ARTICLE pubs.acs.org/JPCC
Evidence of Multiwall Carbon Nanotube Deformation Caused by Poly(3-hexylthiophene) Adhesion Michele Giulianini,*,† Eric R. Waclawik,‡ John M. Bell,† Maurizio De Crescenzi,§ Paola Castrucci,§ Manuela Scarselli,§ Marco Diociauti,|| Stefano Casciardi,^ and Nunzio Motta† School of Engineering Systems and ‡Discipline of Chemistry, Queensland University of Technology, 2 George St., Brisbane (QLD), 4001, Australia § Dipartimento di Fisica, Unita CNISM, Universita di Roma, “Tor Vergata” 00133, Roma, Italy Dipartimento di Tecnologie e Salute, Istituto Superiore di Sanita, 00161, Roma, Italy ^ Dipartimento di Igiene del Lavoro, ISPESL, 00040, Monte Porzio Catone, Italy
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bS Supporting Information ABSTRACT: We show that when a soft polymer like poly(3-hexyl-thiophene) wraps multiwall nanotubes by coiling around the main axis, a localized deformation of the nanotube structure is observed. High resolution transmission electron microscopy shows that radial compressions of about 4% can take place, and could possibly lead to larger interlayer distance between the nanotube inner walls and reduce the innermost nanotube radius. The mechanical stress due to the polymer presence was confirmed by Raman spectroscopic observation of a gradual upshift of the carbon nanotube G-band when the polymer content in the composites was progressively increased. Vibrational spectroscopy also indicates that charge transfer from the polymer to the nanotubes is responsible for a peak frequency relative downshift for high P3HT-content samples. Continuously acquired transmission electron microscopy images at rising temperature show the MWCNT elastic compression and relaxation due to polymer rearrangement on the nanotube surface.
1. INTRODUCTION Carbon nanotubes (CNTs) possess unique physical, mechanical, and electrical properties1-5 and since their discovery, the number of proposed applications for this advanced material has expanded greatly. Some examples include: photovoltaics,6 optics,7 electronics,8,9 photomechanics,10 and gas sensing.11 In addition, carbon nanotubes have frequently been used as fillers and modifiers to the physical properties of other materials and have been incorporated into polymers, 12 epoxies, 13 and ceramics14 for which structural reinforcement and improved mechanical properties have been demonstrated.14-16 Several studies have also shown how long chain molecules wrap around nanotubes17,18 and, as a related effect, the modification of carbon nanotube structural properties after the inclusion in polymers has been explored by various authors with some interesting results. For example, mechanical deformation19,20 and stress-induced fragmentation21 have been reported for multiwall carbon nanotubes (MWCNTs) in polymer matrixes. In this work, we provide compelling evidence of multiwall nanotube deformation due to polymer adhesion in mixtures of poly(3-hexylthiophene) (P3HT) and MWCNTs. First, high resolution transmission electron microscopy (HR-TEM) images of isolated MWCNTs wrapped by P3HT strands provided r 2011 American Chemical Society
evidence of significant MWCNT constriction at pressure points that coincided with P3HT coil locations. Indirect confirmation of these effects was obtained by studying the MWCNT/P3HT mixtures Raman spectra which displayed a considerable increase of the G-band peak frequency when the P3HT content in the compound was raised. Finally, direct HR-TEM imaging under electron beam-induced heating showed the effect of polymer compression on the nanotube structure due to P3HT selfassisted rearrangement.
