Demixing of Polymers under Nanoimprinting Process

Transcription

1 University of Colorado, Boulder CU Scholar Mechanical Engineering Graduate Theses & Dissertations Mechanical Engineering Spring Demixing of Polymers under Nanoimprinting Process Zhen Wang University of Colorado at Boulder, Follow this and additional works at: Part of the Mechanical Engineering Commons, and the Physics Commons Recommended Citation Wang, Zhen, "Demixing of Polymers under Nanoimprinting Process" (2013). Mechanical Engineering Graduate Theses & Dissertations This Dissertation is brought to you for free and open access by Mechanical Engineering at CU Scholar. It has been accepted for inclusion in Mechanical Engineering Graduate Theses & Dissertations by an authorized administrator of CU Scholar. For more information, please contact

2 DEMIXING OF POLYMERS UNDER NANOIMPRINTING PROCESS by ZHEN WANG B.E., Dalian University of Technology, 2008 M.S., University of Colorado, 2010 A thesis submitted to the Faculty of the Graduate School of the University of Colorado in partial fulfillment of the requirement for the degree of Doctor of Philosophy Department of Mechanical Engineering 2013

3 This thesis entitled: Demixing of Polymers under Nanoimprinting Process written by Zhen Wang has been approved for the Department of Mechanical Engineering Prof. Yifu Ding Prof. Scott Bunch Date The final copy of this thesis has been examined by the signatories, and we Find that both the content and the form meet acceptable presentation standards Of scholarly work in the above mentioned discipline.

4 iii Wang, Zhen (Ph.D., Mechanical Engineering) Demixing of Polymers under Nanoimprinting Process Thesis directed by Professor Yifu Ding Polymer blend has been an important area in polymer science for several decades. The knowledge of polymer blend in bulk is well established and technologies based on it have created products ubiquitous in our daily life. More intriguing problem arises when the phase separation process of a polymer blend occurs under physical confinement with length scale less than the correlation length of the intrinsic phase-separated structure. In this thesis, we aim to understand the effect of interfacial interactions between constituent polymers and confinement environment on morphology evolution and to demonstrate viable approaches to fabricating multifunctional materials with unique micro or nanostructures. Phase evolution of free-surfaced polystyrene/poly(methyl methacrylate) (PS/PMMA) blend films on a polymer brush surface (energetically neutral to both components, referred to as neutral surface) was compared with that on a SiOx surface (preferential to PMMA). By tuning the substrate surface energy we can manipulate the pathway as well as the kinetics of the phase-separated domain coarsening process. The conventional formation of a PMMA wetting layer on SiOx surface can be prohibited by the neutral surface and as a result PS domains directly contact with the substrate. In stark contrast, under thermal embossing nanoimprint lithography (TE-NIL) conditions, the physical confinement and/or external pressure prevent the preferential wetting tendency of PMMA, diminishing the influence of the substrate surface energy. When a planar blend film is pressurized between two parallel rigid substrates, the surface-relief structure dominant in the free-surfaced case is completely avoided. The resulted phase evolution process either behaves as if it is frozen or shows anisotropic coarsening depending on the mobility of polymer

5 iv chains. When the confinement is realized by utilizing a template with line-space geometry feature, all the patterned PS/PMMA films show topographically uniform structure, regardless of the surface roughness formed during the initial stage of the phase separation. The phase structures of the PS/PMMA patterns are found to be dictated by the preferential wetting of PMMA onto SiOx substrate despite the pressurization similar to the planar confinement. The interplay between this preferential wetting and the domain coalescence results in a range of complex encapsulated structures. Moreover, step-and-flash NIL (SF-NIL) was applied to NOA65/5CB (4-cyano- 4'-pentylbiphenyl) blend to investigate the phase separation of initially mixed binary composite. Surprisingly, a layered phase-separated structure was observed, which is different from previous reports on UV-induced phase separation in thin films. This is likely due to the preferential wetting of 5CB onto the confining polymer template and the disturbing of the intrinsic randomly distributed 5CB nuclei under confinement.

