thermoplastic elastomer gels. i. effects of composition and processing on morphology and gel...

Thermoplastic Elastomer Gels. I. Effects of Composition andProcessing on Morphology and Gel Behavior

JONATHAN H. LAURER,1* JAMES F. MULLING,1 SAAD A. KHAN,2 RICHARD J. SPONTAK,1,2 RUDY BUKOVNIK3

1 Department of Materials Science & Engineering, North Carolina State University, Raleigh, North Carolina 27695

2 Department of Chemical Engineering, North Carolina State University, Raleigh, North Carolina 27695

3 Telecom and Electrical Products Division, Raychem Corporation, Fuquay-Varina, North Carolina 27526

Received 9 January 1998; revised 16 March 1998; accepted 23 March 1998

ABSTRACT: Thermoplastic elastomer gels (TPEGs) composed of a poly[styrene-b-(eth-ylene-co-butylene)-b-styrene] triblock copolymer and a low-volatility, midblock-compat-ible mineral oil have been investigated at different oil concentrations to ascertain theeffect of composition on TPEG morphology and mechanical properties. The impact ofthermal processing is also examined by comparing gels thermally quenched to 0°C orslowly cooled to ambient temperature. Transmission electron micrographs reveal thatgels with 70 to 90 wt % oil exhibit styrenic micelles measuring ca. 24 nm in diameter.Variation in composition or cooling rate does not have any perceivable effect on micellesize or shape, whereas the rate at which the gels are cooled influences the extent ofmicrostructural order and the time to rupture (tR) at constant strain. Dynamic rheo-logical testing confirms the presence of a physically crosslinked network at TPEGcompositions ranging from 70 to 90 wt % oil, independent of cooling rate. Resultspresented here suggest that the dynamic elastic shear modulus (G9) scales as tR

a , wherea varies from 0.41 to 0.59, depending on cooling rate. © 1998 John Wiley & Sons, Inc. JPolym Sci B: Polym Phys 36: 2379–2391, 1998Keywords: block copolymer; thermoplastic elastomer; physical gel; polymer micelle

INTRODUCTION

As model systems exhibiting molecular self-orga-nization, block copolymers remain a principalsubject of extensive research activity. Comple-mentary theoretical1–6 and experimental7–10 ef-forts are continually directed at enriching the cur-rent understanding of block copolymer phase be-havior and dynamics under a wide variety ofconditions. The ultimate objective of such effortsis to produce microstructured polymeric materi-

als with the tailored properties needed to developnew,11,12 as well as improve existing,13 applica-tions. Addition of a solvent to a microphase-or-dered block copolymer affords a viable route bywhich to probe the order-disorder transition ofhighly segregated copolymers,14 but is seldomemployed15 in morphological studies of orderedblock copolymers. The presence of solvent affectssegmental interactions and may alter interfacialcurvature, thereby influencing morphological de-velopment, especially if the solvent is preferen-tially compatible with one block of the copolymer.A commercially important example of a copoly-mer blended with a block-preferential solvent oflow volatility is the thermoplastic elastomer gel(TPEG), which is ideally suited for comprehensivemorphology-property-processing analysis.

* Present address: Department of Materials Science & En-gineering, University of Pennsylvania, Philadelphia, PA19104

Correspondence to: R. J. SpontakJournal of Polymer Science: Part B: Polymer Physics, Vol. 36, 2379–2391 (1998)© 1998 John Wiley & Sons, Inc. CCC 0887-6266/98/132379-13

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A TPEG is a solvent-rich blend of a thermo-plastic elastomer multiblock copolymer and alow-volatility solvent that is compatible withthe rubbery block(s). While any such copolymercould be used in conjunction with a suitablesolvent to produce a TPEG, only the ABAtriblock architecture is considered here. LinearABA triblock copolymers are similar to their ABdiblock analogs in that they order into a varietyof periodic morphologies,16,17 including complexones such as the cubic bicontinuous gyroid,18,19

if their blocks are sufficiently incompatible. Un-like a diblock molecule, however, each triblockmolecule possesses a B-midblock that mustadopt a looped or bridged conformation, the lat-ter of which serves to connect neighboring A-microdomains. The fraction of bridged mid-blocks in neat triblock copolymers, estimated17

to be 0.40 – 0.45, is responsible for the shape-memory attribute inherent in these materials.Upon addition of a B-compatible homopolymer,ABA copolymer molecules order into the samemorphologies as AB copolymer molecules atcomparable blend compositions, but the fractionof bridged midblocks decreases, resulting in amarked reduction in the elasticity of the blendrelative to that of the parent ABA copolymer.20

Mixing an ABA copolymer with a low-volatil-ity, midblock-selective solvent yields a TPEG atrelatively high copolymer dilution. Under theseconditions, incompatibility between the end-blocks and the midblock-solvent mixture pro-motes micellization21,22 of the A blocks to min-imize interfacial area. As illustrated in Figure1, permissible midblock conformations includeloops (both endblocks of an ABA molecule residein the same micelle), dangling ends (one end-block remains within the matrix), and bridges(each endblock of an ABA molecule resides in adifferent micelle). Bridged midblocks are re-sponsible for the formation of a three-dimen-sional network that binds the solvent and cre-ates a physical gel, stabilized by micelles of theglassy or semicrystalline copolymer endblocks.The TPEG system of interest here utilizes apoly[styrene-b-(ethylene-co-butylene)-b-sty-rene] (SEBS) triblock copolymer and a low-volatility EB-compatible mineral oil.

