Permeability- and Surface-Energy-Tunable Polyurethane Acrylate

Research Article www.acsami.org

Permeability- and Surface-Energy-Tunable Polyurethane Acrylate Molds for Capillary Force Lithography Dongchul Suh,† Hyowon Tak,‡,§ Se-jin Choi,∥ and Tae-il Kim*,‡,§ †

Department of Chemical Engineering, Hoseo University, Asan-si, Chungcheongnam-do 31499, Republic of Korea Center for Neuroscience Imaging Research (CNIR), Institute for Basic Science (IBS), Suwon 440-746, Republic of Korea § School of Chemical Engineering, Sungkyunkwan University (SKKU), Suwon 440-746, Korea ∥ MCNet Company, Limited, Dangjeong-dong, Gunpo-si, Gyeonggi-do 435-030, Korea ‡

S Supporting Information *

ABSTRACT: A permeability- and surface-energy-controllable polyurethane acrylate (PUA) mold, a “capillary-force material (CFM)” mold, is introduced for capillary-force lithography (CFL). In CFL, the surface energy and gas permeability of the mold are crucial. However, the modulation of these two main factors at a time is difficult. Here, we introduce new CFM molds in which the surface energy and permeability can be modified by controlling the degree of cross-linking of the CFM. As the degree of cross-linking of the CFM mold increases, the surface energy and air permeability decrease. The high average functionality of the mold material makes it possible to produce patterns relatively finely and rapidly due to the high rate of capillary rise and stiffness, and the low functionality allows for patterns to form on a curved surface with conformal contact. CFMs with different functionality and controllable-interfacial properties will extend the capabilities of capillary force lithography to overcome the geometric limitations of patterning on a scale below 100 nm and micro- and nanopatterning on the curved region. KEYWORDS: capillary force lithography, degree of cross-linking, permeability, surface energy, micro/nanostructure

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INTRODUCTION

influenced by air permeation when the hydrodynamic pressure builds up as the capillary rise progresses.14 To date, molten polymer in CFL with an impermeable mold has not been able to completely fill a mold cavity.15 Trapped air between the polymer film and the impermeable mold hinders the capillary rise of the polymer, and the polymer cannot fill the void space. With a permeable mold such as polydimethylsiloxane (PDMS), trapped air can escape through the mold so that melted polymer can reach the ceiling of the mold. However, PDMS cannot be used in patterning on a scale below 100 nm due to its low elastic modulus.16 Although cross-linked silicone elastomer D4H−D4V was additionally reported to have several advantages such as UV transparency, low surface energy, and low viscosity precursors, the preparation of the mold took a long time (at least 5 days).17 It also required a special catalyst and careful handling due to the extremely exothermic reaction. Alternatively, a less-permeable mold made of amorphous Teflon (AF 2400, DuPont) was demonstrated.18,19 However, Teflon with a fluorinated functional group has extremely low surface energy (∼15.6 mJ/m2), such that such a low surface

Many researchers have sought to improve unconventional lithography methods to overcome the disadvantages of conventional methods, i.e., high cost, low throughput, and an environmentally unfriendly process.1−6 Among them, capillaryforce lithography (CFL) is a unique technique for one-step patterning by the capillary rise of a polymer thin film heated above Tg (the glass transition temperature).2 Various microand nanostructures on large substrates can be successfully produced using CFL.2,7 These well-defined micro- and nanostructures enable diverse applications in fields such as biology and energy harvesting.8−12 To obtain fine structures with CFL, it is necessary to control several key factors, including the permeability of the mold, capillarity, and the substrate, polymer, and mold interactions.13 Capillarity, which is the driving force for CFL, occurs when a polymer melt wets a cavity wall and moves by reduction of the free energy. With capillaries, a pressure gradient between the gas and liquid phases, which is known as the Laplace pressure, is the driving force for the capillary rise. To understand polymer rising behavior in the cavity of the molds, one should consider the interaction between the permeation of the mold and the Laplace pressure because the capillary rise is generally © XXXX American Chemical Society

