high molecular integration 
 afforded by conveniently processed soft - matter 
 assemblies paves the way toward mechanically responsive systems whose 
 extreme structure  property complexity closely resembles paradigms 
 observed in nature .
this resemblance has fueled the investigation of smart systems in 
 the context of artificial muscles and biomedical implants as well 
 as textiles and nanopositioners among others . in this scenario , not only can mechanical responses and shape memory 
 effects be programmed within native molecular structures of polymers , 
 upon adequate processing , but also they can be suitably combined with 
 additional functionalities such as permeability , dielectricity , or 
 optical transparency . the wealthy 
 synthesis - property space portrayed by mechanically 
 responsive systems
indeed , the inherent molecular architecture flexibility in polymers 
 allows filler addition , which can , in turn , also feature functionalities 
 of their own .
this is the case of carbon nanotubes ( cnts ) , with remarkable properties inclusive of mechanical 
 actuation in ionic solutions and upon suitable bonding 
 within polymeric matrices
. synthesis / property 
 relations are currently the subject of much study in the context of 
 structural and active applications , which 
 extend across the materials space from liquid - crystal elastomers to cnts and graphene polymer composites .
finally , the 
 application space associated with mechanically responsive systems 
 is also augmented by the possibility of multiple stimuli suited to 
 trigger actuation , inclusive of electrical , thermal , and optical , promising engineering flexibility toward 
 device design .
expectation toward deployment of smart devices 
 based on smart composites 
 is mostly hindered by a lack of understanding of the underlying mechanisms .
filler dispersion and alignment are often cited 
 as necessary conditions for effective or enhanced actuation .
precisely , one of the greater challenges in the fabrication of active 
 polymer nanocomposites is filler dispersion , commonly attempted through 
 sonication processes or by addition of 
 chemical dispersants .
in addition , straining 
 procedures have been designed to promote carbon nanotube alignment , 
 possibly leading to mechanical - energy storage , in a variety of polymeric 
 composites .
recently , a photothermal 
 mechanical response has been observed 
 in ethylene vinyl - acetate nanocomposites .
eva is a low - cost polymer that has found a wide range of structural 
 applications and has been successfully used as a reversible shape - memory 
 polymer ( smp ) .
large mechanical responses 
 were favored by the addition of mwcnt fillers to eva , amenable to 
 enhanced light absorption and thermal capacity , yielding strains on 
 the order of 115% , with sharp onset responses , and relaxation 
 times on the order of tens of seconds . intriguingly , shape - memory 
 polymers and composites containing elastomeric matrix with carbon 
 nanotubes can either contract or expand upon stimulation , depending 
 on prestrain .
the extent of prestrain bears 
 important consequences in these material systems , and the mechanism 
 responsible for the bimodal behavior ( contraction or expansion ) with 
 the applied prestrain is not understood at present .
indeed , bimodal and reversible actuation of mwnt / elastomer 
 nanocomposites induced by infrared irradiation was first reported 
 in pdms|mwcnt composites . in this scheme , 

the contracting scenario has been modeled as 
 light being rapidly absorbed by cnts , producing a thermal perturbation 
 to the surrounding polymeric chains , affecting the prestrained distribution . 

overall , a fundamental understanding of photothermal actuation 
 down to the molecular level is still missing , and the proposed models 
 are yet to be verified through direct experimental data .
in particular , 
 the molecular conformation between cnts and polymer chains is the 
 subject of much discussion in nanocomposites , offering insights into 
 physical adsorption mechanisms , conformational 
 configurations , and nucleation dynamics .
however , details of molecular conformation in 
 smart composites have not been explored to the best of our knowledge . in this study , the thermo - active behavior of eva | mwcnt composites 
 was examined by in situ near - edge x - ray absorption fine - structure 
 ( nexafs ) spectroscopy .
nanocomposites described in this 
 article , with prestrains on the order of 50% , belong to the contractive 
 bimodal response scenario . in this scheme , 

nexafs has proved to be an invaluable tool in the study of 
 polymers and cnts , providing a wealth of information from alignment 
 of polymer chains and cnts to surface functional groups on cnts and bond 
 hybridization .
the beamline of 
 choice for these experiments was the national institute 
 of standards and technology ( nist ) u7a beamline at the national synchrotron 
 light source ( nsls ) at brookhaven national laboratory ( bnl ) .
in addition , 
 it is uniquely equipped for in situ experiments , with loading chambers 
 readily available for thermal , electrical , and optical in situ measurements . 

