1” reflection, is a meridianreflection that was only observed in well oriented samples, such asFig. 3(a(x)). This reflection has a 2q position of about 13�, whichcorresponds to a d-spacing of about 6.8 Å, and is attributed toperiodicity within the rigid part of the mesogen. Fig. S1 in Sup-porting Information shows a schematic of the mesogen with foursections labeled. The length of each section was determined bysimulation using Chem3D Pro 8.0 Software (CambridgeSoft Corp.,Cambridge, MA). The segments have lengths ranging from 6.05 Å to6.77 Å, and thus are close to what is experimentally observed forthe M1 reflection. These d-spacings are also quite similar to theperiodic spacing of 6.3 Å that has been observed in oriented liquidcrystalline polymers of 4-hydroxybenzoic acid [17].

The WAXS patterns for E-5H805tB20 recovering from a fixedstate by stepped isothermals show that the smectic layer andmesogen reflections changed as the recovery temperatureincreased (Fig. 3(b)). In the initial state, the samplewas at a strain of164% and had a smectic reflection that was split into four points(Fig. 3(b(i))), an equatorial M2 reflection that was split into twopeaks (Fig. 3(b(ii))), and an M3 equatorial reflection (Fig. 3(b(ii))).As the recovery temperature increased, the four point smecticpattern shifted to the meridian, and those peaks became so close to

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e2820 2815

each other that they appeared to merge to produce just twoopposing meridian reflections (Fig. 3(b(vii))). Once the LCE washeated through the clearing point, orientation was completely lost,as shown by the isotropic ring in the sample relaxed at 100 �C([Fig. 3](b(ix))). The patterns captured at 65.7mm show that theM2and M3 reflections were split in the initial state (164% strain,Fig. 3(b(ii))), and that they remained split until all orientation waslost (Fig. 3(b(x))), although the M2 and M3 reflections becamediffuse and difficult to distinguish as recovery temperatureincreased. Additionally, the spots in theM2 reflection became evenmore split as recovery temperature increased.

The intensity versus two theta plots for the deformation studyare shown in Fig. 4(a,b) for the 236.4 mm (Fig. 4(a)) and the65.7 mm (Fig. 4(b)) sample-to-detector distances. The peak fromthe smectic layer reflection (Fig. 4(a)) shifted to larger two thetapositions as stress and strain were increased in the sample, asobserved by comparing trace (i) to trace (viii). The siloxane reflec-tion (Fig. 4(b)) was broad and it was thus difficult to determine howmuch, if at all, this reflection shifted with increasing stress. TheM1reflection appears as a shoulder at higher two theta values of thesiloxane reflection in trace (ii) and becomes more evident as stressis increased. It was also difficult to tell how the peaks from the M2and M3 reflections varied in the deformation study (Fig. 4(b)). It

Fig. 4. Intensity versus 2q from WAXS for sample to detector distances of: (a,c) 236.4 mm adifferent strains by deformation at the stresses noted in parentheses. Traces are numbered se(12.5 kPa) (iv), 40.2% (15 kPa) (v), 63.3% (25 kPa) (vi), 84.1% (50 kPa) (vii), and 179% (100 kPtemperatures from an initial state of 169% strain. Traces numbered sequentially by strain and(iv), 71.7% (49 �C) (v), 60.0% (51 �C) (vi), 57.3% (52 �C) (vii), 15.2% (53 �C) (viii), 6.59% (57 �

should be noted that all traces are normalized by area under thecurve, and because sample thickness decreased as strain increased,the signal to noise ratio in the traces worsens with increasingstrain. The intensity versus two theta plots for selected tempera-tures from the recovery study are shown in Fig. 4 (c,d) for the236.4 mm (Fig. 4(c)) and the 65.7 mm (Fig. 4(d)) sample-to-detector distances. The two theta position of the smectic layerreflection decreased and the reflection broadened with increasingrecovery temperature. This is the same trend (in reverse order) aswas observed in the deformation study. Specifically, the smecticpeak position in the deformation study shifted to larger two thetavalues (finer d-spacing) and sharpened as strain of the LCEincreased. Similar to the deformation study, the positions of thesiloxane and mesogen reflections did not change in a detectablemanner.

