Article pubs.acs.org/JPCC
Interface Dynamics in Strained Polymer Nanocomposites: Stick−Slip Wrapping as a Prelude to Mechanical Backbone Twisting Derived from Sonication-Induced Amorphization Allen D. Winter,† Klaudia Czaniková,‡ Eduardo Larios,§,# Vladimir Vishnyakov,∥ Cherno Jaye,⊥ Daniel A. Fischer,⊥ Mária Omastová,‡ and Eva M. Campo*,†,§ †
School of Electronic Engineering, Bangor University, Bangor, LL57 1UT, U.K. Polymer Institute, Slovak Academy of Sciences, Bratislava, SK 84541, Slovak Republic § Department of Physics and Astronomy, University of Texas at San Antonio, San Antonio, Texas, United States ∥ School of Computing and Engineering, University of Huddersfield, Huddersfield, HD1 3DH, U.K. ⊥ National Institute of Standards and Technology, Gaithersburg, Maryland 20899, United States ‡
ABSTRACT: In this paper, we examine the effects of excessive sonication during surfactant-assisted multiwall carbon nanotubes (MWCNT) dispersion in ethylene vinyl acetate (EVA) by way of monitoring molecular arrangements upon progressive straining. Aberration-corrected transmission electron microscopy confirms the structural damage on the graphitic layers upon prolonged sonication. The resulting lack on MWCNT alignment is shown by atomic force microscopy. Further, molecular interface dynamics in progressively strained EVA|MWCNT composites have been studied through Raman and NEXAFS spectroscopies. NEXAFS spectra have identified graphitic amorphization and further Cvacancy rehybridization by way of hydrogen passivation as the damage mechanism to the graphitic structure upon sonication. In this scheme, MWCNTs did not align despite the range of composite strains discussed due to stick and slip dynamics of surrounding EVA. Ultimately, damaged MWCNTs rendered the necessary dispersant π−π interactions suboptimal and insufficient for nanomechanically interlocked polymer−filler interactions. Remarkably, upon large strains, polymer chains are seen to unlatch from the MWCNT and undergo mechanically induced backbone twisting. The possibility of mechanically induced backbone twisting might offer alternative processing routes in photovoltaic systems, where chemically induced conjugated backbone twisting yields increased power conversion efficiency.
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INTRODUCTION The possibility of propagating CNTs extraordinary mechanical, thermal, electronic and optical properties to bulk systems has encouraged innovative applications in the realms of both physical and biological sciences.1−6 In a nanocomposite, the propagative nature of such highly sought-after properties requires of intricate relationships at the polymer-nanofiller interface to promote overall system responses.7 Despite lower bond energies, it is worth emphasizing the crucial role that noncovalent bonding plays in the physical properties of soft matter; with recent reports on hydrogenbond engineering of polymer blends yielding increased thermal conductivities of an order of magnitude.8 However, few studies address processing and chemistry on noncovalent bonding dynamics at the molecular level in nanocomposites.9 Moreover, strain and deformation of CNTs has been extensively modeled and simulated but, the combination of local atomistics linking chemistry and processing with bonding structure, is a complex mesoscale problem that remains unsolved.10,11 It is crucial to correlate all three elements to provide a structure−property paradigm toward the optimization of nanocomposites. © 2015 American Chemical Society