Previous studies23–26 of SEBS copolymer/oilblends have shown that the oil is completelybound within the midblock network at oil massfractions (woil) of up to 0.90. Polizzi et al.27 alsofind this to be true for solvated poly(styrene-b-butadiene-b-styrene) (SBS) triblock copoly-

mers. Prior morphological studies28 –30 of SEBS/oil TPEGs have relied almost exclusively onsmall-angle scattering for microstructural anal-ysis. Mischenko et al.28 have examined TPEGssimilar to the one studied here and report theformation of S-rich micelles in dilute systems(woil $ 0.88). Mechanical deformation is foundto affect both the midblock network and localordering of S domains, and is reversible up to ashigh as 150% strain.29 The mode of deformationappears to be concentration-dependent: affinein copolymer-rich systems, but nonaffine whenwoil $ 0.88.29,31 It is therefore expected that themechanical response of TPEGs reflects not onlythe (physically) crosslinked micelles, but alsothe population of midblock bridges and entan-glements comprising the gel network. Directimages of these TPEGs from transmission elec-tron microscopy (TEM) reveal32 that the sizes ofthe S-rich micelles in mechanically mixed andsolvent-cast gels compare favorably with thosereported by Mischenko et al.,29 suggesting thatthe gross morphological features of TPEGs, butnot necessarily their properties, may be rela-tively insensitive to processing. To ascertainhere the relationship between microstructure,mechanical properties, and processing, we ex-

Figure 1. Schematic illustration of the micellar net-work formed by the SEBS/oil gels examined in thisstudy. Self-association of the styrene endblocks due tothermodynamic incompatibility with the oil-rich ma-trix results in the formation of spheroidal micelles. Thecopolymer midblocks adopt either bridged (a) or looped(b) conformations. A fraction of the styrene endblocksmay also reside in the matrix and remain unassociatedwith a micelle, thereby forming a dangling end (c).

2380 LAURER ET AL.

amine the dependence of TPEG morphology andmechanical behavior on blend composition.

EXPERIMENTAL

Materials

The SEBS copolymer employed here was KratonG1654, produced by the Shell Development Co.(Houston, TX). Its number-average molecularweight (M# n) and polydispersity index were deter-mined from GPC to be 161,000 g/mol and 1.09,respectively, and its composition was reported33

by the manufacturer to be 31 wt % styrene. TheEB-compatible extender oil was a saturated ali-phatic white mineral oil (Hydrobrite 380 PO) sup-plied by the Witco Corp. (Greenwich, CT). Accord-ing to the ASTM D2502 testing protocol, the mo-lecular weight of this oil was 468 g/mol.

Methods

Bulk films of the neat copolymer were preparedby casting a concentrated (5% w/v) toluene solu-tion in a Teflon mold and removing the solvent atambient temperature. To ascertain the effect ofsolvent evaporation on copolymer morphology,the solvent was removed both quickly (3 days)and slowly (3 weeks). A series of TPEGs with woilranging from 0.70 to 0.90 (9.3–3.1 wt % styrene)was prepared by mixing predetermined amountsof copolymer and oil for 1 h in a Ross LDM 1-qtdouble-planetary mixer under vacuum at 180°C.Portions of the mixed gels were pressed at 180°Cinto films measuring approximately 3-mm thick.The films were then subjected to two differentthermal treatments. They were either rapidlyquenched to 0°C in an ice/water slurry, or cooledslowly over the course of about 4 h. During slowcooling, samples remained above 170°C for about10 min and above 125°C for about 40 min. Oil lossat these temperatures was deemed negligible,since the mass flux of oil evaporating from a 4 mmdiameter aluminum pan was measured to beabout 9.5 3 10210 g/mm2-s at 180°C. No loss wasmeasured at temperatures below 125°C. Similartests were performed on TPEGs varying in com-position and revealed a very slight reduction inthe oil flux. With the flux of pure oil constitutingan upper limit, changes in TPEG composition dueto oil loss during sample processing were esti-mated to be less than 1 wt %.

