Autophobicity-Driven Surface Segregation and ... - NCSU COE People

energy at the PMMA/air interface by interfacial energy at the PMMA/PS interface, which reduces the energy barrier by an order of magnitude, thereby pe...
0 downloads 0 Views 2MB Size
Download PDF
NANO LETTERS

Autophobicity-Driven Surface Segregation and Patterning of Core-Shell Microgel Nanoparticles

2008 Vol. 8, No. 9 3010-3016

Bin Wei,†,‡ Paul A. Gurr,§ Arif O. Gozen,† Anton Blencowe,§ David H. Solomon,§ Greg G. Qiao,*,§ Richard J. Spontak,†,| and Jan Genzer*,† Department of Chemical and Biomolecular Engineering, Department of Materials Science and Engineering, North Carolina State UniVersity, Raleigh, North Carolina 27695, and Department of Chemical and Biomolecular Engineering, UniVersity of Melbourne, Victoria 3010, Australia Received July 16, 2008

ABSTRACT Core-shell microgel (CSMG) nanoparticles, also referred to as core-cross-linked star (CCS) polymers, can be envisaged as permanently cross-linked block copolymer micelles and, as such, afford novel opportunities for chemical functionalization, templating, and encapsulation. In this study, we explore the behavior of CSMG nanoparticles comprising a poly(methyl methacrylate) (PMMA) shell in molten PMMA thin films. Because of the autophobicity between the densely packed, short PMMA arms of the CSMG shell and the long PMMA chains in the matrix, the nanoparticles migrate to the film surface. They cannot, however, break through the surface because of the inherently high surface energy of PMMA. Similar thermal treatment of CSMG-containing PMMA thin films with a polystyrene (PS) capping layer replaces surface energy at the PMMA/air interface by interfacial energy at the PMMA/PS interface, which reduces the energy barrier by an order of magnitude, thereby permitting the nanoparticles to emerge out of the PMMA bulk. This nanoscale process is reversible and can be captured at intermediate degrees of completion. Moreover, it is fundamentally general and can be exploited as an alternative means by which to reversibly pattern or functionalize polymer surfaces for applications requiring responsive nanolithography.

Nanoscale core-shell microgel (CSMG) particles with unique three-dimensional (3D) structures constitute a class of branched macromolecules1-4 that generally consist of an organic (A) core surrounded by a dense brush of chemically dissimilar (B) arms.5-7 If the B shell is chemically crosslinked, the A core of such nanoparticles can be selectively removed and subsequently used as a hollow compartment for encapsulation.8-11 In this and our previous investigation,5 though, the core is cross-linked so that the B arms can freely interact with species in the local environment. As physical models of permanent AB diblock copolymer micelles12,13 and alternatives to organically modified metal or metal oxide nanoparticles,14-17 CSMG nanoparticles are of growing fundamental and technological interest because of their tailorable chemistry and unique physicochemical properties.18 Depending on their composition, such CSMGs can likewise * To whom correspondence should be addressed. E-mail: jan_genzer@ ncsu.edu (J.G.) and [email protected] (G.G.Q.). † Department of Chemical and Biomolecular Engineering, North Carolina State University. ‡ Current address: Corporate Research, National Starch and Chemical Company, Bridgewater, NJ 08502. § University of Melbourne. | Department of Materials Science and Engineering, North Carolina State University. 10.1021/nl802109x CCC: $40.75 Published on Web 08/08/2008

 2008 American Chemical Society

be made to respond volumetrically to a wide range of external stimuli19-22 or mimic the behavior of nonlinear biomacromolecular assemblies.23,24 Unlike micelles that (i) develop by molecular self-organization to minimize A-B repulsive interactions and (ii) structurally transform in response to environmental changes, however, the spherical core-shell morphology of CSMG nanoparticles remains locked-in from the time of synthesis and thus permits examination of scenarios wherein micelles would otherwise break apart. In this work, we consider CSMG nanoparticles dispersed in a thin film of a homopolymer that is chemically identical to, but of substantially higher molecular weight than, the arms comprising the CSMG shell. Generally speaking, if a homopolymer penetrates a dense brush, the brush is said to be wetted by the homopolymer. Conversely, a nonwetted (dry) brush results from either enthalpic incompatibility between the brush and a chemically dissimilar homopolymer or, as in the present case, the entropic penalty associated with limited mixing of highmolecular weight homopolymers inside a surface-anchored brush.25-28 Specifically, incorporation of high-molecularweight homopolymer molecules into the brush would require entropically unfavorable stretching of the chains comprising

