pubs.acs.org/MacromoleculesPublished on Web 05/07/2010r 2010 American Chemical Society

Macromolecules 2010, 43, 5043–5051 5043

DOI: 10.1021/ma100470w

Investigating the Molecular Origins of Responsiveness in FunctionalSilicone Elastomer Networks

Julie A. Crowe-Willoughby,†,§ Derrick R. Stevens,‡ Jan Genzer,*,† and Laura I. Clarke*,‡

†Department of Chemical & Biomolecular Engineering, NC State University, Raleigh,North Carolina 27695-7905, and ‡Department of Physics, NC State University, Raleigh,North Carolina 27695-8202. §Present address: MeadWestvaco, Center for Packaging Innovation,1021 Main Campus Drive, Raleigh, North Carolina 27606

Received February 28, 2010; Revised Manuscript Received April 25, 2010

ABSTRACT: Dielectric, calorimetric, and dynamic mechanical measurements were performed to delineatethe types and dynamic rates of molecular scale motion in modified poly(vinylmethyl siloxane) (PVMS)stimuli-responsive networks, where pendent groups of the form -S-(CH2)n-OH were chemically attachedto the vinyl moiety of PVMS. The glass transition temperature (Tg) for the unsubstituted PVMS networkmatches that previously reported for linear PVMS indicating that the flexibility of the polymer chainsis unaffected by the network cross-linking. In contrast, Tg increases with the introduction of pendent groupsof the type -S-(CH2)n-CH3 or -S-(CH2)n-OH, where n is 2, 6, or 11, as the different groups constrainthe siloxane backbone to differing degrees. The macroscopic response time and amplitude, as previouslymeasured by dynamic contact angle, are correlated with the observed glass transition temperatures. One con-clusion is that the flexibility of the network and the interactions between pendent groups affect responsiveness.

I. Introduction

Responsive materials alter their physical or chemical proper-ties spontaneously when exposed to an external stimulus (e.g., anelectromagnetic field, mechanical strain, temperature variation,exposure to a fluid, or change of pH or ionic strength).1-4 Whilethe response is primarily due to structural alterations occurring atthe microscopic level, i.e., changes in molecular conformation,these microscopic adjustments result in variations in measurablemacroscopic characteristics, such as material shape, size, color,or wetting properties. When considering applications of thesestimuli-responsive materials, three quantities are important: (1)the magnitude of the response, (2) the speed of the response, and(3) any hysteresis associated with the switching, such as a longreset time after stimulation. Speed of response should be tunablefor specific applications. For instance, an ideal material forsensing possesses a large, fast response with little hysteresis.However, all three of these parameters can be interlinked; e.g.,alterations to increase the response magnitude could also unin-tentionally increase response time.

In this work, we explore the relationship between structuralparameters, microscopic dynamics, and macroscopic responsetime for silicone elastomer networks acting as stimuli-responsivesurfaces. We employ differential scanning calorimetry (DSC),dynamic mechanical thermal analysis (DMTA), and dielectricrelaxation spectroscopy (DRS) to examine thoroughly variousmolecular motions in responsive surfaces made from modifiedpoly(vinylmethyl siloxane) (PVMS) networks. We then correlatethese molecular motions to the measured macroscopic response(as determined by dynamic contact angle (DCA)) as a function ofmaterial structure.

In the PVMS network system, the cross-links of the networkprovide a solid-like nature, despite the low Tg of linear PVMS

(≈150 K) which allows for a high degree of flexibility at roomtemperature.5,6 Due to the difference in surface energies betweenthe vinyl and methyl groups, PVMS is moderately amphi-philic. Because the network is flexible, the side-chains canrearrange readily when placed in contact with another medium,thus delivering the component resulting in the lowest interfacialenergy (dependent on the surroundings) to the interface. Forinstance, the hydrophobic methyl groups are exposed to thesurface when the material is placed in air. When in contact with ahydrophilic environment, i.e., water, the more hydrophilic vinylgroupsmigrate to the surface (Figure 1). As the energy differencebetween the vinyl and methyl moieties is small, we have graftedmercaptoalkanols at the vinyl group via the thiol-ene reactionthereby increasing the contrast between the opposing termini andenhancing the switching property.5,6 We previously studied thekinetics of surface rearrangement of the substituted networks(with pendent -S-(CH2)n-OH groups) by monitoring wett-ability changes using both static and dynamic contact anglemeasurements. In particular, we established that both the rateand magnitude of switching decreased with increasing length ofthe methylene spacer (n) separating the -OH terminus from thesiloxane network.5,6

The aim of this work is to observe and characterize themolecular-level dynamics in PVMS-based networks and to de-termine how thesemolecularmotions are altered by (1) the lengthof the methylene spacer in the mercaptoalkanol pendent groupand (2) the end-terminus of the pendent group (-OH vs-CH3).We perform new analysis of data on the macroscopic networkresponse and discuss how the dynamics at a molecular levelinfluence the macroscopic properties. In general, the responsetime of a stimuli-responsive material will depend not only on themagnitude of the interfacial driving force (i.e., the difference ininterfacial energy between the two switching states) but also onthe innate dynamics within the material. In particular, in thesubstituted PVMS networks, the pendent groups must movethrough the network in order to emerge at the surface; thus the

*Corresponding authors: (L.I.C.) laura_clarke@ncsu.edu; (J.G.)jan_genzer@ncsu.edu.

