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Polymer 55 (2014) 1883e1892

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Polymer

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New insight into microstructure-mediated segmental dynamicsin select model poly(urethane urea) elastomers

Alex J. Hsieh a,e,*, Tanya L. Chantawansri a, Weiguo Hu b, Kenneth E. Strawhecker a,Daniel T. Casem c, Jeffrey K. Eliason d, Keith A. Nelson d,e, Ethan M. Parsons e

aU.S. Army Research Laboratory, RDRL-WMM-G, Aberdeen Proving Ground, MD 21005-5069, USAbUniversity of Massachusetts, Department of Polymer Science & Engineering, Amherst, MA 01003, USAcU.S. Army Research Laboratory, RDRL-WMP-B, Aberdeen Proving Ground, MD 21005-5069, USAdDepartment of Chemistry, Massachusetts Institute of Technology, Cambridge, MA 02139, USAe Institute for Soldier Nanotechnologies, Massachusetts Institute of Technology, 500 Technology Square, Cambridge, MA 02139, USA

a r t i c l e i n f o

Article history:Received 11 December 2013Received in revised form5 February 2014Accepted 9 February 2014Available online 18 February 2014

Keywords:Poly(urethane urea) elastomersSegmental dynamicsDynamic strain-rate sensitivity

* Corresponding author. U.S. Army Research Labodeen Proving Ground, MD 21005-5069, USA. Tel.: þ1 4676.

E-mail address: alex.j.hsieh.civ@mail.mil (A.J. Hsie

http://dx.doi.org/10.1016/j.polymer.2014.02.0370032-3861/Published by Elsevier Ltd.

a b s t r a c t

Segmental dynamics in a series of 4,40-dicyclohexylmethane diisocyanateediethyltoluenediamineepo-ly(tetramethylene oxide) based poly(urethane urea) (PUU) elastomers have been investigated through amulti-scale characterization approach. This includes broadband dielectric analysis, solid-state nuclearmagnetic resonance (NMR), plate impact, and impulsive stimulated scattering. Dielectric relaxationmeasurements applicable at frequencies up to 106 Hz are useful for interpreting the high strain-ratedeformation response; i.e. at the moment of target interaction with an accelerating impact or MHzstress wave excitation. Additionally, the capability of solid-state NMR to differentiate the microstructure-mediated segmental dynamics; correspondingly, the presence of a rigid phase (those in the phase-mixedregions) and a mobile phase associated with the soft-segment domains is demonstrated. These newinsights not only further elucidate the microstructure details discerned through atomic force microscopy,but also enable the prediction of the macroscopically dynamic response in these model PUUs, particu-larly on the temporal scale over the range of msens.

Published by Elsevier Ltd.

1. Introduction

Polymers are pervasive and have become indispensible in everyaspect of our lives; their robust mechanical properties and multi-functionalities have afforded these lightweight materials the po-tential for high performance at low cost and at sustainable use [1].Yet, many of the system designs and integration depend solely onperformance specifications available from manufacturers, but noton the microstructure details. Traditionally, strength, toughness,and thermal stability are among the key material parameters ofchoice for a broad range of engineering polymeric materials.Polymers are viscoelastic in nature, yet a ductile polymer capablefor use as a monolith or a spall layer may not be sufficient to serveas an interlayer with strong energy dissipation in a compositesystem [2,3]. Thus, a multifaceted approach is required for the

ratory, RDRL-WMM-G, Aber-10 306 2292; fax: þ1 410 306

h).

development of enabling polymeric materials technology withsignificantly enhanced survivability against extreme dynamic en-vironments. Specifically, better understanding of molecular mech-anisms that can lead to desired polymer dynamics at themoment oftarget interactionwith either accelerating projectiles or shockwaveexcitation is essential for the design of next-generation polymer-based protective systems.

Amongst the class of thermoplastic elastomers, styreneebuta-dieneestyrene and styreneeisopreneestyrene, for example, havebeen used in a variety of engineering applications; however, thesetri-block copolymers lack specific intermolecular interaction suchas hydrogen bonding. Conversely, extensive hydrogen bonding canoccur in the family of poly(urethane urea) (PUU), polyurea andpolyurethane elastomers. The cohesive strength of the bidentatehydrogen-bonding interaction in urea is much stronger than themonodentate interaction in urethane, thus favoring mechanicalstiffening in polyureas and PUUs over polyurethanes. There havebeen a significant number of studies on compositionemorphologyeproperty characterization of these elastomers [4e18],predominantly focused on the influence of microstructure forma-tion on thermal and mechanical properties. Yilgor et al. in a

A.J. Hsieh et al. / Polymer 55 (2014) 1883e18921884

comprehensive evaluation of non-chain extended polyurethanesand polyureas further revealed that polyureas exhibited greatertendency towards microphase separation than the correspondingpolyurethane analogs, and stronger hydrogen bonding between theurea groups was regarded as an important attribute [13]. It has alsobeen noted, through in-situ microstructure/mechanical character-ization using synchrotron radiation along with small-angle X-rayscattering (SAXS), that disruption of the domain morphology andreorganization of hard segments could be induced in response tostress at slow or moderate strain rates, and as a result the SAXSscattering pattern was found distinctly anisotropic [8,19,20].However, Pathak et al. indicated that at higher strain rates thisbehavior was not apparent in polyurea, as the morphologicalresponse of the hard segments was too slow and an isotropicscattering pattern was observed in SAXS [20].

High performance polyurea and poly(urethane urea) (PUU)elastomers have recently gained considerable interest throughoutthe Department of Defense (DoD), particularly for their potential inballistic impact protection and blast mitigation capabilities [21e30]. This class of elastomeric materials, as pointed out by Rolandand co-workers [21], can be tailored to transition from the rubbery-like into the leathery-like regime or even into the glassy regimewith increasing strain rate, where stress levels may be greatlyenhanced and large energy dissipation can be realized. Other po-tential mechanisms including shock impedance mismatch, shock-wave dispersion and strain delocalization have also been consid-ered effective in blast-wave mitigation and projectile impact en-ergy absorption [21].