6 v ACKNOWLEDGEMENTS This thesis is completed under the guidance of my advisor Dr. Yifu Ding. First and foremost, I would like to express my gratitude to him for his continuous encouragement, patience and support. From him I learned immense factual knowledge of polymer physics and how to think critically. I feel lucky to be his first Ph.D. student. I would also like to thank Dr. Dae Up Ahn who showed me how to carry out experiments perfectly and professionally. I want to acknowledge all the other group members Liang Wang, Sajjad Maruf, Charlie Zhang, Lewis Cox, and Dr. Devid Maniglio for helpful discussion and collaboration. Thanks also go to my thesis committee members Dr. Scott Bunch, Dr. Jerry Qi, Dr. Mark Stoykovich, and Dr. Wei Tan for their insightful suggestions and corrections. Last but not least, I sincerely thank my parents and friends for their help and support.

7 vi CONTENTS CHAPTER I. INTRODUCTION...1 II. THEORY Thermodynamics of polymer blends Wetting on surface Cahn-Hilliard theory III. EXPERIMENTAL Nanoimprint lithography (NIL) Thermal embossing NIL (TE-NIL) Step-and-flash NIL (SF-NIL) Substrate treatment Neutral surface Fabrication of chemical pattern by NIL Characterization of polymer blend morphology Optical microscopy (OM) Atomic force microscopy (AFM) IV. PHASE EVOLUTION OF BINARY POLYMER THIN FILMS WITH FREE SURFACES Introduction Samples and techniques... 39

8 vii 4.3 Phase evolution of PS/PMMA films on SiOx surface Phase evolution of S60/M40 films Phase evolution of S67/M33 films Phase evolution of S80/M20 films Influence of neutral surface of morphology evolution Phase evolution of S60/M40 films Phase evolution of S67/M33 films Phase evolution of S80/M20 films Effect of chemically patterned substrate on structure formation of spun-cast blend films Conclusions V. PHASE EVOLUTION OF PLANAR THIN FILMS CONFINED BETWEEN TWO PARALLEL RIGID WALLS (1-D CONFINEMENT) Introduction Samples and techniques Confinement with symmetric boundary condition (Si wafers with native SiOx surfaces) Confined phase evolution with fixed film thickness Thickness effect on phase evolution Molecular weight effect on phase evolution Effect of substrate surface energy Conclusions VI. COARSENING PROCESS OF POLYMER BLENDS UNDER TE-NIL CONDITIONS (2-D CONFINEMENT) Introduction... 97

9 viii 6.2 Samples and techniques Structure formation of PS(190K)/PMMA(94K) films Patterned morphology of 150-nm PS/PMMA films Morphology of patterned S30/M70 films Morphology of patterned S50/M50 films Morphology of patterned S70/M30 films Substrate surface energy effect on morphology Film thickness dependence of morphology A comparison between phase evolutions in thin films under 1-D confinement and 2-D confinement Conclusions VII. CONFINED REACTION-INDUCED PHASE SEPARATION UNDER SF-NIL CONDITIONS Introduction Samples and techniques Morphology characterization Electro-optic response of ultra-thin confined phase-separated composite film (C-PSCOF) Conclusions VIII. SUMMARY BIBLIOGRAPHY

10 ix FIGURES Figure 2.1 Schematic phase diagram for a binary polymer mixture A schematic illustration of phase-separated structures initialized by nucleation and growth (left) and spinodal decomposition (right) Phase diagram of PS190K/PMMA94K based on compressible regular solution model Contact angle and forces associated with a liquid droplet on solid substrate Final possible equilibrium structure for partial wetting (a) and complete wetting (b) of wettable phase An example of simulation result based on Cahn-Hilliard equation. A, the morphology developed at an early stage of phase separation. B, the morphology after significant coarsening has occurred. C, the simulated morphologies within high-aspect ratio domains at different times D simulation results based on Cahn-Hilliard equation in bulk and confinement situations for two different blend compositions Simulation results for phase separation with asymmetric boundary condition. The bottom of the simulated region has interaction strength = 0.5. A, simulated bulk-like phase-separated structure. B, simulated thin film cases with aspect ratio of 8 and different coarsening times... 21

