Organization of Nanoparticles on Soft Polymer Surfaces - American

Organization of Nanoparticles on Soft. Polymer Surfaces. Zhen Liu, Kristen Pappacena, Jane Cerise, Jaeup Kim, Christopher J. Durning,. Ben O'Shaughnes...
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NANO LETTERS

Organization of Nanoparticles on Soft Polymer Surfaces

2002 Vol. 2, No. 3 219-224

Zhen Liu, Kristen Pappacena, Jane Cerise, Jaeup Kim, Christopher J. Durning, Ben O’Shaughnessy, and Rastislav Levicky* Department of Chemical Engineering, Columbia UniVersity, New York, New York 10027 Received September 7, 2001; Revised Manuscript Received December 14, 2001

ABSTRACT We investigate the organization of gold nanocrystals on thin, fluid films consisting of polymer chains tethered by one end to an underlying substrate in a polymer brush configuration. The thickness of the polymer brush is comparable to the nanocrystal size. Thinner polymer brushes are found to suppress aggregation of the nanoparticles, leading to stable, elongated particle-rich domains. The results suggest new approaches for modification of macroscopic surfaces with nanoscopic particles.

A number of promising fabrication routes to nanostructures have been demonstrated that rely on the order and interactions of “soft” synthetic or biological materials. Of close relation to the present work, several groups have developed methods for templating spatial distributions of nanometer sized particles via in situ synthesis of nanoparticles in ordered block copolymer matrices,1 by introduction of premade nanoparticles into block copolymers or by deposition on their surfaces,2 or via other routes that exploit the soft structure of synthetic polymers (e.g., micelles,3 multilayered structures4). Complementing experimental efforts, theoretical5,6 and computer simulation7 investigations of nanoparticle/ polymer composites have begun to further detail mechanisms responsible for development of structure and physical properties. Prior approaches have typically relied on chemical information to direct the organization of nanoparticles within the polymer environment; for instance, to localize nanoparticles in domains of a particular block in a block copolymer melt due to preference for that chemical environment. This letter reports on organization of nanoparticles in polymer films whose thickness is comparable to the particle diameter. The polymers are attached by one end to an underlying surface, in a “polymer brush” motif.8 A key motivation for the study was to determine how the constraint of end attachment affects distribution of the nanoparticles. For example, one expectation was that it may suppress macroscopic phase separation of particles and polymer because lateral migration of the brush chains is prevented. The ability to control aggregation of nanoparticles should be useful in applications of these nanomaterials to modification of surface properties; for example, to tune uniformity * Corresponding author. Tel: +1-212-854-2869. E-mail: RL268@ columbia.edu. 10.1021/nl015625p CCC: $22.00 Published on Web 01/09/2002

© 2002 American Chemical Society

of coverage of magnetic or luminescent particles on large area substrates. The investigations used polydisperse gold nanocrystals covered by a shell of dodecanethiol ligands. The nanocrystals were synthesized according to previously detailed protocols.9 Transmission electron microscopy (TEM, Philips EM 430) revealed the average gold core diameter to be 2.5 ( 0.7 nm. The dodecanethiol shell on the particles is used to stabilize the gold cores against aggregation; in addition, it serves to prevent direct adsorption of polymer segments to the nanoparticle cores, a process implicated in anomalously slow diffusion kinetics of bare gold nanoparticles in polymers such as poly(tert-butyl acrylate).10 Composite films of nanocrystals and polymers were prepared by sequential spin coating of polymer and nanoparticles onto circular glass coverslips that were precleaned in a 7:3 solution of concentrated sulfuric acid and 30% hydrogen peroxide in water. WARNING: This solution is extremely corrosiVe and should not be stored in tightly sealed containers due to eVolution of gas. First, an adhesion monolayer of a diblock copolymer poly(styrene-b 2-vinylpyridine) (PS-P2VP; 248 000 dalton total mass; 119 000 dalton PS block) was spin coated from 9:1 toluene/ methanol to prevent dewetting of the glass surface by the polymer film. Second, a layer of PS homopolymer was deposited from toluene; this layer served to provide mechanical support for subsequent transfer of films to Formvar coated carbon TEM grids or CaF2 windows. Third, a monolayer of a diblock copolymer poly(styrene-b ethylenepropylene) (PS-PEP; 110 000 dalton total mass; 70 000 dalton PEP block; 0.67 PEP volume fraction; polydispersity 1.04), of variable thickness, was coated from 9:1 octane/ toluene to provide material for the polymer brush. The polymer films were annealed (150 °C, 24 h, vacuum) to order