Recent work based on a single chemistry of select model PUUsconsisting of 4,40-dicyclohexylmethane diisocyanate, dieth-yltoluenediamine, and poly(tetramethylene oxide), PTMO, has alsorevealed that greater phase-mixing between hard and soft seg-ments has a large effect on the viscoelastic relaxation behavior, andcorrespondingly key to property enhancement in both chemicalbarrier and dynamic strain-rate hardening [31e34]. It was notedthat the choice of the molecular weight (MW) of PTMO drasticallyaffected the extent of phase mixing, and the resulting tunablemicrostructure was clearly evidenced by atomic force microscopy(AFM) [33]. The AFM-phase images revealed significant phasemixing in PUUs as theMWof PTMO being 650 or 1000 g/mol. In thecase of PUU with PTMO MWof 1000 g/mol, the AFM-phase imagesrevealed microstructure features, which included (1) a matrixconsisting of predominantly phase-mixed hard and soft segments,wherein the hard segments mostly form fibrillar-like structures;and (2) hard domain regions consisting of rod-like lamellar hardsegments oriented in parallel, which grew in long stacks or lay nextto each other to form irregularly shaped hard domains [33]. Thefibrillar-like feature was postulated to consist of hard segmentsinteracting to form non-lamellar stiff regions oriented parallel tothe polymer chains, whereas the rod-like lamellar hard segmentswere presumably formed through the self-assembly of hard seg-ments with the domain axis (or the long period) orientedperpendicular to the polymer chains [33]. Measurements of typicalneedle-like individual lamellar hard segments resulted in charac-teristic dimensions of w10 nm in width, with a spacing of w5e10 nm [33]. As the MW of PTMO decreased to 650 g/mol, greaterphase mixing was evidenced where the intermolecular hydrogenbonding may presumably facilitated the formation of co-continuous hard and soft segments network, whereby the overallphase contrast between the hard and soft segments regions wasnoted insignificant in both the AFM-phase images [33] and SAXSdata [Beyer, F.L.; Hsieh, A.J. unpublished results].

On the other hand, AFM results displayed significantmicrophaseseparation in PUU with PTMO MW of 2000 g/mol. Coarse-grainedsimulation results further revealed that the flexibility of an

interface between the hard and soft segments was able to affect theoverall morphology including promoting the formation of elon-gated micellar structures. Results indicated that flexibility in theinterface would allow it to coil around the hard segment, pro-moting disordering and phase mixing [34].

It is well recognized that the temperature and pressure de-pendences of a polymer are strongly influenced by its viscoelasticmodes that are probed with respect to the corresponding highstrain-rate, high-pressure deformation behavior [29]. Segmentalmotion, which arises from intermolecular, cooperative conforma-tion changes, typically has a stronger response to thermodynamicvariables than chain reptation dynamics [29]. Thus, segmental dy-namics can be critical to the material response particularly underextreme dynamic environments. However, for the rational designof responsive high performance hierarchical elastomeric materials,it is necessary to elucidate key physical events, which occur onvarious temporal and spatial scales. In this work, we focus oninvestigation of the microstructure influence on segmental dy-namics in a series of 4,40-dicyclohexylmethane diisocyanateediethyltoluenediamineepoly(tetramethylene oxide) based PUUelastomers. A comprehensive study based on a well-coupled multi-scale characterization approach was carried out, which includessolid-state nuclear magnetic resonance (NMR), broadband dielec-tric analysis, impulsive stimulated scattering, and plate impactmeasurements.

NMR is better known as a powerful tool for structural elucida-tion in the liquid state; however, solid-state NMR is also capable ofhigh-resolution detection of the dynamics associated with variousphases in a polymer [35e40]. In solid-state, strong dipolar couplingbetween neighboring protons broadens the wideline NMR spectra.Recent results clearly indicated large difference in the 1H widelinespectra at ambient temperature for PUUs with varying MWs ofPTMO [41]. In the case of PUU with PTMO MW of 650 g/mol, thespectrum exhibited a full-width-at-half-height (FWHH) of 32 kHz,indicative of materials in proximity to the glassy state, where nosignificant narrowing was evident [41]. This glassy-like behaviorwas consistent with the DMA tan d data, where a broadened Tgspanned over room temperature was also observed [33,41]. As theMW of PTMO increased, the FWHH became 10.5 kHz and 3.7 kHzfor PUUs with PTMO MW of 1000 and 2000 g/mol, respectively[41]. It was noted that the trend in FWHH correlated well with theextent of phase mixing observed by AFM [33].

Furthermore, a single exponential fit was shown to adequatelymatch the wideline 1H free induction decay (FID) data acquired forthese PUUs at temperatures of 293, 313, and 333 K [41]. The 1Hdipolar dephasing time (Tdd), a decay constant determined basedon the curve-fitting of the 1H FID data, was inversely proportionalto the peak width of FWHH, and appeared to be a good represen-tation of the rate of molecular motion, particularly with respect tothe propensity to becoming glassy. A linear relationship betweenTdd and the reciprocal of temperature was reported, where theslope increases suggesting greater dephasing with decreasing SSMW as temperature reaches near the vicinity of ambienttemperature.

Dielectric spectroscopy is an alternative to mechanical mea-surements, such as dynamic mechanical analysis, of viscoelasticrelaxation associated with glass transition, though it is more ad-vantageous because dielectric data can be routinely obtained athigh frequencies [29,42] on the range of relevance to the ballisticimpact deformation conditions. In dielectric measurements, polargroups on the polymer can respond to alternating electric fieldbetween two parallel plate capacitors, and absorption maximaoccur when dipole relaxation times s due to molecular motionsmatch the electric field frequency u, i.e. when us z 1 [42]. Studieshave shown good correlation between relaxation data obtained by

Table 1Composition and the calculated values of hard segment content of select modelPUUs.