2. EXPERIMENTAL METHODS Poly(3-hexylthiophene) was employed in this work because its π-conjugated structure strongly interacts with the carbon nanotube outer-walls which provide an electron-rich surface for the polymer adhesion. Regioregular P3HT was purchased from Aldrich (p.n. 445703, Mn ≈ 64 000 g mol-1, head-to-tail regiospecific conformation higher than 98.5%). Multiwall carbon nanotubes have been purchased from Aldrich (p.n. 636 495, carbon content reported higher than 95%) and have been Received: January 2, 2011 Revised: February 12, 2011 Published: March 07, 2011 6324
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The Journal of Physical Chemistry C produced by manufacturer by chemical vapor deposition (CVD) method. Thermogravimetric analysis, available as Supporting Information, was used to estimate the carbon content to be higher than 98%. Initially, MWCNTs were suspended in chloroform (CF) or dichlorobenzene (DCB) and dispersed by using ultrasonic agitation. Poly(3-hexylthiophene) was dispersed in the same solvents but in higher concentrations (10-20 mg mL-1). The relative content of the two materials in the composites, measured by the weight/weight content (w/w), was achieved by controlling the volume of solution added. High resolution transmission electron microscopy was performed with an FEI TECNAI 12 G2 Twin operated at an accelerating voltage of 120 kV (FEI Company, Hillsboro, Oregon, USA). Samples were deposited by solution drop casting on a gold grid (mesh 1000) without using any supporting amorphous carbon film. Scanning tunneling microscopy (STM) analysis was performed by using an Omicron STM apparatus (Omicron, Taunusstein, Germany) and acquiring the data with Omicron software (Omicron Scala Pro). Samples were deposited by solution drop casting on freshly cleaved highly oriented pyrolitic graphite (HOPG). Experiments were performed in constant current mode at room temperature under ultrahigh vacuum conditions (base pressure 8 10-11 mbar) using electrochemically etched tungsten tips (99.9% purity, Goodfellow GmbH, Friedberg, Germany). Raman spectroscopy samples were deposited by drop casting. Samples were excited with a few milliwatts of 532 nm diode laser light, with circular polarization and slightly defocused to minimize laser-induced sample degradation. Spectra were collected from 200 to 4000 cm-1 calibrated against the 520.5 cm-1 peak of a silicon wafer. Thermal annealing was performed in a temperature-controlled oven, for one hour at 120 C. Samples were then slowly cooled to room temperature before performing the Raman analysis. Experimental curve fittings have been performed by using the Gauss-Lorentz cross product function, minimizing the number of oscillators used.
3. RESULTS AND DISCUSSION Figure 1 reports a collection of HR-TEM images, showing the P3HT adhesion onto the multiwall and singlewall carbon nanotube surface. Figure 1(a,b) illustrates HR-TEM images of P3HT strands coiled around narrow multiwall nanotubes of 4-walls and 5-walls having diameters of 4.4 and 5.4 nm, respectively. Thin polymer layer coiling of nanotubes, which was proposed by O’Connell et al.22 in 2001, is now a well-known phenomenon with several examples in literature involving HRTEM23-25 and resulting very similar to the cases proposed here. The polymer self-assembled structure, which appears like a bulky ribbon coiling around the nanotube, is almost transparent, being visible only through the darker edges of the polymer strands. This effect, which is very similar to the way nanotube are imaged under TEM,26 could be due to the different electron transmission coefficient when crossing vertical or horizontal portion of the polymer ribbon surface. The polymer coiling is more evident in slightly larger scale images, i.e., Figure 1(e) reported in the Supporting Information. In both images of Figure 1(a,b), the polymer coiling is irregular with a variable coiling angle, between 60 to nearly 90, measured with respect to the nanotube axis. Consequently, the polymer coil-to-coil distance varied from a few Angstroms to more than 3 nm. In Figure 1(b), blue arrows highlight the MWCNT’s shell interlayer distance of 0.34 nm which is in
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Figure 1. (a, b) High resolution transmission electron microscopy image of P3HT coiled around multiwall carbon nanotubes and, (c, d) around singlewall carbon nanotubes. (e) Slightly inclined MWCNT showing uniform P3HT coiling fading out of the focus plane.
agreement with values reported by other authors that range from 0.34 to 0.39 nm.27 The measured polymer stacking distance is 0.36 nm which matches the π-π interaction stacking distance we reported for adhesion on carbon nanotubes.28 Singlewall carbon nanotubes (SWCNTs) around which are coiled P3HT ribbons are reported in Figure 1(c,d). As can be observed, the polymer’s tendency to wrap the nanotubes by coiling is not affected by the small diameter of these nanotubes, in this case 2.0 and 1.1 nm respectively. This highlights the flexibility of Poly(3-hexylthiophene) structure and demonstrates the bending of the physically adsorbed P3HT around carbon nanotubes with a relatively small curvature radii. We note that Figure 1(c), in particular, is in very good agreement of recently proposed molecular dynamic self-assembly arrangements of P3HT on singlewall carbon nanotubes.29 Therefore, it can be assumed that P3HT coil-wrapping of carbon nanotubes is a general phenomenon that extends over variable lengths, yet may exhibit different arrangements from tube-to-tube. The wrapping is supposed to occur between these two species during the interaction in solution phase and persists after the solvent evaporation following the solution cast of the composite material. Figure 1(e) illustrates another example of P3HT coiling around a 4-wall carbon nanotube slightly inclined. Even if the 6325
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Figure 2. HR-TEM images of polymer compressed multiwall carbon nanotubes. Short (blue) arrows indicate pressure points whereas long (green) arrows highlight the sidewall profile peaks. Black arrow-tips indicate the wider interlayer distance consequence of the surrounding pressure points.