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dynamics of the larger system (e.g., the segmental dynamicsrelated to the Tg of the substituted network) may affect respon-siveness. Beyond the structure-response correlations pertainingto this specific system, these observations are also beneficial fordeveloping rational design strategies for the larger class ofstimulus-response surfaces consisting of polymer chains an-chored to an underlying substrate.2

We will examine six polymer networks whose molecularparameters are summarized in Table 1. Those include: PVMS,PVMS-S-(CH2)2-OH, PVMS-S-(CH2)2-CH3, PVMS-S-(CH2)6-OH,PVMS-S-(CH2)11-OH,andPVMS-S-(CH2)11-CH3. Data from poly(dimethylsiloxane) (PDMS) networks areshown for comparison. The PVMS-S-(CH2)n-CH3 speciesrepresent a special case of amphiphiles whose wettability de-creases with increasing alkyl chain length. In this work, -CH3

terminated networks will serve as a control to understand thedifferences in side-chain interactions due to the terminal -OHgroup.

II. Experimental Methods

Substrate preparation of modified PVMS networks and theexperimental details of the static and dynamic contact angleshave been previously reported.5,6 Elastomeric PVMS network

substrates were synthesized by following an alkoxy-cure methodof PVMS (Mn≈ 40 kDa) through end-group cross-linking asdescribed elsewhere.7 The storage modulus of the PVMS net-works was measured to be ≈8�104 Pa after Soxhlet extractionusing toluene for at least 48 h followed by drying at 75 �C under30 mmHg vacuum. The vinyl moieties on PVMS were subse-quently saturated via azobis(isobutyronitrile)- (AIBN)- initiatedthiol-ene radical addition of HS-(CH2)n-OH (n = 2, 6, and11).8 After the reaction, the PVMS- S-(CH2)n-OH sampleswere re-extracted, dried, and stored in a low-humidity chamber.Experiments using infrared spectroscopy in the attenuated reflec-tionmode confirmed that the grafting density of S-(CH2)n-OHwas approximately equal in each sample.

DynamicMechanical Thermal Analysis (DMTA).DMTAwasperformed using a TA Instrument Q800 operating with singlecantilever geometry. The dried polymers films were cut intoapproximately 25�10�1.5 mm3 rectangles. Storage modulus(E0), lossmodulus (E0 0), and loss tangent (tan δ) were recorded ata heating rate of 3 �C/min from -150 to þ130 �C and at a testfrequency of 1 Hz. The amplitude was set to be within the linearviscoelastic regime.

Differential Scanning Calorimetry (DSC). DSC was per-formed using a TA Instrument Q100. Two runs were carriedout for each sample in the range of -80 to þ130 �C with aheating rate of 3 �C/min under a nitrogen atmosphere. The glass

Figure 1. Schematics depicting the reorientation kinetics of PVMS-S-(CH2)n-OH samples (n = 2, 6, and 11) between hydrophobic (left) andhydrophilic (right) states.

Table 1. Properties of Bare and Modified Silicone Elastomer Networks Assessed with Differential Scanning Calorimetry (DCS), DynamicalMechanical Thermal Analysis (DMTA), and Dielectric Relaxation Spectroscopy (DRS)

Tg (K)a Tm (K)b Local mode (K)

material DSC DMTAc DRSd DSC DMTAc DRSe DMTAc DRSf

PDMS 156( 5 146-166 230 224 245-250PVMS 149 ( 3 140 ( 5PVMS-S-(CH2)2-OH 215 236( 3 230 ( 10 N/A N/A N/A 189( 2 181( 3PVMS-S-(CH2)6-OH 208 228( 12 210 ( 20 N/A N/A N/A 188( 2 180( 10PVMS-S-(CH2)11-OH 284.5 240( 40 270-290g 306 ( 7 303 ( 9 300-320 189( 8 201( 5PVMS-S-(CH2)2-CH3 190( 2 162( 2PVMS-S-(CH2)11-CH3 210 265 ( 5 290-300 140( 5

aGlass transition. bMelting point. cFrom raw data at 1 Hz. dDetermined by extrapolating to τ=100 s. eFrom raw data at 100Hz to 10 kHz. fFromraw data at 100 Hz. gThe dominant feature in the DRS spectra was the peak at Tm. Tg was most clearly observed in aged samples.

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transition (Tg) values were determined as the onset points fromthe second scans if applicable. Themelting point (Tm) was takenas the maximum of the endothermic peak.

Dynamic Relaxation Spectroscopy (DRS). DRS utilized ahighly sensitive Andeen-Hagerling capacitance bridge operatingat 50Hz-20 kHz. Samples weremounted onto planar interdigita-ted electrodes, which enabled dielectric measurements on smallamounts ofmaterial. This technique has been utilized previouslyfor dielectric measurements on self-assembled monolayers9 andpolymers.10 We specifically chose the narrow-band, highly sensi-tive technique, rather than conventional broad-band dielectricspectroscopy, so that we could measure small volumes of mater-ial from the same samples used for wettability measurements.