Molar ratioHMDI:DETA:PTMO

MW of PTMO (g/mol) HS (wt.%)

HSu HSuu

211-2000 2:1:1 2000 16 26211-1000 2:1:1 1000 26 41211-650 2:1:1 650 33 52431-2000 4:3:1 2000 37 44532-1000 5:3:2 1000 34 48

A.J. Hsieh et al. / Polymer 55 (2014) 1883e1892 1885

both DMA and dielectric measurements; the local dynamicsmeasured dielectrically are reported to be nearly identical to thesegmental motions underlying the dynamic mechanical propertiesin the glass transitions [29,42].

Both the broadband dielectric analysis and solid-state NMRmeasurements are effective for characterization and differentiationof molecular dynamics with respect to the tunable microstructures.The goal of this work is to provide a better scale-bridging fordetermination of molecular mobility across the atomistic andmesoscale, with an attempt to better corroborate with micro-structure details discerned from atomic forcemicroscopy, as well asto highlight the macroscale dynamic deformation responseincluding under plate impact and MHz frequency excitation inselect model PUUs.

2. Experimental

2.1. Materials

Select model PUU elastomers composed of 4,40-dicyclohexyl-methane diisocyanate (HMDI), diethyltoluenediamine (DETA), andpoly(tetramethylene oxide) (PTMO), with three different molecularweights (MW) of the PTMO soft segment (SS), 650, 1000 and2000 g/mol, were chosen for this study. The PUU elastomers wereprepared via a two-step, pre-polymer synthesis method wheredetails can be found elsewhere [31]. Fig. 1 is a schematic illustrationof the PUU repeat unit, 4,40-dicyclohexylmethane diisocyanateediethyltoluenediamineepoly(tetramethylene oxide). In the samplenomenclature, the numerals “xyz” refer to the molar ratio ofHMDI:DETA:PTMO, and the succeeding “650”, “1000”, and “2000”refer to the MW of PTMO as 650, 1000, and 2000 g/mol, respec-tively. Throughout this work, 211-650, 211-1000, and 211-2000were used for systematic comparison of the molecular influence onthe segmental dynamics measurements, whereas 532-1000 and431-2000 were also included in the mechanical propertiesmeasurements.

Two different approaches for calculation of hard segment (HS)content are undertaken as follows [7]:

%HSu ¼ 100ðR� 1ÞðMdi þMdaÞ�Mg þ RðMdiÞ þ ðR� 1ÞðMdaÞ

� (1)

%HSuu ¼ 100ððRÞðMdiÞ þ ðR� 1ÞðMdaÞÞ�Mg þ RðMdiÞ þ ðR� 1ÞðMdaÞ

� (2)

where R is themolar ratio of the diisocyanate to PTMO, andMdi,Mdaand Mg are the number average molecular weights of the diiso-cyanate (HMDI), diamine (DETA) and PTMO, respectively. Equation(1) corresponds to the Flory’s formula and only accounts for theportion of diisocyanate that reacts with diamine, while equation (2)also accounts for the additional portion of diisocyanate that reactswith the polyether diol. Table 1 lists the composition alongwith theweight percentage (wt.%) values for both HSu and HSuu for the

Fig. 1. A schematic illustration of repeat unit of the 4,40-dicyclohexylmethane di

select PUUs. These hard segment wt.% values will be differentiatedwhen utilized to interpret the molecular influence on high strain-rate in comparison with the corresponding quasi-state mechani-cal properties.

2.2. Dielectric analysis

Broadband dielectric spectroscopy measurements of selectmodel PUUs were performed at various temperatures using aNovocontrol Concept system in the frequency range of 10�1 to106 Hz. A film sample of w1 mm thick was placed between twogold-plated electrodes.

2.3. Solid-state nuclear magnetic resonance (solid-state NMR)

All NMR experiments were conducted on a Bruker DSX300spectrometer, operating at a 1H frequency of 300.12 MHz and 13Cfrequency of 75.47 MHz, in 4 mm CP/MAS probes. Both 1H wideline(non-spinning) experiments and magic angle spinning (MAS) ex-periments were carried out. 1H spin diffusion experiments wereconducted with a proton dipolar dephasing filter, during which thetransverse 1H magnetization undergoes free evolution and the partfrom the rigid phases loses coherence due to strong dipolarcoupling. The filter length was 0.2 ms (a full rotor cycle at thespinning speed of 5 kHz) such that only the signal from the mobile-SS fraction survives the filter. More details regarding the widelineand the TD-WISE (time-domain wideline separation) experimentsas well as data analysis can be found in Ref. [41].

2.4. Plate impact

Measurements were conducted in the configuration shown inFig. 2. The disk-shaped PUU specimens, nominally 3.4 mm inthickness and 40 mm in diameter, were accelerated in a light gasgun and impacted against stationary soda lime glass target plates(nominally 2.0-mm thick, 40-mm diameter) at impact speeds of298 m/s and 998 m/s. The impact velocities are such that the sodalime glass remains elastic throughout the experiments. The particlevelocities at the center of the free surface of the target plates weremonitored with a Velocity Interferometer System for Any Reflector(VISAR) [43]. These measurements are accurate to within �1% and

isocyanateediethyltoluenediamineepoly(tetramethylene oxide) based PUU.

Fig. 2. A schematic of shot configuration used in plate impact measurements of PUUagainst stationary soda lime glass (SLG) target, in the impact direction as indicated bythe red arrow. Green V points to the locality where a Velocity Interferometer Systemfor Any Reflector (VISAR) is used to monitor the particle velocity. (For interpretation ofthe references to color in this figure legend, the reader is referred to the web version ofthis article.)

A.J. Hsieh et al. / Polymer 55 (2014) 1883e18921886

have a 0.5 ns time resolution. Impact speeds were measured justprior to impact with a series of electrically conducting pins, to anaccuracy of �1%.

The basic concept of the plate impact experiment is to use themeasured behavior of the soda lime glass along with its knownproperties to infer the behavior of the PUU flyer. The impact pro-duces shock waves that transmit into both the flyer and the target.The approach for the experiment is to ensure that one dimensionalstrain conditions are maintained for all relevant data. More detailsof the experimental and data analysis can be found in Ref. [43].