11 x 2.9 Simulation results for phase separation on chemical pattern, which has alternating = 0.5 and = -0.5 with periodicity of 32. A, simulated bulk-like phase-separated structure. B, simulated thin film cases with different film thicknesses Simulated phase-separated morphologies (top view) with different interaction strength values. is 0, 0.05 and 0.5 in A, B and C, respectively. All simulated domains are and time t is Chemical pattern periodicity is Simulation results for phase separation with symmetric boundary condition. Both top and bottom of the simulated region have interaction strength = 0.5. A) simulated bulk-like phase-separated structure. B) simulated thin film cases with different coarsening times Schematic configurations of phase-separated films for different types of boundary conditions, as predicted by Binder s model. A, equilibrium state under non-symmetric boundary conditions; B and C, equilibrium and non-equilibrium states under symmetric boundary conditions. B shows the situation when temperature T is lower than wetting transition temperature Tw, or wetting is less dominant. C shows the situation when T is larger than Tw and wetting of the wettable component plays an important role in the formation of phase separated structure Phenomenological phase diagram for binary liquid mixture under the confinement of a cylindrical pore. The formation of the three configurations of plug, capsule and tube depends on the quenching depth t = ( )/ and, where is molecule length and is the pore diameter Steps in a traditional TE-NIL process. Cast polymer film on a substrate, imprint at elevated temperature with pressure for a certain amount of time, and demold at temperature below the glass transition temperature of the polymer resist... 29

12 xi 3.2 Steps in SF-NIL process. Cast photoactive monomer resist onto a substrate, press and shine UV light to cross-link the resist, and separate the mold Schematic drawing of fabrication process of chemical pattern by NIL AFM cross-section profile of the chemical pattern surface Topographic AFM images (a-d) and optical micrograph (OM) (e) of S60M40 on native SiOx, after annealing at 160 C for different durations as labeled. The lower right insets in (a-e) are the topographic AFM images of the sample after selective removal of either PS (a-d) or PMMA (e). The lower left insets in (c-e) are the FFT images of the AFM image or OM (not to scale). (f) Schematic illustration of the early stages and evolution of the blend morphology (a) Lateral domain correlation lengths, and (b) RMS surface roughness, of thin PS/PMMA films as a function of annealing duration, for a varying PS/PMMA composition (a) Topographic AFM images of as-cast S67M33 on a native SiOx layer. The inset is the topographic AFM image of the as-cast film after the selective removal of PS. (b) Optical micrograph of the S67M33 after annealing at 160 C for 100 min. The inset is the FFT image of the optical image (a) Topographic AFM image of as-cast S80M20 film on a SiOx surface. The inset in (a) is the topographic AFM image of the as-cast film after the selective removal of PS. (b-d) Optical micrographs of S80M20 after annealing at 160 C for different tas, showing: (b) random nucleation of holes in the top PS layer, (c) growth and impingement of the holes, and (d) formation of isolated PS droplets from the breakup of the impinged PS lines. Inset in (d) is the FFT of the optical image. The sizes of these optical images are 400 m 400 m... 46

13 xii 4.5 Topographic (a) and phase (b-f) AFM images of S60M40 on the random copolymer-treated substrate after annealing at 160 C for different tas as labeled. The lower right insets are the corresponding topographic AFM images after the selective removal of either PS (a-e) or PMMA (f). The lower left insets in (c-e) are the FFTs of the corresponding lower right insets (not to the scale) Topographic AFM images of S67M33 on the random copolymertreated substrate after annealing at 160 C for different tas as labeled. The lower right insets are the corresponding topographic AFM images after the selective removal of either PS (a-e) or PMMA (f). The lower left insets in (c and d) are the FFT images of the corresponding lower right insets (not to the scale) Topographic (a-e) and phase (b-d and f) AFM images of S80M20 on the random copolymer-treated substrate after annealing at 160 C for different tas as labeled. The lower right insets are the topographic AFM images after the selective removal of either PS (a-c, e, and f) or PMMA (d). (g) and (h) are the optical images of the film after 200 min and 1000 min, correspondingly Schematics of the later stage morphology after annealing on RCP surfaces for all three blend compositions A comparison of structure formation of phase-separated PS/P2VP films spun-cast from 3 wt.% solution on chemical pattern with that on ODT surface. A and B are topographic AFM images of morphology formed on ODT surface with size 10 μm 10 μm and 50 μm 50 μm, respectively. C and D are corresponding topographic AFM images of morphologies on chemical pattern A comparison of structure formation of phase-separated PS/P2VP films spun-cast from 0.75 wt. % solution on chemical pattern with that on ODT surface. A and B are topographic AFM images of morphologies of ascast film and cyclohexane-etched film formed on ODT surface,