Specimens suitable for TEM were preparedfrom the neat copolymer, as well as from theas-mixed, quenched and slow-cooled TPEGs, bycryosectioning in a Reichert-Jung Ultracut-S cryo-ultramicrotome maintained at 2110°C. Thin sec-tions, 70–100 nm thick, were subjected to thevapor of a 0.5% RuO4(aq) solution for 5 min tostain any existing styrene-rich microstructure.Electron micrographs were obtained on a ZeissEM902 electron spectroscopic microscope oper-ated at an accelerating voltage of 80 kV and anenergy loss (DE) range of 0–25 eV. Dynamic rheo-logical measurements were conducted on a Rheo-metrics Mechanical Spectrometer (RMS800) with25 mm parallel plates and a 3 mm gap. Dynamicstrain sweeps were performed at ambient temper-ature from 0.5 to 150% strain amplitude (go) at anoscillatory frequency (v) of 10 rad/s on all of theTPEGs prepared here to determine the linear vis-coelastic regime of each material. The dynamicstorage (elastic) and loss (viscous) moduli (G9 andG0, respectively) were deduced from the measuredshear stress (t) according to the relationship t5 goG9 sin vt 1 goG0 cos vt, where t denotes time.Moduli were measured at v range 1021–102 rad/sat a go of 1%. Time-dependent experiments con-ducted immediately after sample loading showedno change in moduli, indicating the absence ofhandling or preshear effects. Step-stress relax-ation following a step strain of 3% was also per-formed on all samples as a function of compositionand thermal processing. On the basis of theseexperiments, the relaxation modulus, G(t), wasobtained from t(t,g)/g, where g is the appliedstrain (3%) and t is the measured stress. Time-to-rupture (tR) tests were performed on die-cut gelrings measuring 28.9 mm in outer diameter, 22.5mm in inner diameter, and approximately 3 mmthick. Ten rings of each TPEG composition exam-ined were stretched to a constant elongation of500% between stationary posts and were main-tained at 24 6 3°C until failure.

RESULTS AND DISCUSSION

Morphological Characteristics

It is helpful to examine first the morphology of theneat SEBS copolymer in the absence of extenderoil before discussing the TPEG morphologies.Since the copolymer is compositionally asymmet-ric (31 wt % styrene), the equilibrium morphologyis expected to be cylindrical or lamellar, if the

THERMOPLASTIC ELASTOMER GELS. I 2381

copolymer is microphase-ordered. The ability ofthis block copolymer to microphase-order dependson the magnitude of the thermodynamic incom-patibility (xN) between the S and EB blocks andthe volume-fraction composition of the copoly-mer.7,8,16,17 Here, x denotes the Flory–Hugginsinteraction parameter, and N is the number ofstatistical units along the copolymer backbone.On the basis of the mass densities of polysty-rene34 (1.01 g/cm3) and the SEBS triblock copol-ymer23 (0.91 g/cm3) at the upper glass transitiontemperature of polystyrene (Tg 5 100°C), we es-timate the volume fraction of styrene in the co-polymer to be about 0.28. Since a similar SEBScopolymer (Kraton G1652) of lower molecularweight (M# n 5 87000) and styrene content (29 wt%) is reported35 to be microphase-ordered (exhib-iting hexagonally packed styrenic cylinders in anEB-rich matrix) at the Tg of the copolymer, itfollows that the copolymer examined here mustlikewise be ordered, due to its higher xN. Assess-ment of the copolymer morphology is describedbelow.

Microscopical examination of solvent-cast filmsof the neat SEBS copolymer reveals that the mor-phology depends strongly on the rate of solventevaporation. When the solvent is removed slowlyover the course of three weeks, the copolymer mor-phology consists of numerous lamellar grains, asseen in Figure 2. The presence of numerous grain

boundaries evident in Figure 2 is in marked con-trast to the highly oriented lamellae that are oftenobtained in poly(styrene-b-isoprene) (SI) diblock co-polymers and poly(styrene-b-isoprene-b-styrene)(SIS) triblock copolymers after comparable film-drying conditions. Due to phenyl-specific RuO4staining,36 the S-rich microdomains appear elec-tron-opaque (dark) in Figure 2 and in all other mi-crographs presented in this study. When subjectedto more rapid solvent evaporation (3 days, as op-posed to 3 weeks), a well-defined morphology failedto develop, indicating that the copolymer chainsbecome kinetically trapped during rapid evapora-tion, thereby preventing equilibration. This appar-ent dependence of morphology on the rate of solventevaporation is commonly attributed to morphologymetastability.37 Since kinetic limitations are ex-pected to be more prevalent in the film subjected torelatively rapid solvent removal (over 3 days), thelamellar grains observed in Figure 2 are presumedto be representative of the near-equilibrium copol-ymer morphology in the limit of zero oil content.