2.5. Impulsive stimulated scattering (ISS)

The transient grating technique for ISS is well documented inthe literature and more details can be found in Ref. [44]. In brevity,two w60-ps duration excitation laser pulses (515 nm wavelength)are crossed in the sample, where they interfere to create a sinu-soidal intensity profile. Absorption of the excitation light results insample heating and fast thermal expansion generates a coherentacoustic wave, the wavelength of which is set by the period of theinterference pattern. The temperature profile and stress caused bythe acoustic wave induces a modulation of the complex refractiveindex, which is monitored by diffraction of a 532-nm continuouswave probe beam [45]. Heterodyne detection of the diffractedsignal is employed to cancel out extraneous noise sources. Thesignal intensity is measured by a fast detector and read out on anoscilloscope. From ISS, the sound speed and attenuation rate weredetermined and frequency-dependent modulus values werecalculated.

Fig. 3. AFM-phase images (left: 5 � 5 mm and middle: 1 � 1 mm; 30� light-to-dark) of 211regions is seen [after Ref. [33]] and highlighted within the dashed-blue circles in the middphase-separated microstructure seen in 211-2000 [after Ref. [33]], wherein SS-rich region isthis figure legend, the reader is referred to the web version of this article.)

3. Results and discussion

3.1. Influence of microstructure on segmental dynamics

In accordance with AFM, the molecular weight of PTMO signif-icantly affected the self assembly of hard segments as well as theoverall phase mixing between the soft and hard segments. Fig. 3highlights the tunable microstructure characteristics, particularlya matrix consisting of predominantly phase-mixed hard and softsegments was evidenced in 211-1000, where the hard segment isbright in color (indicative of being stiff) in AFM-phase image(shown on left) and also highlighted within the dashed-blue circlesin the middle. In Fig. 3, AFM-modulus images reveal the presence ofphase-mixed regions in 211-2000, which was found not discerniblein the AFM-phase images [33], in addition to the irregularly-shapedrod-like features and a compliant matrix of SS-rich regions. Thepresence of phase-mixed regions observed in AFM also corroboratewell with the tan d data obtained via dynamic mechanical analysis(DMA), reported in Refs. [31e33].

In DMA tan d at 1 Hz, there was a very distinct and intenserelaxation at about 210 K along with a weak shoulder for 211-2000.The strong relaxation peak presumably corresponds to the SS-richamorphous matrix, whereas the weak shoulder may be associ-ated with the sparsely populated phase-mixed regions. In additionto a w34 K increase in glass transition temperature (Tg), 211-1000also exhibited broadening of the glass transition, indicative ofphase mixing [31,32]. Broadening of the SS relaxation and greaterphase mixing were also evidenced in 211-650, wherein the glasstransition spanned over ambient temperature [33].

Although AFM results clearly reveal the molecular influence onthe tunable microstructure, this work will further discern anddifferentiate the mobility of the SS associated in either the SS-richregions (highlighted with a dashed-green circle shown in bothFig. 3 and the schematic (left) in Fig. 4) or in various phase-mixedregions (highlighted with a dashed-blue circle in both Fig. 3(middle) and in the schematic (right) in Fig. 4).

3.2. Solid-state NMR measurements

To better understand the molecular influence on polymer dy-namics, we exploit solid-state NMRmeasurements on select modelPUUs. Since time-domain 1H NMR does not have the resolution todistinguish the dynamical behaviors of various chemical structuresin a complex material such as PUU, time-domain wideline separa-tion (TD-WISE) measurements were performed to better differen-tiate details on the molecular level. Through utilizing the WISEpulse sequence [38] followed by treating the data in the timedomain rather than in the frequency domain, we can achieve

-1000, wherein a matrix of predominantly phased-mixed fibrillar-like microstructurele; AFM-modulus image (right: 10 kPae1 GPa light-to-dark, log scale) highlights thehighlighted by the dashed green circle. (For interpretation of the references to color in

Fig. 4. Schematic highlighting the soft segments located within a dashed-circle which corresponds to (but not to scale with) either the SS-rich region highlighted in Fig. 3 (left) orthe phase mixed region highlighted in Fig. 3 (middle).

A.J. Hsieh et al. / Polymer 55 (2014) 1883e1892 1887

chemical resolution. The phase components can also be modeledthrough the 1H time-domain line shapes. As a result, the mobility ofSS associated in either SS-rich regions or in various phase-mixedregions as highlighted in Fig. 4 was characterized. More detailsregarding the TD-WISE experimental methods and analysis can befound in Refs. [38,41].

It was found that the soft segments were comprised of onemobile and one rigid fraction, viz. mobile-SS and rigid-SS,respectively, where the corresponding dynamics vary as a resultof the local morphology. The dephasing behavior of the SS, ob-tained from the TD-WISE measurements based on the time-decayof ether carbon signal of PTMO via magic angle spinning (MAS) isshown in Fig. 5 [41]. It is apparent that each of these SS-dephasing data can be curve-fitted by a two-exponential func-tion. The fitting results are listed in Table 2, where the fractionswere normalized to 100% for total SS. For these PUUs, the dy-namics and relative portions of rigid vs. mobile fractions amongthe SS exhibit large contrast. The rigid-SS dynamics is at least anorder of magnitude slower than that of the correspondingmobile-SS, and this is consistent for all the three PUUs regardlessof the SS MW. Additionally, the dynamics in the rigid-SS de-creases with a decrease in the SS MW, meanwhile the populationof the rigid-SS becoming glassy-like also increases withdecreasing SS MW [41]. These results are consistent with the 1Hwideline spectra; where 211-650 is also of most temperature-sensitive toward the glass transition, indicative of greater

Fig. 5. TD-WISE data (the symbols are for the area of the 71 ppm PTMO ether carbonpeak) associated with SS obtained for 211-650, 211-1000, and 211-2000. The lines arethe corresponding best fits (two exponential components) [after Ref. [41]].

strain-rate sensitivity. Thus, solid-state NMR measurements arecapable of revealing the microstructure dependence of segmentaldynamics, highlighting the role of molecular mechanism on themacroscopic deformation response.

