Orientation Control of Block Copolymer Thin Films Placed on Ordered

Oct 3, 2013 - School of Chemical and Biological Engineering, The National Creative Research Initiative Center for Intelligent Hybrids, Seoul National ...
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Orientation Control of Block Copolymer Thin Films Placed on Ordered Nanoparticle Monolayers Taehee Kim,†,‡ Sanghyuk Wooh,† Jeong Gon Son,*,†,‡ and Kookheon Char*,† †

School of Chemical and Biological Engineering, The National Creative Research Initiative Center for Intelligent Hybrids, Seoul National University, Seoul 151-744, Korea ‡ Photo-Electronic Hybrid Research Center, Korea Institute of Science and Technology, Seoul 136-791, Korea S Supporting Information *

ABSTRACT: We investigate orientation and lateral ordering of poly(styrene-block-methyl methacrylate) (PS-b-PMMA) diblock copolymer (diBCP) thin films placed on ordered nanoparticle (NP) monolayers. The densely packed NP monolayers were prepared on silicon substrates with the Langmuir−Blodgett (LB) deposition technique. The perpendicular domain orientation of BCP thin films is obtained on the ordered NP monolayers due to the nanoscale regular roughness which exerts the elastic deformation on the BCP nanodomains and suppresses the substrate-induced parallel orientation. The effect of BCP film thickness as well as the NP size on the orientation of BCP nanodomains was systematically investigated. We also demonstrate the defect-tolerant ordering of the perpendicular orientation of BCP thin films on the NP-vacant sites.

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(NP) monolayers with the tunable NP size and the BCP film thickness. The ordered NP monolayer provides a good model system of well-defined surface roughness. We synthesized monodisperse iron oxide NPs via the thermal decomposition method 22 and employed the Langmuir−Blodgett (LB) technique23 to prepare large area ordered arrays of densely packed NP monolayers on silicon wafer substrates. The size of NPs in ordered NP monolayers can control the period of the regular nanoscale surface roughness. The periodicity of the surface roughness of the NP monolayers and the BCP domains were characterized with the grazing incidence small-angle X-ray scattering (GISAXS) technique. We found that a correlation between the characteristic length scales of the BCP domains and the surface roughness from NP monolayers is a critical factor for the orientation of the BCP films. We also demonstrate the defect-tolerant ordering of domain orientation of BCP thin films on the NP-vacant site.

INTRODUCTION Block copolymers (BCPs) consist of two or more chemically different polymer chains causing self-assembly into periodic nanostructures with controllable feature size and morphology.1,2 An ability to control domain orientation and lateral order of BCP thin films has received great interest for potential applications such as nanolithography3−6 and optoelectronic devices.7,8 In order to control the orientation of BCP thin films, a variety of techniques including electric field,9 solvent vapor annealing,10 surface neutrality,11−14 and chemically patterned substrates15,16 have been developed. However, the vast majority of the previous strategies were mainly centered on the control of the interfacial energy. The driving force for microphase separation of BCP thin films in equilibrium is described through a phenomenological free energy model which is expressed with the elastic free energy of the stretched and compressed polymer chains as well as the interfacial energy at the interfaces.17,18 In case of BCP films confined at the rough interfaces, the effect of elastic deformation dominantly contributes to the domain orientation.19,20 Sivaniah et al. reported the effect of surface roughness on the orientation of BCP domains.21 Their result showed that the elastic deformation could induce the change on the orientation of BCP nanodomains. Unfortunately, the roughness of the substrate used in their study was irregular, and it was difficult to characterize the roughness quantitatively. In this paper, we demonstrate the domain orientation of poly(styrene-block-methyl methacrylate) (PS-b-PMMA) thin films can be controlled using underlying ordered nanoparticle © XXXX American Chemical Society

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EXPERIMENTAL SECTION

Synthesis of Iron Oxide NPs. Oleic acid-covered iron oxide nanoparticles (NPs) with controlled sizes and narrow size distributions used in this work were chemically synthesized via thermal decomposition of iron−oleate precursors in high boiling solvents (octyl ether for 6 nm NPs and 1-octadecene for 22 nm NPs). The iron−oleate complex was prepared by reacting iron chloride (FeCl3· Received: July 31, 2013 Revised: September 23, 2013