DRS Analysis. Dielectric relaxation spectroscopy utilizes thesample as a dielectric within a capacitor and measures thecomplex dielectric constant (ε*=ε0 - iε0 0) under application ofanAC electric field11 at frequencyω. The output of the precisionbridge used in these experiments is capacitance (C) and tanδmeasured as a function of frequency and temperature.C=Aε0,where A is a geometrical factor related to the size and shape ofthe capacitor, whereas tan δ= ε0 0/ε0. Ionic backgrounds weremodeled as tan δ=(A/T) exp(-Ei/kT)/ω

s, whereA is a constantrelated to the DC conductivity, the exponential term reflects thetemperature dependence of ionic carriers with a characteristicbarrierEi, and s is dependent on the nature of ionic diffusion andcan vary from 0 to 1.12,13

The step in the capacitance or the more easily observed peakin tan δ at 1/τ=ω allows for determination of τ (the characteristicrelaxation time or average time between reorientation events) asa function of temperature for a particular applied field fre-quencyω.11,12 This temperature dependence can then be fitted tothe Arrhenius form, 1/τ=ωo exp(-ΔU/kT) to determineωo, theattempt or libration frequency and ΔU the barrier to reorienta-tion. For independent or noninteracting relaxations, where theArrhenius form is appropriate, this analysis yields physicallyinterpretable values (e.g., ω0 ≈ 1013-1015 rad/s); however, forinteracting objects where the barrier (ΔU) changes with tem-perature, particularly for glassy relaxations, fitting with theArrhenius form produces artificially high apparent values forωo andΔU. In this case, a more appropriate analysis utilizes theempirical Vogel-Fulcher-Tammann (VFT) form,14-16 1/τ=ωο exp[-B/(T - To)], where the argument of the exponent pro-vides a barrier that increases with decreasing temperature (B is aconstant.). To is a temperature below which all motion has ceasedand in our fits,ωo is held at 1014 rad/s to reduce the number of freeparameters. Glassy relaxations are quantified by defining Tg as thetemperatureatwhich τ=100s,17 determinedby fitting themeasuredτ versus T with the VFT form and extrapolating to large τ.

In this work, we plotted ln(1/τ) versus 1/T and fit with eitherthe Arrhenius or VFT forms as appropriate to determine eitherΔU, the barrier to rotation, or Tg for glassy relaxations. Ionicbackgrounds were removed from the data for PVMS-S-(CH2)2-CH3 and PVMS-S-(CH2)6-OH samples. For relaxa-tions within interacting side-chains, which are not a glassysystem per se, but which display an elevated apparent ωo value,we utilized a Starkweather analysis. Starkweather argued thatthere are significant entropy changes in cooperative relaxations,in contrast with ΔS=0 for a simple independent motion, andthat this value could be used to quantify the level of interactionin different systems.18 In this analysis, the Arrhenius form isused to fit the data and the observed activation energy ΔU andpeak position T0 are transformed into aΔS value as:ΔS=ΔU-k[1þ ln(k/2πh)þ ln(T0/f )],where T0 is the temperature at whichthe relaxation is f Hz. In our case, we utilized ΔU values fromArrhenius fits, and performed the Starkweather analysis at 1 kHz,in the middle of our experimental frequency range.

III. Results

A. Macroscopic Response. The macroscopic response ofthe substituted PVMS networks was characterized previously

by studying the cycling between hydrophobic (dry) andhydrophilic states (wet) states by dynamic contact angle(DCA). The degree of wettability change and the rate ofresponse, i.e., the variation of the water contact angle (θ)with time (τ), were shown qualitatively to decrease withincreasing the number of methylene units (n) in the mercap-toalkanol hydrocarbon spacer. Moreover, long linkers wereseen to cocrystallize locally thus freezing the conformationuntil it was relieved by heating the sample above the meltingpoint of the alkane crystallites.5,6 The open symbols inFigure 2 depict representative examples of such wettabilitychanges recorded for PVMS (squares), PVMS-S-(CH2)2-OH (circles), PVMS-S-(CH2)6-OH (up-triangles), andPVMS-S-(CH2)11-OH (down-triangles). These data re-veal that while PVMS exhibits a relatively modest, but fast,wettability change, PVMS-S-(CH2)2-OH and PVMS-S-(CH2)6-OH display similar reorganization kinetics,along with a larger change in wetting, upon transitioningfrom a hydrophobic to a hydrophilic environment. PVMS-S-(CH2)11-OH does not exhibit this large change in wett-ability during the cycling behavior even before the side-chains crystallize.

We extend these observations by quantitatively interpret-ing the DCA data.We first analyzed the DCA data by fittingeach cycle individually to the first order exponential decay:

- cosðθÞ ¼ AþB1 expð-β1tÞ ð1ÞIn eq 1, A, B1, and β1 were used as fitting parameters; β1represents an inverse of a characteristic relaxation time. Weaveraged the fitting exponents from runs 2-10 (excludingrun number 1 because of unknown sample history duringpreparation) for PVMS, PVMS-S-(CH2)2-OH, andPVMS-S-(CH2)6-OH, and cycles 2-4 for PVMS-S-(CH2)11-OH (after the fourth cycle the structure frozebecause of local cocrystallization among -(CH2)11-OH).The dotted lines in Figure 2 denote the values of-cos(θ) as afunction of time obtained by the best fits to eq 1. While thefits to PVMSandPVMS-S-(CH2)11-OHwere satisfactory(R2>0.98), the first order exponential decays did not captureaccurately the trends observed in the PVMS-S-(CH2)2-OH

Figure 2. Time dependence of the negative cosine of the contact angle(θ) measured by DCA (second cycle) for (a) PVMS (black squares),(b) PVMS-S-(CH2)2-OH (red circles), (c) PVMS-S-(CH2)6-OH(green up-triangles), and (d) PVMS-S-(CH2)11-OH (blue down-triangles). The dotted and solid lines represent the best fit to eqs 1and 2, respectively, as described in the text.