the goal of this approach was to use nexas for the exploration of 
 mechanisms behind thermal actuation by probing into the fine structure 
 of eva|cnt composites .
it is worth emphasizing that clarification 
 on the nature of filler  matrix bonding in the structural and 
 smart composites will be of paramount importance to both communities 
 given the scarcity of nexafs interfacial studies in polymer nanocomposites . to our knowledge , this is the first attempt to measure fine electronic 
 structure in the context of smart composites .
commercial ethylene vinyl 
 acetate copolymer eva ( evatane 28 - 25 , arkema , france ) containing 28 
 wt % of vinyl acetate and chloroform ( chcl3 p.a . ,
the purity of mwcnt exceeds 95 wt % , 
 the outside diameter distribution is between 60 and 100 nm , the cnt 
 lengths are between 5 and 15 m , and the surface area is 64 
 m / g .
cholesteryl 1-pyrenecarboxylate ( pychol ) was used 
 as a compatibilizer , and synthesis details are provided elsewhere .
following 
 previously reported procedures , a mwcnt / pychol 
 weight ratio of 1:5 was used ( 0.7 wt % ; 70 mg mwcnts and 350 mg of 
 cholesteryl pyrenecarboxylate ( pychol ) ) . upon dispersion in 100 ml 
 of chloroform ,
the solution was sonicated for 1 h under magnetic stirring 
 with a hielscher 400 s sonicator at an amplitude of 20% ( 35 
 m , 60 w / cm ) and a duty cycle of 100% .
the noncovalent 
 modification of mwcnts implies the promotion of  
 interactions between pyrene groups and graphitic walls of mwcnts . 

cholesteryl groups were designed to prevent  
 interactions among adjacent mwcnts , avoiding agglomeration and assisting 
 dispersion .
schematic featuring molecular structures of ( a ) eva and ( b ) proposed 
 mwcnt | pychol interaction . upon sonication , 10 g of eva ( evatane 28 - 25 ) was added , and 
 the 
 final solution was magnetically stirred for several hours .
the solution 
 was subsequently poured into a teflon - coated petri dish and dried 
 at laboratory temperature for 12 h prior to gradual oven drying at 
 40 , 60 , and 70 c for several hours .
finally , composites were 
 placed in a vacuum oven for 6 h at 70 c .
composite films were 
 then compression - molded in a laboratory press ( fontijne sra-100 , the 
 netherlands ) for 15 min under a pressure of 2.4 mpa and a temperature 
 of 80 c .
the pressed composite samples were strained up to 50% 
 at 50 c for 20 min using a custom - made stretching apparatus . 

subsequently , the stretched samples were cooled in ice water to fix 
 the orientation of cnts .
two eva|mwcnt ( 0.7 wt % ) were prepared following 
 this procedure : one with no prestrain and one with 50% prestrain . 