For PUUs with varying MWof PTMO, it is noted that rigid-SS canbe associated with SS directly connected to the hard segments orthose facilitated by the various extent of intermolecular mono-dentate or bidentate hydrogen bonding. Thus, rigid-SS can be in thevicinity near the interphase between the hard segments and softsegments domains for PUUs with higher MW of PTMO, i.e. as seenin 211-2000, within the phase mixed regions or even associatedwithin co-continuous, highly intermolecular hydrogen bonded softand hard segments network. The latter is plausible within themicrostructure of 211-650 which consisting of lower MWof PTMO.1H spin diffusion experiments with a proton dipolar dephasingfilter were also conducted to further discern the phase-mixedmicrostructure, wherein the transverse 1H magnetization un-dergoes free evolution and the part from the rigid phases losescoherence due to strong dipolar coupling. At the filter length of0.2 ms, only the signal from the mobile-SS fraction survives thefilter. During the diffusion period, the HS signal quickly rises atshort diffusion times (�1 ms), and levels off at 50e100 ms.Assuming a dimensionality of 2 (fibrillar) for HS domains, asobserved by AFM (Section 3.1), and using commonly accepted spindiffusion coefficients for these systems [46e48], the average HSdomain diameter (shortest dimension of domains) is estimated tobe about 7 nm. This length scale is consistent with AFM observa-tions. However, the diffusion curve does not follow a well-behavedlinear behavior against the square root of diffusion time, even atshort diffusion times. This is likely because the materials do nothave a well-defined two-phase structure, but rather have fourdynamical components, including rigid-HS, mobile-HS, rigid-SS,and mobile-SS [41], with their relative spatial arrangements un-known. The random distribution of hard- and soft-segmentsequence length would also result in a broad distribution ofdomain sizes, which adds difficulty to data interpretation. Moreaccurate domain size determination by spin diffusion in thesecomplex systems would require much more attention and will be asubject of future studies.

Table 2The fractions of mobile components as well as their corresponding MAS 1H dipolardephasing time, Tdd,mas, for SS as obtained by fitting of TD-WISE data [after Ref. [41]].

PUU Rigid component Mobile component

Tdd,mas (ms) Fraction Tdd,mas (ms) Fraction

211-650 17 � 1 (78 � 2)% 217 � 18 (22 � 2)%211-1000 20 � 2 (53 � 2)% 340 � 10 (47 � 2)%211-2000 43 � 6 (34 � 2)% 1315 � 28 (66 � 2)%

Fig. 7. Influence of SS MW on dielectric loss spectra obtained for 211-650, 211-1000,and 211-2000 at 293 K (solid) and 298 K (dashed), respectively.

A.J. Hsieh et al. / Polymer 55 (2014) 1883e18921888

3.3. Broadband dielectric spectroscopy (DES) measurements

In this work, we attempt to further elucidate the essence ofsegmental dynamics for use to interpret the dynamic materialresponse as a function of strain rate. Dielectric measurements werecarried out in the frequency range of 10�1 to 106 Hz, where the dataspans temperatures from 173 K to slightly above ambient temper-ature. We observed that the dispersion in the dielectric loss shiftstoward lower frequency with decreasing temperature. From thedata, we determined the frequency of the maximum in thedielectric loss, fmax, at each temperature [see Supplementary data];the value of fmax at 298 K, w21,166 Hz, defines the frequencywhereby 211-1000 can transition to becoming glassy-like underambient conditions.

The local relaxation time, s, representing the most probablevalue, is typically defined as s ¼ (2p fmax)�1. Fig. 6(a) is a plotshowing the temperature sensitivity of s along with relaxation timeavailable from the DMA loss modulus measurements [31] obtainedfor 211-2000. For the DMA data, the Tg values associated withrelaxation of both the SS-rich and phase-mixed regions areincluded for comparison. For 211-2000, the dielectric data appearto be much closer to the phase-mixed relaxation data. These resultsindicate that segmental dynamics associated with the glass tran-sition is manifested by a corresponding dielectric and mechanicalrelaxation processes measured by dielectric analysis and DMA,respectively. For 211-2000, curve-fitting of the corresponding fmaxvs. 1000/T data was also performed with the VogeleFulchereTammann (VFT) function [21], and the plot of the VFT fit is shownalong with the fmax vs. 1000/T data in Fig. 6(b),

log fmax ¼ log f0 �B logðeÞT � T0

(3)

In equation (3), f0 is the asymptotic, high temperature value offmax, B is related to the dynamic fragility, and T0 is the Vogel tem-perature, a constant [21,26]. Based on the VFT fit, calculation of thetemperature at a relaxation time of 100 s results in an estimate ofcalorimetric Tg [26], approximately 198 K, which is close to the SS Tgof 211-2000 observed by differential scanning calorimetry [31].

As the SS MW decreases, the segmental dynamics changedrastically among these PUUs. Fig. 7 clearly reveals that for 211-650at 298 K, there appears to be a strong relaxation at w8 Hz alongwith a second loss maximum atw4600 Hz, the former presumablycorresponding to the most phase-mixed regions, where the for-mation of co-continuous, highly intermolecular hydrogen bonded

Fig. 6. (a) (left) Plot of temperature dependence of relaxation times obtained from broadbandenote data associated with the phase-mixed and SS-rich regions, respectively) [31] for 211-fit (solid line), where log f0 is 15.6/s, B is 4041 and T0 is 102.9 K.

HS/SS network is plausible. For 211-1000, a strong relaxation atw21,166 Hz is observed, and additionally there is a small shoulderat around w8e10 Hz, though it is not as significant as the strongrelaxation seen in 211-650.