the brush (loss of translational entropy), which consequently remains segregated from the free homopolymer chains. Such self-segregation, or autophobicity, has been extensively studied29-34 with regard to dense polymer brushes anchored to flat solid surfaces. It has not, however, been addressed to the best of our knowledge in the context of dense polymer brushes grafted onto discrete, permanently shaped polymer particles, such as the CSMG nanoparticles under investigation here. Autophobicity between the CSMG nanoparticles and the homopolymer molecules comprising the thin film causes the nanoparticles to migrate to the film surface where they remain intact, unlike their micellar counterparts. By judiciously altering the surface energy of the film, we can trigger the CSMG nanoparticles to either emerge from or submerge into the film in switchable fashion, thereby providing an alternative route by which to pattern or functionalize polymer surfaces for emerging nanotechnologies. The CSMG nanoparticles primarily employed in this study were synthesized with a cross-linked divinylbenzene (DVB) core and poly(methyl methacrylate) (PMMA) arms via the “arms-first” route by atom transfer radical polymerization previously detailed.4,6 Size exclusion chromatography (SEC) measurements were performed in THF using a Waters 717 Plus autosampler operated in series with a Wyatt Dawn F laser photometer operated at 90° and in parallel with a Waters 410 differential refractometer and a Viscotek T50A differential viscometer. Results from the Viscotek TriSEC software package calibrated with PMMA standards yielded a number-average molecular weight (Mn) and polydispersity index (PI) of 335 kDa and 1.10, respectively, and a gyration diameter (Dg) of 26 nm. The number of PMMA arms (each with Mn ≈ 11 kDa and PI < 1.10) on a single CSMG nanoparticle was 28. With these molecular parameters known and by recognizing that the PMMA arms comprising the shell are analogous to a PMMA brush grafted to the DVB core, the equivalent value of σN, where σ is the grafting density and N is the number of repeat units per chain in the brush, was estimated to be about 1.43. Note that this estimate is highly conservative, since σ ≈ 0.013 PMMA arm/nm2 is based on the outer CSMG surface area (given by πDg2), which is considerably larger than the core-shell interfacial area (recent efforts6 using selectively degradable nanoparticles indicate that the core size ranges from 1/3 to 1/2 the size of the entire nanoparticle). Since σN exceeds unity, a high brush density can be reasonably presumed wherein the PMMA arms compact laterally and stretch outward, normal to the core-shell interface, in similar fashion to brushes grafted to a flat surface.35 The CSMG nanoparticles were mixed (20% w/w) with PMMA (Mn ) 215 kDa, PI ) 1.05) in toluene to form a 2.5 wt % polymer solution, which was subsequently spin-coated onto Si wafer to yield films measuring approximately 110 nm thick, as discerned from ellipsometry. Polystyrene (PS, Mn ) 47 kDa, PI e 1.06) was similarly dissolved in toluene (2% w/w) and spin-coated on glass slides to produce films measuring approximately 67 nm thick. Each PS film was floated on the surface of deionized water and then transferred Nano Lett., Vol. 8, No. 9, 2008

Figure 1. Schematic illustrations showing the distribution of CSMG nanoparticles incorporated into (a) a portion of the PMMA bottom layer and (b) the PS capping layer5 in double-layer thin-film arrangements (see text) prior to annealing. Redistribution of the CSMG nanoparticles after annealing at 180 °C is portrayed in (c), wherein the PS layer dewets into discrete islands. Irrespective of the starting point (a or b), the CSMG nanoparticles and their aggregates protrude from the PMMA surface under the PS islands, whereas the exposed PMMA surface remains relatively smooth.