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Figure 1. TEM (a, b) and SEM (c, d) images of iron oxide NPs. TEM images show as-synthesized NPs with the average particle diameter of (a) 6 nm and (b) 22 nm with the standard deviation of 0.5 and 1.7 nm, respectively. Insets in TEM images are histograms showing the paricle size distributions. SEM images show the Langmuir−Blodgett (LB) monolayers with (c) 6 nm and (d) 22 nm iron oxide NPs transferred to silicon wafers at the surface pressure of 11 mN/m. Insets in SEM images are the magnified images. Source (Quebec, Canada). BCP thin films were produced by spincoating from toluene solution. The film thickness was varied from about 40 to 250 nm by adjusting the solution concentration (from 1 to 5 wt %) and spin speed (from 1700 to 6000 rpm). With the film thickness less than 30 nm, PS-b-PMMA (35.5K−12.2K) thin films dewetted on the NP monolayers. BCP thin films were dried under vacuum for 6 h and thermally annealed at 180 °C for 4 days. During the thermal anneal process, the stable binding of ligands onto NPs was indirectly confirmed from almost constant water contact angle (∼100°) of the NP monolayer even after thermal treatment. For the SEM observation, PMMA cylindrical domains of BCP thin films were removed with UV irradiation (Sankyo Denki, Japan, λ = 254 nm) for 90 min under vacuum, followed by immersion in acetic acid for 1 h at room temperature, and then rinsed with deionized water. Characterization. The film thickness of BCP thin films was measured with a variable-angle multiwavelength ellipsometer (Gaertner L2W16C830) before thermal annealing. The surface morphology of BCP thin films was observed with an AFM (Digital Instrument, Nanoscope IIIA) in tapping mode. A field emission scanning electron microscope (FE-SEM, JEOL, JSM-7401F) operated at 5 kV was used to observe the NP arrangement and the morphology of BCP thin films. GISAXS measurements were carried out at the 4C2 beamline of the Pohang Light Source to analyze both the internal structure of the BCP films and the arrangement of underlying NPs. The incidence angle of the X-ray beam was 0.19°, which is above and below the critical angle of the BCP films and silicon substrates, respectively. Scattering vectors were corrected with a poly(styrene-b-(ethylenerandom-butylene)-b-styrene) (SEBS) BCP standard. The sample-todetector distance was 2.25 m, and the wavelength of the synchrotron X-ray beam was 0.138 nm. A two-dimensional charge-coupled device

6H2O) and sodium oleate. Reactions were carried out in a nitrogen atmosphere. FeCl3·6H2O (Aldrich, 98%), sodium oleate (TCI, 95%), oleic acid (Aldrich, 90%), octyl ether (Aldrich, 99%), and 1-octadecene (Aldrich, 90%) were used without further purification. Synthesized NPs were precipitated with ethanol/acetone (1:1 volume ratio) mixture, centrifuged, and redissolved in hexane. This procedure was repeated three times in order to remove excess surfactants from the solution. Finally, we prepared the NP suspensions in toluene with a concentration of 0.35 wt %. Preparation of LB Films with NPs. The LB instrument used in this work was a KSV 3000 trough. 180 and 500 μL of 6 and 22 nm NP suspensions, respectively, were cast dropwise onto the air−water interface with a Hamilton microliter syringe at the intervals of 30 s to ensure the NPs spread thoroughly. Compression of the Langmuir films was carried out at a speed of the 5 mm/min by moving two barriers symmetrically after the toluene evaporated (40 min). The surface pressure was monitored with a Wilhelmy plate. LB films were prepared at different surface pressures and transferred onto the piranha-treated silicon substrates, which was immersed vertically in pure water before the NP suspensions were cast. The NP monolayers transferred on silicon substrates were heated at 200 °C for 2 h to anchor the NPs to the substrate robustly. After the heat treatment, the arrays of NPs were not disturbed by washing with organic solvents and a prolonged sonication. The detailed description of this test is given in the Supporting Information. Before BCP thin films were deposited, NPcoated substrates were washed with toluene to remove the unbounded surfactants and dried with nitrogen flow. Preparation of BCP Thin Films. Cylinder-forming PS-b-PMMA (Mn,PS = 35.5 kg/mol, Mn,PMMA = 12.2 kg/mol, PDI = 1.04) and Lamella-forming PS-b-PMMA (Mn,PS = 37 kg/mol, Mn,PMMA = 37 kg/ mol, PDI = 1.04) diblock copolymers were purchased from Polymer B

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(CCD) detector with 1024 pixels × 1024 pixels was used to collect the scattering data.

