Reversible Shape Memory of Nanoscale Deformations in Inherently

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Reversible Shape Memory of Nanoscale Deformations in Inherently Conducting Polymers without Reprogramming Michael J. Higgins,* Willo Grosse, Klaudia Wagner, Paul J. Molino, and Gordon G. Wallace* ARC Centre of Excellence for Electromaterials Science (ACES), Intelligent Polymer Research Institute (IPRI), AIIM Facility, Innovation Campus, University of Wollongong, Squires Way, Fairy Meadow, NSW, 2519, Australia

bS Supporting Information ABSTRACT: By using inherently conducting polymers, we introduce new shape memory functionality for stimuli-responsive polymers. The shape memory process is unique in that it utilizes electrochemical control of the polymer redox state to conceal, and temporarily store, preformed nanoscale surface patterns, which can later be recalled. Unlike classical thermoset and thermoplastic shape memory polymers, the electrochemical control does not completely perturb the low entropy state of the deformed polymer chains, thus enabling the concept of reversible transition between the permanent and temporary shapes. This is demonstrated using electrochemical-atomic force microscopy/quartz crystal microbalance to characterize the modulation of nanoscale deformations in electroactive polybithiophene films. Experimental results reveal that cation/solvent exchange with the electrolyte and its effect on reconfiguration of the film structure is the mechanism behind the process. In addition to incorporating conductive properties into shape-memory polymers, the ability to reversibly modulate surface nanopatterns in a liquid environment is also of significant interest in tribology and biointerface applications.

’ INTRODUCTION Inherently conducting polymers (ICP) are plastic materials that are electrically conductive. When operated as an electrode within an electrochemical device, changes in the redox state of these polymers induce electrical, chemical, and mechanical property changes that can elicit color, volume, and morphological responses. These properties can be reversibly switched and form the basis of numerous applications and technologies centered around these materials, including actuators, electrochromics, organic solar cells, and bionics.1 Changes in their properties are caused by an electrochemical doping/dedoping process and can be explained using eqs 1 and 2, where Pþ and P0 are the doped and dedoped state of the polymer, respectively. Ais the dopant anion incorporated into the polymer upon oxidation of the polymer and Cþ is the cation in the surrounding electrolyte medium. Pþ ðA - Þ þ Cþ þ e- T P0 þ A - þ Cþ

ð1Þ

Pþ ðA - Þ þ Cþ þ e- T P0 ðACÞ

ð2Þ

Typically, upon reduction of the polymer, eq 1, the ejection of anions causes the polymer to contract; however, if the anion is immobile, eq 2, cations will instead be inserted into the polymer to balance the loss in charge, causing it to swell. This process of swelling and contraction is generally reversible in both situations and can be used to actuate the polymer for applications such as organic microelectromechanical systems and artificial muscles.1,2 r 2011 American Chemical Society

To date, studies on ICP actuation and its applications have been limited to the macro- and microscale for free-standing films3 and fibers4 and as components such as laminates in actuating devices.5 Therefore, the exploration of ICP actuation in the nanoscale domain presents an exciting dimension. Studying nanoscale actuation in these polymer systems can be thought of as measuring the actuation of autonomous, nanoscale-sized structures whereby the effect of the nanoscale dimensions on the process is of interest. This has been demonstrated for the actuation of individual polypyrrole nanowires embedded in polycarbonate membranes,6 though the difficulty in isolating and configuring nanometre structures is limiting. Another approach is to measure nanoscale displacements even though the actuation is from a macroscale polymer structure. Electrochemicalatomic force microscopy (EC-AFM) has commonly been used for this by measuring displacements or height changes of thin conducting films, as a simultaneous potential is applied to switch the redox properties.7-9 By controlling the potential waveform or film thickness, height changes down to a few nanometers can be measured.10 In this study, we introduce and develop a novel approach that involves the controllable actuation of preformed nanostructures within an ICP film. The approach has uncovered new “shape-memory” functionality in polymers, Received: December 19, 2010 Revised: February 2, 2011 Published: March 14, 2011 3371