Mechanical actuation is one of the most intriguing responses in nanocomposites, paving the way to a number of applications from the micro to the macro scales.12,13 In this context, we have pioneered in situ near-edge X-ray absorption fine structure (NEXAFS) spectroscopy studies of thermo-active ethylene− vinyl acetate multiwall carbon nanotube (EVA | MWCNT) composites, yielding an actuation model that validated the early proposal of active CNT torsion upon excitation.14,15 Further, nanocomposite relaxation elucidated the metastable character of energy landscapes in π−π states. This instability was promoted by both sonication and strain; yielding aged composites that were no longer mechanically active. Remarkably, all processing steps contributed to π−π degradation, ultimately leading to failure.16 Indeed, straining effectively alters interactions between filler, dispersant, and matrix in the context of noncovalent bonding dynamics,8,14,16 and our earlier work highlighted the need to understand how the combination of Received: May 19, 2015 Revised: August 9, 2015 Published: August 10, 2015 20091
DOI: 10.1021/acs.jpcc.5b04773 J. Phys. Chem. C 2015, 119, 20091−20099
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made apparatus was used to produce composites stretched to 10%, 20%, 40%, 50%, and 100%, which were then placed in an oven at 50 °C for 20 min. Composites were finally cooled in ice water to fix the orientation of CNTs in the polymeric matrix. Characterization. NEXAFS spectra were collected at the Carbon K-edge in partial electron yield mode shortly after composite fabrication, at U7A beamline, NSLS, BNL, using a horizontally polarized beam of 2 mm in diameter. Charging of the polymeric samples was prevented by using a flood gun during acquisition. Energy resolution for NEXAFS acquisition was 0.1 eV. NEXAFS spectra were then normalized using the Athena software. Angle-resolved NEXAFS spectra included acquisitions at 20°, or “glancing”; 55°, “magic”; and 90°, “normal”. Aberration-corrected TEM was performed 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 annular darkfield (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 μm (convergence angle 25 mrad), and the HAADF collection angle ranged from 45 to 180 rad. The spot size used was ∼35 pA. Composites were also studied under Raman spectroscopy and AFM microscopy. Raman characterization was performed using a Renishaw Ramascope with a RL633 HeNe laser and Renishaw inVia microscope and an argon laser. Spectra were acquired with strain direction parallel and perpendicular to beam polarization vector. Spectra were later backgroundsubtracted, and the D* band positions were found by leastsquares fitting Voigt functions (WiRE 2.0 Renishaw, Origin Pro 2015). AFM characterization was achieved with a Digital Instruments Dimension 3100 atomic force microscope. Height images were acquired simultaneously while operating in tapping mode under ambient conditions.
sonication and straining affected molecular interactions at the interface between filler and matrix. Indeed, under adequate sonication conditions, nanocomposite straining prior to full cross-linking, lead to the preferential alignment of EVA carboxyl groups along CNTs,14 flagging CNT torsion upon opto-thermal actuation. In this paper, we further our initial investigations of the EVA| MWCNT system by examining the effects of excessive sonication during MWCNT dispersion in EVA, which rendered a nonactuating nanocomposite. This examination was conducted by way of monitoring molecular arrangements during progressive straining. We propose that the resulting molecular map describes the CNT−polymer conformation as it evolves from stick−slip wrapping to backbone twisting upon polymer desorption. This molecular arrangement clarifies the importance of processing and, in particular, the impact of sonication at polymer/CNT interfaces as well as establishing a structure/ property correlation. This study offers a pioneer qualitative description at the molecular level on the effects of sonication along graphitic nanostructures and the subsequent frustrated polymer interactions, paving the way toward an adequate processing-property tandem.