14 xiii respectively. C and D are corresponding topographic AFM images of morphologies on chemical pattern. All AFM images are 10 μm 10 μm A transition from spinodal decomposition-based hierarchical structure to broken lines as the solution concentration decreases. A, B, C, and D are topographic AFM images corresponding to morphologies of ascast films from solutions with concentration 3 wt.%, 1.5 wt.%, 0.75 wt.%, and wt.%, respectively, on chemical pattern. All AFM images are 10 μm 10 μm Topographic AFM images of morphologies of PS48K/PMMA15K films with free surfaces. A, morphology of as-cast film. B, morphology of film annealed in air at 160 C for 30 min. C and D, morphologies of films after annealing in imprinter at 110 C for 2 min and then in air without external pressure at 160 C for 30 min. C is the morphology of the film left on the mold side and D is on the substrate side. All the AFM images have the same size of 50 m 50 m. The color (z-) scale is 40 nm in A and 100 nm in B, C, and D. Insets in B and D are FFT images obtained from corresponding optical microscope images A schematic illustration of the formation process of the symmetric surface-relief structure. (A) as-cast morphology after spin-coating. (B) morphology formed after pressurization at 110 C for 2 min. Arrows indicate the direction of materials flow tendency upon further annealing. (C) thermally induced capillary wave occurred within the film during the annealing without external pressure. (D) final morphology of symmetric surface-relief structure, corresponding to the cross-sectional profiles of Figures 5.1C and 5.1D Topographic AFM images for phase evolution of PS48K/PMMA15K films between two parallel plates. AFM images were taken at film surfaces revealed at the bottom interface after selective removal of PS. Imprinting temperatures are 160 C, 190 C, 210 C, and

15 xiv 230 C, respectively, as marked on each image. All the AFM images have the same size of 50 m 50 m and color (z-) scale of 200 nm Schematic for the driven force of the wetting layer formation. A, wetting occurred in free surface. B, the prevention of wetting layer growth under confinement Schematic of thinning of PMMA layers. A, thinning of intrinsic PMMA wetting layer. B, thinning of artificial PMMA cushion layer. C, topographic AFM image corresponding to the coarsened morphology in B A plot of correlation length, PMMA domain area, and PMMA domain width as a function of imprinting temperature SEM image for film pressurized at 230 C for 30 minutes, which was taken from the same sample as Figure 5.3D. Dark domains in the image are PMMA Histograms of PMMA domain width Lmin. Density defined as normalized frequency is used as y-axis for better visualization. A, B, C and D are for imprinting temperature 160 C, 190 C, 210 C, and 230 C, respectively Histograms of PMMA domain aspect ratio AR. Density defined as normalized frequency is used as y-axis for better visualization. A, B, C and D are for imprinting temperature 160 C, 190 C, 210 C, and 230 C, respectively Phase evolution of PS48K/PMMA15K films with different film thicknesses between two parallel plates. A-E, topographic AFM images of film surfaces revealed at the bottom interface after selective removal of PS. Film thicknesses are ~50 nm, ~100 nm, ~200 nm, ~600 nm and ~960 nm, respectively. F is the plot showing the relationship between film thickness and blend solution concentration SEM images for the corresponding AFM images in Figure A, B, C and D correspond to A, C, D, E in Figure 5.10, respectively... 85

16 xv 5.12 Scatter plots of Lmin as a function of domain area for films with different thicknesses. A, B, and C are scatter plots for films with thickness ~50 nm, ~100nm, and ~600 nm, respectively. D is the normalized deviation of scatter points from ideal isotropic coarsening curve y = Morphology of confined PS190K/PMMA94K film between two flat plates imprinted at 180 C. A, B and C are AFM topographic images of interface between polymer film and template wafer. A is the morphology of film surface. B is the corresponding leftover on the template wafer. C is the film morphology after selective removal of PS. D, E and F are AFM topographic images of interface between polymer film and substrate, corresponding to A, B and C, respectively. All AFM images are 10 m 10 m with color (z-) scale 200 nm Morphology of confined PS190K/PMMA94K film between two flat plates imprinted at 210 C. A, B and C are AFM topographic images of interface between polymer film and template wafer. A is the morphology of film surface. B is the corresponding leftover on the template wafer. C is the film morphology after selective remove of PS. D, E and F are AFM topographic images of interface between polymer film and substrate, corresponding A, B and C, respectively. All AFM images are 10 m 10 m with color (z-) scale 200 nm Morphologies of ~100-nm PS48K/PMMA15K films confined between two plates with neutral surfaces at 160 C and 230 C for 30 min. (A) topographic AFM image for film pressurized at 160 C. The color (z-) scale of 200 nm. (B) SEM image for film pressurized at 230 C. A and B are very similar to their SiOx counterparts Figure 5.3A and Figure 5.7, respectively A schematic of the geometry of the nanoimprinting mold used in Chapter 6, which has a periodicity of 834 nm, a line-to-space ratio of 1:1, and a cavity depth of ~195 nm... 99

