Capillary Force Lithography Pattern-Directed Self-Assembly (CFL-PDSA) of Phase-Separating Polymer Blend Thin Films.

We report capillary force lithography pattern-directed self-assembly (CFL-PDSA), a facile technique for patterning immiscible polymer blend films of polystyrene (PS)/poly(methyl methacrylate) (PMMA), resulting in a highly ordered phase-separated morphology. The pattern replication is achieved by capillary force lithography (CFL), by annealing the film beyond the glass transition temperature of both the constituent polymers, while confining it between a patterned cross-linked poly(dimethyl siloxane) (PDMS) stamp and the silicon substrate. As the pattern replication takes place because of rise of the polymer meniscus along the confining stamp walls, higher affinity of PMMA toward the oxide-coated silicon substrate and of PS toward cross-linked PDMS leads to well-controlled vertically patterned phase separation of the two constituent polymers during thermal annealing. Although a perfect negative replica of the stamp pattern is obtained in all cases, the phase-separated morphology of the films under pattern confinement is strongly influenced by the blend composition and annealing time. The phase-separated domains coarsen with time because of migration of the two components into specific areas, PS into an elevated mesa region and PMMA toward the substrate, because of preferential wetting. We show that a well-controlled, phase-separated morphology is achieved when the blend ratio matches the volume ratio of the elevated region to the base region in the patterned films. The proposed top-down imprint patterning of blends can be easily made roll-to-roll-compatible for industrial adoption.

Introduction

Ever since the development of soft lithography and embossing-based patterning techniques specific to soft films in mid-1990s, there has been tremendous focus on nanopatterning of polymer thin films.[1] Nanopatterned films find wide applications in superhydrophobic surfaces,[2] microfluidic devices,[3] structural color,[4] flexible electronic devices,[5] nanobiotechnology,[6] templates for guiding self-organization, etc.[7] Among the various available techniques,[8−10] capillary force lithography (CFL) is a prominent method, which is capable of creating sub-100 nm features, utilizing capillary-driven rise of a soften polymer meniscus along the contours of a soft stamp during pattern replication.[10] Ease of implementation without any need of major hardware, possible use of a flexible stamp that allows easy peeling after pattern formation, and ability of patterning nonplanar surfaces, etc.[12] make CFL superior over other techniques like nano imprint lithography (NIL), which requires application of high pressure and is often encompassed with difficulties during detachment of a rigid mold. To date, films of various types of polymers,[1,8−11] as well as inorganic sol gel films, have been very successfully patterned by CFL and other embossing-based methods, resulting in purely topographic structures.[12]

Although nanopatterned surfaces find wide applications, polymeric nanostructures with both chemical and topographic features are likely to result in surfaces with multifunctional capabilities that might be useful in many emerging technologies. Patterning of a polymer blend thin film is a viable strategy toward this end, as it is well known that two immiscible polymers cast from a common solvent phase separate and form random domains of both the polymers.[13−21] The exact morphology of the phase-separated domains depends on multiple parameters such as the molecular weight of the two phases, their relative solubility in the common solvent, relative affinity toward the substrate, film thickness and other coating conditions, relative humidity of the surrounding, etc.[14,15] Polymer blend thin films are also important from the fundamental standpoint as they act as model systems for studying and understanding equilibrium and nonequilibrium properties in statistical physics, such as phase transition, composition fluctuation, effect of chain conformation, dynamics of the diffusion process, etc.[19] On the other hand, blending chemically different polymers is a useful approach to optimizing the properties of industrial products[20] and possible fabrication of functional coatings with exciting properties, such as antireflection and self-cleaning capability.[21] Although the phase-separated domains in a polymer blend thin film are isotropic, they have been aligned by “bottom-up” approaches by guiding the phase separation process on a chemically or topographically patterned substrate where the two constituent phases migrate to respective surface domains with preferential affinity toward that particular phase.[20−28] However, such ordering requires a prepatterned template.