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EXPERIMENTAL METHODS Composite Fabrication. A dispersion of 0.0025g MWCNTs (Nanostructured & Amorphous Materials, Inc., Houston, TX, USA, purity higher than 95%, a surface area of 64 m2/g, and an outside diameter in the range of 60−100 nm with lengths between 5 and 15 μm) and 0.0125 g of cholesteryl pyrenecarboxylate (PyChol) compatibilizer in 25 mL of chloroform (versus 100 mL priorly used)14 was sonicated for 1 h with a Hielscher 400 S sonicator at amplitude of 20% (∼35 μm, ∼ 60 W/cm3) with 100% duty cycle, under magnetic stirring. The original acoustic energy density (“mild” sonication), that had been priorly conducive to nanomechanical interlocking through π−π interactions and thermo-mechanical behavior (Figure 1),14,16 was increased 4-fold for the purpose of this study to yield an aggressive sonication scheme. Synthesis of PyChol compatibilizer can be found elsewhere.17
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RESULTS AND DISCUSSION The effects of sonication in CNTs have been thoroughly studied, and it is well established that excessive sonication will yield scission18−20 and amorphization21 through complex dislocation dynamics propagating throughout the graphitic structure.22,23 Both untreated and sonicated MWCNTs are shown in Figure 2, parts a and b, respectively. Habitual carbonaceous impurities from growth are observed in untreated MWCNTs (highlighted by arrows in Figure 2a) covering the outer graphitic layer,24,25 as well as dislocations (Figure 2a).22 Upon sonication, amorphous impurities are removed and sidewall damage is visible on the outer graphitic layer (arrows in Figure 2b). Remarkably, damage is also observed on the internal walls upon sonication.23 The presence of internal damage suggests that cavitation effects propagate through the multiple graphitic layers. Similarly, internal oxidation had been observed by in situ environmental TEM,24 as well as glideclimb dislocation interactions upon thermal treatment.22 With all, structural damage will condition π-systems in CNTs, clearly altering subsequent PyChol π−π functionalization as well as bonding dynamics to EVA.23 Effects of evolving chemical and structural environments,25 as well as tensile strains exhorted upon embedded MWCNTs, can be conveniently monitored by Raman spectroscopy.6,7,26,27 Raman spectrum of MWCNT is characterized by typical D, G
Figure 1. (a) Ethylene−vinyl acetate (EVA) polymer; (b) PyChol molecule (blue) appended to CNT wall through π−π interactions.
After sonication, 10 g of ethylene−vinyl acetate (EVA) with 28 wt % vinyl acetate content (Evatane 28−25) were added to the solution, which was then stirred for 3 h at 1200 rpm. The solution was subsequently poured into a Teflon-coated Petri dish and dried at room temperature for 12 h. The sample was then placed in an oven and gradually heated to 40, 60, and 70 °C for several hours. Additional drying was performed in a vacuum oven for 6 h at 70 °C. The composite foil was prepared by compression molding in a laboratory press (Fontijne SRA100, The Netherlands) for 15 min under a pressure of 2.4 MPa and a temperature of 80 °C. Mechanical stretching is typically used to enhance conformational arrangements at CNT−polymer interfaces.14 A custom20092
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Figure 2. TEM micrographs of (a) nonsonicated and (b) sonicated MWCNTs; (c) Raman spectra of unsonicated and aggressively sonicated MWCNTs; (d) NEXAFS spectra of unsonicated, moderately sonicated, and aggressively sonicated MWCNTs.
and D* bands at 1349.9, 1576.5, and 2695.7 cm−1 respectively. D bands give a measure of “disorder” in the graphitic lattice, though the community has not yet agreed on the precise origin; Lehman and co-workers, for instance, believe the signal arises purely from amorphous carbon, while others believe it originates from structural defects (e.g., heteroatoms, vacancies, heptagon-pentagon pairs, kinks) as well.28,29 Raman G bands correspond to tangential vibrations and give an indication of graphitization.28 This band in fact consists of two features: G+ corresponds to atomic vibrations along the tube axis, and Grelates to circumferential vibrations.28 The D* band is a second-order harmonic of the D band.28,30 Both Raman and NEXAFS spectra of pristine and sonicated MWCNTs have been provided. Raman spectra reveals an increased ID/IG ratio upon sonication, confirming the damaged graphitic character, as shown in TEM. Similarly, the evolution of the 1s → σ*C−O 1s → σ*C−H signals upon progressive sonication, reveals, first, a decreased intensity due to the detaching of carbonaceous residuals from growth, followed by an increased intensity upon further sonication, suggesting an increase on the H and O content. These findings are consistent with the proposed rehybridization of the previously graphitic, now amorphized surface, upon sonication, as seen in Figure 2b. Similarly, the initial increase of 1s → π*CC is attributed to the disentanglement of CNTs, that is rendering more CC groups available to the incident beam. Further decrease of this intensity, confirms the damaged nature of the π systems.