The d-spacing of the smectic layer peak was calculated using thetwo theta position of the peaks from the intensity versus two thetascans and Bragg’s law (Eq. (1)), and these are plotted versus thestrain of the LCE in Fig. 5 for the deformation (Fig. 5(a)) and re-covery (Fig. 5(b)) studies. In the deformation study, the d-spacing ofthe smectic layer was initially 40.3 Å and slightly decreased withincreasing strain to a value of 39.4 Å at 179% strain and a value of34.7 Å at 204% strain. This indicates that the smectic layer was

nd (b,d) 65.7 mm. For the deformation study (a,b), the traces are E-5H805tB20 fixed atquentially by strain and stress: 0% (0 kPa) (i), 10.1% (5 kPa) (ii), 27.1% (10 kPa) (iii), 34.9%a) (viii). For the recovery study (c,d), the traces are E-5H805tB20 recovered at differentrecovery temperature: 169% (25 �C) (i), 165% (31 �C) (ii), 141% (41 �C) (iii), 106% (44 �C)C) (ix), and 0% (100 �C) (�).

Fig. 5. d-spacing of the smectic layer reflection from the deformation (a) and recovery(b) studies of in E-5H805tB20. (a) The d-spacing of the smectic layers is plotted againstthe strain resulting from one-way shape memory cycles run at different stresses. (b)The d-spacing of the smectic layers is plotted against the strain of the LCE (initial strain169%) after isothermal recovery at different temperatures.

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e28202816

thinner for samples fixed at larger strains. The recovery studyshowed that the smectic layer thickness reached intermediate d-spacing values that were not observed in the deformation study.Initially, the sample was at 164% strain and had a smectic layerspacing of 34.7 Å, but as the sample recovered, the layers becamethicker, where the spacing measured 40.3 Å for a completelyrelaxed sample. A recent report of a liquid crystalline network wasshown to increase in smectic layer thickness when annealed at70 �C, which is between its Tg (5 �C) and isotropization transition(w100 �C) [18]. It is possible that the increase in layer thickness inthe recovery series samples, which are held isothermally atdifferent temperatures, some of which are between the mesogen Tgand isotropization transitions, could be due to the same effect.Kinetically, however, this is less favorable in the compositiondescribed in this work than in the referenced study due to thejuxtaposition of themesogen glass transition and the isotropizationtransitions in this LCE.

Fig. 5(b) can be roughly divided into three regions that havedifferent slopes. In Region III, which corresponds to samplesrecovered at the lowest temperatures (25 �Ce41 �C), strain recov-ered without much change in layer thickness. In Region II,which corresponds to samples recovered at higher temperatures(42 �Ce53 �C), layer thickness increased with strain recovery.Finally, in Region I, which corresponds to samples recovered at thehighest temperatures (54 �Ce100 �C), layer thickness dramaticallyincreased with strain recovery. Comparing Fig. 5(b) with the LCEthermal transitions determined by DSC, the three regions roughlycorrespond to temperatures below and through the mesogenglass transition temperature (Region III), temperatures above the

mesogen glass transition and nearing the clearing point (Region II),and temperatures located in and above the clearing point (Region I).These thermal transitions are thought to play an important role inshape recovery by controlling mesogen mobility within the elas-tomeric network. Specifically, below the mesogen Tg (Region III), themesogens are vitrified within the smectic layers, and they pin thepolymer chains and prevent the elastomer from recovering strain.Above the mesogen Tg but below the clearing point (Region II), themesogens are no longer vitrified. As a result, the polymer chainsrecover some strain through rubber elasticity, though liquid crys-talline order is maintained and prevents full recovery. Finally,heating through the clearing transition (Region I) causes loss ofliquid crystalline order and permits full recovery of strain.