Upon addition of mineral oil (woil $ 0.70), theSEBS triblock copolymer molecules self-assembleinto S-rich micelles either at the mixing temper-ature (180°C) or upon cooling. A recent small-angle neutron scattering (SANS) study30 of simi-lar TPEGs has identified a high-temperaturemetastability regime in which the local order ofTPEG micelles increases dramatically, resultingin a clearly discernible transition from randompositioning to bcc packing. Such pronounced long-range order is apparently lost at temperaturesabove or below this regime. Since the effect oftemperature on TPEG behavior is the subject ofthe second part of this series, it is not discussedfurther here. Direct verification of the micellarmorphology in SEBS-based TPEGs through TEMimaging has been provided earlier32 for severalsolvent-cast and mechanically mixed gels. Thediameters of the micelles in those TPEGs mea-sure between 21 to 28 nm, depending on copoly-mer molecular weight, blend composition and pro-cess history, and compare favorably with small-angle X-ray scattering (SAXS) measurements29

obtained from similar systems.TEM images (Figs. 3–5) have been acquired

from TPEGs varying in oil fraction (woil) 2 0.70(Fig. 3), 0.80 (Fig. 4), and 0.90 (Fig. 5)—and sub-jected to the two thermal treatments describedearlier. Note that the TPEGs presented in thesefigures differ from those reported32 previously inthat they have been mechanically mixed and com-pression-molded at 180°C. It is apparent from

Figure 2. Transmission electron micrograph of theneat SEBS copolymer used in this study. The sampleseen here has been solvent-cast from toluene and al-lowed to dry slowly for a period of 3 weeks. The styrenemicrodomains appear electron-opaque (dark) due topreferential RuO4 staining.

2382 LAURER ET AL.

these two series of micrographs that S-rich mi-celles form in each TPEG, irrespective of the ther-mal treatments employed here. The sizes of themicelles, as discerned from these images, arelisted in Table I, with the average micelle diam-eter being about 24 nm. Comparison of the micro-graphs in Figures 3–5 with those presented else-where32 reveals that the spherical shape and sizeof the styrene micelles in these TPEGs are notstrongly dependent on either woil or processinghistory. This observation implies that the corre-sponding micelle aggregation number (n), ex-pressed in terms of the number of styrene-end-blocks within each micelle, is not significantlyaltered by mixing/pressing at 180°C or solutioncasting from a relatively neutral solvent.

If we assume that no oil resides within the mi-celles and if the micellar radius does not change

appreciably upon exposure to the electron beam,then n can be estimated from 4pr3NArS/3MS, wherer denotes the mean micellar radius, NA is Avo-gadro’s number, rS is the mass density of polysty-rene, and MS is the mass of the copolymer endblock.From the values listed in Table I, we find that n isabout 180 6 50. While the first assumption madeabove is difficult to confirm, the second seems rea-sonable based on the quantitative agreement be-tween TEM and SAXS29 data. It must be recognizedthat this estimate of n should be interpreted as aminimum number of endblocks per micelle, sincethe relationship provided above implicitly assumesthat the endblocks adopt a gaussian conformation.Due to the thermodynamic repulsion responsible formicellization, the endblocks are expected to bestretched along the interfacial normal upon aggre-gation. In addition, the aggregation number must

Figure 4. TEM micrographs of the TPEG with woil

5 0.80 quenched from 180 to 0°C (a) and slow-cooledfrom 180°C to ambient temperature (b).




be placed on an endblock, and not molecular, basisdue to lack of information regarding the populationsof looped, bridged, and dangling midblocks. Whilethe micellar radius and, hence, the number of S-endblocks comprising each micelle remain rela-tively invariant with increasing woil, the fraction ofbridged midblocks connecting adjacent micelles, forexample, is expected to decrease since the spacingbetween micelles increases. This increase in spac-ing is most easily seen by comparing the TPEGmicrographs at woil 5 0.70 (Fig. 3) to those at woil 50.90 (Fig. 5). Thus, the S-endblocks within eachmicelle cannot be assigned to copolymer moleculeswithout knowledge of the looped, bridged, and dan-gling midblock populations.

Application of the theory proposed by Kao andOlvera de la Cruz21 for block copolymer micelleformation to the micelle radii measured here can

provide an estimate of the magnitude of x betweenthe styrene endblocks comprising the micelles andthe EB 1 oil matrix. Although this equilibrium the-ory is developed for AB/hA copolymer/homopolymerblends, it should be equally applicable in thepresent study since the copolymer midblock is dis-solved in a low-molar-mass oil, as opposed to ahighly entangled homopolymer. In fact, the oil-in-duced increase in chain diffusion during micelliza-tion is expected to promote equilibration more rap-idly in a TPEG than in a comparable copolymer/homopolymer blend. According to this theoreticaltreatment, the standard free energy of block copol-ymer micellization (F) can be written in units of kT(where k is the Boltzmann constant and T is abso-lute temperature) as

F 5 Fint 1 Fdef 1 Fx (1)

Here, the subscripts denote formation of the core-corona interface (int), deformation of the micelle-forming block (def), and contact loss between co-polymer blocks (x). Each contribution is given by

Fint 54pr2

a2 Sx

6D1/2

(2a)

Fdef 5 nS 3r2

2NSa2 1p2NSa2

6r2 2 3.1449D (2b)

Fx 5 2 nxNS (2c)

Table I. Dimensions of Styrene Micelles in theTPEGs Examined in the Present Studya

woil

Thermalroute

Meandiameter

(nm)

Prior to annealing0.70 Quenched 25.00.70 Slow-cooled 24.80.80 Quenched 22.50.80 Slow-cooled 25.00.90 Quenched 21.30.90 Slow-cooled 21.5After annealing for 16 h at 180°C0.80 As-mixed 28.00.80 Quenched 25.00.80 Slow-cooled 25.0

Average 24.2 (62.1)

a Measured from TEM micrographs through the use of NIHImage software.