polymer coiling is present on the whole nanotube structure, it is difficult to see (almost invisible) on the right part of the figure, while is easily visible and resolved on the left part. This difference is due to the non-planar nature of the whole structure which passes through the focus plane and evidences different features on the MWCNT surface. Another interesting feature of the P3HTcoiled nanotube is the two clear deformation points, indicated with dark dots which are visible on the upper part of the structure, in correspondence with polymer coils. Two more deformation points can be individuated on the lower part (indicated empty circles) but in this case, it is not possible to exclude that the deformation is due to the buckling of the nanotube. Nevertheless, the overall effect that can be observed is the diameter reduction of the P3HT coiled carbon nanotube. In order to further investigate the effects of the polymer coiling on the underlying nanotube structure, we imaged MWCNTs possessing a relatively large number of inner shells (12-16) that have been wrapped by P3HT. In Figure 2, it can be observed that both nanotubes exhibit compressed regions and display sizable deformations of their outer radius. (Other HR-TEM images showing nanotube deformation of MWCNTs are reported in the Supporting Information.) The deformations occur in straight nanotube segments, far away from any bends or kinks, ruling out the possibility that the structural defects imaged originate from long-range nanotube deformation. Radial deformations30 up to a complete nanotube collapse31 have been observed in the absence of any covering material and have been attributed to pure van der Waals interaction among nanotube shells. The same interaction, exerted by the polymer wrapping is likely to be the reason (or one of the reasons) for the targeted MWCNT sidewall deformation. By observing Figure 2(a), it can be noted that on one MWCNT sidewall a harmonic deformation (peaks and valleys) of the MWCNT outermost shells radii is present. Remarkably, as highlighted by the short-blue arrows, at each compression point (valley) it is possible to observe the polymer coil cross section, whereas the peaks appear to coincide with polymer-free regions of the nanotube’s external surface. As a consequence of the polymer-induced surrounding radial pressure, larger interlayer distances between the MWCNT’s inner shells are also observed which appear to coincide with the sidewalls peaks. This result was unexpected since it is often assumed that the radial stress is transferred through the nanotube layers causing the MWCNT
interlayer distance and/or internal radius reduction. This latter phenomenon can be observed in Figure 2(a) highlighted by the red arrowhead which is coincident with a polymer-induced pressure point. In this case, it is observed that the external stress is transferred elastically to the inner shells causing eventually the innermost nanotube radius reduction. It should be noted that Figure 2(a) displays an asymmetric multiwall nanotube, where the sidewalls exhibit a different number of inner graphene shells and therefore different thicknesses. (During HR-TEM experiments, a large number of carbon nanotubes exhibiting a different number of shells between the sidewall have been imaged. Some examples are included in the Supporting Information.) The upper and the lower sidewalls of the tube in Figure 2(a) are formed by 12 and 16 shells, respectively, with a constant interlayer thickness of 0.38 nm. The measured difference between the minimum and maximum diameter is 1.46 nm, whereas the nanotube average diameter is 17.54 nm, resulting in a deformation of 4.2%. From these considerations, the distance between the P3HT coils can be estimated by considering the “polymer compressed zones” of the MWCNT structure and results in a polymer repeat-distance of 8-9 nm. Considering the multiwall carbon nanotube reported in Figure 2(b), it is observed that it exhibits a 12-shells structure, with an interlayer thickness 0.33-0.35 nm. The average diameter is 13.01 nm, whereas the diameter variation is 0.98 nm, yielding to deformation of 3.8%. In this case, the increased distance between the inner nanotube walls is strongly evident. In addition, although this nanotube’s cross-section is symmetric, the polymer compression