to the top of a PMMA/CSMG film in such fashion that half of the PMMA/CSMG film was covered by PS while the other half-remained exposed (cf. Figure 1a), thereby permitting the effect of the PS film to be directly discerned under a given set of treatment conditions. After drying for at least 24 h in air at ambient temperature, each layered specimen was annealed in a Mettler-Toledo hot stage at 180 °C for various times under continuous N2 flow. Visual observation of the postannealed specimens was performed with an Olympus BX60 optical microscope, whereas atomic force microscopy (AFM) of the PMMA/ CSMG surfaces (washed with cyclohexane36 to remove PS as required) was conducted with a Digital Instruments 3000 instrument, equipped with a Si3N4 tip and operated in tapping mode. According to optical microscopy, the top PS layer dewets from the PMMA/CSMG substrate and eventually forms discrete islands under the conditions examined here. This observation is consistent with our previous reports,37,38 including one wherein CSMG nanoparticles were incorporated into the PS (rather than the PMMA) layer as a stabilizing agent (cf. Figure 1b).5 Results from that study indicate that bare PMMA regions from which PS fully dewets appear smooth, whereas regions below PS-rich islands become rough because of CSMG aggregation along the PS/ PMMA interface, as depicted in Figure 1c. The fact that the CSMG nanoparticles did not sink into the PMMA substrate below the PS islands suggests the existence of a nonenthalpic barrier, since the CSMG nanoparticles possess a PMMA shell. A plausible explanation offered for this observation is autophobicity between the CSMG nanoparticles and the highmolecular-weight PMMA chains comprising the bottom layer. In order to test this hypothesis, we have incorporated the CSMG nanoparticles into the PMMA matrix in the present study, and the results upon comparable annealing are remarkably similar to Figure 1c. For instance, the surface of the exposed portion of the PMMA/CSMG film without the added PS capping layer remains smooth, with a measured 3011

Figure 2. Series of AFM height images showing the PMMA surface after different exposure conditions at 180 °C: (a) complete dewetting of the PS capping layer after 80 h; formation and subsequent solvent removal of PS islands after (b) 80 and (c) 20 h; and (d) further annealing of (b) without the PS capping layer for 40 h.

root-mean-square roughness (R) of 0.78 nm, after annealing for 80 h. Figure 2a confirms that the dry regions from which the PS layer dewetted are likewise featureless (R ) 0.83 nm). In both cases, no evidence indicates that CSMG nanoparticles, at least in the form of discernible aggregates, locate atop the PMMA layer. In marked contrast, the PMMA/CSMG surface directly below the dewetted PS islands after 80 h at 180 °C is observed to significantly roughen (R ) 3.51 nm), as illustrated by the AFM height image displayed in Figure 2b. In this case, the PS has been thoroughly removed by a selective solvent (cyclohexane). Bright features measuring ≈30 nm in height are attributed to protruding aggregates of CSMG nanoparticles, which not only segregated from the PMMA homopolymer matrix but also emerged from the PMMA matrix along the PMMA/PS interface. (It is important to recognize that our previous study5 of PS dewetting from PMMA verified that the PMMA/PS interface is featureless with R ) 0.57 nm in the absence of CSMG nanoparticles in either polymer layer.) At shorter annealing times (e.g., 20 h), fewer, discrete features measuring on the order of a few tens of nanometers laterally and approximately 25 nm in height become evident along the PMMA/CSMG surface below the dewetted PS islands (cf. Figure 2c), revealing that the occurrence, as well as extent, of CSMG protrusion is a timedependent and most likely temperature-dependent process. Further annealing of the PS-capped PMMA/CSMG specimen heated for 80 h at 180 °C for 40 h after the PS layer was removed results in total disappearance of the CSMG nanoparticles. At shorter anneal times before and after PS removal (20 h before and 24 h after, as in Figure 2d), the CSMG nanoparticles largely, but not completely, disappear as they sink back into the PMMA matrix, thereby confirming that 3012