Figure 1. Structure of the investigated samples. A layer of PSPEP diblock copolymer is topmost, with the PEP blocks forming a polymer brush at the air interface. The PS blocks anchor the polymer brush to an underlying PS homopolymer film, and the bottommost layer of PS-P2VP serves to promote adhesion to the glass substrate.

them into the structure depicted in Figure 1, in which the PS-PEP chains have their PS blocks anchored in the underlying PS homopolymer and their PEP blocks segregated to the film-air interface, forming a PEP brush. When thickness of the PEP brush surpassed 20 nm, optical and atomic force microscopy revealed that extra PS-PEP layers formed on top of the brush, leading to the formation of multilayers. The range of brush thickness investigated (2 nm to 13 nm) was deliberately kept below the 20 nm limit so that all of the PS-PEP polymer could be incorporated into a single brush monolayer. Segregation of PEP to the free surface of the film was checked by water contact angle measurements (PS reference film: 80°; PEP reference film: 98°; polymer brush: 95°), by neutron and X-ray reflectivity experiments (see below), and by angle-dependent X-ray photoelectron spectroscopy (XPS; Perkin-Elmer PHI 5500). For a 5 nm thick brush, changing the XPS takeoff angle from 17° to 45° to 70° (respective sampling depths of 3, 7, and 9.4 nm) resulted in an increased ratio of a shake-up satellite C1s signal at 291 eV, originating uniquely from the PS benzene rings, to the main C1s signal at 284 eV, originating from PEP, dodecanethiol ligands, as well as PS. For these takeoff angles, the satellite-to-main ratio increased from ∼0.0 to 0.009 to 0.015. The ratio was 0.073 for a pure PS film annealed at 150 °C. The improved “visibility” of the PS satellite signal at more perpendicular takeoff angles indicates that the PS is indeed buried underneath an overlayer of PEP. Nanoparticles were spin coated from octane directly on top of the ordered polymer films. The nanocrystals were allowed to diffuse and organize on the PEP brush for various times, either at room temperature (∼22 °C) or at 60 °C under a nitrogen atmosphere. Under these conditions, the entire polymer film is glassy save for the topmost PEP brush layer, for which the glass transition temperature is -60 °C.11 The brush chains cannot migrate along the surface because they are tethered to the glassy PS underlayer. In addition, the annealing temperatures are above the melting transition for the dodecanethiol shell, which is therefore expected to be disordered and liquid-like (differential scanning calorimetry placed the melting transition 220