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Figure 1. Shape memory effect in polymers. (A) One-way process requiring the reprogramming of the temporary shape through mechanical deformation and/or heat. (B) New process with reversible recovery of the temporary shape without reprogramming.

particularly in the recovery of polymers after nanoscale deformation using atomic force microscopy (AFM). Most often, the shape-memory effect in classic systems is a one-way process shifting the material from a permanent f temporary f permanent shape (Figure 1A). Two-way shape memory is also possible, though in each case repeated programming of the temporary shape through mechanical force and/or heat is required. Here, we assessed the possibility of using ICP to implement a nanoscale process that allows reversible shifting between the temporary and permanent structures without reprogramming (Figure 1B). Furthermore, the use of ICP, which are not considered to be shape-memory polymers due to their nonthermoplastic/set properties, imparts additional functionality such as conductivity.

’ MATERIALS AND METHODS Electrochemical Polymerization of Polybithiophene (PBth) Films. PBth doped with perchlorate was grown as films

galvanostatically at a current density of 1 mA/cm2 on a gold Mylar working electrode in 0.1 M bithiophene monomer/tetrabutylammonium in acetonitrile. Polymer growth was performed for 20 s corresponding to a polymerization charge of 0.02 C/cm2 in a three-electrode electrochemical cell consisting of a platinum mesh counter and Ag/AgCl reference electrode using an eDAQ potentiostat/recorder. The growing solution was degassed with nitrogen for approximately 15 min prior to polymer growth. Analysis of Mechanical Properties. AFM images of the conducting polymer films were obtained in air using an MFP3D Asylum Research AFM (Asylum Research, Santa Barbara, CA) operated in ac mode with a silicon nitride cantilever (k = ≈40 N/m). Using the same cantilever, a series of nanoindentations were performed by carrying out force measurements using different maximum loads at selected x-y positions in the images. The force measurements were converted to force versus

indentation curves and analyzed using Oliver-Pharr model within the Asylum Research AFM software and IGOR PRO (Wavemetrics). There are several assumptions and errors associated with the application of Oliver-Pharr analysis to AFMbased indentation measurements; thus, Young’s modulus is given as an estimate. Electrochemical AFM (EC-AFM). EC-AFM measurements were done using a Biowizard II AFM (JPK Instruments, Germany) with integrated electrochemical cell consisting of a pseudo silver wire reference electrode, platinum wire counter electrode and the polybithiophene films as the working electrode substrate. The measurements were carried out in 0.1 M tetrabutylammonium perchlorate (TBAP) in propylene carbonate (PC) for several different polybithiophene film samples. The ability to plastically deform the sonicated polythiophene films at low forces, particularly when immersed in PC electrolyte, facilitated the use of lower spring constant cantilevers (Nanoworld PNP, k = ≈0.25 N/m). These cantilevers had the advantage of first doing a grid-array of indentations, typically at load forces smaller than those observed under dry conditions (≈100 nN), then imaging the nanoindentations before and after electrochemical switching at low imaging forces so as to avoid modifying or damaging the surface. Electrochemical-Quartz Crystal Microbalance Measurements. EC-QCM measurements were performed using a Stanford Research Systems (SRS) QCM200 with SRS AT-cut 5 MHz Au-coated crystals. The EC-QCM cell consisted of a platinum mesh counter electrode and Ag/AgCl reference electrode, with the Au electrode on the sensor crystal acting as the working electrode. Polybithiophene films were initially electrochemically polymerized in situ using the same growth conditions described in the Introduction. The sensor was then removed, sonicated in deionized water, and dried under a stream of Nitrogen gas. Thereafter, the sensor was replaced in the E-QCM and immersed 3372

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Figure 2. (A) AFM height image in air of as-grown PBth film showing aggregates (black asterisk) and underlying smooth layer (red asterisk). 25 μm AFM scan. Inset: optical image of the same film. (B) AFM height image of same film sonicated in water for ≈30 s showing removal of aggregates. Inset: optical image for comparison with inset in (A). (C) Cyclic voltammagram (CV) of sonicated film in (B). The CV is started at 0 V and arrows show the direction of the applied voltage. (D) AFM height image of the sonicated film after applying a series of nanoindentations at different load forces. 750 nm AFM scan.

in a solution of 0.1 M tetrabutylammonium perchlorate (TBAP) in polypropylene carbonate (PC). An eDAQ potentiostat/recorder was used to perform the electrochemical measurements as well as record the QCM output data.