emergence of the nanoparticles along the PMMA/PS interface is reversible. While ongoing studies seek to elucidate the temperaturedependent kinetics associated with the CSMG nanoparticles as they emerge and submerge, the mechanism by which the nanoparticles sink back into the PMMA matrix in the absence of PS capping is illustrated in Figure 3a, which shows the periphery of where a PS island resided prior to its dissolution. Numerous CSMG nanoparticles are clearly observed where PS contacted the PMMA/CSMG after 6 h at 180 °C. Just outside this region, CSMG nanoparticles at different levels of submersion are evident, which can be explained by the dynamic nature of the experiment. As schematically depicted in Figure 3b, the CSMG nanoparticles emerge as before along the PMMA/PS interface, but the front of the PS island recedes as dewetting proceeds. The area in which CSMG nanoparticles only partially protrude corresponds to an initially PS-wetted region that ultimately dewetted during the course of thermal treatment. Exposure of the CSMG nanoparticles along the recently exposed (dry) surface forces the nanoparticles back into the matrix. This sequence of events, in conjunction with the results described in Figure 2, indicates that the surface segregation and emergence of CSMG nanoparticles from the PMMA matrix depends on the subtle interplay between entropically driven autophobicity and enthalpically governed surface/interfacial energy. While autophobicity between the CSMG PMMA arms and the PMMA homopolymer drives the nanoparticles to segregate from the PMMA matrix and form protruding aggregates along the surface, the surface energy due solely to PMMA (γPMMA) serves to suppress CSMG protrusion, as evidenced by the absence (or disappearance) of nanoscale surface features. Two structural aspects contribute to autophobicity between free and surface-grafted chains. The first is the high grafting density of the brush, which comprises the outer layer of the CSMG nanoparticles as discussed above. The second factor is the disparity in chain length between the grafted brush and free homopolymer chains. In the present case, the CSMG arms are much shorter (N ) 110) than the PMMA homopolymer (with a degree of polymerization, P, of 2260 repeat units). Pioneering the concept of wet and dry polymer brushes, Leibler and co-workers25,39 have proposed a scaling model to predict the onset of autophobicity. Despite differences in system details, their model can be invoked here for strictly qualitative analysis. According to the model of Leibler and co-workers, free homopolymer chains that are longer than a critical value (P*) given by (Nσ/a2)2/3 are unable to wet a chemically identical brush composed of chains wherein the repeat units are of size a. If the statistical length of a single MMA repeat unit is taken as 0.67 nm, P* ≈ 708 in the present PMMA/CSMG system so that P . P*. In addition, the PMMA chains are also sufficiently long to satisfy the condition of P > 5N required by Liu et al.30 for autophobic dewetting to occur. In synergistic fashion, the high grafting density of CSMG arms and the long chain length of the PMMA homopolymer promote autophobic segregation of the nanoparticles from the PMMA matrix, Nano Lett., Vol. 8, No. 9, 2008

Figure 3. (a) AFM height image showing the PMMA surface after annealing for 6 h at 180 °C. The dashed line marks the periphery where a dewetted PS island resided prior to selective dissolution, and the circle identifies several representative partially protruding CSMG nanoparticles that have begun to sink back into the PMMA matrix upon dewetting-induced surface exposure. A schematic series of top and side views are provided in (b) to illustrate this mechanism of surface energy-induced CSMG submersion as a function of annealing (PS dewetting) time.