In Raman spectra of EVA|MWCNT nanocomposites, G and D* bands are observed at similar wavenumbers to pristine MWCNTs, but the D band is not present, or at least observable. Indeed, Jorio and co-workers reported D bands of 100 times weaker intensity than G bands in 50% of Raman spectra from a large number of characterized CNTs.29 It is possible that the already weak D band signal is attenuated as the Raman spectra in the present study were acquired from already prepared composites i.e. CNTs embedded in EVA polymer. Figure 3a shows the downshifted progression of D* emissions upon increased strain (up to 40%) associated with axial lengthening of C−C bond lengths in MWCNTs; indicative of a certain degree of load transfer.26,31,32 The trend also shows an oscillatory behavior (Figure 3b), attributed to “stick and slip”.6 In an optimized load-transfer configuration 1 cm−1 D* shifts result from 1% strain.27 However, 10−40% strain is needed here to produce 1 cm−1 D* shifts; suggesting load transfer is not optimal in the current EVA | MWCNT systems. On those lines, 50% strain yields a strongly blueshifted D* emission. Although typically associated with compression in load-conducive composites,27 in these suboptimal ensembles the blue-shifted D* emission suggests that a new chemical environment around the MWCNTs is unfolding at 50% strain, followed by a hydrostatic effect at 100%.6 Indeed, Schadler et al. had identified mechanisms through which tensile strain could overcome interactions of noncovalent 20093
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Figure 4. Polarized Raman spectra of G bands.
Abbasi and co-workers observed G∥/G⊥ ratios of 1 for randomly oriented CNTs and up to 1.75 for oriented filler arising from processing methods. With the exception of 10% strain, ratios (Table 1) do not show significant variation, suggesting no alignment progression with increasing strain. Table 1. G∥/G⊥ Ratios of Composites with Increasing Strain
Figure 3. (a) Raman D* band of strained composites, with fits shown by dashed lines; (b) shifts in D* band versus strain. The standard deviation is noted on the graphs, with the noted oscillatory behavior being reflective of stick and slip dynamics.
nature, such as those reported in the EVA|PyChol|MWCNT ensemble, i.e., CH−π between EVA and PyChol and π−π between PyChol and MWCNT.16,27 In this scheme, deterioration of noncovalent interactions would break EVA-PyCholMWCNT interconnectivity and hence compromise load transfer. Polarized Raman spectroscopy has been used to determine alignment of CNT filler within host materials.14,33 Raman spectra are measured with polarization vector parallel (G∥) and perpendicular (G⊥) to strain direction. Ratios of G∥/G⊥ of unity are indicative of isotropy, while increasing ratios indicate increasing alignment. Since the D band is not observed in these composites only the G band is used. Figure 4 shows polarized Raman spectra of composites in the region of the G band. Interestingly, only the G+ component is consistently observed. As mentioned above, the G+ band arises from atomic displacements along the CNT axis and is independent of diameter, whereas the G− band is the result of circumferential displacements and its frequency decreases with decreasing diameter.29 The intermittence of G- signal in composites could therefore be suggestive of CNT exfoliation due to sonication, where damage to the outer graphitic layer interrupts the effective CNT diameter, but still allows tangential G+ vibrations from inner layers. This sonication-induced damage was not observed in our earlier study, but the sonication volume here was reduced by a factor of 4.