2384 LAURER ET AL.

where r is the micellar core radius, a is the sty-rene segment length (0.68 nm), and NS is thenumber of segments in the micelle-forming S-block (' 240). Substitution of eq. (2) into eq. (1)and minimization of F with respect to r at con-stant n and x, followed by algebraic rearrange-ment, yields

x 5 6n2Fp2NS2~a/r!4 2 924pNS

G 2

(3)

Evaluation of x between styrene and the EB 1 oilmatrix on the basis of the molecular propertiesprovided above and the micelle radii listed inTable I yields a range of values lying on averagebetween 0.01 and 0.02 at the temperature of mi-celle vitrification (i.e., at the styrene Tg). To putthis value into context, the temperature depen-dence of x between styrene and isoprene seg-ments in SI diblock copolymers can be ex-pressed14 as 33.0/T 2 0.0228, in which case x ispredicted to lie between 0.088 at 25°C and 0.066at 100°C. These x values representing S/I inter-actions are higher than that calculated from eq.(3) for S/(EB 1 oil) interactions. Due to enhancedchemical dissimilarity between the S and EBblocks, however, xS/EB is expected to be signifi-cantly higher (ranging from 0.37 to 0.61, depend-ing on the E/B ratio38) than xS/I. This dilemmacan be resolved if some of the mineral oil residesin, plasticizes, and slightly swells the S-rich mi-celles.32 Intramicellar oil entrapment may there-fore constitute another potential source of uncer-tainty in the above calculations, and should beconsidered in the design, analysis, and commer-cial manufacture of TPEGs.

A comment regarding the images in Figures3–5 is warranted at this juncture. Images such asthese are two-dimensional projections of three-dimensional features. Since the sections nomi-nally measure between 70 and 100 nm thick,these images provide information from severallayers of micelles, since the average micellar di-ameter is only 24 nm. Thus, the images mayappear to possess a higher concentration of mi-celles than the bulk compositions would imply.Moreover, while precautions are exercised to min-imize specimen damage due to electron beam ir-radiation and the extent of such damage is ex-pected to be tolerable within the context of thisstudy, it must always be remembered that sectionshrinkage and thinning, which may induce signif-icant spatial distortion, readily occur in irradi-

ated organic specimens due to specimen heatingand mass loss.36

While morphological differences betweenquenched and slow-cooled TPEGs after compres-sion-molding at 180°C appear minimal, a subtledifference between the two microstructures is ev-ident. The quenched TPEGs displayed in Figures3a, 4a, and 5a are rapidly cooled to a temperaturebelow the styrene Tg, thereby preserving the size,shape, and distribution of the S-rich micelles im-mediately following the pressing process. Sincethe quench conditions employed here are not ex-pected to instantaneously immobilize the mole-cules comprising the TPEGs (as would a quenchinto, for example, a liquid cryogen39), it remainsquestionable as to whether the micelles exhibitevidence of long-range order at 180°C or undergomicrostructural refinement during quenching.Due to the extended period spent above the sty-rene Tg (ca. 90 min), the slowly cooled TPEGsshown in Figures 3b, 4b, and 5b consist of reason-ably well-ordered S-rich micelles, many of whichappear to sit on a cubic lattice. Such long-rangeorder is similar to that obtained from high-tem-perature annealing of high-viscosity copolymermelts, and is more clearly evident in the low-magnification images presented in Figure 6. Re-gions of localized order are indicated in micro-graphs of both the quenched (Fig. 6a) and slow-cooled (Fig. 6b) TPEGs with woil 5 0.70. Suchregions are more abundant in the slow-cooledsample, although neither material possesses sig-nificant long-range order. As with other block co-polymer systems,37 time above the styrene Tg ap-pears to be a crucial consideration in refining theextent of long-range microstructural order inthese materials.

To verify this apparent trend, TPEGs from theseries with woil 5 0.80 have been annealed at180°C for 16 h under argon. As mentioned earlier,a small fraction of oil loss occurs due to evapora-tion at this temperature. On the basis of the mea-sured flux of oil from these samples at 180°C, theuncertainty in the composition of these specimensis ' 2.5 wt % after annealing. As anticipated fromthe micrograph pairs presented in Figures 3–6,annealing at this temperature clearly promotesalignment of the S-rich micelles after 16 h. [Whilethis time period would not constitute a lengthyanneal for a copolymer melt, annealing is consid-erably more efficient in TPEGs due to their lowerviscosity and higher polymer diffusion rate.] Elec-tron micrographs of annealed samples with woil' 0.80 appear qualitatively identical, regardless