pristine pychol , eva , and methanol - dispersed cnts were also used for 
 comparison .
untreated mwcnts were characterized using 
 a jeol jem-2010f field - emission operated at 200 kv and a jeol jem - arm200f 
 electron microscope .
stem images were simultaneously recorded in both 
 the high - angle anular dark - field ( haadf ) and bright - field ( bf ) modes 
 at 80 kv .
probe correction was performed with a ceos corrector obtaining 
 a 12-fold ronchigram with a flat area of 40 mrad .
images were 
 registered with a condenser lens aperture of 30 m ( convergence 
 angle 25 mrad ) , and the haadf collection angle ranged from 45 to 180 
 mrad .
the spot size used was 35 pa . polarized raman 
 spectroscopy was conducted on a 161b renishaw system equipped with 
 a 632 he 
near - edge x - ray absorption fine - structure spectroscopy 
 probes transitions of excited 1s core electrons to unoccupied ( bound 
 or continuum ) states , emitting a photoelectron . upon core  hole 
 decay ,
released energy promotes two processes : fluorescence emission 
 ( bulk sensitive ) or auger electron emission ( surface sensitive , partial 
 electron yield mode ) , as seen in figure 2 . 
 on observing these processes , information on the electronic structure , 
 chemical environment , and orientation of molecules
is derived . when compared with other techniques , such as 
 infrared spectroscopy ( ir ) , nexafs places fewer demands on the physical 
 characteristics of substrates ( such as transparency ) , which is advantageous 
 for in situ analysis .
carbon k - edge spectra were collected 
 at u7a ( nsls - bnl ) in partial 
 electron yield ( pey ) mode using a horizontally polarized beam . a toroidal 
 spherical grating monochromator with 600 lines / mm and slits opening 
 of 30 m  30 m provided an energy resolution of 
 0.1 ev at the carbon edge .
the resulting beam 
 size is 2 mm  2 mm , and given the chemistry of the composite , 
 we find that 1.5 in a million monomeric units are vinyl acetate groups 
 in proximity to the mwcnt , which add up to 3  10 vinyl acetate groups surrounding cnts within the sampled volume .
for in situ temperature studies , samples were secured with a vacuum - safe 
 epoxy ( torr - seal ) on a tantalum metal plate that was inserted on a 
 customized heating stage .
the stage angle was controlled by a goniometer , 
 and the temperature was controlled by a voltage supply .
genzer and co - workers published more exhaustive details on the in 
 situ capabilities of the beamline .
the 
 prestrained composite was mounted with the strain direction perpendicular 
 to the beam polarization .
spectra were acquired at magic angle by 
 sequentially running macros that controlled the coordinates of the 
 stage and specified the energy parameters .
the temperature sequence 
 was 30 , 35 , 38 , 42 , and 48 c .
the pey signals were normalized 
 to the incident beam intensity 
 using the photoemission signal from a freshly evaporated au mesh located 
 along the incident beam path .
the spectra were energy - calibrated using 
 the photoemission current from an amourphous carbon mesh also located 
 along the path of the incident beam .
the 
 nexafs spectrum of pristine eva ( figure 3 ) 
 presents three prominent features with an additional low - intensity 
 feature at 284.5 ev , indicative of c = c groups .
this is taken 
 to be contamination , as no such group exists in pristine eva .
the 
 emission at 287.5 ev is assigned to 1s  *c  h , 288.4 ev corresponds to 1s  *c = o from vinyl acetate , and the broader resonance at 292.4 ev arises 
 from 1s  *c  c .
carbon k - edge nexafs 
 spectra acquired at magic angle and room temperature 
 of ( bottom to top ) pristine eva , pychol , eva/0.7 wt % mwcnt composite , 
 and prestrained eva/0.7 wt % mwcnt composite .
the nexafs spectrum of mwcnts shows prominent emissions at 
 285 , 
 287.6 , and 292 ev , attributed to c = c  * , c  h  * , 
 and c  c  * , respectively , with c = o  * and 
 c  o  * in the range of 288 to 289 ev . the cnt spectra are in good agreement with reported spectra 
 of unpurified mwcnts , showing c = o 
  * , possibly associated with native defects along graphitic 
 walls , as well as an unexpectedly large 
 contribution to the c  h  * and c  o  * emissions , 
 similar to those reported in ozone - treated cnts .
the possibility of residual species on untreated cnts was 
 investigated by stem ( figure 4 ) .
low - magnification 
 stem images show undulated external layers of 
 cnts ( figure 4 , left ) .
this pattern transfers 
 throughout at least a few graphitic layers through the inner tube , 
 suggesting structural defects .
local high crystallinity of graphitic 
 walls ( middle section in figure 4 , right ) coexists 
 with structural defects , as seen in the upper left region in figure 4 , right .
the sharp atomic contrast is convoluted 
 by the presence of an uneven layer that we tentatively attribute to 
 residual species as discussed above .
recently , in situ aberration - corrected 
 tem studies explored chemistry dynamics during cnt oxidation , where 
 a similar carbonaceous layer has been identified .
low - magnification stem image showing uneven external graphitic 
 layers of mwcnts with bends along the wall ( left ) and residual species 
 at higher magnification ( right ) .
similar features to eva are observed in the pychol nexafs 
 spectrum 
 ( figure 3 ) given the similar chemical species 
 in both pychol and eva .
the 1s  *c = c transitions are responsible for the peaks at 284.2 and 285 ev .
the 
 different energies are indicative of different chemical environments 
 for c = c groups .
the 1s  *c  h transitions are present at 287.7 ev , and at 289.1 ev , emissions 
 are observed from the 1s  *c = o transitions , 
 which are less intense that in eva because there is only one c = o 
 group per pychol molecule .
intensities associated with the 1s 
  *c  h , 1s  *c = o , and
1s  *c  c transitions are present 
 in the unstrained composite 
 ( figure 3 ) at the same photon energies and 
 with similar intensities as in pristine eva
 