strain (%)
G∥/G⊥
0 10 20 40 50 100
1.00 0.58 1.02 0.96 1.02 0.97
Although the ratio at 10% would suggest alignment perpendicular to the strain direction, this is unfeasible: another possibility may be poor graphitic quality. Lack of increased alignment may be due to poor π−π interactions between filler and compatibilizer as a result of CNT exfoliation. Damage to CNT walls would alter the pristine π systems around their circumference, preventing a thorough PyChol coverage. In this scheme, residual strain is communicated to CNTs, as observed in D* shifts, but communication between polymer matrix and filler is not conducive to an anisotropic configuration of CNTs. These findings support the notion of excessive damage to π− systems in CNTs as a result of sonication, ultimately rendering suboptimal π−π interactions between PyChol and MWCNT. Topographic atomic force microscopy (AFM) images are shown in Figure 5. The maps show needle-like structures with diameters in the micron range, suggesting MWCNT bundling. AFM maps do not show alignment regardless of strain, and in fact, MWCNTs remain bundled upon straining; resonating with the notion of frustrated MWCNTs and PyChol π−π interactions (Figure 1). Frustrated π−π interactions is a byproduct of excessive sonication-induced damage (Figure 2b), leading to poor load transfer (Figure 3a). 20094
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Figure 6. Angle-resolved C K-edge NEXAFS spectra of nanocomposites under various strains.
pyrene. Note that π*2 in PyChol appears at the same energy as 1s → π*CC in CNTs, as the chemical environments of carbon atoms in both CNT walls and pyrene groups are spectroscopically identical. The spectrum of EVA has three distinct peaks at 287.6, 288.6, and 293 eV, attributed to 1s → σ*C−H, 1s → π*CO, and 1s → σ*C−C respectively.14 Table 2 below correlates the observed emissions with the electronic transition as well as their molecular origin. This Table 2. Assignment of NEXAFS Electronic Transitions and Corresponding Molecular Origins Figure 5. AFM height profiles (40 μm regions) of composites with (i) 0%, (ii) 10%, (iii) 20%, (iv) 40%, (v) 50%, and (vi) 100% prestrains, where the strain direction is highlighted.
energy (eV)
Ultimately, damaged MWCNT prevent PyChol coverage and perpetuate bundling and misalignment despite large strains. This is in great contrast to the use of adequate sonication parameters that successfully promote dispersion and alignment and favor actuation.14,17 Raman spectroscopy has provided a description of strainpromoted variations on MWCNTs vibrational modes. In this section, the impact of strain on the conformational matrix will be addressed. NEXAFS spectroscopy offers fine detail of polymer conformation dynamics through high beam polarization and high chemical specificity,14,34,35 and has been widely deployed in the study of both nanocomposites and of CNTs, often in the context of energy materials.36,37 Figure 6 shows the C K-edge spectra of composites at glancing, magic and normal incidences. NEXAFS emissions from the individual components in EVA | MWCNTs are provided here toward the building block discussion. Spectra of CNTs typically consist of a pre-edge resonance at 285 eV arising from 1s → π*CC transitions from graphitic lattice, and a broader feature centered at ∼290 eV, attributed to 1s → σ*C−C. A series of intensities in the range of 287−289 eV are usually present, indicating hydrogen and oxygen impurities.38 PyChol is spectrally similar to MWCNTs, with additional 1s → π*CC transition at 284.3 eV, as the molecule has CC groups in two distinct chemical environments.14 In this study, these are denoted by π*1 at 284.3 eV, arising from cholesteryl groups, and π*2 at 285 eV, from
electronic transition π*CC π*CC σ*C−H π*CO
284.3 285.0 287.6 288.7
1s→ 1s→ 1s→ 1s→
290.0 293.0
1s→ σ*C−O 1s→ σ*C−C
molecular origin PyChol (cholesteryl) MWCNT, PyChol (pyrene) EVA, MWCNT (impurities), PyChol EVA (main contribution), MWCNT (impurities), PyChol EVA, MWCNT (impurities), PyChol EVA, MWCNT, PyChol
identification justifies the evolution of emission at 288.7 eV as the result of conforming polymer chains (those appended to CNTs) upon strain, by monitoring CO orientation. Remarkably, and as an exclusive result from processing, several spectral differences are found between aggressively sonicated composites (Figure 6) and the previously reported mildly sonicated systems.14 First, the intense 1s→ π*2 CC transitions in mild sonication, indicative of pristine π-systems, are significantly reduced as a consequence of aggressive sonication (Figure 6), in good agreement with weak Raman G bands as discussed above. Second, despite identical chemistries, aggressive sonication yields intense 1s → σ*C−O transitions from PyChol. Mild sonication, however, had yielded reduced σ* C−O intensities, suggesting more PyChol molecules become accessible to the incident beam upon aggressive sonication, i.e., are not latched to graphitic MWCNT.14,16 Third, a drastic increase of σ* C−H intensities in aggressively versus mildly sonicated systems strongly suggests an increase on the overall hydrogen content. Arguably, this increase is related to the passivation of dangling bonds resulting from carbon vacancies originated from excessive 20095