5046 Macromolecules, Vol. 43, No. 11, 2010 Crowe-Willoughby et al.

and PVMS-S-(CH2)6-OH specimens.We then reanalyzedthe DCA data by fitting to the second order exponentialdecay:

- cosðθÞ¼ AþB1 expð-β1τÞþB2 expð-β2τÞ ð2ÞThe solid lines in Figure 2 represent the best fits to eq 2 withA, B1, β1, B2, and β2 being used as fitting parameters. Thequality of the fits improved considerably for all data sets. InFigure 3we plot the values of the decay coefficients β1 (closedsymbols) and β2 (open symbols) for all substrates. From thedata in Figure 3, it is apparent that β1 decreases system-atically with increasing n, while β2 remains relatively con-stant for all specimens.Recalling thatβj is inversely related to

a relaxation time and assuming that the second order ex-ponential decay is a physically reasonable functional form,there are two superimposed processes that governwettabilitychanges in the specimens: a fast one (given by β1) and aslower one (given by β2).Wewill return to the issue of secondorder exponential decay later in the Discussion.

B. . Characterization of Molecular-Level Motion. DSCscans for PVMS, PVMS-S-(CH2)2-OH, PVMS-S-(CH2)6-OH, and PVMS-S-(CH2)11-OH are shown inFigure 4; the results are also summarized in Table 1. Thefirst three compounds exhibit a clear step-change in the heat

Figure 3. Coefficients in the 2nd order exponential decay function (β1:solid symbols, β2: open symbols) for PVMS (black squares),PVMS-S-(CH2)2-OH (red circles), PVMS-S-(CH2)6-OH (greenup-triangles), and PVMS-S-(CH2)11-OH (blue down-triangles) as afunction of the number of methylene units in PVMS-S-(CH2)n-OH.

Figure 4. DSC measurements for PVMS and PVMS-modified sub-strates. The heating rate was 3 �C/min, see text for details.

Figure 5. DMTA data for PVMS and PVMS-mercaptanol modified substrates depicting the storage (left ordinate, closed blue symbols) and loss(right ordinate, open red symbols) moduli, E0 and E0 0, respectively, for PVMS (a), PVMS-S-(CH2)2-OH (b), PVMS-S-(CH2)6-OH (c), andPVMS-S-(CH2)11-OH (d). Temperature ramps were performed at a heating rate of 3 �C/min and at 1 Hz. Amplitude was set within the linearviscoelastic regime. The inset graphs depict tan δ, i.e, tan(E0/E0 0).

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flow output at the glassy transition state plateau. Theresultant Tg values are given in Table 1. Data from PVMS-S-(CH2)11-OH reveal a melting transition that results in adistinct endothermic peak superimposed near the glass tran-sition feature. We have previously reported on crystalliza-tion of PVMS-S-(CH2)11-OH.5,6

DMTA data for PVMS, PVMS-S-(CH2)2-OH, PVMS-S-(CH2)6-OH, and PVMS-S-(CH2)11-OH are shown inFigure 5. The largest step change in elastic modulus E0,maxima in loss modulus E00, and maxima in tan δ (E0 =3G0,E00 =3G00, and tan δ=E00/E0) indicates the position of theR transition (TR), which is assigned as the glass transitiontemperature (Tg).

19,20 PVMS exhibits a clear transition at≈150 K associated with the Tg. In PVMS-S-(CH2)2-OHand PVMS-S-(CH2)6-OH, the most dramatic changecorresponds to the Tg; these values match well with thosefrom DSC (cf. Table 1). In addition, a smaller secondaryrelaxation appearing at a temperature below Tg is present.The locations of these features (labeled as β) are alsosummarized in Table 1. The presence of crystallinity com-plicates the spectrum of PVMS-S-(CH2)11-OH. In thiscase, the highest temperature feature is due to melting. Inaddition, there is a feature associated with the glass transi-tion as well as a secondary (β) relaxation below Tg. Inaddition to single frequency data, DMTA measurementswere taken at a range of frequencies for the PVMS andPVMS-S-(CH2)11-OH samples. The R and β relaxationsin PVMS-S-(CH2)2-CH3 were also studied byDMTA (cf.Table 1).

DRS spectra at 10 kHz for PVMS, PVMS-S-(CH2)2-OH, PVMS-S-(CH2)6-OH, and PVMS-S-(CH2)11-OHare plotted in Figure 6. The trends in the DRS data followthose in the DMTA results. For PVMS a single relaxation isobserved which corresponds to the R relaxation in DMTA;whereas both R and sub-Tg (β) relaxations are observed inthe PVMS-S-(CH2)n-OH samples. At the frequency uti-lized for the data depicted inFigure 6, the β andR relaxationsaremerged; however, they can be seenmore clearly in Figure 7,whichdisplays tanδ for PVMS-S-(CH2)2-OHat frequencies

varying from100Hz to 20 kHz. Figure 8 summarizes the sub-Tg (β) relaxations for PVMS-S-(CH2)2-OH, PVMS-S-(CH2)6-OH, and PVMS-S-(CH2)11-OH, and PVMS-S-(CH2)11-CH3. The R peak for PVMS-S-(CH2)2-OHand PVMS-S-(CH2)6-OH is shown for comparison.

For PVMS-S-(CH2)2-OH, the lower temperature (β)DRS peak is the smaller in amplitude, and appears at 180 to210 K for applied frequencies ranging from 100 Hz to 5 kHz.At higher frequencies, the smaller β peak shifts into the largeramplitude, higher temperature R peak, and is no longerclearly distinguishable. The β peak positions form a straightline on an Arrhenius plot (see Experimental Methods), with

Figure 6. Capacitance (left ordinate, blue solid symbols) and tan δ (right ordinate, red open symbols) fromDRSmeasurements at 10 kHz for PVMS(a), PVMS-S-(CH2)2-OH (b), PVMS-S-(CH2)6-OH (c), and PVMS-S-(CH2)11-OH (d).