’ RESULTS AND DISCUSSION Electrochemically polymerized polybithiophene (PBth) films had dark aggregates (Figure 2A, inset) that formed part of a dual layer morphology comprising a smoother, underlying layer (Figure 2A). Dual layer growth has been attributed to initial deposition of less soluble, higher molecular weight fractions followed by lower molecular weight fractions,11 the latter occurring as the concentration of oligomers reacting in solution decreases and polymerization slows. Adsorption spectra following PBth film growth has indicated polymer chains are well ordered up to 50 Å but have a degree of disorder as the film thickness increases.12 Brittle films or powdery deposits can also form with continued growth, most likely from low cohesion of the observed aggregates, and leave behind residue when physically handled.12 Even at our early stage of film growth, the aggregates were easily removed by scanning with the AFM tip or sonication in water for ≈30 s. Sonication of the films was valuable for completely removing the aggregates to render smoother films (cf. Figure 2B and inset) suitable for nanoindentation experiments.

The sonicated films remained electroactive and had the characteristic redox profile of PBth (Figure 2C), which is important as this was to be the control mechanism for the “shape-memory” functionality. The additional requirement of having a deformable polymer was assessed by performing a series of nanoindentations at different load forces using AFM under dry conditions. After indenting with the AFM tip, the polymer underwent plastic deformation (Figure 1D), with a minimum force of ≈300 nN required to permanently deform the polymer by ≈3.5 nm. The deformation process had a linear force versus indentation relationship and estimation of the effective Young’s modulus using Oliver-Pharr analysis yielded a value of 493 MPa. Polymer chain effects during the nanoscale plastic deformation of ICP have recently been studied with poly(3-hexylthiophene-2,5diyl) (P3HT). Incident X-ray diffraction shows induced chain ordering within the nanopatterns of nanoimprinted P3HT films.13 AFM mechanical lithography of regioregular P3HT films results in patterns significantly larger than the AFM tip due to “cutting”, as the polymer under tension pulls away due to strain relaxation.14 However, this does not appear to be the case here as our nanoindentations represent the AFM tip dimensions. We assessed and analyzed other types of conducting polymer films, including polypyrrole, poly(3,4-ethylenedioxythiophene), and poly(2-methoxyaniline-5-sulfonic acid), using the same approach. In contrast to PBth, these films did not undergo plastic deformation except for polypyrrole at the highest load of 1 μN 3373

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Figure 3. (A) AFM height image of a PBth film after applying an array of nanoindentations and prior to application of electrochemical switching. 1 μm AFM scan. (B) AFM height image of the same scan area in (A) after the application of -1 V showing erasure of the nanoindentations and polymer recovery. 1 μm AFM scan. (C) Enlarged regions of AFM height images following changes in nanoindentations (numbers 7, 8, 9) after the application of different applied potentials. No Stim refers to no applied voltage. Below graphs showing the change in indentation depth (D) and diameter (E) for each indentation number as a function of the applied potential. Black asterisk indicates no measurable depth/diameter and full polymer recovery.

(see Supporting Information, Figure 1). Due to the lack of plastic deformation, these polymers were not examined further in this study. Having demonstrated suitable properties for the PBth films, the main aim was to investigate the effect of electrochemically switching the polymer redox state on modulating the topography