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RESULTS AND DISCUSSION Figures 1a,b show the transmission electron microscope (TEM) images of as-synthesized iron oxide NPs. The NPs were covered with oleic acid as a surfactant and their average diameters of 6 and 22 nm and the standard deviations of 0.5 and 1.7 nm, respectively. The histograms show the narrow particle size distributions of the corresponding NPs. Figures 1c and d show the scanning electron microscope (SEM) images of ordered NP monolayers prepared by the LB technique with 6 and 22 nm iron oxide NPs transferred onto the silicon substrates at the surface pressure of 11 mN/m. Although the NPs are not perfectly ordered, they exhibit a high degree of hexagonal short-range order and sufficient long-range surface coverage on the silicon substrates. The detailed SEM images of LB films transferred at the different surface pressures are shown in Figure S1. Figure 2 shows representative SEM images for the cylinderforming PS-b-PMMA (35.5K−12.2K) films with a film thickness of 240 nm placed on a bare silicon substrate and the NP monolayers. The BCP films were prepared by spincoating using 3 wt % solution in toluene and thermally annealed at 180 °C for 4 days under vacuum to achieve the ordered morphologies. The parallel orientation of BCP films with respect to the substrate was preferred on the bare silicon and the 6 nm NP monolayer, while the perpendicular orientation was observed on the 22 nm NP monolayer. The grazing incident small-angle X-ray scattering (GISAXS) measurement supports the result on the domain orientation of the BCP films as shown in Figure 3. The GISAXS experiment was carried out at incident angle of 0.19°, which is between the critical angles of the BCP films and the silicon substrate, so that X-ray beam fully penetrates the BCP films. The diffraction peaks appeared at qy = 0.745 nm−1 in Figure 3a and qy = 0.257 nm−1 in Figure 3b correspond to dNP = 8.43 nm and dNP = 24.4 nm, respectively, which are the repeat period of the NPs in the ordered monolayers. The diffraction spots in Figure 3c show the parallel orientation of the cylinder-forming BCP film deposited on a bare silicon substrate.24 The GISAXS pattern shown in Figure 3d can be understood as a superposition of scatterings from the parallel-oriented cylindrical BCP nanodomains and the underlying 6 nm NP monolayer. This result confirms that BCP nanodomains orient parallel to the substrate on the 6 nm NP monolayers while underlying NP monolayers maintain their ordered arrangement. The GISAXS pattern shown in Figure 3e can be considered as the superposition of diffraction peaks from the perpendicularly oriented cylindrical BCP nanodomains and the underlying 22 nm NP monolayer. In Figure S2, we show the in-plane profile of the GISAXS pattern of the BCP film on the 22 nm NP monolayer in comparison with that of the pristine NP monolayer along the qz = 0.151 nm−1. The typical rodlike diffraction peaks appearing in the in-plane profile, with a relative q y ratio of 1:√3:√4:√7:√9, indicate that cylindrical BCP domains oriented perpendicular to the substrate.25 The main driving force for the parallel orientation of BCP domains on the bare silicon substrates is the interfacial energy minimization between PMMA/SiOx and PS/SiOx. The perpendicular domain orientation of the BCP films placed on the NP monolayer is attributable to the surface roughness that can exert the elastic deformation on the BCP domains. The

Figure 2. Plan-view SEM images for cylinder-forming PS-b-PMMA (35.5K−12.2K) thin films with a film thickness of 240 nm showing the change of BCP domain orientations on the different substrates. The BCP thin films were deposited on (a) a bare silicon wafer, (b) a 6 nm NP monolayer, and (c) a 22 nm NP monolayer. PMMA cylindrical nanodomains were removed with UV irradiation followed by rinsing with acetic acid after thermal annealing at 180 °C for 4 days. Inset in (c) is the FFT of perpendicularly oriented cylindrical domains. All the scale bars shown on the figures are 100 nm.

theoretical consideration of the orientation of BCP domains on rough surfaces was reported by Tsori and Andelman.19 They calculated the relative energy difference between the parallel and perpendicular lamellar domains placed on a corrugated surface. We employed their calculation formulas to our experimental result. According to their calculation, the bulk free energy of the parallel orientation is given as C