which decreases the system free energy (∆Fauto) by an amount proportional to kBT/N1/2a2, where kB is the Boltzmann constant and T denotes absolute temperature. The results presented in Figure 2 and Figure 3 indicate that the propensity for CSMG nanoparticles to emerge alone or as aggregates from the PMMA matrix due to autophobic segregation is regulated by surface (γPMMA) or interfacial (γPMMA/PS) energy. (It is important to recognize that block copolymer micelles would break apart under similar conditions without chemical treatment to induce shape retention.40) Segregation to and subsequent emergence along the PMMA surface would induce surface roughening and incur an energy penalty (∆Fsurf) on the order of kBT/a2, assuming that γPMMA is unaffected by the presence of CSMG nanoparticles with chemically identical arms. This energy increase is an order of magnitude larger than ∆Fauto due to the CSMG arms (N ) 110), in which case γPMMA (reported41 to be ≈30 dyn/ cm) affords strong resistance against CSMG protrusion from the PMMA surface and thereby ensures that the PMMA surface remains (or becomes) relatively smooth. Addition of the PS capping layer reduces the energy barrier from γPMMA to γPMMA/PS, which is about an order of magnitude smaller42 (at ≈1 dyn/cm) than γPMMA, again assuming that the presence of CSMG nanoparticles again has negligible influence. In this case, the interfacial energy is of the same magnitude as ∆Fauto, and our observations in Figure 2 and Figure 3 reveal that ∆Fauto is sufficiently large to offset the increase in free energy due to interfacial roughening (∆Finter) caused by CSMG emergence along the PMMA/PS boundary below isolated PS islands. A similar sequence of events, this time governed again by γPMMA, likewise explains the submersion of CSMG nanoparticles back into the PMMA matrix when the PS islands are removed and the specimens with an exposed PMMA surface are further annealed. The scenario of CSMG surface segregation envisaged in Figure 3 could, in principle, be directly confirmed by crosssectional transmission electron microscopy (TEM) insofar Nano Lett., Vol. 8, No. 9, 2008

as suitable specimen preparation protocols could be developed. Technical difficulties arising from (i) insufficient mechanical integrity of the thin-film assemblies along the PMMA/PS interface (which results in delamination during ultramicrotomy), and (ii) insufficient vapor- or liquid-phase permeation of electron-dense species (e.g., Ru and Br) through the PMMA shell of the CSMG nanoparticles to selectively stain, and hence enhance the contrast of, the PS core (which is needed to identify the spatial location of the nanoparticles) preclude use of the present CSMG nanoparticles for this purpose. Thus, a new species-reversed nanoparticle, hereafter designated CSMG-R, has been synthesized with an acrylic core derived from ethylene glycol dimethacrylate (EGDMA) and PS arms. In this case, the PS arms can be selectively stained, thereby permitting direct visualization of the unstained CSMG-R cores. According to the experimental procedures provided in Supporting Information, the Mn and PI of the CSMG-R nanoparticles are similar to those of the CSMG nanoparticles, namely, 376 kDa and 1.16, respectively. The corresponding characteristics of the PS arms on the CSMG-R nanoparticles are likewise comparable to those of the PMMA arms on the CSMG nanoparticles: Mn ) 10.6 kDa, PI ) 1.01 and 32 arms/nanoparticle. Although attempts to cross-section PS/PMMA bilayered assemblies containing CSMG-R nanoparticles in the top PS film still failed because of the weak PS/PMMA interface, the CSMG-R nanoparticles can alternatively be incorporated into a thin, spin-coated film of PS (Mn ) 216 kDa, PI ) 1.06) welded between two thick solvent-cast films of highmolecular-weight PS (Mn ) 900 kDa, PI ) 1.10), which is considerably more viscous than the 216 kDa PS. Homopolymer entanglements induced during annealing at 180 °C for 24 h keep the trilayered assembly intact during sectioning, whereas the viscosity difference between the PS layers ensures that the CSMG-R nanoparticles do not quickly migrate from the sandwiched thin film because of the 3013

Figure 4. Energy-filtered TEM images of CSMG-R nanoparticles acquired from (a) the interior of a thin PS film and (b,c) the interface between 216 kDa PS and 900 kDa PS after annealing at 180 °C for 24 h. The PS homopolymer matrices and arms of the CSMG-R nanoparticles are selectively stained (see Supporting Information for details) and appear electron-dense (dark), whereas the light features correspond to the unstained acrylic cores of the nanoparticles. The regions separated by dashed lines in (b) are discussed in the text. The circled feature in (a) and the arrow in (b) highlight examples of large nanoparticle aggregates.