Received: April 13, 2015 Accepted: September 28, 2015

A

DOI: 10.1021/acsami.5b06975 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

the CFM molds, we pressed a transparent PDMS slab or glass substrate (used as a temporary supporting backplane) lightly on a liquid drop of the CFM precursor on the master. After the polymerization of CFM by UV flood exposure through the transparent backplane, the CFM replica mold was manually removed from the master and the supporting backplane. CFM molds replicated at least 50 times could be obtained from a single master mold without any contamination or damage. Figure S1 displays SEM images of master mold (2 μm line, 1 μm space, and 700 nm height) and the 1st, 10th, and 50th replicated molds. The detailed scheme for fabricating CFM molds is shown in Figure S2. For the CFL experiment (Figure 1), commercially available polystyrene (PS; Sigma-Aldrich, molecular weight 3.5 × 104 and 1.92 × 105) and a silicon wafer as the substrate were used. For PS films, a precleaned Si substrate was spin-coated with PS solution (3−10 wt %) diluted in toluene. To exclude the solvent effect and enhance adhesion, we annealed each sample at 100 °C in an oven for 30 min. The CFM mold was placed on the smooth surface of the PS film at room temperature and pressed lightly (negligible pressure, ∼0.01 N) with a thin PDMS backplane for conformal contact. After contact, PDMS backplane was removed and PS film with CFM mold was heated above the glass transition temperature of the PS polymer. After being annealed above Tg and cooling to ambient temperature, the mold was mechanically removed by sharp tweezers. The patterned polymer structure by CFL was examined by atomic-force microscopy (AFM) and scanning electron microscopy (SEM). Figure 2 illustrates the chemical formulas of the prepolymers used for tunability of the mechanical properties and surface energy of the CFM molds. In the CFM molds, modified PUA belongs to a segmented polymer tipped with acrylic functionality. The average functionality of the polymer makes it possible to adjust the degree of cross-linking, and it improves the mechanical properties of the CFM. As the average functionality of the CFM increases, the cross-linking density increases. In the case of CFM-110S, low acrylate functionality produces rubber elasticity, leading to low mechanical modulus. The cycloaliphatic backbone notably produces a high degree of mechanical stiffness, and it can tune the mechanical moduli of various CFM molds. More interestingly, the surface energy is substantially dependent on the polarity of the polymer, which is determined by the type and number of functional groups. Nonpolar moieties, such as vinyl groups of aliphatic acrylated prepolymers, reduce the surface tension of the polymer. As a result, CFM-130H, which contains the highest number of vinyl groups per molecule, has the lowest surface tension among the CFM materials. The mechanical and chemical properties of the CFM molds are displayed in Table 1. Materials for CFM Molds. First, 30 wt % 1,6-hexanediol diacrylate (M200, MIWON Co., Korea), which is a low-viscosity acrylate monomer used as a diluent with respect to the prepolymer, was prepared. Subsequently, Irgacure 184 (1-hydroxy-cyclohexyl-phenylketone) and Darocur 1173 (2-hydro-2-methyl-1-phenyl-1-propane) from Ciba Specialty Chemicals (Switzerland), were added as photoinitiators at 1.5 wt % with respect to the total amount of the blend of the prepolymer and reactive monomer diluent. In addition, the releasing agent, Rad 2200N from TEGO Chemie Service, was blended at a loading level of 1.0 wt % with the prepared mixture for easy detachment. The mixture was thoroughly stirred until the powder-type photoinitiator (Irgacure 184) was completely dissolved. Fourier-Transform Infrared Spectroscopy. Infrared measurements were obtained with a Bruker TENSOR27 (German, detector loaded: MIR_ATR(ZnSe), scanned 32 times, resolution of 4 cm−1). Scanning Electron Microscopy. Samples were sputter-coated with a thin (∼2 nm thick) platinum layer and imaged on a Merlin Compact high-resolution, field-emission scanning electron microscope (Carl Zeiss, Germany). Atomic Force Microscopy. Measurements were performed in tapping mode in air on a XE-100 Instrument (PSIA Inc., Korea) using silicon cantilevers (SSS-NCHR, Park Systems, Korea) with a spring constant of 4.2 × 104 mJ/m2 and a resonance frequency of 330 kHz (manufacturer’s specifications).