In addition to changes in the layer thickness, deformation andrecovery affect the orientation of the smectic layers and themesogens. As such, azimuthal scans of the smectic layer reflection(Fig. 6 (a,b)), the M2 reflection (Fig. 6 (c,d)), and the M3 reflection(Fig. 6 (e,f)) are shown for the deformation (Fig. 6(a,c,e)) and re-covery (Fig. 6(b,d,f)) studies. In this figure, the strain axis is locatedat 90� and 270�. At the start of the deformation study, the azimuthalscan of the smectic layer reflection (Fig. 6(a), trace (i)) has no peaks,indicating that the polydomain unstretched sample had no overallorientation. As stress was increased, two broad peaks appeared inthe azimuth plot (Fig. 6(a), trace (ii)) from the orientation of thesmectic layer reflections at the meridian of the sample. These broadpeaks each split into two peaks, which became sharper and moreintense, but did not shift in position, as the fixed strain increased. Atthe start of the recovery study (LCE is at 164% strain), the smecticlayer reflection (Fig. 6(b)) was split into four sharp peaks (trace (i)).As the recovery temperature increased, the peaks decreasedin intensity, increased in breadth, and e importantly e shifted inazimuthal position, indicating rotation of the layer normal. Inparticular, the peaks shifted towards the meridian and are so closein trace (viii) that they appear merged together, although fourpeaks can still be discerned. Once the LCE was heated through theisotropization transition and relaxed at 100 �C, all orientation of thesmectic layer reflection was lost (Fig. 6(b), trace (x)).

The azimuthal scans of the M2 reflection measured from thedeformation study are shown in Fig. 6(c). Like the smectic layerreflection, there was no orientation of the M2 reflection in theunstretched sample. As the strain increased, however, two equa-torial reflections appeared and increased in magnitude more andmore as strain continued to increase. The M2 reflection split withincreasing deformation, with two peaks evident in trace (vii) andtrace (viii), though the flat top of the peaks in all the traces mayindicate two peaks remained present but that they were too closeto resolve. The peaks did not shift in position, as shown in Fig. 6(c).In the recovery study (Fig. 6(d)), however, the M2 peaks wereinitially split and, in addition to becoming less intense, the peaksmoved apart and approached the meridian as the recovery tem-perature increased. The azimuthal scans of the M3 reflection areshown in Fig. 6(e,f) for the deformation (e) and recovery (f) studies.Fig. 6(e) shows that, for the deformation study, equatorial re-flections appeared and intensified as strain increased. The peaks inthe recovery study have the same trend: as strain was reduced, thepeaks became less intense. It should be noted that the two peaks oneach side of the equatorial reflection in the recovery study (Fig. 6(f))are due to theM2 reflection, which overlaps with theM3 reflection.

The azimuthal scans in Fig. 6 show that orientation of the re-flections changes with deformation and recovery. The angular po-sitions of the peaks within the azimuthal scans can be used todetermine the microstructure of the LCE (see SI for details). Threeangles were calculated for the patterns that displayed orientation:F, the angle between the smectic layer normal and the strain axis,4, the angle between the mesogen major axis and the strain axis,

Fig. 6. Intensity versus azimuth of the smectic (a, b), M2 (c, d), and M3 (e, f) reflections for the deformation (a, c, e) and recovery (b, d, f) studies. (a, c, e) LCE was fixed at differentstrains by deformation at different stresses. Traces are numbered sequentially by strain and stress: 0% (0 kPa) (i), 10.1% (5 kPa) (ii), 27.1% (10 kPa) (iii), 34.9% (12.5 kPa) (iv), 40.2%(15 kPa) (v), 63.3% (25 kPa) (vi), 84.1% (50 kPa) (vii), and 179% (100 kPa) (viii). (b, d, f) LCE was recovered at different temperatures from an initial state of 169% strain, with tracesnumbered sequentially by strain and recovery temperature: 169% (25 �C) (i), 165% (31 �C) (ii), 141% (41 �C) (iii), 106% (44 �C) (iv), 71.7% (49 �C) (v), 60.0% (51 �C) (vi), 57.3% (52 �C)(vii), 15.2% (53 �C) (viii), 6.59% (57 �C) (ix), and 0% (100 �C) (�).