at around 5 °C in similarly sized gold nanoparticles;12 molecular dynamics simulations have reported the melting transition increasing from 7 to 21 °C with particle size13). Infrared absorption measurements on ∼20 nm thick control films of pure nanoparticles, prepared in parallel with each set of samples, verified that desorption of thiol ligands during annealing remained below 5%. Films were separated from the glass coverslip supports by floating on deionized water, following a 10 s 1 M NaOH etch, and picked up on TEM grids or CaF2 plates for subsequent characterization. The samples were stored in a class 1000 cleanroom when not used. Specimen composition was analyzed by decomposing infrared absorption spectra, measured in transmission from films on CaF2 plates, into a linear superposition of calibrated spectra of the individual PS, PEP, and P2VP components. These measurements determined a thickness of 22 ( 1 nm for the total PS content, and 3.2 ( 1 nm for P2VP (an unknown amount of PS-P2VP polymer may have been retained on the original glass support during floating). Thickness values for the PEP brush, which were adjusted from experiment to experiment, are reported below. Lateral organization of the nanoparticles in the PEP film was imaged with a JEOL JEM-100C TEM (100 keV). Particle coverages were extracted from TEM negatives using ImageJ software (http://rsb.info.nih.gov/ij/). Figures 2 and 3 show TEM images of gold nanocrystals in PEP brushes of two different thicknesses, 13 ( 2 nm and 5 ( 1.5 nm. For each brush thickness, three images corresponding to progressively longer annealing times are shown. The size scale is much larger than that of an individual nanoparticle; for example, the circled “island” in Figure 2a contains about 500 particles. At higher magnifications, individual particles are distinguished (inset in Figure 2c). For all samples, surface coverage of particles was between 4% and 6%. In this paper, “coverage” refers to surface fraction occupied by nanoparticles’ gold cores only, as seen in the TEM image (inclusion of the particles’ ∼1 nm thick dodecanethiol shell would translate the surface coverage to about 15%). Vertical stacking of particles, in a multilayer arrangement, was not observed in any of the experiments. The following observations are evident from Figures 2 and 3. (i) Phase Separation. PEP polymer and the gold nanoparticles phase separate into particle-rich and particle-poor domains. Consistent with this observation, nanoparticles also did not disperse in bulk PEP homopolymer “solvent” of the same 70 000 dalton molecular weight as the brush-forming PEP chains. (ii) Effect of Annealing Time. Reorganization of nanoparticles occurs upon increasing annealing time from 25 h at room temperature (RT) to 50 h at RT plus 15 h at 60 °C (Figures 2a and 2b; 3a and 3b). An increase in the average size of particle-rich domains is observed in both the 13 nm and the 5 nm brush; however, aggregates in the thicker brush are much larger. Additional annealing for 15 h at RT plus 10 more hours at 60 °C resulted in only modest additional change (Figures 2b and 2c; 3b and 3c). Nano Lett., Vol. 2, No. 3, 2002

Figure 2. TEM images of nanoparticle organization on a 13 nm thick brush. Annealing times increase from (a) to (c) as specified in the lower left of each image. In the upper left, areal percentage occupied by the particles’ gold cores is indicated. The circled nanocrystal aggregate in (a) consists of about 500 particles. The inset in (c) is an enlarged view.

Figure 3. TEM images of nanoparticle organization on a 5 nm thick brush. Annealing times increase from (a) to (c) as shown. Areal percentage occupied by the particles’ gold cores is displayed in the upper left of each image.

(iii) Effect of Film Thickness. In the thicker, 13 nm brush, extensive particle-rich domains can be seen, containing around 1000 nanoparticles (Figures 2b and 2c). These large aggregates are interspersed with smaller ones. Irrespective of size, domain shape tends to be rounded. Similar observations were obtained on a 9 nm thick brush. In contrast, aggregates in the 5 nm brush (Figures 3b and 3c) are ∼200 particles or less, and their shape tends to be asymmetric and ribbon-like above an aggregate size of about 30 nanoparticles. The width of these ribbons around the narrower dimension corresponds to about 5 nanoparticles, or ∼20-30 nm. A number of smaller, roughly circular domains are also present. Typically, the diameter of these circular domains does not exceed 30 nm. An even thinner, ∼2 nm thick brush, yielded qualitatively similar observations (asymmetric and circular

aggregates) as the 5 nm brush, albeit aggregate dimensions and shapes were more broadly distributed. Before discussing the above observations, we note that nanoparticles spin coated directly on PS homopolymer at similar coverages did not resemble any of the brush specimens. Rather, a random arrangement of many isolated particles and small, few particle aggregates persisted even after annealing at 60 °C. The particle distribution on these “brushless” samples is likely far from equilibrium, locked in by low mobility of the nanoparticles on the PS film (though large, multilayered aggregates could be realized if the particles were deposited by precipitation from an evaporating but otherwise quiescent solution). These control experiments demonstrate that the presence of the fluid PEP brush enhances particle mobility along the surface.

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Figure 4. Electron density profile (Fe) as a function of depth (z) for a sample structure similar to that illustrated in Figure 1. z ) 0 at the air-film interface. The PEP brush is located between 0 and 12 nm and corresponds to a 10 nm thick PEP layer with 20% coverage of Au. The profile shows an accumulation of gold nanoparticles at the film-air interface. Inset: Specular X-ray reflectivity data (points) and calculated fit (line) based on the plotted Fe profile. The term qZ, the z-component momentum transfer, is given by 4π sinθ/λ (θ: angle of reflection; λ: wavelength).