of whether the preannealed specimens weremixed only; mixed, pressed, and quenched; ormixed, pressed, and slow-cooled. Shown for illus-tration in Figure 7a is a quenched TPEG afterannealing. Such similarity in the extent of micro-structural order suggests that processing historycan be effectively erased through high-tempera-ture annealing. A corresponding two-dimensionalFourier transform40 of the image in Figure 7a (aswell as of other images of woil ' 0.80 samplessubjected to comparable annealing) is provided inFigure 7b, from which the interdomain spacing(i.e., the distance between micelle centers) can bederived. Measuring from the center of the patternto the intensity maxima in Figure 7b yields aspacing of ca. 40 nm, independent of the thermal

history prior to annealing. Although this mea-surement lies within the range of periodicitiesfrom SAXS for similar TPEG systems,29 onlysemiquantitative assessments can be made heredue to the image-related limitations discussedearlier. An illustration of the Fourier transformin Figure 7b is presented in Figure 7c and, cou-pled with existing data,41 indicates that the mi-celles comprising this TPEG most likely exhibittwinned bcc symmetry.

Comparison of Figures 6 and 7 reveals thathigh-temperature annealing improves long-rangemicellar order in these TPEGs relative to even theslow-cooled TPEGs, which remain above the sty-rene Tg for an extended period of time. If thelong-range order evident in annealed TPEGs (Fig.7) is representative of near-equilibrium, then itfollows that the morphologies observed in theslowly-cooled specimens (Figs. 3b–5b) are closerto thermodynamic equilibrium than are those inthe quenched TPEGs (Figs. 3a–5a), even thoughthe sizes of the micelles in these systems arevirtually identical. Therefore, it is reasonable to

Figure 7. TEM micrograph (top) of a quenched TPEG(woil 5 0.80) after a 16-h anneal at 180°C under argon.The styrene micelles exhibit relatively long-range or-der, demonstrated by the maxima (bright spots) in theFourier transform of the image. A schematic showingthe arrangement of the Fourier transform maxima re-veals that the microstructure exhibits twinned bccsymmetry. The scale bar in the transform is equal to0.025 nm21.

Figure 6. Low-magnification TEM micrographs of (a)quenched and (b) slow-cooled TPEGs with woil 5 0.70,illustrating the differences in local order due to thermalhistory. Note that the slow-cooled material exhibitsnumerous locally ordered regions (arrows).

2386 LAURER ET AL.

conclude from these data that the coordination,rather than size, of the micelles provides a mea-sure of microstructural refinement, in which caserefinement necessarily involves the conforma-tions of the copolymer midblocks residing in theoil-rich matrix. As mentioned earlier and depictedin Figure 1, the three conformation-based mid-block populations (bridged, looped, and dangling)govern the extent of micelle connectivity, in whichcase they must be at least partially responsiblefor any measured differences in the mechanicalproperties of TPEGs within the range 0.70 # woil# 0.90. In the next section, results obtained frommechanical tests capable of probing the elasticnetwork of TPEGs are presented and discussed.

Rheological Behavior

Figure 8 shows the frequency spectra of the dy-namic storage modulus (G9) for TPEGs producedby thermal quenching and slow cooling. (Mea-sured values of G0 from only one sample are in-cluded in this figure to maintain clarity.) Twonoteworthy features of Figure 8 support the iden-tification of these triblock copolymer/oil blends asthermoplastic elastomer gels42–44: (1) G9 consis-tently exceeds G0 over the entire v spectrum forall compositions, and (2) G9 is virtually indepen-

dent of v over three orders of magnitude in v. Wealso find that tan d (AG0/G9) from these frequencyspectra scale remarkably well with v« (data notshown), where « depends on gel composition. Ac-cording to Figure 8, G9 decreases monotonicallywith increasing woil. This characteristic, as dis-cussed in detail later, reflects an increasing frac-tion of mineral oil (a viscous liquid), as well as acorresponding reduction in the elastic copolymernetwork. Since the distance between micelles in-creases with increasing woil,

31 bridged midblocksbecome more highly stretched. The energy pen-alty associated with the formation of a midblockbridge45 therefore increases as woil increases, inwhich case the populations of looped and danglingmidblocks both increase.

The effect of strain amplitude on the dynamicmoduli of representative samples is shown in Fig-ure 9a. Several features are evident from thisfigure, which displays G9 as a function of g andwoil for quenched and slow cool TPEGs. (As in Fig.8, only a single G0 data set is included for the sakeof clarity.) The first is that G9 consistently exceedsG0 by at least an order of magnitude over 2-1

2decades in g, adding support to Figure 8 thatthese materials behave in an elastic (solid-like)fashion at typical application conditions. More-over, G9 is again seen to decrease monotonicallywith increasing woil. This behavior is consistentwith the tactile observation that the TPEGs be-come more pliable as the fraction of oil increases.Lastly, the linear viscoelastic (LVE) regime ex-tends to higher strain levels upon increasing woil,yet exhibits no distinct dependence on thermalhistory. Values of g identifying the departurefrom LVE behavior (gLVE), as discerned from theabrupt increase in tan d(g) (data not shown), areprovided for several values of woil in Figure 9b,which reveals that gLVE depends moderately onTPEG composition and thermal history. Since thequenched samples may possess a higher bridgingfraction than the slow-cooled TPEGs due to ki-netic entrapment, the results in Figure 9 suggestthat (1) any such processing-related difference isminimal, and (2) midblock bridging occurs duringhigh-temperature mixing and not upon cooling.