 *c = c of sp - hybridized carbon 
 atoms in nanotubes appear at 284.7 ev . at slightly lower energies , 
 the 1s  *c = c transitions from pychol 

the main contributor to the 1s  
 *c = o transitions is eva , whereas pychol and 
 cnts dominate the 1s  *c = c transitions . 

all of the components contribute to the 1s  *c  h and 1s  *c  c transitions .
the dual 
 functionality of the pychol structure is conducive to the dispersion 
 of cnts ( prevented through the action of cholesteryl groups ) and to 
 the bonding of eva polymers to cnts .
indeed , pyrene groups in pychol 
 molecules attached to graphitic walls through  
 interactions provide a similar environment to pristine cnts for eva 
 molecules to bond .
remarkably , upon stretching , transitions 
 from the prestrained composite 
 produce a significantly different spectrum ( figure 3 ) .
the intensity at 288.4 ev from c = o groups shows 
 a 50% increase , causing the 287.7 ev feature to appear as a shoulder . 

straining is likely to promote alignment of polymer chains , conferring 
 a preferred distribution of the c = o groups .
the distinct double 1s  *c = c transitions observed prior to stretching now appear convoluted as 
 one feature .
we hypothesize that straining has promoted 
 alignment of cnts , yielding a higher amount of  * vectors ( normal 
 to the cnt surface ) parallel to the incoming electric field ( whose 
 polarization is parallel to the cnt surface ) .
the resulting overlap 
 between the polarized electric field and the  system is decreased 
 owing to the orbital geometry of the  system , thus diminishing 
 the 1s  *c = c intensity observed by 
 nexafs . by conducting polarized raman spectroscopy along the 
 main axis 
 orientation before and after strain , an indication of achieved alignment 
 can be provided .
indeed , monitoring of 
 d and g band intensities from cnts in a matrix whose straining direction 
 is oriented parallel (  ) and perpendicular (  ) to the 
 laser polarization is commonly performed to address the degree of 
 filler alignment upon straining .
figure 5 shows 
 raman spectra from pristine drop - casted mwcnts as well as unstrained 
 and strained eva|mwcnt composites .
spectra from composites were acquired 
 along parallel and perpendicular orientations , with d and g bands 
 at 1336 and 1585 cm originating from mwcnts , as 
 seen in figure 5a , and commonly attributed 
 to disorder and graphitic crystallinity in cnts , respectively .
additional emissions are observed in spectra 
 from composites at 1302 and 1439 cm , which have 
 been attributed to ch2 twisting and scissoring vibrational 
 modes on eva .
raman spectrum from drop - casted 
 mwcnts on si ( a ) . polarized raman 
 spectra acquired at parallel ( solid line ) and perpendicular ( dashed 
 line ) configurations , as discussed in the text , from unstrained ( b ) 
 and strained ( c ) eva|mwcnt composites .
ratios of id/id and 
 ig/ig were both 1.1 in unstrained 
 composites , suggesting isotropy .
this degree of gained alignment 
 upon processing is on the same order of magnitude as that reported 
 by abbasi et al . when comparing cnt alignment derived from simple 
 compression molding versus microinjection compression molding .
furthermore , 
 ratios of ( id / ig) in unstrained and prestrained 
 composites were 1.35 and 1.32 , respectively , suggesting that the graphitic 
 structure of cnts has not been damaged upon straining .
spectra acquired at increasing temperature 
 registered increased 
 emissions from c = c  * and decreased emissions from c 
h 
  * , c = o  * , and c  o  * for all samples 
 ( figure 6 ) . in all cases ,
c  c  * 
 emissions experienced small variations with temperature . every material 
 in the system ( cnt , phycol , and eva ) contributes to c  c *. 
 the actual c  c  * bonds are at many large angles to the 
 average chain anisotropy , and chains are still flexible , so this orientational 
 bias averages out .
in addition , spherical symmetry from sigma bonds 
 affords a higher tolerance toward conformational perturbations .
this 
 is not the case for c = c  * , which is highly directional 
 and less accommodative of nonaxial deformations , often triggering 
 concerns on pyramidalization .
carbon k - edge 
 nexafs spectra of pristine eva as well as unstrained 
 and prestrained composites at various temperatures . decreased intensities of c  h and c = o 
 groups in pristine 
 eva suggest rearrangement of polymeric chains .
figure 6 also shows increasing intensities at 284.5 ev with increasing 
 temperature , possibly because of enhanced cross - linking of adsorbed 
 contaminants on the surface of pristine eva .
both  * c  h 
 and  * c = o intensities from the unstrained composite 
 ( figure 6 ) decrease with increasing temperature , 
 as was observed in pristine eva , although to a much lesser extent . 
 as cnts are assumed to be randomly ordered if no strain is applied 
 upon fabrication , an explanation for this scenario could revolve around 
 cnts acting as anchors for polymeric chains , preventing their rearrangement . 