energy cannot allow easy surface modification with a selfassembly monolayer (SAM) for controlling the interfacial property of the CFL.19,20 To overcome the above shortcomings, we introduced a UVcurable polyurethane acrylate, which has been widely used for sub-100 nm lithography.16,21 Suh et al. first presented the formation of several intriguing nanostructures using CFL with a PUA mold.22 In another study, they also demonstrated high aspect ratio polymer nanohairs formed by two sequential steps, molding and drawing.23 Micro- and nanoscale hierarchical structures fabricated by a two-step UV-assisted capillarymolding technique have also been presented.24,25 Nevertheless, all approaches have focused on how to effectively form nanostructures without the modulation of surface energy and permeability, even when these two properties are key issues in CFL. In this work, we introduce permeability and surface-energytunable polyurethane acrylate capillary-force material (PUA CFM) molds by changing the chemical composition and degrees of cross-linking. We endeavor to study correlations between two crucial factors: permeability and surface energy. Moreover, we show good agreement between the experimental results and a simple kinetic model for the capillary rise of polymer melts in which the capillary rise in permeable molds is proportional to the square root of the annealing time. Additionally, the unique properties of the mold materials make it possible to modulate the capillarity and mechanical moduli for nanoscale patterning.

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EXPERIMENTAL SECTION

The procedure for the CFL with CFM is illustrated in Figure 1. Thin (20 μm thick) CFM molds was fabricated by replica molding from silicon master molds prepared by e-beam lithography.27 To fabricate

Figure 1. (a) Schematic of deterministic nanostructure assembly by CFL on a substrate coated with a polymer layer. (b) Photograph of the master for replica molding. (c) CFM mold replicated from the master (shown in panel b). (d) Polymer pattern formed by CFL. B


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Gas Permeability Test. Air permeability was tested through a VVolumetric ASTM D1434 procedure using a BT-1 by TOYOSEIKI, Inc. The test was performed by measuring the pressure difference between the high and low cells. Air was introduced into the high cell in the vacuum state and moved to the low cell after penetrating a 600 μm thick sample. The gas pressure sensor detects a change of the pressure at the low cell and then displays gas permeability. The test temperature was controlled to 25 (±1) °C in air. ASTM D3985 (O2) and ASTM F1249 (water) were V-Volumetric procedures used to measure oxygen and water permeability. OX-TRAN Model 2/21 (O2) and PermatranW 3/33 MA (water) instruments of Mocon, Inc. were used. The sample was placed in a test cell composed of two chambers. The inner chamber was filled with carrier gas (N2), and the outer chamber was filled with oxygen or water vapor. Molecules of the test gas diffused through the sample and moved to the sensor via the carrier gas. The sensor detects the concentration of the test gas in the carrier gas and calculates the permeability. The test temperature was controlled to 23 (±2) °C for oxygen and 38 (±2) °C for water.

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RESULTS AND DISCUSSION The molecular structure of all synthesized copolymers shown in Figure 3 was characterized via Fourier transform infrared (FT-

Figure 2. Chemical structures of various CFM mold materials consisting of (a) CFM-110S, (b) CFM-120M, and (c) CFM-130H.

Figure 3. FT-IR spectra of various CFM mold materials.