A second point concerns penetration of nanocrystals into the brush interior. The vertical particle distribution was investigated using X-ray reflectivity. Figure 4 shows an electron density profile determined for a 10 nm thick PEP brush with a 20% surface coverage of Au. The higher particle loading was used to enhance sensitivity to the presence of the nanoparticles. The experimental reflectivity data, together with a calculated fit based on the plotted density profile, are shown in the inset. The sharp increase in electron density near z ) 0, where the brush-air interface is located, is attributed to an enhanced concentration of gold from the nanoparticles. Therefore, reflectivity data indicate that nanoclusters preferentially segregate to the film-air interface. The segregation would be expected to reflect contributions both from particle-PEP incompatibility as well as differing surface tensions against air, though the relative magnitude and sign of these effects is not known at present. A detailed discussion of the reflectivity measurements and surface segregation behavior will be reported separately. Segregation of nanoparticles to the brush-air interface was also consistent with angle-resolved XPS measurements of the Au4f signals. Furthermore, coarsening of the particle domains did not lead to changes in the relative C1s, Au4f, and S2p signal strengths in XPS spectra measured at a 70° takeoff angle. At this nearnormal takeoff condition and for the thinner 5 nm brush, the 291 eV C1s shake-up satellite peak of the buried PS layer was also observable, albeit weakly. Reproducibility of relative signal strengths before and after coarsening of the particle domains suggests that, within the sensitivity of XPS, the extent of particle aggregation does not perturb the vertical arrangement of PEP, PS, and nanoparticles. We now return to considering the lateral particle organization (Figures 2 and 3). The observed aggregation of nanocrystals points to the existence of attractive interactions between the particles. The strength of these attractions will be affected by the (unknown) extent of particle immersion in the brush. Assuming that particles are fully immersed, it 222

is straightforward to estimate a lower limit of ∼0.1 kT on the attraction at contact between two particles14,15 (k: Boltzmann constant; T: absolute temperature; room temperature conditions). On the other hand, if the nanoparticles are fully separate from the brush, then their attractive interactions should be of the same order of magnitude as the energy difference (the cohesive energy) between an isolated, dodecanethiol passivated gold nanoparticle and one embedded in a superlattice aggregate of such particles. In this case, particle-particle attractions are expected to reach ∼100 kT16 (computer simulations of ref 16 found a 15 eV, or ∼600 kT, cohesive energy per particle of 140 Au atoms passivated by 62 dodecanethiol ligands, located in a distorted fcc lattice of such particles). The cohesive energy arises chiefly from interactions between dodecanethiol ligands which, however, should be strongly screened by contact with the chemically similar, hydrocarbon PEP medium. Albeit the range of possible interaction magnitudes (0.1 kT to 100 kT) quoted above is very broad, the actual value is suspected to fall between 1 kT and 10 kT. This is because isolated particles can be seen in the TEM images, often at a fraction of about one percent of the total particle number; for interactions weaker than kT there should be many more isolated particles, while for interactions stronger than 10 kT there should be virtually none. The pronounced changes brought about by increasing annealing time (Figure 2a,b and 3a,b) from 1 to 2.5 days clearly indicate that, over this time scale, particles are able to move and rearrange in both the 13 and 5 nm films. In the thicker brush, as time progresses, large nanocrystal domains grow at the expense of smaller ones (Figure 2a,b). Such a trend may reflect an Ostwald ripening mechanism, in which coarsening of a phase separated system occurs by diffusion of individual particles from smaller to larger domains.17 In three dimensions, Ostwald ripening is driven by a decrease in surface free energy as larger domains grow, since fewer particles find themselves in energetically unfavorable surface sites. In two dimensions, a situation more similar to the present experiments, line tension plays the role of surface energy. Theory predicts the average domain size to increase with time to the 1/3 power.18 While this weak, sublinear dependence is qualitatively consistent with the slowing of aggregate growth with time observed from Figure 2a to Figure 2c, investigations of large numbers of aggregates under thermodynamically constant conditions would be needed for a more detailed analysis. In contrast to the thick brush, in the thinner 5 nm film only a microphase (local) phase separation is observed, with aggregate size smaller by about an order of magnitude compared to the thicker film. One possible cause for the stifled phase separation is suppression of particle mobility, so that aggregate growth is slowed. However, changes in particle organization from Figure 3a to 3b clearly indicate that rounded aggregates of ∼10 nanocrystals are transformed to larger, elongated aggregates of up to ∼200 nanocrystals over the duration of about two and a half days. This observation argues against a drastic reduction in particle mobility compared to the thicker specimen, though does not Nano Lett., Vol. 2, No. 3, 2002