The dielectric loss data indicate that 211-650 can readily becomeglassy at low frequency near ambient temperatures in comparisonwith both 211-1000 and 211-2000. Thus, for 211-650 the dynamics-induced glass transition, presumably as a result of the well phase-mixed regions, occurs at a frequency nearly three and five orders ofmagnitude sooner than that for 211-1000 and 211-2000, respec-tively. In fact, 211-2000 may still be at most in the leathery stateeven as the frequency reaches 105 Hz at ambient temperatures.Table 3 lists the values of fmax, s and the segmental mobility, definedas 1/s, determined at 298 K for select model PUU elastomers. For211-2000, the s value, 5 � 10�8 s, was determined based on theextrapolation of the curve-fit of the s vs. 1000/T data to 298 K. Thus,the corresponding segmental mobility of 211-2000 is w1.9 � 107/s,indicative of being mostly rubbery-like characteristic among thethree PUUs. It is also noteworthy that the log(fmax) at 298 K for 211-1000 isw4.3, much lower than the reported value, 5.95, at 296 K fora commercial polyurea, despite similar SS with MW of 1000 g/mol[21]. We postulate that this difference is the result of a variation inthe extent of phase mixing; the lower log(fmax) value correspondsto greater phase mixing based on our dielectric analysis results andmicrostructure analysis via AFM. Thus, the extent of phasemixing is

d dielectric analysis (filled circle) as well as from DMA (open circle and inverted triangle2000; (b) (right) Plot of the corresponding fmax vs. 1000/T data shown in (a) with a VFT

Table 3Measurements of fmax and the calculated values of s and 1/s of 211-650, 211-1000,and 211-2000 PUUs.

211-650 211-1000 211-2000

fmax (Hz) 8 21,166 e

s (s) 0.02 7.5 � 10�6 a5 � 10�8

1/s (1/s) 50 1.3 � 105 1.9 � 107

a Calculated through extrapolation of the curve-fit of the dielectric s vs 1000/Tdata (Fig. 6(a)).

A.J. Hsieh et al. / Polymer 55 (2014) 1883e1892 1889

presumably to be greater in 211-1000 than the correspondingcommercial polyurea.

3.4. High strain-rate mechanical deformation

The segmental dynamics as indicated above is a strong functionof SS MW, thus one can expect its influence on high strain-ratemechanical deformation, particularly as the strain-rate reachesthe segmental mobility of these PUU elastomers. We first comparethe quasi-static mechanical response including both loading andunloading of 211-650, 211-1000 and 211-2000; Fig. 8 is a plot of thetrue stress vs. true strain data obtained at a true strain rate of 0.001and 0.01 s�1. It appears that 211-650 exhibits greater strain-ratesensitivity in addition to greater strain hardening than 211-1000and 211-2000. For 211-2000, the flow stress level shows a verycompliant behavior and the least amount of the inelastic strainfollowing recovery (permanent set deformation), indicative ofrubbery-like responsewhich is consistent with the aforementionedsegmental dynamics. The flow stress values are higher for 211-650and 211-1000 than 211-2000, corresponding to their difference inthe hard segment content, HSu. The calculation of HSu only ac-counts for the HMDI that reacts with the chain extender DETA toform the urea linkages but not those reactedwith PTMO to form theurethane linkages. Thus, keeping the stoichiometric ratio constantthe HSu values vary from 16 to 33 wt.% among these PUUs (Table 1).It is noted that 211-650, in addition to a moderate increase in flowstress as the strain rate increases, exhibits considerably higherextent of strain-hardening. The glassy-like stressestrain behaviorin 211-650 also corroborates well with the DMA data, where the Tg,particularly for the phase-mixed domains of 211-650, was seenspanning across the ambient temperature.

Fig. 8. Compressive stressestrain curves for 211-650, 211-1000, and 211-2000 ob-tained at a strain rate of 0.001 (dashed) and 0.01 (solid) s�1.

In this work, we focus on the high strain-rate material responseincluding deformation under the split-Hopkinson pressure bar(SHPB) impact loading. For the SHPBmeasurements the strain ratesare in the range of 1000e2500 s�1; Fig. 9 displays the true stress vs.true strain obtained at a strain rate of 2150� 100 s�1 in comparisonwith those obtained at 0.001 s�1 for the three PUUs. At high strainrates, the relative increase in flow stress corresponding to 0.2 truestrain varies from w4.7 MPa for 211-2000 to w16.3 MPa andw41.3 MPa for 211-1000 and 211-650, respectively. As indicated bythe segmental dynamics measurements, the 1/s values for 211-1000 and 211-2000 are much higher than the strain rates underSHPB, thus both 211-1000 and 211-2000 most likely do not exhibitthe deformation-induced glass transition. On the other hand, the svalue for 211-650 is w0.02 s, thus the corresponding segmentalmobility, 1/s, is w50 s�1, which is much lower than the strain ratesunder SHPB. As a result, the material response of 211-650 is ex-pected to become glassy-like under SHPB and therefore will exhibitmuch higher stiffness than 211-1000 and 211-2000.

To better highlight the segmental dynamics influence on dy-namic mechanical deformation response, we also compare thetrend in strain rate sensitivity of select model PUU elastomers thathave similar HSu contents but varying in SS MW. The 532-1000 and431-2000 PUUs are the materials of choice; 532-1000 has slightlyhigher HSu content, 34 wt.%, than 211-1000 (26 wt.%), whereas37 wt.% for 431-2000 (versus 16 wt.% for 211-2000), both aresimilar to that of 211-650, 33 wt.%. In Fig. 10, the flow stress valuesfor these PUUs are plotted as a function of strain rate on a loga-rithmic scale, where the flow stress was calculated based on 0.2true strain from the stressestrain data. At low strain rates, all threePUUs, 431-2000, 532-1000 and 211-650, exhibit similar flow stressvalues (Fig. 10); this is consistent with the hard segment contentwhere the flow stress can approximately be a function of HSu. Athigher strain rates under SHPB, 211-650 exhibits significantlygreater strengthening than 532-1000 and 432-2000, presumablydue to the fact that these strain rates are greater than the segmentalmobility of 211-650, where the latter exhibits a glassy-like defor-mation response in contrast to deformation observed in 532-1000and 431-2000. The rate of segmental mobility for 532-1000 and431-2000 is higher than the strain rates considered. More discus-sion regarding the segmental relaxation of 532-1000 is given inSection 3.6.