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K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e28202818

and b, the tilt angle of the mesogen with respect to the smecticlayer normal. Only two of these angles are independent quantities,given the inter-relation of Eqn. (2). The angles are illustratedin a schematic of the smectic layers in the chevron patterns inScheme 2. In addition to these angles, discussion regarding thechevron structure belowwill also refer to the chevron’s angle, whichis the acute angle of the smectic layer zigzag. Another way to thinkof this angle would be to consider rotating the chevron pattern inScheme 2 by 90� to give a “W” type pattern. The chevron angle is theangle of the peaks and valleys of this “W”. A sharper chevronwith alarger pitch would correspond to a taller, narrower “W”, while abroader chevron with a shallower pitch would correspond to ashorter, wider “W”. Using geometry, it is possible to show that onehalf of the chevron’s angle is complementary to F. “Sharpening” ofthe chevron or an increase in the chevron’s pitch both correspond toa decrease in the chevron’s angle and an increase in F.

Fig. 7 shows the dependence ofF,4, and b on deformation stress(Fig. 7(a)) and recovery temperature (Fig. 7(b)). Fig. 7(a) shows thatFwas initially small (27� at 5 kPa stress), but then this increased to45� at 25 kPa stress, and then slightly decreased to reach 42� at140 kPa. The increase in F corresponds to the sharpening of thechevron. The M2 reflection was split only in samples deformed atstresses that exceed 25 kPa, thus 4 was zero for samples deformedat stresses less than this limit. Because 4 was zero and F whileincreased, beta appears to increase to a value of 53�. It should benoted that it is possible that the mesogens were tilted with respectto the strain axis (to yield a nonzero 4), but that this could not beresolved in the patterns. In such a case, the mesogen tilt angle withrespect to the smectic layer normal, b, may be constant throughoutthe deformation series. Regardless of the mesogen tilt angle b atlow strains, the deformation study shows that the deformed sam-ple reaches about the same mesogen tilt b (Fig. 7(a)) and smecticlayer spacing (Fig. 5(a)) when it is deformed at stresses that exceed

Scheme 2. Schematic of structural changes in E-5H805tB20 during deformation (a) and reoriented along the strain axis and the smectic layers have a broad pitch. As deformation strerespect to the strain axis and smectic normal. (b) As the LCE recovers during interrupted isconstant.

25 kPa, which suggests that this structure is favorable energeticallyand is not highly stress or strain dependent.

In contrast, the recovery experiments show much more varia-tion in F and 4 thanwas observed in the deformation experiments.Fig. 7(b) shows that, as recovery temperature increases (move fromleft to right on the plot), F decreases and 4 increases. The decreasein F means that the pitch of the chevron decreased and its anglesbecame less sharp as recovery temperature increased. Because thedecrease inF occurred with an increase in 4, the mesogen tilt anglewith respect to the smectic layers (b) remained relatively constant:b decreased only from 53� to 49� when recovery temperatureincreased from 25 �C to 52 �C. When the recovery temperaturereached and exceeded 53 �C, b decreased to 46�, changes that arethought to be possible due to the proximity of the clearing transi-tion. The tilt angle could not be determined above 55 �C becauseorientation of the M2 reflection could not be detected in theazimuthal scans. The recovery study shows that the LCE recoversstrain as the pitch of the chevron decreases, and this is accompa-nied with mesogen rotation with the layers (b; Fig. 7(b)) andthickening of the smectic layers (Fig. 5(b)). The mesogen tilt anglewith respect to the smectic layers (b) remains relatively constantthroughout the process. The tilt angle of a mesogen within asmectic layer has been studied theoretically by relating the tilt tothe Z-like molecular shape that occurs when the rigid mesogen isattached to two chain ends [19,20]. The mesogen tilt has also beenrelated to a permanent dipolar interaction that breaks the sym-metry of rod-like uniaxial molecules to give a biaxial tilted phase[21,22]. A weakness of these studies, however, is that they did notallow free rotation of the molecules, which experimental studieshave shown to be important. More recently, simulations haveshown that a balance of dipole interactions and GayeBerne in-teractions can predict smectic-C phase formation, and that chang-ing these interactions leads to different tilt angles [23]. It is thus