Figure 7. Dissipation factor (tan δ) as a function of temperature forPVMS-S-(CH2)2-OH. The logarithmic scale enables tracking of thelow temperature relaxation. An arbitrary y-axis offset was applied toeach data set.

5048 Macromolecules, Vol. 43, No. 11, 2010 Crowe-Willoughby et al.

slightly elevated attempt frequency of ≈1016 rad/s. and abarrier of 10.9 ( 0.8 kcal/mol. For the 100 Hz to 20 kHzfrequency range, the higher amplitude and temperature Rpeak occurred at 240-250Kwith only a small peak shift as afunction of frequency. Such small peak shifts with frequencychange are consistent with “glassy” or strongly interactingrelaxations, which reflect the very rapid change in the rate ofthe molecular motion over very small changes in tempera-ture. Thus, we assign this relaxation as associated withsegmental motion and the glass transition.

TheRDRSpeak for the-(CH2)6-OHsamples appears atsimilar temperatures as the-(CH2)2-OH equivalent, rangingfrom 245 to 270 K, for 100 Hz to 20 kHz. As in PVMS-S-(CH2)2-OH, the lower temperature peak in PVMS-S-(CH2)6-OH DRS data was smaller in amplitude than theR-relaxation. At 100 Hz, this β peak occurred at the samelocation as the -(CH2)2-OH (180 K), before shiftinginto the higher temperature R feature at frequencies above500 Hz.

For PVMS-S-(CH2)11-OH, two peaks correspondingto the β and melting relaxations were observed in mostspecimens. The data in Figure 5d are representative. Thelower temperature (β) peak had the larger amplitude andoccurred at 200-230 K, for 100 Hz to 20 kHz. For agedsamples of-(CH2)11-OH, the melting peak was not presentand an R relaxation associated with the Tg was observed.

PVMS-S-(CH2)11-CH3 showed a high temperaturepeak in the dielectric relaxation data occurring at ≈290-300 K, which had characteristics associated with a meltingtransition (including a discontinuous change in capacitanceat Tm). It is not uncommon for side-chains with n=11 tocrystallize, even without a terminal -OH group. A relaxa-tion due to Tg was observed in the capacitance (enhanced bytaking a derivative with respect to temperature). Tg, Tm, andβ peak locations fromDRS for the C11 substituted networksappear in Table 1.

IV. Discussion

A. Comparison of Tg as a Function of Structural Changes.DMTA and DRS data for PVMS networks were consistentin displaying a sharp dispersing peak varying in location

from 145 to 155 K from 0.1 Hz to 20 kHz. Analysis of thedata (see Experimental Methods) resulted in an estimatedTg=140( 5K,which is consistentwith theTg value recordedby DSC. These values match the Tg of 143 K reported forlinear PVMS.21 Thus, the cross-linking employed to form thenetwork does not appear to alterTg and the PVMSbackboneretains its highly flexible properties. This observation alsoholds for PDMS (see Table 1; for comparison linear Tg forPDMS: 145-150 K, depending on crystalline fraction22).

From the PVMS-S-(CH2)2-OH DRS data, Tg=230 (10 K, a nearly 100 K increase from the PVMS. This Tg isconsistent with the DMTA value of 236 ( 3 K at 1 Hz. TheDSC Tg of 215 K also displays a similar increase over thePVMS network value (148 K). This indicates that the Tg hasbeen increased significantly due to the presence of the -S-(CH2)2-OH side-chains. Dramatic changes in PVMS glasstransition temperatures (≈50-200 K) have been observedpreviously when substituting bulky side-chains, which pre-sumably hinder motion of the flexible siloxane backbone.21

The increase in Tg indicates that there are sufficientsteric interactions between the pendent groups or betweenthe pendent groups and backbone to alter the backbonedynamics.

The RDRS peak for the-(CH2)6-OH samples appearedat similar temperatures as the -(CH2)2-OH equivalent, witha Tg of 210 ( 20 K. The value from DSC was 208 K. TheseTg values overlap with the corresponding results for-(CH2)2-OH. The DMTA data are also similar for thetwo materials with a primary peak centered at≈230-240 K (1Hz).While this lack of change inTg when transitioning to thelonger chain may be attributed to the well-known plasticiz-ing effect of alkyl chains of moderate length,23 it couldalternatively be caused by a slight variation of the -S-(CH2)n-OH loading in the PVMS-S-(CH2)n-OH net-work. Though, as we note in the Experimental Methods,the loading of -S-(CH2)n-OH inside the PVMS-S-(CH2)n-OH network appeared to be approximately thesame, as revealed by IR spectroscopy.

PVMS-S-(CH2)11-OH displayed Tm features in therange 299-311K (DSC) and 294-312K (DMTA). Evidenceof melting (often associated with a large discontinuouschange in the ε00 or, equivalently, capacitance) was observedin the dielectric data at 300-320 K. For most DRS samples,this was the dominant high temperature relaxation and nopeak associated with the glass transition was observed. Forsome samples, particularly those that had been aged atambient conditions for a year or so after fabrication, a peakranging from 282 to 287 K (320Hz to 20 kHz) was observed.This results in an estimated Tg in the 270-290 K range,which is consistent with that observed in the DSC (284.5 K).Features inDMTAassociatedwith theR relaxation varied inposition from 240 to 270 K. These Tg values are significantlyhigher than for the networks with shorter pendant groups,indicating that the backbone dynamics are further con-strained by the longer chains. Generally the DMTA andDS data from -(CH2)11-OH samples exhibited more com-plex structure and sample to sample differences than theother materials, which is consistent with a semicrystallinematerial.