17 xvi 6.2 Optical micrographs (OM) and topographic AFM images of PS/PMMA blend films with film thickness around 150 nm and three different compositions. The first, second and third rows show morphologies of S30/M70, S50/M50 and S70/M30, respectively, as marked on the left. The images in the first column are optical micrographs of ascast films. The second, third and fourth columns are topographic AFM images corresponding to the morphologies of as-cast films, films after PS was selectively dissolved by cyclohexane (CLE), and films after PMMA was selectively removed by acetic acid (AA), respectively. All AFM images are 10 µm 10 µm in size with color scale Optical micrographs (OMs) and topographic AFM images of 150- nm S30/M70 blend films annealed at three different temperatures as marked on the left. The images in the first column are optical micrographs of surface morphologies. The second, third columns are topographic AFM images corresponding to surface morphologies, and interfacial morphologies after PS was selectively dissolved by cyclohexane (CLE). All AFM images are 10 µm 10 µm in size Images of S30/M70 patterns imprinted at three different temperatures, as marked on the left. From left to right, different columns correspond to optical micrographs of films after imprinting, corresponding topographic AFM images of the as-imprinted films, after the removal of PS with CLE, and after the removal of PMMA with AA. All AFM images are 10 µm 10 µm. Insets are the FFT images of the optical micrographs Optical micrographs (OMs) and topographic AFM images of 150- nm S50/M50 blend films annealed at three different temperatures as marked on the left. Listed in the first column are the optical micrographs of surface morphologies (A1, B1, and C1). The dark and bright domains are the PS and PMMA, respectively in all three OM images. The second and third columns are topographic AFM images corresponding to surface morphologies, and interface morphologies after PS was selectively

18 xvii dissolved by cyclohexane (CLE). All AFM images are 10 µm 10 µm in size Images of the S50/M50 patterns imprinted at three different temperatures, as marked on the left. From left to right, different columns corresponding to optical micrographs of films after imprinting, corresponding topographic AFM images of the as-imprinted film, after the removal of PS with CLE, and after the removal of PMMA with AA. All AFM images are 10 µm 10 µm. Insets are the FFT images of the optical images. The dark and bright regions in the optical images correspond to the PS and PMMA phases, respectively Images of the S70/M30 patterns imprinted at three different temperatures, as marked on the left. From left to right, different columns corresponding to optical images of films after imprinting, corresponding topographic AFM images of the as-imprinted film, after the removal of PS with CLE, and after the removal of PMMA with AA. All AFM images are 10 µm 10 µm. Insets are the FFT images of the optical images. D, The cross-sectional profiles of B2, B4, and C4, as marked in the AFM images Phase evolution of 150-nm free-surfaced planar S50/M50 films annealed for 30 min at different temperatures on neutral surface. The first column is AFM topographic images for as-cast films. The second, third and fourth columns are AFM topographic images for films annealed at 150 C, 180 C and 210 C, respectively. The first row reveals the surface morphology while the second row shows the interfacial morphology after selective removal of PS by cyclohexane. The third and fourth rows are for planar films on SiOx. All AFM images are 10 µm 10 µm in size Surface and interface roughnesses of S50/M50 planar films annealed at different temperatures on SiOx surface and neutral surface Morphology of 150-nm S50/M50 films nano-imprinted at 180 C for 30 min on neutral surface. A1, B1 and C1 are topographic AFM images