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of the nanoindentations. We hypothesized that, by oxidizing the films, their swelling due to the uptake of the dopant anions (as described by eq 1 and in this case perchlorate) from the electrolyte would erase the nanoindentations and recover the polymer (permanent shape) due to smoothing out of the topography. Smoothing out of the polymer surface due to swelling has previously been reported to result in a decrease in roughness.15 Furthermore, it was expected that upon electrochemical switching to a reduction potential, if the polymer retained the “memory” of the nanoindentations (temporary shape), then induced ion expulsion and subsequent contraction of the film would result in the recovery of the nanoindentations. The entire process is therefore said to be a reversible process without reprogramming. To test our hypothesis, a series of nanoindentations were investigated using a combined electrochemical-AFM setup. First, to achieve stability in polymer electroactivity, an initial cyclic voltammogram (CV) was run from 0 to 1.2 V for a total of 10 cycles and then terminated at 1.2 V to leave the film in the oxidized state. A grid array of nanoindentations (no. 1-9) was made in the film (Figure 3A) and a series of constant potentials (0.68 V f -0.5 V f 1 V f -1 V) subsequently applied for 60 s. By following the effect of each applied potential on the nanoindentations (no. 7-9) using AFM imaging (Figure 3C), it was observed that the topography of the nanoindentations for 0.68, -0.5, and 1 V did not significantly differ from the starting film. Many of the topographic features surrounding the nanoindentations were also unchanged at these potentials. These observations at the potential of 0.68 V and more negative potential of -0.5 V agreed with our hypothesis that the nanoindentations should remain unchanged under these electrochemical conditions. However, a lack of erasure of the nanoindentations as the film was oxidized at 1 V did not agree with our hypothesis that at these positive potentials the insertion of perchlorate anions should smooth out the film. When attempting to switch the potential back to -1 V, the nanoindentations were, however, completely erased to the point where they were deemed unrecognizable and the polymer had recovered. Larger topographical features remained similar, though smaller features such as the nanoindentations and other film undulations appeared to be smoothed out. Erasure of the nanoindentations, or henceforth referred to as “polymer recovery”, occurred for the entire grid array (cf. Figure 3B), except partially for nanoindentations 3 and 6. Most of the nanoindentations were erased or differed significantly from their original form. Therefore, in contrast to that expected at oxidation potentials, cation insertion (i.e., tetrabutylammonium, TBAþ) and/or associated solvent uptake is most likely responsible for film swelling and polymer recovery, as the polymer is reduced at -1 V. Quantitative analysis of the nanoindentations showed that their depths ranged from 5 to 7 nm, which is within the statistical variation of acquiring each height profile, and remained relatively constant after applying each potential (Figure 3D). With the exception of nanoindentations 3 and 6, no residual depth of the nanoindentations could be recorded for the applied voltage of -1 V (Figure 3D). Similarly, no significant change was observed for indentation diameters ranging from 20 to 45 nm until the polymer recovered at -1 V (Figure 3E). In developing the above experiments, we found that the polymer recovery at -1 V was often better initiated by performing an initial switching cycle of -1 V f 1 V prior to commencing, suggesting that a “break-in” behavior16 facilitated the polymer swelling 3374

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Figure 4. AFM height images showing an array of nanoindentations before electrochemical switching (A), after polymer recovery at -1 V (B), and after indentation recall at 1 V (C). (D) Height profiles taken across each row of nanoindentations in the array before electrochemical switching (red trace), after polymer recovery at -1 V (black trace), and after indentation recall at 1 V (blue trace). Below graphs showing the change in indentation depth (E) and diameter (F) for each indentation number as a function of the applied potential. Each indentation number has three bars corresponding to before electrochemical switching (blue), after polymer recovery (red), and after indentation recall (green). Black solid line is linear fit to the percentage recovery. Inset in (E) shows an indentation profile where the observed change in dimensions in (D) can be described by a modulation of the parabolic function y = a(x2) þ b(x) þ c, as described in the text.

in the subsequent cycle. Break-in behaviors typically involve nonsteady-state redox processes at high overpotentials that can rearrange the film structure and enhance electroactivity. We also found that 30 s was the minimum time period required for recovery, though we typically applied longer times to ensure completion.