IR) spectroscopy. In the functional group region of these spectra, there are characteristic peaks at 1114 cm−1 (C−O−C stretching vibration), 1249 cm−1 (C−O stretching vibration), 1453 cm−1 (bending vibration of aliphatic CH), 1734 cm−1 (CO stretching vibration), 2873 and 2958 cm−1 (stretching vibration of aliphatic CH), and 3349 cm−1 (NH stretching vibration of urethane).28 The functional group region (1500 cm−1 to 1700 cm−1) related to the double bond (CC) characteristic stretching vibrations was relatively clear in CFM130H, which indicates more utilization of unsaturated sites during copolymerization, resulting in higher cross-linking density.28 To show the replication capability of the mold for features under 100 nm, we used a 60 nm line-and-space pattern. The SEM images in Figure 4a,b for the UV-cured CFM mold with CFM-120-M material and the resultant PS patterns formed by the CFL process, respectively, demonstrate that a dense pattern under 100 nm was successfully transferred to the polymer film. However, CFM-110S and -130H were not suitable for nanoscale pattern replication. The CFM-110S mold with a low mechanical modulus, 52.19 MPa, exhibited pattern pairing between the two adjacent structures. The CFM-130H mold with a relatively high degree of cross-linking additionally produced high shrinkage during polymerization, and it did not

Table 1. Material Properties of Various CFM Molds surface energy [mJ/m2]

material CFM110S CFM120M CFM130H

elastic modulus [MPa]

crosslinking density

γ

d

γ

2

low

5.2

31.0

36.2

52.19

3

middle

8.8

21.3

30.1

606.09

6

high

15.3

12.7

28.0

1042.76

average functionality

p

γ

total

Contact-Angle Determination. The contact angle was obtained with EasyDrop from KRÜ SS GmbH (Germany). Water and ethylene glycol drops were placed on the sample, and images were taken a camera (monochrome CMOS sensor 752 × 480 px). Contact-angle evaluations were performed with SW23 (DSA2) using drop-shape analysis. The accuracy of the contact-angle measurement was ±0.1°. Elastic Modulus Measurement. Moduli of were measured by Instron 3342 (Instron Inc.) under tension. Samples with each CFM material were prepared with 250 μm, 0.5 cm, and 3.5 cm for thickness, width, and length, respectively. The samples hung on the instrument were stretched by 0.1 mm/min velocity. Elastic moduli are defined as the slope of the samples stress−strain curve in the elastic deformation region. C


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standing film without a supporting layer such as PET, and all CFL experiments and analysis can be successfully taken without consideration of the effect of PET. For the interfacial properties of the CFM molds, the surface energy was calculated by contact angle measurements. The measurement of contact angles of pure liquids with known surface tension on a given solid surface is a common way of obtaining the surface energy of the solid material. Figure S5a shows contact angle measurements performed with deionized water and ethylene glycol on polymer surfaces. To ensure reproducibility, we performed at least five experiments for each liquid on a fresh polymer surface. The contact angle of CFM110S was 73.9° for deionized water, which was slightly higher than the 70.1° for ethylene glycol. The contact angle on CFM130H increased to 86.2° for deionized water and decreased to 61.5° for ethylene glycol, corresponding to a 16.6% increase and a 12.3% reduction, respectively, compared to the results for CFM-110S. The contact angle of deionized water on CFM130H was greater than that of ethylene glycol and is a consequence of the stronger dispersive interactions of the conjugated CFM surfaces with the ethylene moiety of ethylene glycol. CFM-110S and CFM-120M, which are relatively dominated by the polar component, are hydrophilic, and CFM-110S is more hydrophilic due to the lower value of γsd/ γsp, where γsd and γsp are the dispersion and the polar components of the solid surface tension, respectively. However, CFM-130H that is dominated by the dispersive component is hydrophobic because γsd is larger than γsp (Figure S5b). The hydrophobic polystyrene melt is therefore likely to flow into the channel of CFM-130H without additional pressure. Additionally, CFM-130H has the lowest overall surface tension, enabling it to be removed easily and cleanly from the polymer pattern after CFL. The surface tension γ and Laplace pressure ΔPL of the CFM film were calculated according to the Young’s and geometricmean equations suitable for a low-energy system such as an organic liquid, water, and a polymer. These equations are expressed as