nanoparticle concentration gradient. A series of energyfiltered TEM images of stained cross sections of the annealed trilayered assembly is presented in Figure 4 and reveals 3014

several key morphological characteristics. The first such feature evident in Figure 4a is the presence of light dispersions, attributed to the unstained acrylic cores of the CSMG-R nanoparticles, uniformly distributed in an electrondense (stained) PS matrix. The smallest dispersions measure about 5-15 nm in diameter, whereas the average dispersion size is closer to 20-25 nm, although these and larger dispersions (such as the circled one in Figure 4a) most likely represent nanoparticle aggregates and not individual nanoparticles. Note that the image width greatly exceeds the thickness of the thin film, which implies that the sectioning/ viewing direction is not orthogonal to the polymer/polymer interfaces. In Figure 4b, discrete dispersions are visible in the lefthand region (I) of the image but disappear for the most part in the region labeled II. This apparent lack of dispersions in Figure 4b and the enlargement in Figure 4c may reflect a reduction in CSMG-R nanoparticles in region II, which is unlikely due to the uniform distribution seen in images such as the one displayed in Figure 4a. Conversely, it may be a consequence of a substantial increase in nanoparticles, which would explain why the diffuse features in Figure 4a are absent, while large, irregularly shaped aggregates (cf. Figure 4b) exist. If, however, the specimen was tilted during sectioning, as alluded to above, then the polymer/polymer interface would also appear expanded in projection. Just beyond the polymer/polymer interface in region III of Figure 4b, CSMG-R nanoparticle aggregates are seen to extend from the 216 kDa PS thin film into the 900 kDa PS thick film. It is important to recognize that such nanoparticle extension from the host film into an immiscible polymer, such as PMMA, is expected to be considerably less pronounced because of enthalpically repulsive interactions. These observations of a similar, but not identical, layered assembly containing species-reversed CSMG nanoparticles are generally consistent with the postulated picture of nanoparticle enrichment/emergence along the polymer/polymer interface illustrated in Figure 3. Our results and explanation more broadly suggest that the autophobic segregation/emergence of CSMG nanoparticles, as well as similarly interacting nanoscale objects, along a polymer surface can be reversibly controlled (i.e., switched) by manipulating the surface/interfacial energy along the surface. The generality of this approach has been tested by replacing the PS capping layer that stochastically dewets upon annealing with a poly(dimethylsiloxane) (PDMS, Dow Corning Sylgard 184) film that is permanently and periodically corrugated (“wrinkled”) due to chemically cross-linking under mechanical deformation.43 This stamp allows contact only between the PMMA/CSMG film and the ridges of the PDMS wrinkles, while leaving other areas of the PMMA surface exposed. After annealing the system for 77 h at 180 °C and removal of the corrugated PDMS by peeling (no solvent was required), stamp transfer is visibly observed on the PMMA/CSMG film as a pattern of alternating bright stripes that signify regions of contact, as seen in Figure 5a. Subsequent AFM analysis (cf. Figure 5b) indicates that CSMG aggregates emerge from the PMMA film only in areas Nano Lett., Vol. 8, No. 9, 2008

and the principles by which CSMG nanoparticles undergo surface segregation and patterning are general enough to suggest that other nanoparticulate systems14 may be designed with similar capability. Acknowledgment. This work was supported by the Kenan Institute for Engineering, Technology and Science at North Carolina State University and a Linkage-International project sponsored by the Australian Research Council (G.G.Q. and R.J.S.). We thank Dr. K. Efimenko for providing the PDMS stamp. Supporting Information Available: Details of materials, equipment, synthesis, and microscopy. This material is available free of charge via the Internet at http://pubs.acs.org. References Figure 5. Soft stamping as a route to spatially pattern CSMG nanoparticles protruding from the bulk PMMA. Periodic surface patterning on the PMMA/CSMG film (a) is achieved by pressing a corrugated PDMS elastomer43 (a scanning electron micrograph is included in the inset and the periodicity λ is labeled) against the film for 77 h at 180 °C, followed by peeling. Within contact regions (white lines), CSMG nanoparticles emerge from the film, as seen in the AFM image (5 µm × 5 µm scan area, 60 nm height marker) displayed in (b). Illustrations of this patterning process, provided in (c), depict the cross-sectional spatial distribution of CSMG nanoparticles relative to the PDMS stamp.