effect is visible only on one side of the nanostructure as it was for the previous asymmetric cross-sectioned nanotube. Remarkably, in both cases, the buckling side of the nanotube occurs where the polymer coils, while the nanotube flat side shows a more uniform P3HT coverage. It is important to observe that under induced stress, the radial deformation of singlewall nanotubes increases with the radius, but decreases if at the same radius a higher number of inner shells is considered.32,33 Therefore, for small radii nanotubes like in the case of SWCNTs or DWCNTs, we expect very rigid structures and no visible deformation while for large radii, thin-wall MWCNTs we consider deformations more likely to be observed. In first instance, we investigate the P3HT mechanical capability to compress the nanotube, excluding any other mechanism 6326
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(for example charge transfer) that could possibly aid the nanotube deformation. We start from the MWCNT elastic radial modulus reported values, measured experimentally with atomic force microscope nanoindentation. The radial compressive elastic modulus, under asymmetric normal force, can be defined as follows: ERAD ¼ ðF=AÞ=ðΔD=DÞ
ð1Þ
where F is the applied force, A is the contact area and ΔD/D is the strain measured on the diameter. Due to the nonlinearity of the phenomena involved, a better estimation of ERAD is obtained by differentiating eq 1 or reducing the extent of the compression. A recent independent study by Palaci et al.34 provides an accurate estimation of the MWCNT radial elastic modulus. The study was conducted on nanotubes with radii very similar to our case, reducing the extent of the induced deformation and, as for our case, utilizing MWCNTs synthesized by chemical vapor deposition. (The nanotubes reported in Palaci’s study had external radii ranging from 0.2 to 12 nm, while in our case, the external radii are 6.5 and 8.75 nm. Moreover, the referenced work was conducted and MWCNTs with constant Rext/Rint ratio of 2.2 ( 0.2 nm. In our case, the Rext/Rint ratio is 2.23 and 2.46, confirming the good agreement between the two studies.) For multiwall nanotubes with external radius larger than 4 nm, the value of ERAD saturates to 30 ( 10 GPa, close to the value of graphite of 36 GPa.35 Assuming 30 ( 10 GPa as a valid value, the estimation of the applied pressure is 1.2 ( 0.4 GPa. Regarding the polymer, the experimental value for the P3HT’s Young modulus has been recently measured by using the buckling-based method36 and the value reported is 1.33 ( 0.01 GPa. (The value provided for the Young's modulus of P3HT has been measured for polymer films. The authors believe that the tensile strength of the single polymer backbone can be higher, confirming the possibility of carbon nanotube deformation by P3HT strands.) Therefore, we can affirm that the calculated pressure exerted on the MWCNT by P3HT coiling can be structurally sustained, and hence applied, by the polymer. We notice that it is significant that the deformation observed is close to the maximum that the polymer can withstand and further compression cannot occur. From our HR-TEM analysis, it is hence possible to assess that one effect of the polymer coiling is to induce structural defects on the nanotube sidewalls, such as diameter reduction and interlayer displacements. As shown in several studies, radial deformation is able to modify the nanotube electronic37,38 and transport properties,39 with expected significant repercussions on device performance. It must also be noted that the electron transfer from the polymer to the nanotube is suspected to play a fundamental role in the nanotube deformation as observed for potassium doped doublewall carbon nanotubes.40 Electron transfer from P3HT to carbon nanotube has been proposed,41 modeled,42 and recently experimentally observed by our group with scanning tunneling spectroscopy.43 The local charge transfer on the nanotube surface is expected to weaken the C-C bond on the nanotube surface and to soften the vibrational modes.44 In order to support the previous results and provide further evidence of polymer adhesion-induced nanotube compression, MWCNT/P3HT samples were investigated with vibrational spectroscopy. Raman spectroscopy studies have demonstrated that frequency displacement of CNTs' first-order vibrational modes often occurs when the carbon nanotubes are incorporated in composites. Specifically, G-band peak frequency upshifts and