Figure 4. SEM images of (a) the CFM-120-M mold (60 nm line and space and height of 120 nm) and (b) the replicated PS pattern formed by the CFL process. The scale bar indicates 100 nm. (c) Perspective and (d) magnified images of the PS dot patterns replicated by the thinfilm CFM-110S mold on the curved surface of a 16 mm tube that has a radius of 3 mm.

exhibit high-fidelity pattern replication for patterns under 100 nm. The master mold and its replicated CFM-110S and -130H molds are shown in Figure S3. However, CFM-110S mold with low modulus and sufficient flexibility is useful to generate microscale patterns on nonflat substrate. Figure 4c,d show SEM images of PS dot patterns on the curved surface (3 mm radius of curvature), which was fabricated by CFM-110S. As previously mentioned, air permeability and interfacial properties are major issues of CFL. First, the air permeability coefficients of the CFM molds were measured according to the variable-volume method, as shown in Figure 5. The gas

ΔPL =

Figure 5. Air permeability of prepared CFM films.

γl cos θ (1)

L

γsl = γs − γs cos θ

(2)

γsl = γs + γl − 2(γsd·γld)1/2 − 2(γsp·γlp)1/2

(3)

where θ is the contact angle, L is the half-channel width, γsl is the interfacial tension between the solid and liquid, γs is the surface tension of the solid, γl is the surface tension of the liquid, and superscripts d and p are the dispersion and polar components of the surface tension, respectively. γ represents the overall surface tension equivalent to the sum of the dispersive and polar components where the dispersive component is contributed to by van der Waals interaction, and the hydrogen bonds and dipole−dipole interaction contribute to the polar component. Figure S5b shows the resultant surface energy of each CFM material as summarized in Table 1. The surface energy of the mold can be tuned to the range of 28−36 mJ/m2 by modulating the content nonpolar groups. Because the polymer molecular forces do not change appreciably within the temperature range of 100−200 °C, the polarity (γp) value is constant and independent of the temperature. Based on the polarity value (∼0.17) measured

permeability of the polymer films is related to the gas diffusivity. A decrease in the average functionality of the CFM films results in a lower cross-linking density and high diffusivity, while the air permeability increases. CFM-110S has a relatively higher air permeability (1.54 × 10−10 cm3·(cm/cm2)· s·cmHg), which is attributed to its large free volume from the low cross-linking density. The more-cross-linked structure of the CFM-120-M and -130H can shorten the gap between the macromolecular chains that have lower permeability, 1.16× 10−10 and 1.19 × 10−10 cm3·(cm/cm2)·s·cmHg, respectively. However, such a consistent trend is clearly shown in the oxygen and H2O permeability results shown in Figure S4. To date, a PUA mold with PET film as a supporting layer produced poor air permeability due to the PET layer.13−16,22 Generally, the permeability of the PUA mold is mostly governed by the PET film.14 In this work, all CFM materials were fabricated as a freeD


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ACS Applied Materials & Interfaces at 140 °C,29 the components of the surface tensions (γd and γp) of the PS polymer were calculated using the γtotal measured at 100 °C. As a result, PS had a γd of 33.7 mJ/m2, a γp of 6.9 mJ/ m2, and a γtotal of 40.7 mJ/m2. Based on the PS surface tension and contact angle between melted PS (liquid) and CFM molds (solid), the Laplace pressure was estimated from Young’s equation, excluding the effect of a slightly slanted wall and nonuniformly curved meniscus. As the surface tension of the CFM mold decreases, the interfacial tension with PS also decreases, resulting in increasing Laplace pressure, as shown in Table 2. The actual Laplace pressure would be somewhat

Figure 6. Capillary rise of PS polymer (molecular weight 1.92 × 105) as a function of annealing time comparing the experimental data and theory with the effective viscosity value for three different CFM materials.