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regardless of synthetic approaches.9,14,45 Clearly, the dynamic intensity range at each isostrain correlates with the angular dependence of 1s → π*CO transitions, and is inversely proportional to the spread in the angular frequency distributions (Figure 8); that will enable the prediction of
sonication, in combination with available hydrogen from cavitation.18,39 With all, the notion that PyChol molecules are in fact not latched to the MWCNTs, is solidifying. Through our studies of EVA|MWCNT systems, we have found the angular dependence of CO groups (1s → π*CO at 288.6 eV) particularly revealing of conformational EVA behavior. These findings show the high capability of NEXAFS, efficiently probing into bond directionality. Angle-resolved intensity variations from these groups are attributed to conformation from polymeric chains surrounding MWCNTs exclusively, as intensities from the disordered matrix will average out.14 Furthermore, the statistical contribution of PyChol to this intensity is almost negligible compared to that of EVA, as there is only one CO group per PyChol molecule. Because of the mainly amorphous character of EVA,40 and the lack of phenyl groups in the molecular structure, straining of the pristine matrix would not provide appreciable alignment, and the lack of fillers as molecular anchors, would render the polymeric chains unconforming.41−43 Figure 7 shows the angle-resolved intensity interplay of 1s → π*CO transitions with applied strain. Here, the original 18 data
Figure 8. Angular frequency distributions of CO polymer groups from different composites.
Figure 7. Angle-strain-resolved π* CO resonance intensity map.
points were assembled into a 3D map linking strain (x axis) and incidence angle (y axis) with the recorded CO intensities (z axis, with values from Figure 6). The interpolations on both axis were done on a linear fashion, to complete a physically meaningful 3D plot. Indeed, different incident angles from the polarization vector in NEXAFS probe on the geometry of the bond,44 and the data represented here provides a reasonable description of the angular continuum (that has only been probed in three locations). The interpolation on the x axis is also satisfactory, since strain was highly probed at regular strain intervals between 0% and 50%, as well as 100%. The interpolation does not apply between 50 and 100% and is represented by the black box. Higher intensities are observed at 90° and lower at 20° for all strains; and given the π symmetry of the bond, this is indicative of CO groups predominantly positioned normal to the surface.44 This is unsurprising, as CO groups have demonstrated a preference to be oriented perpendicular to the surface as a result of intense carbonyl π* transitions in a number of systems, inclusive of EVA, PMMA and other blends,
Figure 9. Polymer unwrapping through “stick and slip” mechanism (0−40% strain), followed by backbone twisting (50−100% strain) in cross-sectional (left) and lateral view (right).
molecular orientations (Figure 9). In this scheme, a perfectly aligned system, for example a scenario where 100% of the population is aligned at 90°, would show maximum dynamic range in the 3-dimensional plot and its angular frequency distribution would be a delta function. Conversely, a perfectly random oriented system would show no intensity variation across the incidences (a dynamic range of 0), and the angular frequency distribution would be represented with a constant function. Figure 8 describes the isostrain cross-sectional cut-offs at the three measured angles in Figure 7 that have now been fitted to 20096
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Figure 10. Molecular interactions model. (a) Aggressive sonication induces damage to nanotube walls, hindering noncovalent π−π interactions with PyChol; (b) PyChol-MWCNT ensembles introduced to EVA matrix, but EVA chains mostly interact directly with nanotubes through CH-π interactions; (c) under small strains CH−π interactions begin to break, some residual strain is communicated to MWCNTs, but not enough to align them; (d) higher strains result in polymeric chains unlatching from nanotubes as well as backbone twisting.