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where S is the surface area and σ is the surface field that is taken to be a constant throughout the surface which is chemically homogeneous. ϕ0 is the amplitude of the order parameter, and K is the bending modulus. qBCP (= 2π/dBCP) and qsub (= 2π/ dNP) are the wavenumbers of repeat period of the BCP domains and the substrate roughness, respectively. dBCP and dNP are repeat period of the BCP domains and the NPs. R is the amplitude of the substrate roughness, and we assumed the R is equal to the radius of the NPs in our system. The bulk free energy of the perpendicular orientation is similarly derived: Fperp S

∼

σ 2ϕ0 2 1 K qBCP

(2)

They simply assumed perpendicular orientation is not affected by surface roughness owing to their geometry while parallel orientation is significantly affect by the period and amplitude of roughness. The difference between the two free energies on two different orientations is 2 ⎧ ⎫ ⎪ K 1 ⎪⎛ qBCP ⎞ 2 ⎜ ⎟ ⎨ (F − Fperp) = (qBCPR ) − 1⎬ ⎜ ⎟ 2 2 para ⎪ qBCP ⎪ Sσ ϕ0 ⎩⎝ qNP ⎠ ⎭

(3)

When Fpara − Fperp > 0, the perpendicular orientation is favored. The calculation result based on eq 3 is shown in Figure 4.

Figure 4. Calculation result based on eq 3 showing preferred orientations of BCP microdomains on the ordered NP monolayers. The black symbols indicate our experimental results: the circles and squares represent the cylindrical and lamellar BCPs, respectively.

We can consider qBCP/qsub (= dNP/dBCP) as a relative periodicity of the substrate roughness to the domain spacing of the BCP film. The increase on qBCP/qsub implies that the BCP domains undergo deformation with the long-range undulation induced by the large periodicity of the substrate roughness while qsubR is an independent value with NP size. In eq 3, the large qBCP/qsub increases the energy penalty of the parallel ordering, and consequently the perpendicular orientation of BCP nanodomains is preferred. In the present case, the largesized NP monolayer imparts the long-range elastic deformation to the parallel-oriented BCP nanodomains and induces the perpendicular domain orientation of the BCP film on the NP

Figure 3. Two-dimensional GISAXS patterns of (a) the 6 nm NP monolayer and (b) the 22 nm NP monolayer; the cylinder-forming PS-b-PMMA thin films placed on (c) the bare silicon wafer, (d) the 6 nm NP monolayer, and (e) the 22 nm NP monolayer.

Fpara S

2 σ 2ϕ0 2 1 ⎛ qBCP ⎞ ⎜ ⎟ (q )2 ∼ K qBCP ⎜⎝ qsub ⎟⎠ BCP

(1) D

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Figure 5. AFM height (a) and phase (b, c) images for the cylinder-forming PS-b-PMMA (35.5K−12.2K) thin films deposited on (a) the bare silicon wafer, (b) the 6 nm NP monolayer, and (c) the 22 nm NP monolayer at different film thicknesses after thermal annealing at 180 °C for 3 days.

Figure 6. AFM height (a) and phase (b, c) images for the lamella-forming PS-b-PMMA (37K−37K) thin films deposited on (a) the bare silicon wafer, (b) the 6 nm NP monolayer, and (c) the 22 nm NP monolayer at different film thicknesses after thermal annealing at 180 °C for 3 days.

E

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Figure 7. SEM images and a Voronoi diagram for the cylinder-forming PS-b-PMMA (35.5K−12.2K) thin film deposited on a 22 nm NP monolayer: (a) a cross-sectional SEM image showing the perpendicularly oriented cylindrical domains; (b) a plan-view SEM image at the fractured edge; (c) a large-area SEM image showing the long-range order of the perpendicular domain orientation; (d) a Voronoi diagram of (c). Inset in (c) is the FFT of the corresponding SEM image. In the Voronoi diagram, PMMA cylindrical domains having five and seven nearest neighbors are marked in blue and red, respectively.