To assess whether the films retained the “memory” of the nanoindentations, grid arrays were performed across several different samples (total of 43 indentations) and their topographies imaged after the application of -1 V and then after switching back to 1 V. Nanoindentations were clearly visible prior to electrochemical switching (Figure 4A), and the polymer 3375

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The Journal of Physical Chemistry B recovery subsequently occurred after the application of -1 V (Figure 4B), though in some cases only partially (see nanoindentations (3, 4, 5, 6, 9). Importantly, the nanoindentations completely reappeared after switching back to 1 V (Figure 4C), thus supporting a reversible recovery process shifting from a temporary f permanent f temporary shape without reprogramming, as outlined in Figure 1B. We henceforth refer to the reappearance of the nanoindentations as “indentation recall”. Height profiles taken across each of the nanoindentations before electrochemical switching (red trace) and after polymer recovery (black trace) and indentation recall (blue trace) show the extent to which the dimensions of the nanoindentations change during the switching process (Figure 4D). In particular, the relatively good agreement between the red and blue traces highlights the good ability of the polymer to retain the memory of the nanoindentations. Quantitative analysis of the total nanoindentations was done to summarize the overall success rate of the reversible recovery process (Figure 4, E and F). For each indentation number, the depth and width values before switching (blue), after polymer recovery (red) and indentation recall (green) were plotted adjacent in sequence. In 79% of the indentations, the polymer showed complete recovery. The remainder that partially recovered occurred more frequently as the depth and width values increased, suggesting limitations in the bulk film strain required for polymer recovery. In terms of the reverse process of indentation recall, 83% of the nanoindentations were recalled to varying percentages of their original dimensions. Small nanoindentations of ≈2 nm were recalled to 100% of their original depth, which decreased linearly to 6070% for nanoindentations of depths >6 nm (Figure 4C). In contrast, their diameters remained relatively constant after indentation recall, suggesting that a change in the depth dominated the process. Further indication of this can be seen in height profiles of the partially erased nanoindentations listed above and shown in Figure 4D. By approximating the shape of the nanoindentations as a simple paraboloid, the changes in dimensions during polymer recovery (with the reverse for indentation recall) are predominately described by a decrease and increase in both the quadratic a(x2) (width) and (c) coefficients (vertex height), respectively. This is opposed to only a decrease in the a(x2) coefficient that would result in widening and apparent shallowing of the indentation. We suspect that this change in indentation geometry is related to the response of the residual stress distribution of the nanoindentations to the bulk film strain. In testing the number of reversible cycles, we found that we could only repeat the recovery process 2-3 times and that the nanoindentations were eventually unrecoverable, as shown and discussed further below. We investigated the possibility of the “memory” effect occurring for a trench geometry with different dimensions. In this case, a load force was applied using the AFM tip so as to form a trench in the film with depths and widths ranging from 10 to 15 nm and from 25 to 35 nm, respectively, and also with lengths of 125200 nm (Figure 5). Upon electrochemical switching for up to three cycles (i.e., 3 times -1 V f 1 V), these trenches showed a reversible recovery effect (Figure 5, images a-g and below graph). The first cycle (Figure 5, a-c) showed partial polymer recovery followed by close to full trench recall, while the second cycle (Figure 5, c-e) showed both full polymer recovery and trench recall. In contrast, the last cycle (e-g) showed a full polymer recovery but the trenches could not be recalled, indicating that reversibility was limited. In contrast to the

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Figure 5. (a-g) AFM height images following a series of trenches as a function of the applied potential for three temporary f permanent shape cycles (first cycle, a-c; second cycle, c-e; third cycle, e-g). Below graph showing depth, width, and length values for each trench at each applied potential over the three cycles. Depth and width showed a cyclic change while the length remained constant. Inset shows change in trench geometry during the process. It is noted that, as the width decreases, the depth of the trench is not known due to the limitations associated with the AFM tip dimensions.

indentation patterns above, the width of the trenches changed dramatically (Figure 5, graph and inset), suggesting that a change from a radial to bilateral stress distribution may account for differences in the changing trench geometry during the process. Electrochemical-quartz crystal microbalance (EC-QCM) was used to assist in elucidating the electrochemical swelling mechanism of the reversible recovery effect. Initial cyclic voltammetry (0 f 1.2 V) showed an expected mass increase (frequency decrease) and decrease (frequency increase) associated with 3376

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Figure 6. Electrochemical-quartz crystal microbalance measurements for sonicated PBth films: current (red), voltage (blue), and frequency shift (black). (A) CV measurements run between 0 and 1.2 V at 50 mV/s. Starting voltage is 0 V. (B) Simulation of two temporary-to-permanent shape cycles by applying 2(-1 V f 1 V). (C) CV measurements run between 1 and -1 V at 50 mV/s. Starting voltage is 0 V.