Careful review of the published literature reveals that very few articles have reported direct “top-down” patterning of a thin polymer blend film,[29,30] arguably because of scientifically intriguing nature of the problem arising out of the spatial variation in the interaction between the phase-separated domains of the film and the stamp used for patterning. The presence of multiple confined interfaces between the two phases with random orientation leads to complexities during pattern replication. Ding et al. successfully patterned a polystyrene (PS)/poly(methyl methacrylate) PMMA blend film by nano imprint lithography, with an applied pressure of 4 MPa at different temperatures.[29] Huck et al. used NIL to pattern an active polymer blend layer and fabricated photovoltaic devices with better power conversion efficiencies.[30] In contrast, CFL has never been explored for patterning a polymer blend thin film, despite its wide success in patterning homopolymer thin films.[10,11] As already mentioned, CFL has several advantages as compared to those of NIL including reduced accumulation of stresses within the imprinted features and easy implementation without any dedicated instruments. In this article, we report a new technique termed as capillary force lithography pattern-directed self-assembly (CFL-PDSA), which utilizes CFL for patterning PS/PMMA blend thin films of different compositions. Although an overall perfect negative replica of the stamp is obtained for all of the blending ratios, the phase-separated morphology is found to be strongly dependent on the composition of the film as well as the time given for patterning. The organization of the phase-separated domains further evolves with progressive thermal annealing of the film while in conformal contact with the stamp because of confined capillary dynamics of the molten polymer above the TG of both the phases. We show that under certain conditions of controlled film thickness, blend composition, and annealing temperature and time, the blend films exhibit vertical, anisotropic patterned phase separation as PMMA has stronger affinity toward the silicon substrate and PS, toward the cross-linked PDMS stamp, resulting in a vertically phase separated, ordered morphology.

Results and Discussion

As-Cast Morphology and Temporal Evolution of Unconfined Blend Films

Films of four different blend compositions; (i) PS 33%/PMMA 67% (designated as S33/M67), (ii) PS 50%/PMMA 50% (S50/M50), (iii) PS 67%/PMMA 33% (S67/M33), and (iv) PS 75%/PMMA 25% (identified as S75/M25); were used in our experiments. To understand the dynamics and alignment of the phase-separated domains in the patterned blend films via CFL-PDSA, it is important to observe the as-cast morphology for each blend composition on a flat substrate and also to understand their evolution with time during thermal annealing. Although there is wide literature available on the morphology of PS/PMMA blends on silicon wafers, almost none of them report the morphology of PS/PMMA blend films prepared by flow coating. The as-cast morphology of a film with composition S50/M50 is shown as inset A2 of Figure A. The morphology comprises a continuous PMMA matrix layer with isolated PS domains, which is attributed to preferential wetting of the oxide-coated silicon substrate by PMMA.[33,34] The fractional coverage of PS (AF-PS) in percentage is found to be ≈37.8%. The as-cast morphology of films with other compositions is shown in Figure S1 of Supporting Information, where it can be seen that PMMA forms a continuous matrix even in films with composition S67/M33 because of the enhanced affinity of PMMA toward the silicon substrate, though the size of the PS domains continuously grows with an increase in PS fraction. Only in films with composition S75/M25, the matrix formation of PMMA is suppressed and a bicontinuous morphology is observed.

Figure 1

Atomic force microscopy (AFM) images of 3% S50/M50 blend films annealed for (A) 2 min, (B) 30 min, and (C) 8 h, at 180 °C without PDMS pattern confinement. Inset A2 shows the as-cast morphology. The cross-sectional profiles are shown under respective frames. (D–F) Corresponding images after selective extraction of PS. Scale bar in all cases is 10 μm.

The morphological evolution of a S50/M50 film after 2 min, 30 min, and 8 h of thermal annealing at 180 °C is shown in Figure A–C, along with respective solvent-extracted morphologies (Figure D–F) after preferential removal of PS. A comparison between Figure A and inset A shows that the isolated PS domains start to reorganize almost immediately as thermal annealing starts. With time, the PS domains start to coalesce, forming larger hemispherical droplets. The size of the rounded PS domains increases with annealing because of coarsening until an equilibrium configuration is achieved, which is driven by minimization of the interfacial area of PS with both PMMA and air. It can be observed that AF-PS reduces from 37.8% in Figure A to ≈27.8% after 30 min of annealing in Figure B, which clearly highlights that PMMA present in the system preferentially wets the substrate. Interestingly, AF-PS remains nearly unaltered during annealing beyond 30 min. As a consequence, the morphologies of the films shown in Figure B,C are nearly identical, including the cross-sectional profiles, feature height, and spacing of the PS domain (11.76 ± 0.37 and 11.77 ± 0.29 μm in Figure B,C respectively). The near-identical morphology clearly highlights that the domain reorganization as well as morphological evolution in the films is complete after 30 min of annealing itself. This pinned state is considered as a long-lived metastable state, as found in recent simulations of phase separation in capillaries.[16,17] The rapid reorganization dynamics of the domains can be attributed to very low molecular weight of the constituent polymers. This can be verified in Figure S2 of Supporting Information, where it can be seen that in a thin film of a symmetric blend (S50/M50) comprising high-molecular-weight PS (280 K) and PMMA (350 K) there is virtually no evolution after 8 h of annealing. Only partial morphological evolution is observed after 24 h of annealing. Similar coalescence of PS domains is also observed in blend films of compositions S67/M33 and S33/M67. For S75/M25 blend films with bicontinuous morphology, the PS threads initially take well-rounded configuration, before they disintegrate into isolated droplets. With time, the PS droplets coalesce into larger droplets. The evolution sequence for films with bicontinuous morphology is shown in Figure S3 of Supporting Information.