MWCNTs with a finite number of points (Figure 9). Remarkably, higher overall intensities were recorded at 50% strain; highlighting the fact that more CO groups are becoming accessible to the beam, and populating all possible orientations. This observation can be explained by polymer chains unlatching from the CNTs, rendering CO groups available to the beam as backbones twist, also in agreement with a new chemical environment predicted by Raman spectroscopy at 50% strain. Upon further strains (100%) orientational order of pendant CO groups increases again (Figures 8 and 9), attributed to further mechanically induced backbone twisting. Interestingly, the twisted conformational arrangement, a byproduct on the quest to CNT alignment, has been the motor behind extensive pendant group functionalization in the context of photovoltaic devices, where ionization potentials are seen to increase upon molecular orbital overlap of conjugated backbones.48 It is worth highlighting that polymer chains have detached at strains above 40%, as a result, no stick and slip communication will develop between 50 and 100% strain. Regardless of the intermediate scenarios of polymeric chain configuration as a result of progressive strain beyond 50%, the final configuration at 100% strain reveal polymeric chains with some level of alignment of the CO groups, as seen in Figure 9. However, polymeric chains remain disconnected from the CNTs in this
Gaussians and mirror-symmetrized to enable statistical estimates on the angular distribution of the CO population. The resulting sigmas, specified in Figure 8, indicate that 17%, 19%, 16%, 26%, 12%, and 17% of the beam-accessible CO groups lie within 85° to 95° for strains of 0, 10, 20, 40, 50, and 100%. This progression is highly informative on the possible conformation scenarios between EVA and MWCNTs to accommodate increasing strain, and the proposed models are depicted in Figure 9. According to Baskaran, polymer wrapping of CNTs is a universal phenomenon, promoted by CH−π interactions.46 In fact, Monte Carlo simulations by Gurevitch and Srebnik showed that in equilibrium, polymer wrapping around CNTs resulted from blending.47 This situation in depicted in Figure 9 at 0% strain. Unsurprisingly, the oscillatory behavior on the CO population between 85° and 95° for strains up to 40% keeps a strong parallelism to the oscillatory behavior in D* band shifts (Figure 3b), attributed to “stick and slip” mechanics, unfolding a model where EVA chains are progressively unwrapping from CNTs, as shown in Figure 9. The oscillation dynamics in this process is likely a function of both MWCNT geometry and polymer properties, all favored by the highly damaged graphitic lattice. At 40% strain, the CO distribution is the most anisotropic, with more than 25% of the population heavily distributed around 90°, where EVA is anchored to the 20097
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regime, as shown by D* shifted Raman emissions in Figure 3b, where D* shifts remain positive for both 50% and 100% strains. This miscommunication with surrounding polymeric chains is also shown in AFM images in Figure 5, where CNTs did not align. We now combine key findings from Raman and NEXAFS spectroscopy into a model describing molecular interactions. PyChol molecules are expected to cover CNT walls during sonication, through π−π interactions.14 Excessive sonication energy, however, causes damage to the graphitic lattice ranging from partial exfoliation in outer layers to CNT scission.18,21,23,49 Although sonication damage of outer graphitic layers in MWCNTs will enable P orbitals of inner shells to be potentially accessible to PyChol bonding, π−π stacking is unlikely in such an abrupt graphitic structure (Figure 10a). Additionally, atoms lining extended vacancies in the graphitic structure undergo rehybridization to amorphous sp3 carbon (Figure 10a inset), as supported by decreased G and π*2 intensities from Raman and NEXAFS analysis. Extended vacancies also incorporate hydrogen as a passivation strategy (increased σ* C−H intensities), further breaking continuity of CNTs π-systems.39 These irregularities provide a defective interface for PyChol molecules to append to, ultimately hindering optimal π−π configurations and leading to EVA interacting directly with MWCNTs (Figure 10b). Under small applied strains (