covery (b). (a) When the LCE is deformed at low stresses, the mesogen appears to bess is increased, the pitch of the chevron sharpens, and the mesogens are inclined withothermal heating, the smectic layers thicken, F decreases, 4 increases, and b remains

Fig. 8. Angle between the smectic layer normal and the strain axis (F) (a), angle be-tween the mesogen major axis and the strain axis (f) (b), and the mesogen tilt angle(b) (c) plotted against strain of the LCE from the deformation (black circles) and re-covery (gray triangles) studies.

Fig. 7. Angle between the smectic layer normal and the strain axis (F), angle betweenthe mesogen major axis and the strain axis (4), and the mesogen tilt angle (b) plottedagainst the controlled variables stress (a) and recovery temperature (b) from thedeformation and recovery studies, respectively.

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e2820 2819

hypothesized that the reasonwhy the tilt angle does not vary in thedeformation and recovery studies is because the thermodynamicscontrolling the mesogen environment do not change appreciably.

The unexpected occurrence of a soft material displaying suchexcellent shape memory properties prompted study of the micro-structure of E-5H805tB20 deformed at different stresses andrecovered from an initially fixed large strain (169%). We have pre-viously shown [5] that these main-chain LCEs are soft due to theirsiloxane backbone, yet they have two mesogen-dependent transi-tions, both of which may be detected by DSC and DMA. The higherof these two transitions is the isotropization transition, belowwhich mesogen/siloxane layering exists as a nanoscale, smecticphase. The lower temperature transition is due to the glass tran-sition of the mesogen within the smectic layers of the fluid-likesiloxane network. For the LCEs to remain soft while capable ofshape memory cycling, the mesogenic layers of the smectic phasemust spatially percolate. While a high degree of orientation may beheld fixed by shape fixing, the authors have found that it is possibleto fix other types of deformation featuring very little orientation.Thus, orientation is apparently not required for strain fixing inthese materials. The LCE’s microstructure, including the mesogenorientation, mesogen tilt angle, and spacing between layers ofmesogens (smectic layers), is thus highly relevant to macroscopicproperties of the network. Interestingly, differences were found inthe microstructure of the LCE with respect to strain in the defor-mation vs. recovery studies. These differences are highlightedquantitatively in Fig. 8. Concerning the angle between the smecticlayer normal and the strain axis, F, the recovery study showed thatthis decreased linearly with the recovered strain, but was nonlinear

(increased and then plateaued) in the deformation study. Theplateau in F versus strain in the deformation study shows thatchevron’s pitch is similar to what was observed in the LCE at thestart of the recovery study, and that once the chevron reaches thisconfiguration, it is invariant with increasing stress and strain. Theshallow pitched chevron found in the recovering LCE was observedonly in samples deformed using small (less than 25 kPa) stresses,though it is noted that this may not be the same structure becausethe tilt angles were different between the two cases. The differencein the chevron structure between the deformation and recoverystudies proves that strain itself does not determine the smecticlayer structure. Rather, the structure of the chevron changes basedon the thermomechanical path of the LCE to that strain. Themicrostructural difference between the deformation and recoverysamples can be attributed to the lowest energy structure, which ishypothesized to be limited by the highest recovery temperature.