Table 2. Calculated Interfacial Tension, Contact Angle, and Laplace Pressure on Three CFM Moldsa material CFM110S CFM120M CFM130H a

interfacial tension [mJ/m2]

contact angle with melted PS [°]

Laplace pressure [kPa]

21.1

68.2

15.115

12.1

63.7

18.019

4.3

54.4

27.635

polymer might have an effect on nonuniform and fluctuated pattern heights given by stretch and drawing of the meniscus,23 as shown in Figure S7. Such an interfacial effect is clearly confirmed by using different scale molds and different molecular weight (3.5 × 104) of PS. Figure 7 shows top-view SEM images for a 2 μm line and

The half-channel width, L is 1 μm.

smaller because a small gap can exist between the mold and the polymer that causes air permeation even when pressure is applied to create conformal contact.22 This trend is more significant in the case of CFM-130H due to the high stiffness caused by the high cross-linking density. To evaluate the capillary kinetics of polymer melt in the microcavities, the effect of the permeability and the thickness on polymer films confined to the substrate need to be considered. The permeabilities of the CFM molds in this work are greater than that of the less permeable PET film by 3 orders of magnitude, such that the CFM films can be regarded as permeable media.14 According to a previous work,13 the capillary kinetics for a permeable mold and sufficiently high polymer-film thickness is given by ⎛ 2Lγ cosθ ⎞1/2 t⎟ z=⎜ ⎝ 3η ⎠

(4)

where z is the capillary rise in time t, L is the half-channel width, γ is interfacial tension, θ is the contact angle, and η is the viscosity of the polymer melt. This equation reveals that capillary rise is proportional to the square root of time when the Laplace pressure and viscosity are determined from the mold geometry and material properties. Figure 6 illustrates the experimental data compared to the results from the theoretical model with annealing temperatures of 120 °C and a channel width of 2 μm lines. On the basis of eq 4 by mold size (2 μm lines), the calculated Laplace pressure (shown in Table 2), and time, we could surmise the best fittings by the effective viscosity of the PS polymer at a given temperature. A value of 7 ×109 Pa· s obtained by the interpolation method30 was used as the effective viscosity of PS polymer (molecular weight 1.92 × 105) at 120 °C (Figure S6). The average height of the meniscus corresponds to the capillary rise when the meniscus has uniform curvature in the microcavity. The results of CFM120M and -130H show good agreement with the theory, although a minor difference and broad data scattering in CFM110S data were observed. High adhesion caused by high interfacial tension between the CFM-110S mold and PS

Figure 7. SEM images of PS (molecular weight 3.5 × 104) patterns formed by CFL with (a) CFM-110S, (c) CFM-120M, and (e) CFM130H. AFM profiles of CFM molds and PS patterns formed by (b) CFM-110S, (d) CFM-120M, and (f) CFM-130H molds.

1 μm (700 nm height) pattern formed by three kinds of CFM molds at an annealing temperature of 100 °C for 3 h. When equilibrium was reached after a certain annealing time, the interface adopted a minimum area representing a cylindrical cap, presumably due to the reflow of the low-molecular-weight polymer melt with low viscosity. As shown in the figure, the height or aspect ratio is a minimum for CFM-110S and a maximum for CFM-130H, where the PS film reached the ceiling of the mold and filled the entire cavity. The maximum height was ∼700 nm, which is in good agreement with the E



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height of the original master mold. On the basis of the AFM profiles from Figures 6 and 7, we determined that the CFL patterns by CFM-110s with the highest permeability coefficients have the lowest height. It turns out that Laplace pressure has much more dominant on the capillary rise than permeability of the mold. It is convinced that interfacial property between PS and the mold is more crucial factor in CFL because Laplace pressure is related to polymer−mold interaction.