Patterning of the PS/PMMA Blend Films

Whereas a flat film is in contact only with the substrate, a film during patterning is confined between the stamp and the substrate and therefore spatio-temporal reorganization of the phase-separated polymer domains during thermal annealing is strongly influenced by the wetting behavior of the constituent polymers on both the stamp and the substrate. Using water, ethylene glycol, and toluene as a probing liquid in a contact angle goniometer, the surface energy of the cross-linked Sylgard 182 block is calculated to be γPDMS = 24.5 mJ/m2.[32] The surface energies of PS and PMMA used are γPS = 28.9 mJ/m2 and γPMMA = 29.2 mJ/m2, respectively, both of which are higher than γPDMS. The interfacial tensions of PS and PMMA with cross-linked PDMS are found to be γPS–PDMS = 5.96 mJ/m2 and γPMMA–PDMS = 8.91 mJ/m2, respectively. The surface and interfacial tension values were calculated by dispensing water, ethylene glycol, and diodomethane on spin-coated thin films (roughly 300 nm thick) of PS and PMMA.[35] The values of spreading coefficients of PS and PMMA on PDMS substrates were found to be SPS–PDMS = −10.36 mJ/m2 and SPMMA–PDMS = −13.61 mJ/m2, respectively. Whereas negative values of both the spreading coefficients indicate that neither PS nor PMMA would prefer to wet PDMS, a lower value of SPMMA–PDMS as compared to that of SPS–PDMS implies that under confinement a PDMS surface would prefer to be in contact with PS than with PMMA.

Next, we estimate the time needed by the molten polymer blends to fill the PDMS mold (tfill) from the following equation

where R is the hydraulic radius, half the channel width (205 nm); η is the polymer zero-shear viscosity; z is the mold depth (140 nm); γ is the surface tension of the polymer; and θ is the contact angle at the polymer–mold interface.[10] At 180 °C, ηPS = 11.8 Pa s, ηPMMA = 297.7 Pa s, θ ≈ 84°, and corresponding tfill is in the range of 5.2 × 10–4 s (pure PS) to 0.02 s (pure PMMA), indicating that the mold fill process is extremely fast.[33] However, the calculated value of tfill is based on the assumption that the polymers are in the molten state. In reality, once the samples are subject to annealing, there is a finite time requirement for the polymer to get heated beyond TG and transform to a molten state. Furthermore, we estimated the maximum height of capillary rise using the following equation[10]where ρ is the density of polymer (approximately 1.1 g/cm3), g is the gravitational constant (9.81 m/s2), and L is the channel width full width at half maximum (FWHM) = 410 nm. At 180 °C, hmax is found to be ≈1.55 m. This confirms that the capillary force is strong enough to fill the polymer into the mold, as hP is only 140 nm. Finally, a simple calculation based on the remnant layer thickness ≈30 nm in all cases, it is found that the volume fraction of the elevated region to the base region in a patterned film is 71:29 for the specific stamp used for patterning, as can be seen in Figure S4 of Supporting Information.

Figure A shows the AFM image of a patterned S50/M50 blend film after 30 min annealing at 180 °C. Inset A1 of the figure shows the AFM image of the stamp used for patterning. Figure B shows the superimposed cross-sectional profile of the patterned film and the stamp used for patterning (inset A1). Inset A2 of Figure A presents a detailed scaled model of the pattern cross section based on the actual pattern dimensions obtained from AFM analysis including the measurement of the thickness of the remnant layer by scratching, which is found to be ≈30 nm.