Figure 5. Postulated mechanisms of nanocrystal aggregation. (a) In a thick brush, nanocrystals diffuse and aggregate at the brushair interface. (b) In a thin brush, a constrained lateral phase separation of nanocrystals and polymer occurs.

rule out a modest decrease. Allowing that significant particle mobility must be present, the observation of elongated rather than rounded aggregates is puzzling. Rounding is expected based on minimization of line tension, and could be achieved by migration of particles over relatively short distances along the perimeter of a domain. These considerations led to the postulation that an equilibrium mechanism is at play; in other words, a free energetic contribution is present that favors formation of asymmetric domains in the thinner film. A mechanism that can account for differences in nanocrystal organization between thick and thin samples is depicted in Figure 5. For the thick sample (Figure 5a), nanoparticles float on several nm of liquid PEP medium, diffusing and aggregating at the brush-air interface with relatively little perturbation to the underlying polymer matrix. This kinetic process results in growth of extensive nanocrystal domains as time progresses, with the domains tending to a rounded shape to minimize line tension. In contrast, the thin layer (Figure 5b) cannot support the nanoparticles as effectively and individual chains, forced to stretch around the particle aggregates, are strongly deformed. The deformation cannot be relaxed because the chains are immobilized at one end, preventing their lateral mobility. As a nanocrystal domain grows perturbation to the surrounding polymer matrix increases until, eventually, further growth of the particle aggregate becomes unfavorable. At this stage, a limit on aggregate size is reached. Elongated aggregates would arise naturally above a certain size because their asymmetric shape allows polymer chains to stretch around the narrower (minor) dimension of an aggregate, even as aggregate growth proceeds along the perpendicular (major) direction. The elongated shape would then reflect a compromise between particle attractions driving aggregate growth along the major direction while chain stretching suppresses it in the other. This hypothesis, which argues that deformation of polymer chains is a free energetic force that affects organization of the nanocrystals, is in the same spirit as theoretical models that implicate chain stretching in producing effective particle repulsions in polymer brushes6 and in influencing the compatibility of particles with block copolymer matrices.7a An important check of the above-proposed mechanism for the thin brush can be provided by allowing the brush chains Nano Lett., Vol. 2, No. 3, 2002

Figure 6. TEM image of nanoparticle organization on a 5 nm thick brush after 4 h of annealing at 120 °C.

to diffuse along the surface, since that should allow formation of large particle aggregates without the associated chain stretching penalties. Figure 6 shows particle organization in a 5 nm thick brush that, after deposition of the nanoparticles, was annealed for 4 h at 120 °C. This temperature is above the glass transition of PS (100 °C), so that the PS anchor blocks of the brush chains are expected to be mobile. As is evident from the figure, even for the short 4 h annealing duration, significantly more extensive aggregation of nanoparticles is observed than for a brush with immobile chains (Figure 3). This evidence lends credence to the argument that chain stretching suppresses aggregate growth in immobile brushes. However, some caution in interpreting these results is necessary, as under the harsher 120 °C annealing about 12% of the thiol ligands on the nanoparticles desorbed, implying that particle-particle interactions may also have been perturbed. Stretching of a polymer chain becomes energetically significant (i.e., comparable to kT) once a chain is induced to stretch a distance comparable to its unperturbed radius of gyration RG.19 It is only after aggregate size reaches this characteristic length that chain stretching could be a viable mechanism for preventing further growth of an aggregate. Observation of aggregates stabilized at dimensions much smaller than RG would undermine the proposed model since chain stretching over such short distances is energetically insignificant and therefore should not strongly influence organization of the nanoparticles. On the other hand, costs associated with stretching chains over distances many times RG become increasingly prohibitive. For a 70 000 dalton PEP chain, RG ) 10 nm.20 From the TEM data, the narrower dimension of the aggregates falls in the range 20 to 30 nm, comparable to RG and consistent with above expectations. More generally, the combination of short-range attractive interactions (e.g., interparticle attractions) favoring growth of an aggregated phase with competing contributions that oppose it above a certain domain size (e.g., chain stretching when aggregate size begins to exceed RG) is a hallmark of so-called modulated phases.21 A common observation in many modulated phase systems is that, in two dimensions, circular domains are observed at low densities, with gradual 223