where the film contacted the PDMS ridges. This result, illustrated in Figure 5c, corroborates that autophobicity is sufficient to overcome polymer/polymer interfacial energy (γPMMA/PDMS in this case) and promote CSMG emergence from the substrate, whereas surface energy suppresses autophobic segregation. By further annealing the patterned film for an additional 40 h, the CSMG aggregates within the stripes sink back into the matrix. Thus, by using a simple elastomeric stamp to selectively contact the PMMA/CSMG film, we have demonstrated that protruding CSMG nanoparticles, first noticed under the stochastically dewetted PS capping layer, can be controllably patterned and reversibly switched by further annealing performed with and without (in alternating fashion) the contacting stamp. The results of this study establish that autophobicity between the dense brush shell of CSMG nanoparticles and long homopolymer chains constitutes a sufficiently strong driving force to overcome polymer/ polymer interfacial energy. Conversely, polymer surface energy prevents CSMG nanoparticles from emerging by keeping the homopolymer film surface smooth. By manipulating the energetic nature of the film surface, we can reversibly control the extents to which (i) autophobic segregation occurs and (ii) the spatial distribution of CSMG nanoparticles can be patterned on a polymer surface. We propose that the relevant time scales can be shortened by reducing the matrix viscosity through the use of temperature or plasticizing agents. Controllable transport, surface patternability, and nanoscale size of the particles make them attractive for nanotechnologies requiring, for instance, highdensity memory, sensory capability or switchable optics,44 Nano Lett., Vol. 8, No. 9, 2008

(1) (a) Inoue, K. Prog. Polym. Sci. 2000, 25, 453. (b) Jikei, M.; Kakimoto, M. Prog. Polym. Sci. 2001, 26, 1233. (2) Jikei, M. Kakimoto Prog. Polym. Sci. 2001, 26, 1233. (3) Solomon, D. H.; Qiao, G. G.; Abrol, S. WO 99/58588, The University of Melbourne, 1999. (4) Gurr, P. A.; Qiao, G. G.; Solomon, D. H.; Harton, S. E.; Spontak, R. J. Macromolecules 2003, 36, 5650. (5) Wei, B.; Gurr, P. A.; Genzer, J.; Qiao, G. G.; Solomon, D. H.; Spontak, R. J. Macromolecules 2004, 37, 7857. (6) Wiltshire, J. T.; Qiao, G. G. Macromolecules 2006, 39, 9018. (7) Connal, L. A.; Gurr, P. A.; Qiao, G. G.; Solomon, D. H. J. Mater. Chem. 2005, 15, 1286. (8) Zhang, Q.; Remsen, E. E.; Wooley, K. L. J. Am. Chem. Soc. 2000, 122, 3642. (9) Wooley, K. L. J. Polym. Sci. A: Polym. Chem. 2000, 38, 1397. (10) Sauer, M.; Streich, D.; Meier, W. AdV. Mater. 2001, 13, 1649. (11) Kaneda, I.; Vincent, B. J. Colloid Interface Sci. 2004, 274, 49. (12) Riess, G. Prog. Polym. Sci. 2003, 28, 1107. (13) Joralemon, M. J.; O’Reilly, R. K.; Hawker, C. J.; Wooley, K. L. J. Am. Chem. Soc. 2005, 127, 16892. (14) Caruso, F. AdV. Mater. 2001, 13, 11. (15) Bockstaller, M. R.; Mickiewicz, R. A.; Thomas, E. L. AdV. Mater. 2005, 17, 1331. (16) Kim, B. J.; Chiu, J. J.; Yi, G. R.; Pine, D. J.; Kramer, E. J. AdV. Mater. 2005, 17, 2618. (17) Spontak, R. J.; Shankar, R.; Bowman, M. K.; Krishnan, A. S.; Hamersky, M. W.; Samseth, J.; Bockstaller, M. R.; Rasmussen, K. Ø. Nano Lett. 2006, 6, 2115. (18) (a) Ho, A. K.; Iin, I.; Gurr, P. A.; Mills, M. F.; Qiao, G. G. Polymer 2005, 46, 6727. (b) Connal, L. A.; Qiao, G. G. AdV. Mater. 2006, 18, 3024. (19) Kato, T. F.; K., n/a.; Kawaguchi, H. Polym. Gels Networks 1994, 2, 307. (20) Kuckling, D.; Vo, C. D.; Wohlrab, S. E. Langmuir 2002, 18, 4263. (21) Lapeyre, V.; Gosse, I.; Chevreux, S.; Ravaine, V. Biomacromolecules 2006, 7, 3356. (22) Wei, B.; Spontak, R. J.; Genzer, J., in preparation. (23) Kiser, P. F.; Wilson, G.; Needham, D. Nature 1998, 394, 459. (24) Freedman, K. O.; Lee, J.; Li, Y. G.; Luo, D.; Skobeleva, V. B.; Ke, P. C. J. Phys. Chem. B 2005, 109, 9839. (25) Leibler, L.; Ajdari, A.; Mourran, A.; Coulon, G.; Chatenay, D. In Ordering in Macromolecular Systems; Teramoto, A., Kobayashi, M., Norisue, T., Eds.; Springer-Verlag: Berlin, 1994; p 301. (26) Shull, K. R. Faraday Discuss. 1994, 98, 203. (27) Hamdoun, B.; Ausserre, D.; Cabuil, V.; Joly, S. J. Phys. II 1996, 6, 503. (28) Orso, K. A.; Green, P. F. Macromolecules 1999, 32, 1087. (29) Liu, Y.; Rafailovich, M. H.; Sokolov, J.; Schwarz, S. A.; Zhong, X.; Eisenberg, A.; Kramer, E. J.; Sauer, B. B.; Satija, S. Phys. ReV. Lett. 1994, 73, 440. (30) Kerle, T.; Yerushalmi-Rozen, R.; Klein, J. Macromolecules 1998, 31, 422. (31) Borukhov, I.; Leibler, L. Phys. ReV. E 2000, 62, R41. (32) Reiter, G.; Khanna, R. Phys. ReV. Lett. 2000, 85, 2753. (33) Mu¨ller, M.; MacDowell, L. G. Europhys. Lett. 2001, 55, 221. (34) MacDowell, L. G.; Mu¨ller, M. J. Phys.; Condens. Matter. 2005, 17, S3523. 3015