Figure 2

(A) AFM image of the patterned S50/M50 blend film. Inset A1 shows the image of the grating-patterned PDMS stamp. (B) Cross-sectional profiles of the patterned film and PDMS mold exhibiting the perfect negative replica formation. (C) Optical micrograph of the patterned blend film. The scale bar corresponds to 20 μm.

It is important to highlight that for a specific stamp the thickness of the remnant layer depends on dinitial, which was kept constant at ≈100 nm for films with all different compositions. Both the stamp and the patterned film have λP = 750 nm and FWHM of the height region peak is 410 nm, which confirms that a perfect negative replica of the stamp pattern forms on the blend film, with no loss of fidelity during pattern transfer by CFL. The optical micrograph in Figure C shows that uniform pattern replication has taken place over a large area of the patterned S50/M50 film, which confirms the large area of the film surface.

Identical pattern replication was also observed in blend films with other compositions as well (refer to Figure S5 of Supporting Information) for annealing times of 30 min and above. Interestingly, the large-area optical microscope images of the patterned films, both after 30 min and 8 h of annealing, show no signs of phase separation at all, which is always observed in flat thin films of immiscible blends, such as PS/PMMA, and is shown in Figure . Even in the only earlier article on nano imprinting of PS/PMMA blend thin films by Ding et al., signatures of phase separation on the patterned films over large areas was observed.[29] The disappearance of the phase-separated domain boundaries in Figures and S5 is attributed to flow-driven reorganization of low-molecular-weight PS and PMMA under a confining stamp, clearly highlighting the effectiveness of CFL in patterning a polymer blend film over NIL, particularly in terms of pattern uniformity control. In certain cases, the patterns obtained after 2 min annealing exhibit incomplete replication with signs of phase-separated domains. This is due to the finite time required by the polymers to transform into the molten state under actual experimental conditions.

Evolution of Phase-Separated Morphology of Patterned Films

In Figure , we show the morphology of the PMMA domains after selective extraction of the PS phase, in films annealed for 2 min (column 1), 30 min (column 2), and 8 h (column 3). The main frames of Figure show the solvent-extracted morphology (after preferential removal of PS) of the patterned films with different compositions: S33/M67 (series A); S50/M50 (series B); S67/M33 (series C), and S75/M25 (series D). The insets to the first and second images in each series show the morphology of the patterned film before selective removal of PS. The images illustrate how the PS and PMMA domains reorganize with time under the confining stamp. Column 4 shows the cross-sectional profiles of the samples shown in column 4, and column 5 schematically represents the organization of the phase-separated domains in the patterned films after 8 h of annealing (the images are not to scale).

Figure 3

AFM images of patterned PS/PMMA blend films with varying compositions and annealing times after the extraction of the PS phase. The compositions are S33/M67 (series A), S50/M50 (series B), S67/M33 (series C), and S75/M25 (series D). The patterning times used are 2 min (A1–D1), 30 min (A2–D2), and 8 h (A3–D3). The insets show the images of patterned films before PS phase extraction. (A4–D4) Sections in each case for 8 h annealed samples after phase extraction. (A5–D5) Schematic showing the composite architecture of individual stripes in each case.

The first image in each series shows that the patterns start to appear just after 2 min of annealing, though some signatures of the initial phase-separated domains are still present. By 30 min of annealing, a perfect negative replica of the stamp patterns was obtained on the film surface for each blend composition (the inset to the second figure in each series of Figure ). This happens due to very short mold filling time (5.2 × 10–4 s > tfill > 0.02 s) as well as very shallow depth of the stamp used for patterning (140 nm). However, in many cases, the phase-separated domains are seen to further reorganize beyond 30 min, without affecting the overall replicated pattern morphology. This means that the morphological evolution in a film confined between a stamp and the substrate during patterning is slower than that in a flat film, as we have seen in Figure that the morphological evolution as well as domain reorganization in a flat film is complete within 30 min. We argue that the additional shear resistance to flow at the film–stamp interface is responsible for sluggish domain reorganization in a film sandwiched and confined between two rigid substrates as compared to that in a flat unconfined film, which has a zero-shear air–film interface.