It is noteworthy that the segmental mobility calculated based ondielectric analysis reflects the most probable dynamics associated

Fig. 9. True stressetrue strain data for 211-650, 211-1000, and 211-2000 obtainedat strain rates w2150 � 100 s�1 under SHPB measurements (solid), incomparison with those obtained under quasi-static compression at a strain rate of0.001 s�1 (dashed).

Fig. 10. Flow stress values obtained at 0.2 true strain for 211-650 (green square), 532-1000 (red circle), and 431-2000 (black triangle) under quasi-static compression as wellas from SHPB measurements in the range of 1000e2500 s�1. (For interpretation of thereferences to color in this figure legend, the reader is referred to the web version of thisarticle.)

A.J. Hsieh et al. / Polymer 55 (2014) 1883e18921890

with a given polymer, yet solid-state NMR reveals the presence ofrigid-SS components associated with the phase-mixed regions in211-650, as well as in 211-1000 and 211-2000. Thus, deformation-induced glass transition can presumably occur in the phase-mixedregions in 211-1000 and 211-2000, thoughmay not be as significantas those seen throughout 211-650 which is predominantly phase-mixed. We attempt to interpret the high strain-rate flow stressdata by taking into account the effective hard segment content(HSuu), where HMDI that reacts with PTMO to form the urethanelinkages in addition to those reacted with the chain extender DETAto form the urea linkages are both included in the calculation ofHSuu, according to equation (2). This is consistent as the bidentateureaeurethane hydrogen bonding that contributes to greater phasemixing can also undergo deformation-induced glass transition,leading to the formation of the rigid-SS phase as evidenced in solid-state NMR. The calculated value of HSuu listed in Table 1 is greaterfor 211-650 (52 wt.%) than for 532-1000 (48 wt.%) and 431-2000(44 wt.%).

Fig. 11. Stress vs. strain behavior of 532-1000 obtained from plate impact measure-ments (arrows pointing to data on the Hugoniot).

3.5. Coarse-grained calculation of high strain-rate deformationresponse

As noted from Section 3.3, the strain rates available from theSHPB measurements are not high enough for 211-1000 to exhibitthe deformation induced glass transition on a global level. In thiswork, we also exploit a coarse-grained modeling approach tocalculate the material response since computations are typicallyperformed at very high strain rates of order 106 and larger. In thiswork, we calculate the stressestrain behavior of 211-650, 211-1000and 211-2000 at a strain rate of 107 s�1. At this strain rate, 211-1000is presumably to be glassy as the applied strain rate is much greaterthan the segmental mobility, 1/s, 1.3 � 105/s, whereas 211-2000may be towards the glassy-like deformation response as the strainrates approach its segmental mobility. Preliminary simulation re-sults indicate that the trend in the calculated values of flow stresscorresponds well to the variation in hard segment contents[Chantawansri, T.L.; Hsieh, A.J. unpublished results], where the

deformation-induced glass transition is presumably imposed uponall three PUUs.

3.6. Dynamic material response to the challenges with strain-rateor frequency on the temporal scale over the range of msens

To further validate the influence of segmental mobility on dy-namic material response, particularly for the PUUs with 1000 and2000 g/mol MW PTMO, we have extended the dynamic measure-ments to include those under high strain-rate, high pressure plateimpact and MHz-frequency impulse excitation loading conditions.

3.6.1. Plate impact measurementsUpon plate impact, the PUU target is under uniaxial strain,

whereas uniaxial stress is prevalent under both quasi-static andSHPB loading conditions. Upon exposure to uniaxial strain, bothhigh strain rate and high pressure are critical and need to be takeninto account when considering the material deformation. However,in this work our intent is to highlight the role of segmental dy-namics while interpreting the target response at the moment uponinteraction with an accelerating impact. The plate impact mea-surements were carried out at impact velocity of 298 m/s and998 m/s, correspondingly the strain rates reach approximately1 � 106/s and 7 � 106/s, respectively (estimated from the slopes ofthe shock fronts).

The 532-1000 PUUwas chosen for the plate impact study and itss value calculated is w11 ms, correspondingly 1/s w 9.5 � 104/s.Although this rate is much faster than the strain rates under SHPB,it is slower than the imposed strain rates under plate impact. Thus,532-1000 can presumably undergo deformation-induced glasstransition upon plate impact at strain rates w1e7 � 106/s.

For the experiment performed at 998 m/s, 532-1000 is partiallyunloaded in three successive steps after the initial compression.Details of the calculated values of free-surface particle velocityestimated from the initial shock and two release states can be foundin Ref. [43]. In Fig. 11, the stress versus strain data of 532-1000 areshown, where at each impact velocity the highest stress point is on

Table 4The calculated values of speed of sound and apparent modulus along with the decaytime of 211-650, 211-1000, and 211-2000measured from ISS, also the correspondingstorage modulus values at 1 Hz from DMA are included for comparison.

DMA, 1 Hz ISS, w300 MHz

Storagemodulus(GPa)

Speed ofsound,C (m/s)

Apparentmodulus(C2 x density) (GPa)

Decaytime (ns)

211-2000 0.02 1770 3.2 w8211-1000 0.11 2001 4.2 w8211-650 0.35 2208 5.2 w8

A.J. Hsieh et al. / Polymer 55 (2014) 1883e1892 1891

the Hugoniot (as pointed out by an arrow), and additionally, thestressestrain data along the release path appear to coincide withthose on the Hugoniot obtained for both the 298 and 998 m/s ex-periments [43]. It is nevertheless noted that the high stress, on theorder of GPas, upon plate impact reflects the measurement oflongitudinal stress, which includes a substantial pressure contri-bution; thus, these may not be used for direct comparison with theflow stress measurements based on the uni-axial stress loading.

Futureworkwill explore the pressure-dependence of segmentaldynamics in order to better account for the overall materialresponse under high pressure, high strain-rate deformation.