CONCLUSIONS In summary, we have presented a CFM mold that is permeability- and surface-energy-controllable by modulation of the degree of cross-linking in mold materials. During capillary rise, the polymer properties and surface energy of the mold play a significant role in pattern replication in CFL. Depending on the average functionality of the polymer, the change in the surface tension and cross-linking density was observed for three kinds of PUA molds used in the experiment. Furthermore, the permeability could be adjusted by the degree of cross-linking density. For a given mold geometry, the rate of the capillary rise in the relatively highly permeable CFM mold is proportional to the square root of the annealing time regardless of the permeability and is relatively low in the CFM110S mold but high in CFM-130H. Therefore, the CFM-110S mold has patterning capability on a nonflat surface because of the low modulus for conformal contact and the CFM-120-M with a relatively high stiffness can be used for a fine (under 100 nm) and fast patterning process at a fast rate of capillary rise. A relatively high modulus material, CFM-130H, may be useful for nanoimprint lithography. Therefore, these CFM molds with easily tunable chemical structures could provide valuable tools for controlling and obtaining well-defined nanostructures. ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.5b06975. SEM images of micro- and nanoscale master molds and their replicated CFM molds; fabrication scheme for CFM mold; oxygen and water permeability of CFM molds; contact angle and surface energy of the molds; viscosity value of PM; capillary rise of PS as a function of annealing time. (PDF)

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REFERENCES

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]; Tel: +82 31 290 7312. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is dedicated to the late Kahp-Yang Suh, a pioneer of capillary-force lithography. This work was supported by the Institute of Basic Science (IBS-R015-D1), NRF-2013- 430 R1A1A1061403, the Basic Science Research Program (20090083540), and the Pioneer Research Center Program (NRF2014M3C1A3001208) through the National Research Foundation of Korea funded by the Ministry of Science, ICT, & Future Planning. F


Research Article

ACS Applied Materials & Interfaces (22) Suh, K. Y.; Choi, S. J.; Baek, S. J.; Kim, T. W.; Langer, R. Observation of High-Aspect-Ratio Nanostructures Using Capillary Lithography. Adv. Mater. 2005, 17 (5), 560−564. (23) Jeong, H. E.; Lee, S. H.; Kim, P.; Suh, K. Y. Stretched Polymer Nanohairs by Nanodrawing. Nano Lett. 2006, 6 (7), 1508−1513. (24) Jeong, H. E.; Kwak, R.; Kim, J. K.; Suh, K. Y. Generation and Self-Replication of Monolithic, Dual-Scale Polymer Structures by TwoStep Capillary-Force Lithography. Small 2008, 4 (11), 1913−1918. (25) Jeong, H. E.; Kwak, R.; Khademhosseini, A.; Suh, K. Y. UVassisted Capillary Force Lithography for Engineering Biomimeticmultiscale Hierarchical Structures: From Lotus Leaf to Gecko Foot Hairs. Nanoscale 2009, 1, 331−338. (26) Kwak, R.; Jeong, H. E.; Suh, K. Y. Fabrication of Monolithoic Bridge Structures by Vacuum-Assisted Capillary-Force Lithography. Small 2009, 5 (7), 790−794. (27) Suh, D.; Choi, S. J.; Lee, H. H. Rigiflex Lithography for Nanostructure Transfer. Adv. Mater. 2005, 17, 1554−1560. (28) Pavia, D. L.; Lampman, G. M.; Kriz, G. S.; Vyvyan, J. A Introduction to Spectroscopy, 5th ed; Cengage Learning: Boston, 2014; pp 14−106. (29) Wu, S. Calculation of Interfacial Tension in Polymer Systems. J. Polym. Sci., Part C: Polym. Symp. 1971, 34, 19−30. (30) Kim, S. H.; Teymour, F.; Debling, J. A. A Unifying Approach for Melt Rheology of Linear Polystyrene. J. Appl. Polym. Sci. 2007, 103, 2597−2607.

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