downshifts occur and have been attributed to stress-induced compression,45,46 molecular adsorption on the nanotube surface47 and charge transfer.44 Studies on singlewall nanotubes48-50 proved that the G-band peak position in Raman spectra, related to transversal modes, increases almost linearly with the external pressure applied, before attaining a nonlinear zone with a reduced increase rate that can be related to reversible nanotube collapse. For MWCNTs, the same trend was observed51 with a reported increase rate of 4.3 cm-1 GPa-1. A series of first-order Raman spectra starting from pristine MWCNT samples and moving to increasing P3HT relative weight content up to 1:1 ratio were collected for this work. Figure 3(a) shows the first-order Raman spectra of pristine multiwall nanotubes and 1:1 w/w MWCNT/P3HT samples. The pristine MWCNT spectrum between 1000 and 1800 cm-1 features the defect induced D-band52 centered at 1337 cm-1 and the characteristic tangential modes related G-band53 centered at 1568 cm-1. Conversely, the 1:1 w/w MWCNT/P3HT sample’s spectrum exhibits vibrational bands originating from the polymer’s presence. The highest intensity P3HT-related peak occurs at 1446 cm-1 which is attributed to CRdCβ stretching modes, whereas weak vibrational modes at 1377 cm-1, related to Cβ—Cβ’ stretching and Cβþ—H bending, are barely visible because they overlapped the MWCNT D-band. Other vibrational modes for the polymer occurring at 1090 and 1210 cm-1 are discussed elsewhere.54 The G-band of the 1:1 w/w MWCNT/P3HT composite displays an upshift of 8 cm-1, with the experimental maximum centered at 1576 cm-1. The peak is also broadened, showing a fwhm of 55 cm-1 which results larger than the 41 cm-1 fwhm of the pristine MWCNT sample. Figure 3(b) shows the G-band Raman upshift as the polymer content in the compound is raised. At low polymer content, the contribution of compressed nanotubes is negligible with little or no upshift observed in the spectrum. As the P3HT quantity is raised, the upshift increases up to a maximum of 11 cm-1 for the 0.5 w/w sample. A relative downshift of 3 cm-1 is observed for 1:1 w/w sample reported in Figure 3(b). A similar gradual upshift related to increasing nanotube contents was reported by Baskaran et al.,47 by McNally et al.,55 and by Musumeci et al.46 studying polybutadiene-MWNT, polyethylene-MWNT, and P3HTMWNT composites, respectively. The origin of the relative downshift at high polymer loads can be attributed to charge migration from the P3HT. It has been reported that negative charge-doped nanotubes exhibit a G-band displacement to lower wave numbers.44 We propose that the compression effect is predominant until the nanotubes are completely wrapped or the amount of charge transferred (proportional to the polymer content) equilibrates, and this eventually reverses the peak shift. Considering that both polymer-induced MWCNT compression and doping are related to P3HT adsorption on the nanotube surface, we conclude that the resultant G-band shift is the result of these two phenomena operating in competition. Therefore, the electron migration in our measurements placed a limit on the compression-related G-band peak upshift observed that could have otherwise resulted in a G-band upshift higher than the value measured. (It must be noticed that a charge migration from the nanotube to the polymer could have been able to support upshifts higher than expected, but would have failed to explain the relative downshift at higher polymer contents.) Although we can assess that a completely covered MWCNT is characterized by an upshifted G-band, we must observe that a partially coiled nanotube like the 6327
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Figure 3. (a) Normalized Raman spectra of polymer free multiwall nanotubes sample and 1:1 w/w MWCNT/P3HT sample evidencing the compression-induced G-band peak upshift, indicated with Δ. The G-band peak for pristine MWCNT samples is centered at 1568 cm-1. (b) Wavenumber shift of multiwall carbon nanotubes G-band versus P3HT relative content. (c) P3HT highest intensity peak shift as the compound relative content of the polymer is raised.