The mesogens within the smectic layers adjusted their orien-tation with the strain axis as the smectic layers rotated so that the

K.A. Burke, P.T. Mather / Polymer 54 (2013) 2808e28202820

tilt angle between the mesogen and the smectic layer normal re-mains about constant for the deformation and recovery processes.This led to the construction of schematics that highlight the dif-ferences between the two processes. At low stresses in the defor-mation study, the chevron had a shallow pitch and the tilt of themesogens with respect to the strain axis could not be determined,but as stress increased, the chevron pitch (F) increased, and the tiltangle of the mesogens became detectable (Scheme 2(a)). Once thisstructure was reached, increasing the deformation stress and fixedstrain in the LCE did not change the structure, which suggests thatthis structure is the deformed LCE’s minimum energy state at 25 �C.In the recovery study, the sample is initially organized in thischevron structure, but recovery caused the chevron’s pitch todecrease (Scheme 2(b)). The flattening of the chevron’s pitch, asevidenced by the transition of the smectic layer reflections from asharp, four point pattern to broad and nearly-merged meridianreflections (Fig. 6(b)), does not suggest that the elastomer isbecoming smectic-A in nature, which typically has sharp meridianlayer reflections and diffuse equatorial mesogen reflections. Unlikea smectic-A structure, themesogens in this LCEmaintain a tilt angleranging from 46� to 53� with respect to the smectic layer normal.Thus, the LCE recovered by relaxing the pitch of the layers, butmaintained its smectic-C phase.

One final point about the microstructure of the strained LCEs isthat the M3 reflection is always oriented along the strain axis. Thismesogen peak may be due to the population of mesogens at thepeaks and valleys of the chevron structure. Another possible causeof this reflection is from “pockets” of mesogens oriented in anematic phase. Because the LCE is a polydomain, it is unclear howmesogens orient between the domains. If there are mesogens in anon-smectic liquid crystalline phase, then it is possible that theywill orient along the strain axis and produce these equatorial re-flections. The sharpness of the M3 reflection is maintained evenwhile M2 broadens and the chevron changes and is thought to bean important clue to its yet undetermined origin.

4. Conclusions

We have reported the microstructural characterization of a softshape memory network, a main-chain LCE. The LCE studied in thiswork was prepared by hydrosilylation chemistry, in which dienemesogens were polymerized with a hydride-terminated PDMSspacer and crosslinked with a tetravinyl crosslinker. The particularLCE composition, E-5H805tB20, was selected because both its meso-gen glass transition and isotropization temperatures are super-ambient, which allowed room temperature x-ray studies of theelastomer to be conducted at a temperature below these materialtransitions. Previous work has demonstrated that several composi-tions of 5H-5tB LCEs have excellent shape memory fixing and re-covery behavior, which is highly unusual in a soft material and thusprompted further investigation. It was found that deforming theelastomer above its isotropization transition to different stresses leadto similar microstructures above a critical deformation stress. Uponrecovery, however, intermediate microstructures developed. These

results suggested that bulk shape changes observed in the LCE areindeed tied to microstructural changes, which are, in turn, dictatedby the mobility of the mesogen in the network. It is thought that thisworkmay aid in future studies of the physics of shapememory in thecontext of liquid crystallinity and soft elasticity that is observed insmectic-C elastomers.

Conflict of interest

The authors declare no competing financial interest.

Acknowledgments

The authors are grateful for the support of the National ScienceFoundation. PTM acknowledges support from National ScienceFoundation (DMR-0758631 and DMR-1004807). KAB acknowl-edges the support of a Graduate Research Fellowship from theNational Science Foundation (DGE-0234629).

Appendix A. Supplementary data

Supplementary data related to this article can be found at http://dx.doi.org/10.1016/j.polymer.2013.03.049.

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