in this scheme , cnts would either restrict or cancel mobility within 
 the polymeric system . both c = c  * intensities on the 
 unstrained composites increase with temperature . as discussed earlier , 

 * signal from pristine eva is produced by a 
 surfacial contamination layer , which increases with both irradiation 
 time and temperature , as expected .
the evolution of c = c  * 
 from both unstrained and strained eva is attributed to intrinsic phenomena 
 to the composites , as will be discussed in the next section .
decreasing trends from c = o  * intensities are enhanced 
 on prestrained composites . in this case , the c  h  * emissions 
 experience little variation with temperature .
* 
 intensities at 284.7 ev , contributed to mostly by cnt carbons and 
 pyrenes from pychol appended to graphitic wall , seem to increase with 
 temperature with some correlation .
in fact , pyrene groups at pychol 
 are likely to undergo the same deformations as cnt carbons .
it is 
 unclear if pychol provides a homogeneous coverage of cnts , and eva 
 molecules possibly bond to either pristine graphitic walls or pychol - modified 
 regions .
relative intensities of c = c  * and c = o 
  * 
 were computed to investigate trends of intensities with temperature . 

figure 7 shows relative changes in the c = c 
  * and c = o  * emissions with respect to intensities 
 at room temperature , calculated as ( ii irt)/irt , where the intensities are corrected for 
 the surficial c = c  * contamination contribution observed 
 in pristine eva .
it is observed that trends 
 of intensity variation ( tv ) from strained 
 composites are coupled , described by eq 1 , by 
 a factor of 10 ( xstrained = 10 ) .
furthermore , 
 unstrained composites show weaker coupling by a factor of 5 ( xunstrained = 5 ) .
the implications of this result 
 will be discussed in the next section because both chemical groups 
 could be involved in temperature - driven dynamic effects on eva composites.1 relative change in the c = c and c = o emissions 
 in composites 
 with temperature .
in addition , the prestrained 
 composite shows a greater sensitivity 
 to temperature ( yc = c = 7 and yc = o = 4 in eq 2 ) , which is consistent with prior cnt alignment / actuation assumptions .
results of these relations are summarized in table 1.2 indeed , prestraining is expected to align 
 both cnts and polymeric 
 chains uniaxially with the purpose of enhancing actuation . prestraining 
 above 10% favors contraction along the straining axis ; this has been confirmed in eva|mwcnts composites .
thermal conductivity is expected to increase 
 with aligned cnts , as they can provide more direct conductivity pathways 
 across the entire composite thickness . increased thermal conductivity on the overall system
however , aligned cnt and polymeric chains could also 
 present a different response to phonon coupling .
a number of scenarios could occur 
 that would lead to the trends shown in figure 7 . decreasing intensities of 1s  *c = o could be explained as conformational changes in polymeric chains
. 
 published work has reported on molecular orientation 
 as a function of temperature . to explore conformation further , nexafs 
 spectra acquired at glancing ( 30 ) , magic ( 55 ) , and normal 
 ( 90 ) incident angles , both at room temperature and at 48 c , 
 were compared ( figure 8)
. spectra acquired 
 at room temperature show a slight variation with incidence angle in 
 both the c = c and c = o  * emissions .
although the 
 variation is small , it suggests that prestraining the sample results 
 in some preferred orientation of cnts ( c = c  * 284.7 ev 
 emission ) but more significantly of the polymer chains ( 288.4 ev emission 
 corresponding to the 1s  *c = o transitions ) .
angular 
 dependence of prestrained eva/0.7% cnt nanocomposite at 
 room temperature and 48 c . the angular dependence associated with the 288.4
ev emission 
 is 
 no longer apparent in spectra recorded at 48 c , suggesting that 
 c = o groups no longer have a preferred orientation .
the 284.7 
 ev emission attributed to the 1s  *c = c transitions still shows an angular dependence .
this effect could 
 be explained by cnts experiencing torsional perturbations with increasing 
 temperature while maintaining their overall 
 uniaxial alignment .
polymer chains are likely to be wrapping around 
 cnts through noncovalent ch interactions , resulting in c = o groups from eva aligned at a preferred orientation 
 with respect to the polymer backbone ( figure 9 ) .
torsional straining on cnts could therefore break the uniform 
 orientation of c = o groups , as proposed in the model depicted 
 in figure 9 .
the lack of communication between 
 adjacent pychol molecules coating the outside wall of the mwcnt has 
 important consequences .
it inhibits the propagation of a perturbation 
 from the underlying cnt through the pyrene lattice in pychol .
in this 
 model , an isotropic distribution of c = o for pychol follows , 
 resulting in their omittance from both the discussion of thermal effects 
 as well as the schematic of figure 9 .
beam travels 
 along z axis with electric field 
 vector along x axis ( a ) .
proposed configuration of 
 ( left ) strained eva chain along a cnt side view ( b ) and top view ( c ) 
 where the c = o bonds have a preferred orientation . upon cnt 