(35) Carignano, M. A.; Szleifer, I. J. Chem. Phys. 1994, 100, 3210. (36) Recent results reported by Harton, S. E.; Luning, J.; Betz, H.; Ade, H. Macromolecules 2006, 39, 7729. suggest that selective dissolution of PS with cyclohexane is not as effective as 1-chloropentane. Using NEXAFS spectroscopy, we have recently confirmed this observation and have found evidence of residual PS on the PMMA surface after cyclohexane treatment. Despite this observation, we do not believe that our conclusions would be detrimentally affected, as confirmed by the use of the cross-linked PDMS stamp. (37) Wei, B.; Genzer, J.; Spontak, R. J. Langmuir 2004, 20, 8659. (38) Wei, B.; Lam, P. G.; Braunfeld, M.; Agard, D. A.; Genzer, J.; Spontak, R. J. Langmuir 2006, 22, 8642.

3016

(39) Cloitre, M.; Borrega, R.; Leibler, L. Phys. ReV. Lett. 2000, 85, 4819. (40) Alexandridis, P.; Hatton, T. A. Colloids Surf. A 1995, 96, 1. (41) Wu, S. Polymer Interface and Adhesion; Marcel Dekker: New York, 1982. (42) Carriere, C. J.; Biresaw, G.; Sammler, R. L. Rheol. Acta 2000, 39, 476. (43) Efimenko, K.; Rackaitis, M.; Manias, E.; Vaziri, A.; Mahadevan, L.; Genzer, J. Nat. Mater. 2005, 4, 293. (44) Serpe, M. J.; Kim, J.; Lyon, L. A. AdV. Mater. 2004, 16, 184.

NL802109X

Nano Lett., Vol. 8, No. 9, 2008