ones reported previously in the text is likely to show a Raman spectrum which could be more affected by the polymer doping rather than the compression. Further studies on limited number of nanotube spectra are being pursued at the moment to elucidate this point. Raman spectra of the same samples were next collected after a low temperature annealing (120 C). The annealing step was expected to lower the compressive forces on the carbon nanotubes due to polymer’s reorganization and thermal relaxation. As can be seen in Figure 3(b), after the thermal treatment, the maximum upshift observed is reduced to 9 cm-1. Focusing on the P3HT CRdCβ stretching vibration mode in the Raman spectra, Figure 3(c) shows the measured central frequency of this feature for the different composites. As can be observed, the vibration mode frequency measured at the expected value of 1446 cm-1 is downshifted for the lowest P3HT content sample (1:0.1 w/w). We note that the 6 cm-1 observed downshift occurs for the composite with the maximum carbon nanotube surface available for polymer adhesion. Heller et al.56 reported P3HT CdC stretching modes downshift that coincided with increased values of the polymer’s effective conjugation length. We therefore consider that the electron rich MWCNT surface plays a significant role in the P3HT self-organization, leading to polymer segments with order-enhanced structures. In order to confirm this observation, the same analysis was performed on the samples after the thermal annealing. Even at the highest content ratio, the P3HT main peak is downshifted by a few wave numbers which
we interpret as evidence for the contribution of P3HT segments possessing extended conjugation lengths interacting with nanotubes. (Although thermal annealing can enhance the polymer crystallization order, Raman spectra collected on pristine P3HT samples before and after the thermal treatment did not show any downshift, reported in the Supporting Information. This confirms that the nanotube sidewall structure can act as template to promote a more ordered polymer phase.) Further support of the effect of thermal annealing on P3HT self-assembly onto MWCNTs, is provided by continuously acquired HR-TEM images of a polymer-covered nanotube under high energy electron beam. [With high energy electron beam is intended a radiation able to raise the temperature of the sample under analysis. Although there is no control of the temperature of the sample in the time frame under analysis we can exclude damages to the nanotube. The damage of the sample occurred after a long time exposure (not shown).] Figure 4 shows four images collected at different times clearly showing the polymer reorganization onto the nanotube surface. If Figure 4(a) is taken as reference, then in Figure 4(b), a new polymer induced defect (indicated with R) is observed on the MWCNT structure after more than 11 min electron beam exposure. The following image, part (c), shows two other defective features, indicated with β and γ. In the same figure, the compressed zone R is more evident and generates a sizable, localized inner diameter reduction. After almost 24 min of electron beam exposure, the polymer is still rearranging its 6328
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Figure 4. (a-d) Effect of the electron beam induced heat on the P3HT adhesion onto a MWCNT. The sizable effect of the polymer-induced compression is shown with consecutive images collected at different times on the same structure. Scalebar is 5 nm for all images.
structure on the nanotube surface, while the defect β is barely visible and the induced elastic compression in R is no longer perceptible. Conversely, compression point γ becomes more obvious and the nanotube diameter reduction appears at another location of the structure. The elastic restoration of structural defect R proves that up to 24 min of analysis after the reference time, no irreversible nanotube deformation or structural damage has been induced by the electron beam.
4. CONCLUSIONS We have provided microscopic evidence of the P3HT coilwrapping around MWCNT revealing the nanotube sidewalls deformation caused by the P3HT coils. The value calculated for the observed deformation is consistent with the polymer structural stiffness. Raman scattering spectra for MWCNT/P3HT samples show the G-band peak position increasing upshift which is connected to the relative weight content of the polymer in the composites, confirming the radial compression effect due to the wrapped nanotubes contribution. This effect is reduced in thermally annealed samples due to the relaxed polymer adsorption on the nanotube sidewalls. High resolution transmission electron microscopy images acquired as a function of time during thermal annealing, confirm that the deformations can be the result of the polymer thermally assisted reorganization. Nanotube deformation as a result of the polymer coiling is expected to have major consequences on the MWCNT electronic and
mechanical properties. Therefore, multiwall nanotubes included in P3HT matrix could have a response different from what expected and result in performance lower than projected.
’ ASSOCIATED CONTENT
bS
Supporting Information. TGA analysis of nanotube purity, HRTEM images showing visible and invisible polymer coiling, other nanotubes showing deformation due to polymer wrapping, examples of nanotube sidewalls with different thicknesses, and polymer Raman spectra pre- and post-annealing. This material is available free of charge via the Internet at http://pubs. acs.org.
’ AUTHOR INFORMATION Corresponding Author
*Fax: þ 61 7 3138 8381. E-mail:
[email protected].
’ ACKNOWLEDGMENT This work has been supported by the Queensland Government through the NIRAP project “Solar Powered Nanosensors”. M.G. thanks Dr M. Rybachuk and Dr K.R. Moonoosawmy for the fruitful discussion and M.C. Hales for the help with Raman Spectra. 6329
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