In order to study the role of the terminal group inmodifying dynamics within the network, PVMS-S-(CH2)2-CH3 and PVMS-S-(CH2)11-CH3 were also in-vestigated. In a DMTA study of PVMS-S-(CH2)2-CH3

and PVMS-S-(CH2)2-OH, the Tg dropped significantlyfrom 236( 3K to 190( 3Kwithout the-OHgroup present(Table 1). This value is still larger than the Tg for the PVMSnetwork without pendent groups (≈150 K). The-(CH2)11-

Figure 8. Dissipation factor for (right ordinate) PVMS-S-(CH2)2-OH (filled red circles), PVMS-S-(CH2)6-OH (filled green up-triangles), and (left ordinate) PVMS-S-(CH2)11-OH (filled bluedown-triangles), PVMS-S-(CH2)11-CH3 (open blue down triangles)at 100 Hz. A background due to ionic conductivity has been removedfrom each sweep. This sweep of PVMS-S-(CH2)11-OHdata displaysamelting peak at≈315K. The higher temperature peak for PVMS-S-(CH2)2-OH and PVMS-S-(CH2)6-OH is associated with Tg.

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analogues were studied by DRS. When comparing the-(CH2)11-CH3pendentgroupwith themorepolar-(CH2)11-OH, the Tg only decreased slightly to 265 ( 5 K (comparedwith -(CH2)11-OH estimate of 270-290 K). Interestingly,the DSC measurement of Tg shows a much larger decreasefrom 285 to 210K. These results indicate that both the lengthand termination of the pendent group affect the glass transi-tion temperature; that is, that pendent-pendent or pendent-backbone interactions determine the effective Tg.

B. Sub-Tg Relaxations. The low-temperature (β) peak inthe PVMS-S-(CH2)2-OH DRS data is consistent withrotation within the pendent-group alkyl chains. Relaxationsdue to local motions within side-chains often follow Arrhe-nius dynamics. Side-chain peaks are generally smaller inamplitude than those associated with Tg (R relaxations)and appear at sub-Tg temperatures. There exists a well-known local relaxation in the 120-150 K range (for 100 Hz)that is associated with alkyl side groups24-27and has alsobeen observed in alkyl self-assembled monolayers.9,28 Pre-vious workers have assigned this relaxation to a crankshaft-likemotionwhere two gauche defects alternate in orientationwithout significantly changing the average volume taken upby the chain.29 The presence of the polar-OH group and/orpolarized bonds near -S- would enable dielectric observa-tion of these motions. Although the exact nature of therelaxation is still debated,26,27 it is mostly independent ofthe backbone structure25 and thus a common feature forpolymers with alkyl side-chains. A similar local relaxationoccurs in polyethylene (named the PE-γ relaxation in thatcase).28,29 The barriers observed for this PE-γ relaxation aregenerally ≈6-7 kcal/mol. However, in side-chain polymersthis relaxation has been observed at higher temperature(larger barriers) when interchain interactions are present.30

Our results are consistent with the presence of a PE-γ-likerelaxation, with an elevation in barrier (to 10.9 ( 0.8 kcal/mol) resulting from interchain interactions due to the -OHgroup. Sub-Tg or side-chain motions are not generally ob-servable in DSC, which is consistent with this assignment.

Comparing peak positions for the β relaxation fromDMTA and DRS (190 K at 1 Hz, and 180 K at 100 Hz) itappears that the barrier is lower in the dielectric case. Thishas previously been observed for other local relaxations.26,31

In that case, the authors argued that the coupling betweenthe electric field and the dipole was direct; whereas thecoupling between the applied mechanical stress and theside-chains was more indirect, and therefore associated witha higher barrier. Thus, the two techniquesmeasured differentresponses. In our materials only the terminal-OH group ora polarized bond near the thiolene attachment need respondto provide a dielectric signal, which may have slightlydifferent dynamics than a relaxation in the interior of thechain.

Comparing the response of-(CH2)2-OH and-(CH2)6-OH pendent groups, we find that the DMTA data is similarfor both, with peaks at≈190K. This similarity also holds forthe DRS data (180 K for 100 Hz). Considering the con-straints on the data due to peak merging, we conclude thatthe relaxation is similar for-S-(CH2)2-OHand-S-(CH2)6-OH pendent groups.

In contrast, the β dielectric relaxation in PVMS-S-(CH2)11-OH (DRS) was shifted to higher temperatures(200 K at 100 Hz). DMTA data showed an equivalentβ relaxation in the range of 180-197 K at 1 Hz. These peakpositions are higher in temperature than for the shorter side-chains, indicating slower dynamics. Analysis of the DRSdata revealed that the -(CH2)11-OH pendent groups havesignificant interaction, with an elevated apparent attempt

frequency of ≈1020 rad/s and an effective barrier of 16 (0.8 kcal/mol. A Starkweather analysis (see ExperimentalMethods) showed ΔS=115 ( 32 J/mol/K, compared toΔS=60( 10 J/mol/K for-(CH2)2-OH. Thus, we concludethat the -(CH2)11-OH chains are significantly more inter-acting than the-(CH2)2-OH or-(CH2)6-OH side-chains.This is consistent with previous observations of crystalliza-tion of the-(CH2)11-OH side-chains. Although this relaxa-tion has interacting dynamics, it did not appear in the DSCsweep, which is representative of an interacting side-chainrelaxation rather than a glass or melting transition.