3.6.2. MHz-frequency impulse excitationImpulsive stimulated scattering (ISS) measurements were also

used to further investigate the frequency dependence of dynamicrelaxation behavior in these PUUs. The oscillation frequency de-pends upon the material’s acoustic velocity, and we first compare211-650 with a polydimethylsiloxane (PDMS) for reference evalu-ation to ensure the feasibility of ISS. PDMS is a thermally stableelastomer and is used in many engineering applications. Fig. 12highlights the difference in oscillation decay between 211-650and PDMS. The calculated speed of sound is about two times fasterin 211-650 than in PDMS, 2177m/s vs. 1074 m/s, presumably due tothe fact that 211-650 is much more frequency-dependent thanPDMS and thus 211-650 exhibits a greater glassy-like responseunder ISS.

In addition, we compare the acoustic response among 211-650,211-1000 and 211-2000, and the calculated results from the ISSmeasurements are listed in Table 4. It is noted that the difference inmagnitude of the frequency-dependent (u w 300 MHz) modulusvalues between these three model PUUs is much smaller than themodulus values determined at a frequency of 1 Hz via DMA. Forexample, the apparent modulus of 211-650 from ISS is only about60% larger than that of 211-2000 upon the acoustic oscillation atw300 MHz frequency (w108/s), whereas the corresponding dif-ferencewith respect to the storagemodulus is about 17 times largerwhen measured at 1 Hz via DMA. These results indicate that allthree of these PUUs are presumably at the glassy state at the300 MHz frequency range. This further confirms that deformation-induced glass transition is also present in 211-2 K upon 300 MHzacoustic wave excitation.

Fig. 12. Comparison of ISS data (curve-fits are shown for clarity) for 211-650 (red) andPDMS (black). (For interpretation of the references to color in this figure legend, thereader is referred to the web version of this article.)

4. Conclusion

A comprehensive multi-scale characterization of select modelpoly(urethane urea) elastomers was performed. It is clearlydemonstrated that the propensity towards the formation of a 3-Dinterconnected hydrogen bonded network wherein the tunablemicrostructure in these PUUs and the corresponding segmentalmobility of SS are key to the dynamic response over a broad rangeof the temporal scales, msemsens. The segmental mobility calcu-lated based on broadband dielectric analysis reflects the mostprobable dynamics associated with the respective PUU, which wasutilized to interpret the dynamic material response under split-Hopkinson pressure bar impact as well as in the moment of thetarget interaction with either an accelerating impact or stress waveexcitation. It is expected that deformation-induced glass transitioncan occur in the case when the imposed strain-rate reaches and isgreater than the segmental mobility of a given PUU. The 211-650becomes glassy upon SHPB impact in contrast to 211-1000 and 211-2000, where the segmental mobility of 211-650 is w50 s�1. Uponplate impact at strain rates of about 1e7 � 106/s, 532-1000 maypresumably undergo deformation-induced glass transition as theseimposed strain rates are much higher than its segmental mobilitywhich is about 9.5 � 104/s. However, plate impact measurementsinclude a strong pressure contribution; thus, future work will takeinto account the pressure-dependence of segmental dynamics.Upon the impulsive stimulated scattering measurements, the dy-namic response associated with a 300 MHz acoustic wave oscilla-tion reflects that 211-2000 along with 211-1000 and 211-650 are atthe glassy state, which correlates with their respective segmentalmobility. Thus, the calculated values of apparent modulus from ISSare on the same order, which is in contrast to the storage modulusvalues obtained from DMA at a frequency of 1 Hz.

The microstructure influence on segmental dynamics wasfurther elucidated via solid-state NMR time-domain wideline sep-aration measurements, wherein the dynamics in the rigid-SS is atleast an order of magnitude slower than that in the mobile-SS,which is consistent among 211-650, 211-1000 and 211-2000.Additionally, the dipolar-dephasing time of the rigid-SS fractiondecreases with decreasing SS MW, and the population of the rigid-SS also increases with decreasing SS MW. The presence of rigid-SScomponents, revealed via the solid-state NMR, is presumablyassociated with the phase-mixed regions in 211-650, as well as in211-1000 and 211-2000. These observations corroborate well withthe microstructure details discerned by AFM and the viscoelasticrelaxation data by DMA. Thus, deformation-induced glass transi-tion, when considering in a temporal scale on the order of micro-second, appears to be dominant in the phase-mixed regions,though those occur in 211-1000 and 211-2000 may not be as sig-nificant as those seen throughout the 211-650.

From the segmental dynamics perspective, broadband dielectricanalysis provides coarse-grained dynamics, whereas solid-stateNMR measurements reveal the corresponding dynamics with

A.J. Hsieh et al. / Polymer 55 (2014) 1883e18921892

more details on the atomistic scale. These measurements not onlyare capable for validation of the presence of deformation-inducedglass transition in these select model PUUs, but can be potentiallyused to guide the synthesis of hierarchical elastomers toward theoptimization of multi-functionality.

Acknowledgments

AJH and TLC acknowledge funding support from the ArmyResearch Laboratory (ARL) Director’s Research Initiative program.AJH, JKE, KAN and EMP also acknowledge collaboration opportunitythrough funding support by the U.S. Army through the Institute forSoldier Nanotechnologies (ISN), under Contract W911NF-07-D-0004. WH thanks the Army Research Office for funding, also theUniversity of Massachusetts-Amherst Materials Research Scienceand Engineering Center (MRSEC) and the Materials Research Fa-cilities Network (MRFN) for supporting the use of DSX300 solid-state NMR instrument. AJH acknowledges Dr. Richard C. Becker ofARL for the review and critical comments for this manuscript. AJHalso thank Dr. Norm Rice of Triton Systems, Inc. (Chelmsford, MA)for providing PUU materials, through an ISN project funded by theArmy.

Appendix A. Supplementary data

Supplementary data related to this article can be found online athttp://dx.doi.org/10.1016/j.polymer.2014.02.037.

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