Figure highlights that with progressive annealing the PMMA phase starts to migrate toward the substrate because of its strong affinity to the SiO surface, irrespective of the blend composition. This in turn results in migration of PS present in the films toward the stamp, as it maximizes the contact area of the PDMS stamp with PS, which is obvious from the values of spreading coefficients reported earlier. Furthermore, within the grooves, both the PMMA phase and PS phase coalesce with their neighboring PMMA phase and PS phase, respectively, and subsequently get packed along the replicated stripes. In films with compositions S67/M33 and S75/M25, the PS domains coalesce and accumulate under the entire stamp grooves to form PS threads spanning the grooves over the PMMA layer, as has been schematically shown in Figure C5,D5, respectively. Interestingly, in both the cases, the PS fraction is very close to 71%, which is the volume fraction of a patterned film that lies within the grooves. As the PS fraction is just below 71% in the S67/M33 film, the PMMA layer forms a very shallow arch below every PS thread (Figure C3 and cross section in Figure C4). In contrast, a higher PS fraction in a S75/M25 film leads to a concave curvature of the PMMA meniscus at the PS–PMMA interface below each stripe (Figure D3,D4). PMMA forms a tiny ridge along the edge of the stamp grooves arguably due to pinning. At both the compositions, we can argue that perfect vertical phase separation between PS and PMMA has been achieved, resulting in the formation of aligned isolated PS threads over a continuous PMMA underlayer.

In contrast, in films with low PS fraction (S33/M67 and S50/M50), PS fails to cover the entire span of the stamp groove because of the inadequate amount of PS present in the films. Consequently, PMMA is also present in the protruded areas of the patterned film, under the stamp grooves. It can be seen from Figure A3,B3 that with a decrease in PS fraction, the PMMA domains within the replicated stripes become progressively larger. Figure B2 shows an interesting intermediate morphology, where PS and PMMA threads laterally coexist along each replicated stripe after 30 min of annealing. However, with time, the tiny PS threads break down in favor of alternating PS and PMMA domains along the stripes. Vertical scale bars of ≈140 nm in both Figure A3,B3 imply that the PMMA domains span the depth of the stamp grooves. This reorganization is probably attributed to minimization of the PS–PMMA interfacial area within the grooves. The estimated volume fraction of PMMA within the elevated regions of the patterned films as a function of blend composition and annealing time is shown in Figure . The figure clearly highlights the gradual migration of PMMA from the stripes to the base of the film with progressive annealing and also near-complete stratification of PS within the grooves and PMMA over the entire base in PS-rich films.

Figure 4

Estimated ratio of the volume of PMMA in the elevated region (VPMMA) to the volume of the elevated region (Ve) in films annealed for 2 min (black square), 30 min (red circle), and 8 h (blue triangle) as a function of the (A) volume fraction of PS (ϕPS) and (B) annealing time.

Finally, we would like to demonstrate one specific application of the patterned blend thin films, particularly the ones with ordered vertical phase separation (shown in Figure C3,D3) where strips of PS are formed over a continuous matrix or remnant layer of PMMA. We demonstrate that the ordered PS strips are elegantly used as a self-organized mask to etch the exposed portion of the PMMA remnant layer below it. This is achieved by simple UVO exposure of the patterned blend films. It is well known that UVO exposure preferentially degrades PMMA as compared with PS.[36] Thus, when the patterned film is exposed to UVO in a UV–ozone chamber (Novascan Technologies), the thickness of the PMMA film between the PS strips gradually reduces with time, resulting in a grating structure that has a higher feature height than that of the original stamp. Figure shows the morphology of such a grating where the stripe height is 158 nm obtained after 25 min of UVO of a patterned S75/M25 blend film. It may be recalled that the height of the stamp features is 140 nm, which can be seen in the superimposed cross-sectional profile. Thus, patterning of blend thin films with vertical phase separation, in conjugation with UVO exposure, can be used as a simple method for creating nanopatterns, which are taller than the original stamp patterns and cannot be created by any direct embossing technique, and is considered to be a scientific challenge which is only partially resolved.[37]

Figure 5

AFM image of a patterned blend thin film (S75/M25) UVO-exposed for 25 min. The cross-sectional profile of the UVO-exposed patterned film (continuous line) and the superimposed cross-sectional profile of the stamp (dashed line).