pychol 
 molecules have been omitted for simplicity . cross - sectional view of 
 cnt conformational changes upon actuation ( d ) .
a torsional deformation of carbon nanotubes could stretch 
 bonds , 
 leading to distortion of the  systems and making them less 
 stable .
indeed , increased cnt diameters have been previously proposed 
 as a result of increased bond lengths , leading to more populated antibonding 
 states , and defective regions , as seen 
 in figure 4 , could be especially prone to local 
 twisting and bending .
in addition , suitability 
 of cnts to buckle and twist upon temperature and mechanical deformation 
 opens the door to a number of conformational possibilities conducive 
 to an increase in c = c emissions .
indeed , torsional responses of cnts to stimuli is a recurrent 
 topic 
 in the literature of actuators , with baughman and co - workers having 
 recently proposed torsional cnts toward artificial muscles .
as vaia had previously summarized , cnt 
 buckling is necessary to explain contractions of up to 15% ( beyond 
 reported strains in cnts of 1 to 2% ) .
both a direct response from defective 
 cnts ( responding in an inhomogeneous fashion to light ) and inhomogeneity 
 of local strain within the surrounding matrix have been proposed . 
 in the second scenario
, the prestrained polymeric chains would absorb 
 thermal energy effectively supplied by cnts , and they would contract , 
 returning to a relaxed state .
the mechanisms prone to yielding a response 
 directly from the cnts , as proposed in the first scenario , are still 
 unclear .
the model in figure 9 proposes 
 a cnt upon 
 straining with a preferential uniaxial configuration and with c = o 
 groups in eva also at a preferred orientation in the xy plane ( figure 9a , b ) , consistent with the 
 early argument .
however , synthesis and mechanical prestrain have not 
 fully aligned the cnt either on the xy plane ( as 
 seen by the presence of kinks and bends ) or on the xz plane ( figure 9d , left ) . in this configuration , 

 systems in the highlighted rectangular region ( figure 9d , left ) would not be accessible to the beam and 
 therefore would not contribute to the 1s  *c = c intensity .
upon increased temperature , cnts could experience 
 additional torsion 
 and buckling that could result in increased 
 uniaxial symmetry ( figure 9d , right ) , where 
 a higher density of  systems are accessible to the beam and 
 therefore contribute to the registered 1s  *c = c intensity .
this effect could be promoting the observed increase 
 in c = c emission intensity , as more transitions to *c = c are now accessible to the beam .
intensity 
 variations in 1s  *c = c and 1s  
 *c = o possibly result from 
 conformational effects . because cnts and polymer chains are likely 
 to be connected through ch interactions other than 
 those remaining intact upon conformation ( through the examined temperature 
 range ) , it is not surprising that both emissions are indeed related . 

the implications of a 10-fold correlation in strained composites , 
 however , are unclear , as a correlation between two geometrical conformations 
 involving c = c and c = o bonds would need to be established . 