The amplitude of the DRS β relaxation PVMS-S-(CH2)11-OH increased with increasing temperature (highermeasurement frequency) in contrast to the short side-chainnetwork data where the amplitude was constant. Such anincrease indicates that the number of units participating inthe relaxation is increasing with temperature.32,33 This couldoccur because the side-chain relaxation is localized in theamorphous regions, and the amorphous fraction is increas-ing with temperature. In addition, an asymmetrical potentialsurface could explain this result. In thermally activatedhopping, any reorienting group is hindered by a potentialsurface with more than one accessible minimum, i.e., re-orientation involves transition from oneminimum to anotherover an effective barrier ΔU. However, if the energy differ-ence between the minima, B0, is sufficiently large (comparedto kT), a particular unit will remain frozen in a single well,and thus not reorient. For a distribution of B0 values, as thetemperature increases, more units will be available forreorientation. Thus, in general, an increasing signal withtemperature is associated with asymmetrical potential sur-faces. Such asymmetry can occur when strong local interac-tions are present, which is consistent with the significantinteractions between the -(CH2)11-OH pendent groups.

The β transition was also studied for methyl-terminatedpendent groups. The relaxation in PVMS-S-(CH2)2-CH3

shifts downward from that in PVMS-S-(CH2)2-OH, in-dicating faster dynamics (DMTA measurements, Table 1).This occurs in the -(CH2)11- system as well, as studied byDRS. In that case, the peak shifted downward significantly(140-170 K, from 100 Hz to 20 kHz), and displayed aresultant barrier to rotation of 7.7 ( 0.5 kcal/mol (comparedto 16 kcal/mol for the C11-OH networks). Starkweatheranalysis revealed ΔS = 34 ( 12 J/mol/K, indicating asignificant drop in interaction compared with the -(CH2)11-OH case, and similar or lower interactions as within the side-chains of PVMS-S-(CH2)2-OH. Thus, the presence of the-OHgroup increases the effective barrier (in both theC2 andC11 cases), and in the absence of the-OH group, the barrierto rotation and peak location of the β-motion are similar toγ-PE relaxation in alkyl side-chain polymers. This confirmsour assignment of the relaxation.

C. Relationship between Tg and the Macroscopic Respon-siveness. Generally, the three techniques (DSC, DRS, andDMTA) agree in their determination of Tg within theexperimental error. Cross-linking to form the network doesnot affect the flexibility of the PVMS backbone; howeverupon grafting of the -S-(CH2)2-OH pendent group Tg

increases by >60 K. The Tg of PVMS-S-(CH2)11-OH issignificantly higher than that of the other substituted sys-tems, indicating that the strong interactions among the-S-(CH2)11-OH groups hinder substantially the motionof the siloxane backbone.

The trends inTg match qualitatively with the response ratedetected in the DCA experiments at room temperature.Namely, the response rate at room temperature ofPVMS-S-(CH2)n-OH increases with decreasing Tg. This

5050 Macromolecules, Vol. 43, No. 11, 2010 Crowe-Willoughby et al.

trend is not unexpected as materials farther above their glasstransition temperature should possess faster dynamics. Tothat end, PVMS exhibits the fastest response rate, followedby PVMS-S-(CH2)2-OH and PVMS-S-(CH2)6-OH andfinally, PVMS-S-(CH2)11-OH exhibits the most sluggishresponse.

Recall that earlier in the paper we pointed out that theDCA experimental data could not be fitted satisfactorilywith a simple exponential decay function. In contrast, asecond order exponential decay, i.e., eq 2, provided a betterfit. Interestingly, the decay coefficients β1 and β2 used in eq 2plotted in Figure 3 exhibit two distinct trends. While β1decreases steadily with increasing the length of themethylenespacer, β2 is independent of the size of the pendent group. Ifthe network Tg is well correlated with the response time, asreflected in β1, this would indicate that the segmental dy-namics control the motion of the pendent group through thenetwork. The time scale of segmental fluctuations is given byτ micro, the average time between reorientation events, thevalue of which (at room temperature) can be estimated fromthe DRS data. This τ micro should be related to the macro-scopic response time τmacro=1/β1 from the analysis of thedynamic contact angle (cf.Figures 2 and 3). Thus, in order tofurther correlate microscopic and macroscopic motion, wefit our DRS relaxation data (see Experimental Methods),and then extrapolated this fit to estimate τ micro at roomtemperature where the DCAmeasurements were taken. Thisis similar to the extrapolation utilized to determine Tg. Asalready discussed, these Tg values (listed in Table 1) agreewell with data from DSC and DMTA data taken at 1 Hz,which indicates that the fits are appropriate.

Since many segmental fluctuations would be required fora single -S-(CH2)n-OH group to emerge from the sur-face upon application of water, we expect the microscopicand macroscopic relaxation times to vary by many ordersof magnitude. The τ micro values for -S-(CH2)2-OHand -S-(CH2)6-OH were similar whereas the value for-S-(CH2)11-OH was about ten times larger. These ratiosare similar to those for the equivalent 1/β1 values. Compar-ing the four τ pairs (for the different network types) we findτ macro ≈1�109 τ micro. We conclude that there is a goodqualitative correlation between these two τ estimates, whichresult from completely isolated and independent measure-ments of the samples properties.