transition to ribbon-like elongated domains at higher surface coverages. In agreement with these general considerations, in the present system a greater fraction of elongated domains is observed at the higher 5.4% particle coverage of Figure 3b than at the 4.3% coverage of Figure 3c. Furthermore, TEM images for particle coverage of 8% revealed almost exclusively elongated particle aggregates. This report explored use of polymer brushes to direct organization of nanoparticles on surfaces. Using a system in which the polymer and particles prefer to phase separate, control over the extent of phase separation was obtained by variation of brush thickness. Strong suppression of the growth of particle-rich domains was observed in brushes with thickness comparable to the diameter of the particles, an effect that was rationalized in terms of chain stretching penalties that arise due to perturbation of inner brush structure by the growing particle aggregates. The described approach should be useful in applications in which nanoparticles are employed to modify properties of macroscopic surfaces. Acknowledgment. Dr. Axel Hoffman is gratefully acknowledged for his help with the X-ray reflectivity measurements. The authors are also thankful to Louis Brus, SiuWai Chan, Alex Couzis, Jeff Koberstein, and Feng Zhang for valuable input and assistance. Shell Chemical Company donated the poly(styrene -b ethylenepropylene) diblock copolymer (Kraton G1701). This work was supported primarily by the MRSEC Program of the National Science Foundation under Award Number DMR-9809687. References (1) (a) Chan, Y. N. C.; Craig, G. S. W.; Schrock, R. R.; Cohen, R. E. Chem. Mater. 1992, 4, 885. (b) Ciebien, J. F.; Clay, R. T.; Sohn, B. H.; Cohen, R. E. New J. Chem. 1998, 22, 685. (c) Bronstein, L.; Seregina, M.; Valetsky, P.; Breiner, U.; Abetz, V.; Stadler, R. Polym. Bull. 1997, 39, 361. (2) (a) Fogg, D. E.; Radzilowski, L. H.; Dabbousi, B. O.; Schrock, R. R.; Thomas, E. L.; Bawendi, M. G. Macromolecules 1997, 30, 8433. (b) Lin, B. H.; Morkved, T. L.; Meron, M.; Huang, Z. Q.; Viccaro, P. J.; Jaeger, H. M.; Williams, S. M.; Schlossman, M. L. J. Appl. Phys. 1999, 85, 3180. (c) Zehner, R. W.; Lopes, W. A.; Morkved, T. L.; Jaeger, H.; Sita, L. R. Langmuir 1998, 14, 241. (d) Hamdoun, B.; Ausserre, D.; Joly, S.; Gallot, Y.; Cabuil, V.; Clinard, C. J. Phys. II 1996, 6, 493. (3) (a) Spatz, J. P.; Roescher, A.; Moller, M. AdV. Mater. 1996, 8, 337. (b) Bronstein, L.; Kramer, E.; Berton, B.; Burger, C.; Forster, S.; Antonietti, M. Chem. Mater. 1999, 11, 1402. (4) (a) Kotov, N. A.; Dekany, I.; Fendler, J. H. J. Phys. Chem. 1995, 99, 13065. (b) Schmitt, J.; Decher, G.; Dressick, W. J.; Brandow, S. L.; Geer, R. E.; Shashidhar, R.; Calvert, J. M. AdV. Mater. 1997, 9,

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Nano Lett., Vol. 2, No. 3, 2002