clarification on the chemical origin of the observed 1s  *c = o signal is important to visualize a realistic actuation 
 model .
pychol molecules are latched to cnts by noncovalent  
 interactions , as discussed earlier .
phenyl groups act as a coating 
 layer to cnts whose elements are not interconnected .
the c = o 
 group is bonded directly to the phenyls ; therefore , because of the 
 reduced size of phenyls and their lack of connectivity , any conformational 
 movement of cnts is not going to translate into a correlated conformation 
 of the c = o groups .
variations of  * c = o will be 
 anisotropic , and their evolution with temperature would not correlate 
 with  * c = c , as seen in the nexafs spectra .
the lack 
 of connectivity between adjacent phenyl groups , as discussed above , 
 ultimately renders the c = o groups in pychol inefficient to 
 flag conformation of cnts and does not contribute significantly to 
 the observed variation in the 1s  *c = o signal .
it is worth highlighting that mwcnt in the pdms composite 
 did not use compatibilizers to disperse cnts , yet bimodal photoactuation 
 was reported .
moreover , cnt torsion 
 and expansion effects have been proposed 
 as key mechanisms behind mechanical actuation , and our findings are consistent with those models .
mechanical actuation 
 mechanisms were tentatively modeled first as rigid nanotubes suffering 
 orientational order imposed by uniaxially applied strain .
more recently , smart behavior in hydrated media 
 is seemingly related to polarization effects .
this scheme involved polarization at the polymer / cnt interface because 
 of excitation of cnts by incident light .
resulting electric fields 
 could then promote migration of hydrogen ions and water molecules 
 toward polymer / filler interfaces as well as distortion of the  
 systems in cnts , resulting in mechanical bending .
the question 
 certainly arises pertaining the nature of external 
 stimulus on responses down to the molecular level .
actuation 
 studies at the macroscopic level seem to suggest that all electrical , 
 optical , and thermal external stimuli yield an internally induced 
 thermal actuation .
the basis for this 
 phenomenology is founded on the suitability of cnts to modify thermally 
 their surrounding environments ( i.e. , a thermal perturbation is induced 
 locally on surrounding polymer / elastomer either by melting or by the contraction / relaxation of the polymeric 
 chains , depending on prestraining history ) . further supporting this hypothesis , hu et al .
recently emphasized the importance of internal thermal 
 actuation in polymer nanocomposites , where cnts are believed to function 
 as nanoantennas , converting external stimulus to strategically localized 
 thermal perturbations in the composite .
the availability of simultaneous 
 cooling and heating capabilities for in situ nexafs could offer further 
 insight into the nature of these thermal effects .
if cnts indeed 
 convert external light into localized heat other 
 than actuation kinetics , then dynamic effects observed by nexafs during 
 external heating are likely to be comparable to those during external 
 heating reported by czanikova et al . in 
 which
phonon coupling could be inducing both cnt 
 torsional effects and polymer conformation at the molecular level . 

how this phenomenology would produce a macroscopic response is in 
 need of further investigation .
the thermal - induced effects promoting molecular conformation on 
 eva | cnt composites were studied by nexafs .
our results show inverse 
 variations in intensity between the emissions at 284.8 and 288.5 ev 
 with temperature , associated with the 1s  *c = c and 1s  *c = o transitions .
intensities 
 in prestrained composites are linearly correlated by a factor of 10 
 across the investigated temperature range .
these findings suggest 
 strong conformational coupling between c = c from cnts and c = o 
 from eva , possibly involving cnt torsional effects .
both a direct 
 response from defective cnts ( responding in an inhomogeneous fashion 
 to light ) and inhomogeneity of local strain within the surrounding 
 matrix have been proposed .
the observed dynamics at the molecular 
 level could contribute to the macroscopically observed thermal actuation .
 nexafs 
 spectroscopy was used to investigate the temperature dependence 
 of thermally active ethylene - vinyl acetate | multiwall carbon nanotube 
 ( eva|mwcnt ) films . 
 the data shows systematic variations of intensities 
 with increasing temperature . 
 molecular orbital assignment of interplaying 
 intensities identified the 1s  *c = c and 1s  *c = o transitions as the 
 main actors during temperature variation . 
 furthermore , enhanced near - edge 
 interplay was observed in prestrained composites . because macroscopic 
 observations confirmed enhanced thermal - mechanical actuation in prestrained 
 composites , our findings suggest that the interplay of c = c 
 and c = o  orbitals may be instrumental to actuation . 
Introduction
Experimental
Methods
Results and Discussion
Conclusions