The stress relaxation behavior of the TPEGsinvestigated here (Fig. 10) provides another indi-cation of the presence of an interconnected mid-block network. In this figure, we observe that anon-zero shear modulus (G) plateau is reachedafter 120 s. The magnitude of the plateau G foreach composition agrees reasonably well with theG9 values shown in Figures 8 and 9 (see Table II

Figure 8. Dynamic storage modulus (G9) presentedas a function of oscillatory frequency (v) at a strainamplitude of 1% for quenched (filled symbols) and slow-cooled (open symbols) TPEGs differing in oil fraction(woil): 0.70 (circles), 0.75 (triangles), 0.85 (inverted tri-angles), and 0.90 (diamonds). Illustrative dynamic lossmodulus (G0) data are included for only one composi-tion (labeled) to maintain figure clarity.


for a comparison). As is evident from Figures8–10, the change in rheological properties in-duced by varying the oil fraction (woil) is consid-erable. To illustrate the significance of thischange, the stress-relaxation modulus (G) and thedynamic elastic modulus (G9, from both dynamicfrequency and strain analyses) of each TPEG arepresented as functions of copolymer concentration(C, expressed in g/cm3) in Figure 11, and areclearly seen to increase with increasing C. In

terms of either the dynamic modulus or the relax-ation modulus, this relationship can be accuratelydescribed by G9 ; Cb, where b is equal to 2.43 andis independent of cooling rate within experimen-tal uncertainty. Since b . 1, these moduli dataare not indicative of the Rouse semidilute re-gime,46 in which G9 (or G) scales linearly with Caccording to CRT/M# n (which is shown, for com-parison, as the line with a slope of RT/M# n ' 1.543 105 dyn-cm/g passing through the datum pointat the lowest copolymer concentration). Thus, itcan be concluded from the data presented in Fig-ure 11 that elastic contributions from entangledmidblock loops and dangling ends, as well as from

Figure 9. In (a), the dynamic storage modulus (G9) isdisplayed as a function of strain amplitude (go) at v5 10 rad/s for quenched (filled symbols) and slow-cooled (open symbols) TPEGs differing in oil fraction(woil): 0.70 (circles), 0.75 (triangles), 0.85 (inverted tri-angles), and 0.90 (diamonds). To avoid confusion, lossmoduli (G0) are only included for one composition (la-beled). In (b), the strain identifying departure from thelinear viscoelastic regime (gLVE) is shown as a functionof woil for slow-cooled (‚) and quenched (F) gels.

Figure 10. Stress relaxation results displaying therelaxation modulus (G) as a function of time for slow-cooled TPEGs after application of a 3% step strain.Each material reaches a non-zero plateau value within120 s, indicating the presence of a midblock gel networkat five different oil fractions (woil): 0.70 (E), 0.75 (‚),0.80 (h), 0.85 (ƒ), and 0.90 ({). Stress relaxation datacollected from the corresponding quenched TPEGs arenearly identical to the results shown here and are notincluded for that reason.

Table II. Average Shear Moduli Measured byRheological Methods in the Present Studya

woil G9 (dyn/cm2) G (dyn/cm2)

0.70 7.73 3 105 (612%) 7.73 3 105

0.75 4.15 3 105 (64%) 4.56 3 105

0.80 2.10 3 105 (615%) 1.53 3 105

0.85 1.10 3 105 (612%) 9.36 3 104

0.90 5.21 3 104 (613%) 4.74 3 104

a Values in parentheses represent one standard deviation.

2388 LAURER ET AL.

midblock bridges, contribute to the measuredmoduli.

Despite expected differences in the populationsof copolymer midblocks (due to kinetic entrap-ment), there appears to be no significant changein G9 for the TPEGs subjected to the two differentcooling rates examined here. This observation isconsistent with the earlier conclusions that (1)bridging/looping and micellization occur primar-ily during mixing at elevated temperatures and(2) the cooling rate has little impact on either themidblock bridging fraction or micelle size andshape. Upon decreasing the block copolymer massfraction in these TPEGs, the degree to which thegel network is connected must diminish due to anoverall reduction in the number of midblocksavailable for bridging, which, explains at leastpartially, the observed monotonic increase in G9with increasing C. Another measure of the gelnetwork is tan d, which, as alluded to earlier, isfound to scale with v«. Shown in the inset of

Figure 11 is the composition dependence of « eval-uated from the frequency spectra displayed inFigure 8. An interesting feature of these data isthat « increases almost linearly with increasingwoil for gels slow-cooled from 180°C. The depen-dence of « on woil clearly becomes nonlinear forgels quenched from 180°C.