more hydrophobic PS block due to the hydrophobic alkyl chains at the NP surface. This behavior is the symmetric wetting with the commensurable film thickness of ndBCP. Particularly, the BCP thin films placed on the NP monolayers could relieve the strain at an incommensurable film thickness with mixed orientations (parallel and perpendicular) rather than forming hole or island structures because of near equal energetic conditions between the two different orientations. We observed mainly the parallel orientation of BCP nanodomains at 2dBCP and 3dBCP, while the perpendicular orientation of BCP nanodomains was observed in the vicinity of 2.5dBCP film thickness on the NP monolayers. This transition of the orientation occurs because the parallel orientation of BCP nanodomains is energetically stable at ndBCP thickness on the NP monolayers, while unstable at (n + 1/2)dBCP thickness due to the symmetric wetting. On the relatively large sized NP monolayer, BCP nanodomains tend to orient perpendicular to the substrate due to the contribution of the large elastic energy propagated from the rough surfaces of the NP monolayers. This tendency depending on the NP size is still consistent with the above results from the thick BCP films and the theoretical calculation. We investigated the detailed structure for the perpendicularoriented PS-b-PMMA (35.5K−12.2K) thin films deposited on the 22 nm NP monolayer. A cross-sectional SEM image taken with a tiling angle of 5°, and a plan-view image near the fractured boundary shown in Figure 7a,b confirm the perpendicular orientation of the BCP thin film on the ordered NP monolayer. Long-range lateral ordering is also confirmed with the large-area SEM image in Figure 7c. The boundaries of the perpendicular-oriented cylindrical BCP domains in Figure 7c were analyzed with the Voronoi diagram as shown in Figure 7d. In the Voronoi diagram, four different grains of perpendicularly oriented BCP nanodomains are visible; that is consistent with the fast Fourier transform (FFT) of Figure 7c. The Y-shaped dark trace appeared in Figure 7c indicates the grain boundary of the underlying NP monolayer. We can see the boundaries of the underlying NPs and BCP nanodomains

monolayer. Since this calculation method is more valid in case that the film thickness is much larger than the domain spacing, the calculation and experimental observation showed good consistency. It is generally known that the ordering phenomena of BCP nanodomains in thin films are strongly influenced by the commensurability between the film thickness and the domain spacing of BCP nanodomains (dBCP). Our group has intensively studied about PS-b-PMMA orientation changes with different substate/surface properties according to the film thickness.14,26 In order to investigate the effect of film thickness on the domain orientation of BCP thin films placed on the NP monolayers, surface morphologies of BCP thin films were examined by using atomic force microscope (AFM) with the film thickness range from 2dBCP to 3dBCP. The film thickness was finely tuned with spin-coating by adjusting the concentration of the BCP solutions as well as the spin speed. Figures 5 and 6 demonstrate the orientation behaviors of cylinderforming (35.5K−12.2K, dBCP = 25 nm) and lamella-forming (37K−37K, d BCP = 40 nm) PS-b-PMMA thin films, respectively, depending on the film thickness as well as the NP size. A domain spacing of cylindrical nanodomains is defined as the interplane distance of hexagonally packed cylinders.27,28 When a PS-b-PMMA thin film is deposited on a bare silicon substrate, PS block is preferentially segregated at the air surface while PMMA block preferentially wets the surface of the silicon substrate. This wetting behavior is referred to as the asymmetric wetting. In the case of asymmetric wetting, the commensurable film thickness of the parallelorientd BCP nanodomains is (n + 1/2)dBCP, where n is an integer. This means the parallel orientation of BCP nanodomains is energetically favored at the (n + 1/2)dBCP thickness. At an incommensurable film thickness, the hole and island structures are formed at the free surface in order to relief the strain caused by stretched or compressed polymer chains in the parallel-oriented BCP domains. In contrast, when a PS-bPMMA thin film is deposited on the NP monolayer, both air and the NP surface are preferred by less surface energy and F

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Figure 8. SEM images along with schematics showing the orientation change of cylindrical nanodomains of a PS-b-PMMA (35.5K−12.2K) thin film with a film thickness of 68 nm deposited on a 22 nm NP monolayer with NP-vacant sites. The bright and dark regions in the SEM images indicate the BCP thin films with and without the underlying NPs, respectively. The dotted lines in the figures indicate the width of NP-vacant sites. The perpendicular orientation of the cylindrical nanodomains persists on NP-vacant sites when the width (W) of the NP-vacant sites is less than 290 nm (∼11.6dBCP).

result directly shows a clear evidence of the defect-tolerant ordering of BCP self-assembly and the limitation.