Conclusions

We have utilized capillary force lithography pattern-directed self-assembly to fabricate topographically uniform patterns on phase-separated polymer blend thin films. We show that a perfect negative replica of the stamp pattern can be obtained on blends of any composition after 30 min of annealing. However, the phase-separated domains continue to evolve and reorganize with progressive annealing without altering the overall patterned morphology when the films are annealed beyond 30 min. During this stage, preferential migration of PMMA toward the substrate coupled with coalescence of individual PS and PMMA domains leads to a well-ordered phase-separated morphology. Of particular interest is the vertical phase separation in PS-rich films, where PS accumulates entirely along the stamp grooves, forming PS threads over a continuous PMMA layer. This is the first example of the formation of such well-controlled phase-separated structures due to vertical phase separation in a polymer blend thin film, confined between a stamp and the substrate. On the other hand, alternating PMMA and PS domains are arranged along each replicated stripe in films with a high PMMA fraction. This capillary force lithography pattern-directed self-assembly (CFL-PDSA) has much potential for a wide range of vertical patterning of a large range of polymer blend films ubiquitously used in the thin film industry. The patterns are extremely reproducible and span over large areas (cm2). The proposed method is facile, easy to implement without the need of any major lithography or fabrication facility, and highlights the possibility of fabricating polymeric nanostructures with both chemical and topographic features, which are likely to find application as surfaces with multifunctional capabilities that might be useful in many emerging nanotechnologies.

Materials and Methods

Material

Polystyrene (PS, weight-average molecular weight, Mw = 16 400 g/mol; polydispersity index, PDI = 1.16; Polymer Source Inc.) and poly(methyl methacrylate) (PMMA, Mw = 15 000 g/mol; PDI = 1.12; Polymer Source Inc.) were used as received. TG values for PS and PMMA are 95 and 101 °C, respectively. Poly(dimethyl siloxane), PDMS, Sylgard 182, was purchased from Dow Corning. p-Type silicon substrates were bought from Silicon Quest International, Inc., and digital versatile disc (DVD) was purchased from Sony Corporation, Japan. For some control experiments, PS with Mw 280 kg/mol and PMMA with Mw 350 kg/mol (Sigma) were also used.

Preparation of Stamp for Patterning

The stamp for patterning the blend films was created by replica molding the structures present on the surface of a commercial DVD on a self-standing block of Sylgard 182 (cross-linked PDMS). For this purpose, the polycarbonate layer of DVD was separated and the metallic reflecting layer on the surface of DVD pattern was removed by scotch tape, followed by rinsing in acetone, as has been described in detail elsewhere.[31] The liquid PDMS and curing agent were thoroughly mixed in the weight ratio of 10:1, and the mixture was degassed to remove the entrapped air bubbles. The PDMS prepolymer mixture was then poured onto the DVD master and was cured at 120 °C for 4 h for complete cross-linking. The patterned cross-linked Sylgard stamps were subsequently manually peeled from the DVD master. The stamps had grating geometry, with pattern periodicity (λP) = 750 nm, stripe height (hP) = 140 nm, duty ratio = 1, and FWHM of the height region peak around 410 nm.

Blend Thin Film Preparation

The concentration of the casting solution was maintained at 3% polymer in toluene (w/v) for all different blend compositions. Before coating the films, silicon substrates were rinsed with toluene, followed by 1 h ultraviolet–ozone treatment. Then, PS/PMMA blend films were flow-coated from prepared PS/PMMA blend solutions onto the UVO-treated silicon substrate. The gap and speed of flow coater were adjusted to obtain the predetermined initial film thickness (dinitial) of 100 nm. The film thickness after coating was measured by an interferometer. Prior to the thermal annealing process, the residual solvent was extracted by placing the films in an oven at 50 °C for 6 h under vacuum.

Patterning of the Blend Films

The blend films were patterned by a low-pressure-assisted CFL.[11] The patterned cross-linked PDMS stamp was first brought in contact with the blend films and was subject to a uniform load of 4 kPa to ensure complete conformal contact between the stamp and the film surface. The films sandwiched between the substrate and the stamp were then subjected to annealing in an oven at 180 °C under vacuum. We report the morphology of the patterned films after 2 min, 30 min, and 8 h of annealing.

Selective Solvent Extraction

To study the morphology of the phase-separated domains, PS was extracted by immersing the films in cyclohexane for 30 min at 22 °C, followed by drying in a stream of dry nitrogen. In certain cases, the PMMA layer was also degraded by UVO exposure.

Characterization

The surface topography of the films (as-cast, patterned, and solvent-extracted) were investigated using an atomic force microscope (AFM, Dimension ICON, Veeco at Akron and 5100, Agilent Technologies at Kharagpur) in intermittent contact mode. To measure the thickness of the residual layers and absolute height of the pattern, the scratch test was performed using a sharp blade, followed by measuring the height difference between the silicon substrate and the film surface.