The values of G9 reported in Figure 11 areacquired by dynamic rheology within the LVEregime to ensure that the microstructure is rela-tively unaffected by the measurement. In markedcontrast, the time to rupture (tR) constitutes anultimate (failure) property well outside the LVEregime. As seen in Figure 12, tR is observed todecrease exponentially with woil for bothquenched and slow-cooled TPEGs. Unlike themeasurements conducted within the LVE regime,the values of tR in this figure clearly demonstratethat the cooling rate has an increasingly signifi-cant effect on TPEG failure as woil increases, withquenched TPEGs exhibiting longer tR than slow-cooled materials. This observation is consistentwith the intuitive expectation that quenching re-sults in a kinetically entrapped microstructure.Since both G9 from Figure 11 and tR from Figure12 are highly sensitive to woil, G9 is shown as afunction of tR in the inset of Figure 12. This insetreveals that G9 scales as tR

a , where a is found toequal 0.41 for quenched specimens and 0.59 for

Figure 12. Variation of the time to rupture (tR) withwoil for slow-cooled (‚) and quenched (F) TPEGs. Thesolid lines reveal that tR decreases exponentially withwoil (error bars based on one standard deviation of thedata appear smaller than the symbols). Shown in theinset is G9 as a function of tR, indicating that G9 scalesas tR

a (solid lines) over the range in woil explored.

Figure 11. Mean values of the dynamic storage mod-ulus G9 (triangles, from strain sweeps; circles, fromfrequency sweeps) and stress-relaxation modulus G (di-amonds) as a function of copolymer concentration (C)for slow-cooled and quenched TPEGs (open and filledsymbols, respectively). The solid line is a power-law fitto the data, indicating that G9 (or G) over this concen-tration range scales as Cb (b ' 2.43). The dashed linethrough the data at the lowest copolymer concentrationcorresponds to the linear relationship between G9 (or G)and C in the Rouse semidilute regime,46 in which theslope is equal to RT/M# n (' 1.54 3 105 dyn-cm/g). Thedependence of the exponent « (from tan d ; v«, wheretan d 5 G0/G9) on C is seen in the inset for both slow-cooled (‚) and quenched (F) gels. The line in the insetis a linear fit of « to C for the slow-cooled gels.


slow-cooled gels. This scaling relationship pro-vides a direct correlation between two very differ-ent measures of the midblock-stabilized gel net-work, one within the LVE regime (G9) and theother at more realistic application conditions (tR),over the TPEG concentration range 0.70 # woil# 0.90.

CONCLUSIONS

Transmission electron microscopy and dynamicrheology have been employed in concert to ana-lyze the morphological and property developmentof thermoplastic elastomer gels (TPEGs), whichwere prepared by mechanically blending a SEBStriblock copolymer with a low-molecular-weight,midblock-selective mineral oil. Evidence of sty-rene endblock micellization is provided by TEMwhen the oil fraction lies between 0.70 and 0.90.Compression molding of the TPEGs appears tohave no discernible effect on micelle shape or size,as compared to similar materials that are mixedat elevated temperature (180°C) or cast from aneutral solvent, and demonstrates that these ma-terials are surprisingly robust.32 The TPEGs ex-amined here have been subjected to two thermaltreatments, rapid quenching and slow cooling, af-ter pressing. As a result of the ability of theseTPEGs to recover rapidly from relatively highlevels of imposed stress, materials subjected toboth cooling rates exhibit comparable rheologicalbehavior. Since the slow-cooled TPEGs remain attemperatures above the styrene Tg longer thanquenched TPEGs, they possess greater long-range micellar order than their quenched analogsof equal concentration. Long-range order can beenhanced in these TPEGs, regardless of initialsample preparation, through extended annealingat elevated temperatures. Frequency spectra andstress relaxation experiments confirm the pres-ence of a gel network (formed by bridged copoly-mer midblocks) at oil fractions ranging from 0.70to 0.90, relatively independent of cooling rate. Inall of the rheological tests performed here, theTPEGs immediately recovered their inherentshear moduli, which were found to decrease byover an order of magnitude with increasing oilcontent. Thus, a reduction in the population ofcopolymer midblocks with increasing woil appearsto be responsible for the formation of midblocknetworks that are less effective at promotingTPEG elasticity. Measurement of an ultimate ma-terial property (i.e., the time to rupture) likewise

reveals a marked reduction with an increase inwoil, as well as a power-law correlation with re-spect to G9. While these TPEGs serve as modelblock copolymer systems that contain a block-selective solvent, they are also of commercial im-portance, principally as sealants and shock-ab-sorbing materials.

This work was supported by the Shell DevelopmentCo., the Raychem Corp. and, in part, by the NationalScience Foundation (CMS-941-2361). We thank Ms.E. M. Hawkins and Ms. S. H. Roberts for technicalassistance.

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