are independent. This result shows that the BCP self-assembly does not exactly reproduce the underlying NP arrangement. Additionally, according to our SAXS and GISAXS results, the cylinder-forming BCP in the bulk and the perpendicularly oriented BCP thin film showed the same diffraction peak position (q = 0.254 nm−1), implying that the NP size does not directly affect the periodicity of the BCP cylinders. It is expected that the lateral ordering of BCP thin films depends on the surface coverage of the underlying NP monolayer. If a NP monolayer contained the large vacant site, it might be impossible to achieve the long-range lateral ordering of oriented BCP thin films on the NP monolayer. In order to address this problem, we analyzed both the BCP orientation and the NP arrangement all together with SEM. When the thin BCP films (thinner than 150 nm) placed on the NP monolayers were observed under plan-view SEM, it is possible to detect both the orientation of BCP nanodomains and the indication of the underlying NP monolayer as shown in Figure 8. The bright and dark regions in Figure 8 indicate the BCP thin films with and without the underlying NPs, respectively. The NP monolayer was prepared at the surface pressure of 10 mN/m, which is less than the optimal surface pressure (11 mN/m), so that the NP monolayer contains the more defect sites. Figure 8 shows the perpendicular orientation of cylindrical nanodomains in PS-b-PMMA (35.5K−12.2K) thin films with a film thickness of about 68 nm is persisted on the NP-vacant site, when the width of NP-vacant sites is less than 290 nm (∼11.6dBCP). On the other hand, the parallel orientation of BCP nanodomains was observed on the large NP-vacant sites which are wider than 290 nm. This interesting finding implies that self-assembly of BCP thin films show the defect-tolerant ordering on the small-sized defect of underlying NP monolayers. A few studies for the long-range lateral order of BCP self-assembly and the defect-tolerant ordering of BCP thin films have been reported so far,16,29 but the limitation of the ordering has not been fully elucidated. Our experimental

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CONCLUSION

In conclusion, we have demonstrated that ordered NP monolayer can be served to control the orientation of microphase-separated domains of BCP thin films. The sufficient roughness of the NP monolayer assembled with relatively largesized NPs exerted the elastic deformation on the paralleloriented nanodomains, so that the perpendicular domain orientation was obtained in BCP thin films placed on the ordered NP monolayer. The orientation of BCP thin films placed on the NP monolayers was strongly influenced by the film thickness as well as the NP size. The perpendicular orientation of the cylinder-forming and lamella-forming BCP thin films placed on the NP monolayers was favored at the film thickness in the vicinity of (n + 1/2)dBCP because PS-b-PMMA thin films show the symmetric wetting on the NP monolayers. We observed the perpendicular orientation of BCP nanodomains was persisted on the NP-vacant sites up to the width of NP-vacant sites less than 290 nm (∼11.6dBCP) in our system. This defect-tolerant ordering implies the long-range lateral order of the perpendicular orientation of BCP thin films can be achieved. Our results open up exciting opportunities for controlling the orientation and the lateral order of BCP nanodomains using the topographical pattern comprised of NPs. We expect that the topological pattern-induced orientation transition of (block) copolymer thin films could be utilized for functional device application. For instance, by controlling the orientation of conjugated polymers on semiconducting NPs, the charge mobility of the conjugated polymers could be enhanced in organic field-effect transistor (OFET) or organic photovoltaic (OPV) devices. G

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ASSOCIATED CONTENT

S Supporting Information *

Detailed SEM images of NP arrays at the different surface pressures in the LB experiment and in-plane profiles of GISAXS patterns. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (K.C.). *E-mail [email protected] (J.G.S.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research was financially supported by the WCU (World Class University) Program through the Korea Science and Engineering Foundation funded by the Ministry of Education, Science and Technology (R31-2008-000-20012-0) and the National Research Foundation of Korea Grant funded by the Korean Government (MEST) (The National Creative Research Initiative Program for “Intelligent Hybrids Research Center” (2010-0018290)), the Global Frontier Research Program (2011-0032156), and Korea Institute of Science and Technology (KIST) Internal Project (2E23821). We thank D. Andelman for helpful discussions and Prof. T. Hyeon and Dr. T. Yu for the LB experiments. The experimental support by the staff at the 4C1 and 4C2 Beamline of Pohang Light Source is also gratefully acknowledged.

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REFERENCES

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