We could find no correlations between the slowermacroscopic relaxation time (given by the inverse of β2)and any other physical quantity measured. The slowerprocess is unlikely to originate from the relaxations ofthe side groups because those relaxations are likely fasterthan or have similar dynamics to the segmental motions atroom temperature. Moreover, the process responsible forβ2 should be independent of the molecular architecture ofthe silicone elastomer chain. Unlike other analytical meth-ods employed in this work, which probe the bulk propertiesof the system, DCA is sensitive only to the surface and nearsubsurface region of the network. We thus tentativelyassociate the slower relaxation process observed in theDCA data with the rearrangement of subsurface chainsthat move when the macromolecules present at the surfacereorganize. This process may thus be associated tentativelywith the overall mobility of the networks that woulddepend on the average molecular weight between thecross-link points; those were exactly the same in all speci-mens studied. In order to test this hypothesis, experi-ments similar to the one described here would have to becarried out using networks that possess different cross-linkdensities.

D. Relationship between Pendant Group Interactions andTg. Generally, longer pendent groups and stronger interac-tions between the pendent groups (as exemplified by the-(CH2)11-OH case) leads to an increase in Tg, as the stericand/or dipolar interactions between the chains limit theflexibility of the PVMS backbone. These interchain interac-tions also affect, and are reflected in, the β relaxation. Thus,although the β relaxation within the pendent groups is notdirectly related to the responsiveness of the networks, obser-ving its dynamics enhances our understanding of the pendentgroup interactions that lead to a resultant network Tg, andthus a characteristic response time to applied stimulus.

The R relaxation results indicate that both steric anddipolar interactions can effect interchain or chain-backboneinteractions and thus the stiffening of the PVMS network.For the short chain case, both effects appear distinctlyimportant, as exhibited in the difference in Tg values forthe -(CH2)2-OH and -(CH2)2-CH3 networks. For thelonger chain, the situation is less clear, in part due to thecomplications of measuring a semicrystalline material undera variety of conditions. It may be that steric interactionsbetween the side-chains are the dominant effect for longchains, and thus primarily responsible for the increase in Tg,explaining why the two-(CH2)11- pendent groups result ina similar Tg.

The polarity of the pendent group is important because itdetermines themagnitude of the response. As outlined above,the flexibility of the network, through which the pendentgroup must travel, affects the response time. For this parti-cular set of networks, these two effects are correlated.Because the PVMS backbone is highly flexible, interactionsbetween the pendent groups stiffen the network, raising Tg,and slowing the response time. Increasing the difference inhydophobicity between the two switching groups (the-CH3

and the vinyl or -OH pendent terminus), creates a largerdrive (energy difference) to respond to a change in environ-ment (application of water or air). However, rather thanenhancing the response time (due to the increased drivingforce) the increased polarity of the pendent groups causesinteractions which stiffens the network, thus (counterintui-tively) slowing down the response.

These results reveal the delicate balance needed in thedesign of such responsive networks. For instance, decreasingthe density of pendent groups would significantly decreaseinterchain interactions and thus allow for groups withstronger polarity without significant increases in Tg. How-ever, a lower density of pendent groups would also decreasethe magnitude of the response. An ideal design will utilize anoptimal density and polarity of pendent groups to maximizeboth response size and speed.

V. Conclusions

By combining dielectric, calorimetric, and mechanical mea-surements we observed both intrapendent and PVMS-backbonerelaxations within stimulus-responsive PVMS networks. Tg wasnot affected by network formation but increased by ≈40 K withintroduction of the -S-(CH2)2-CH3 pendent, and another25-40K for the polar analogue, i.e.,-S-(CH2)2-OH.Tg valuesfor PVMS-S-(CH2)2-OH and PVMS-S-(CH2)6-OH weresimilar, but the Tg increased significantly (≈50 K) for the longerpendent group, i.e., -(CH2)11-OH and -(CH2)11-CH3. Thus,as the pendent groups are altered, the network Tg changes, as thedifferent groups constrained the siloxane backbone to differingdegrees. This indicates that network dynamics can be tunedsignificantly (ΔTg= 130 K) by adjusting the size, shape, andpolarity of the pendent groups.

Article Macromolecules, Vol. 43, No. 11, 2010 5051

β relaxations were also observed for all pendent groups, andwere assigned to rotationwithin the pendent alkyl chain. Study ofthe β relaxation enabled an understanding of the changes inTg asresulting from steric or dipolar interchain or chain-backboneinteractions which stiffen the system.

For macroscopic responsiveness the pendent groups musttravel through the PVMS network, thus the flexibility of thenetwork appears to be one limitation to response time. Forsystemswith significant interactions between the pendent groups,Tg is elevated and the response time is slow. These two effectsmayarise from the same source; namely, the increased interactionsbetween pendent groups, which increase the rigidity of the entiresystem. When pendent groups are less interacting, some of theflexibility of the PVMS backbone is retained, and macroscopicresponse times approach that of the unsubstituted PVMS. Takenas a whole, these results indicate that a holistic approach,understanding both the driving forces and the intrinsic molecularmotions within a system, is desirable for effective design ofstimuli-responsive systems.

Acknowledgment. J.A.C.-W. and J.G. acknowledge the sup-port through the Office of Naval Research, Grant No. N00014-08-1-0369.We thank Professor SaadKhan and ProfessorRussellGorga for the generous use of their DSC andDMTA equipment.

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