reversible movement of perchlorate anions during oxidation and reduction, respectively (Figure 6A). Subsequent EC-QCM measurements aimed at monitoring the frequency changes during two simulated polymer recovery and indentation recall cycles involved two cycles of holding the potential constant for 60 s at -1 V and then 1 V (i.e., 2 times -1 V f þ1) (Figure 6B). In the first cycle at -1 V, a very small mass increase confirmed the difficulty of initially incorporating relatively large cations into the polymer. Following this, switching to 1 V showed a mass increase in accordance with perchlorate anion insertion during oxidation. However, the next polymer recovery cycle at -1 V showed an initial instantaneous small mass decrease, presumably due to ejection of the perchlorate anion, followed by an exponential increase in mass, indicating that significant swelling of the film is a driving force for polymer recovery. This subsequent mass increase also confirmed the presence of a “break-in” effect that facilitated swelling in this subsequent cycle. Based on charge balance, the insertion of TBAþ is likely during swelling, though additional solvent uptake may also explain the dramatic mass increase.17 The process is less clear for indentation recall at 1 V, though it is observed that the mass increase (perchlorate anion uptake) is not accompanied by an initial decrease, suggesting that the ions/solvent incorporated during reduction remain trapped in the film. The latter may contribute later to the loss of reversibility of the indentation recall. Finally, a CV of the full potential range highlighted the multiple processes in the mass change (Figure 6C), particularly the increase at a reduction peak of -1 V (Figure 6C) responsible for polymer recovery.

The reversible recovery effect observed here for the inherently conducting polymer films is unique from that of classical shapememory functionality in polymers. We particularly refer to the large class of thermoset or thermoplastic polymers which by definition can be stabilized in the deformed state. In these polymers, the shape-memory effect relies on a processing and programming stage.18 First, the polymer is processed into its permanent shape and then mechanically deformed into a temporary shape that remains fixed. The permanent shape is stored while the temporary shape is on display. The shape-memory effect occurs when the polymer is subsequently heated above its melting (TM) or glass transition temperature (TG) to recover the permanent shape. The molecular basis for the shape-memory effect is that sufficient physical cross-links (or crystalline nodes for a thermoset) are present to form a “memorable” polymer network. During polymer deformation, the strain-induced crystalline regions incorporated in between the physical cross-links prevent the polymer network from immediately re-forming the permanent shape. The induced crystalline regions are in a low state of conformational energy, but can be switched to a higher entropy state and free volume to recover the permanent shape (or stored strain) upon heating above TM. Cooling does not recover the temporary shape, and thus this process in nonreversible and reprogramming is required. For thin thermoset films such as PMMA and epoxy, AFM-based studies have been used to “write” nanoindentations into the polymer which can easily recover after heating.19 For our conducting polymer films, the programming step is similar. The load forces of the AFM tip induces chain order or “crystalline” regions, therefore stabilizing 3377

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The Journal of Physical Chemistry B the indentations. However, the molecular mechanism for the reversible recovery differs by way of electrochemical switching of the redox states. In this case, the intercalation of TBAþ ions and/or associated solvent during swelling increases the distance between polymer chains and volume of the film structure. Such effects predominately underlie the polymer recovery but because the polymer network remains under entropic elastic strain, the memory of the induced “crystalline” regions is retained. For example, the elastic strain energy is not sufficient to perturb their low entropy state. Upon switching back, the induced “crystalline” regions also now act as effective cross-links to restabilize the temporary (indentation) shape as the stored elastic strain is recovered. As the process occurs at low reduction potentials, continual cycling of the film can then remove the induced “crystalline” regions through structural rearrangement of chains,16 or charge trapping/film degradation,20 and reversibility of the system is lost. The reversible recovery effect works by entropic elasticity coupled to the ability to program a deformable polymer. It is not a shape-memory effect in the true sense, as defined by the molecular mechanisms in thermoset/plastic polymers, but is analogous to the observed phenomenon. Although we have detailed the underlying principle for recovering the temporary shape without reprogramming, further optimization of a conducting polymer system is required to improve aspects such as increasing the number of reversible cycles by operating under electrochemical conditions where the redox properties of the polymer are not irreversibly affected. The concept here provides nanolithographic data storage capabilities, particularly in light of the ability to conceal the stored data like a form of steganography. The ability to also modulate a nanopatterned surface of choice in a liquid medium has uses in tribological applications where controlling lubrication, friction, and fluid flow are desired, and also in biological applications for modulating cellular-substrate interactions via dynamic nanotopographies. Furthermore, this new concept of nanoactuation in inherently conducting polymers and “shape-memory” functionality of polymers in general should be applicable to other stimuli-responsive polymer systems which show nanoscale deformation under load.

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(2) Madden, J. D.; Madden, P. G.; Hunter, I. W. Proc. SPIE 2002, 4695, 176–190. (3) Spinks, G. M.; Liu, L.; Wallace, G. G.; Zhou, D. Adv. Funct. Mat. 2002, 12, 437–440. (4) Mazzoldi, A.; Degl’Innocenti, C.; Michelucci, M.; De Rossi, D. Mater. Sci. Eng.: C 1998, 6, 65–72. (5) Jager, E. W. H.; Smela, E; Ingan€as, O. Science 2000, 290, 1540–1545. (6) Lee, A. S.; Peteu, S. F.; Ly, J. V.; Requicha, A. A. G.; Thompson, M. E.; Zhou, C. Nanotechnology 2008, 19, 165501. (7) Suarez, M. F.; Compton, R. G. J. Electroanal. Chem. 1999, 462, 211–221. (8) Smela, E.; Gadegaard, N. J. Phys. Chem. B 2001, 105, 9395–9405. (9) Higgins, M. J.; McGovern, S.; Wallace, G. G. Langmuir 2009, 25, 3627–3633. (10) Gelmi, A.; Higgins, M. J.; Wallace, G. G. Biomaterials 2010, 31, 1974–1983. (11) O’Neil, K. D.; Semenikhin J. Phys. Chem. B 2007, 111, 9253–9269. (12) Roncalli, R. Chem. Rev. 1992, 92, 711–738. (13) Aryal, M.; Trivedi, K.; Hu, W. ACS Nano 2009, 3, 3085–3090. (14) Jones, A. G.; Balocco, C.; King, R.; Song, A. M. App. Phys. Lett. 2006, 89, 013119. (15) Pyo, M. J. Electrochem. Soc. 2005, 152, E90–E93. (16) Tatsuma, T.; Hioki, Y.; Oyama, N. J. Electroanal. Chem. 1995, 396, 371–376. (17) Hillman, A. R.; Swann, M. J.; Bruckenstein, S. J. Electroanal. Chem. 1990, 291, 147–162. (18) Lendlein, A.; Kelch, S. Angew. Chem., Int. Ed. 2002, 41, 2034–2057. (19) Nelson, B. A.; King, W. P. App. Phys. Lett. 2005, 86, 103108. (20) Semenikhin, O. A.; Ovsyannikova, E. V.; Ehrenburg, M. R.; Alpatova, N. M.; Kazarinov J. Electroanal. Chem. 2000, 494, 1–11.

’ ASSOCIATED CONTENT

bS

Supporting Information. Figure showing AFM indentation of different conducting polymers. This material is available free of charge via the Internet at http://pubs.acs.org.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected] (M.J.H.); [email protected] (G.G.W.). Tel: þ61-2-4221-3989 (M.J.H.); þ61-2-4221-3127 (G.G.W.). Fax: þ61-2-4221-3114 (M.J.H., G.G.W.).

’ ACKNOWLEDGMENT We gratefully acknowledge the Australian Research Council (A.R.C.) for funding support. This research was supported by the ARC Federation Fellowship Program (G.G.W.). ’ REFERENCES (1) Wallace, G. G.; Spinks, G. M.; Kane-Maquire, L. A. P.; Teasedale, P. R. Conductive Electroactive Polymers—Dynamic Properties and Intelligent Material Systems; CRC Press: Boca Raton, FL, 2002. 3378

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