19 xviii for as-imprinted film, film etched by cyclohexane and film etched by acetic acid, respectively, as marked on top of the figure. A2, B2 and C2 are taken from Figure 6.6 and are included for comparison. All AFM images are 10 µm 10 µm in size Morphology of as-imprinted ~75-nm S50/M50 film (at 180 C for 30 minutes) on neutral surface and size distribution of elevated domains. A, Topographic AFM image of pattern after NIL; B, Gaussian-like apparent area distribution of elevated domains Topographic AFM images of surface and interface morphologies of patterned ~75-nm S50/M50 film (210 C for 30 minutes). A, as-imprinted morphology; B, morphology of PMMA after PS was selectively dissolved by CLE; C, morphology of PS after PMMA was selectively removed. All AFM images are 10 μm 10 μm Topographic AFM images and cross-sectional profiles of interfacial morphologies of patterned ~300-nm S50/M50 films after PS was selectively dissolved by CLE. A, B and C are topographic AFM images of films imprinted for 30 minutes at 150 C, 180 C and 210 C, respectively. All AFM images are 10 μm 10 μm. D is the cross-sectional profiles corresponding to the marks drawn on A, B and C Topographic AFM images and cross-sectional profiles of interfacial morphologies of patterned ~300-nm S50/M50 films after PMMA was selectively dissolved by AA. A is the topographic AFM image of films imprinted for 30 minutes at 210 C. B is the cross-sectional profiles corresponding to the marks drawn on A A comparison of morphologies between pressurized ~150-nm S50/M50 film and patterned S50/M50 film. A and B are extracted from Figures 5.13 and 5.14, and they are topographic AFM images for films pressurized for 30 minutes at 180 C and 210 C, respectively. C and D are extracted from Figure 6.6, and they are topographic AFM images for

20 xix films imprinted for 30 minutes at 180 C and 210 C, respectively. All AFM images are 10 μm 10 μm Schematic of major flow directions during coarsening under imprinting condition A transition of encapsulated PS domains from blocks to continuous threads as PS concentration increases from 30% to 70%. AMF images are extracted from figures shown before and are morphologies of AA-etched films imprinted at 210 C. All AFM images are 10 μm 10 μm A schematic illustration of confined reaction-induced phase separation process. PAA pattern was first made by TE-NIL. Then while NOA/5CB mixture was confined within PAA trenches mostly, SF-NIL was carried out to induce the phase separation of NOA and 5CB. Finally, after selectively dissolving PAA and 5CB, the NOA-5CB interface was revealed and characterized Phase-separated morphologies of NOA/5CB mixtures confined by 833 nm-width pattern. A, B, and C are topographic AFM images of NOA- 5CB interfacial morphologies after 5CB was selectively removed for NOA87/5CB13, NOA50/5CB50 and NOA30/5CB70 films, respectively. All AFM images are 5-µm wide. D is interfacial cross-section profiles for different compositions Schematic cross-sectional profile of the PSCOF device structure. 4, cross-linked NOA; 5, liquid crystal; 6, UV source Topographical AFM images of surface morphology (A) for pure NOA pattern and NOA-5CB interfacial morphologies (B and C) after selective removal of 5CB for NOA85/5CB15 and NOA50/5CB50 films. All films were confined with 150 nm-width pattern and all AFM images are of size 1 µm wide A schematic illustration of fabricaiton of ultra-thin C-PSCOF (confined phase-separated composite film) cell. With a PAA pattern, pure NOA pattern was replicated by the same method as in Figure 7.1. Then

21 xx NOA pattern was used as substrate to conduct SF-NIL to induce the confined phase seperation of NOA65 and 5CB forming C-PSCOF cell Ion effect in PDLC cell. A, Schematic of a bilayer of ion charges that produce an internal electric field opposite to the external field; B, Ion concentration dependence of ion effect on dynamics of the optical response of PDLC film Experimental setup for the measurement of electro-optic response of C-PSCOF device A typical electro-optic response of C-PSCOF cell corresponding to an applied electric field (AC voltage of 10 V, 100 Hz) along with schematics of alignment directions of the liquid crystal molecules at different stages (A), and a schematic illustration of the ideal alignment states of liquid crystal molecules in voltage-on state and voltage-off state (B) Voltage dependence of electro-optic responses of C-PSCOF cell. A, electro-optic responses of C-PSCOF cell to different applied voltages ranging from 6V to 10V. B, plot of relationship between maximum/minimum output value and the input voltage Frequency dependence of electro-optic responses of C-PSCOF cell. A, a comparison of electro-optic responses under 200-Hz 10-V applied voltage with 100-Hz 10-V applied voltage. B, plot of relationship between maximum/minimum output value and input frequency Schematic summary of the thesis including thin films on chemically homogenous substrates (SiOx surface, neutral surface), on chemical pattern, within two parallel rigid substrates, and under traditional TE-NIL/SF-NIL conditions

22 1 CHAPTER I INTRODUCTION Nanoimprint lithography (NIL) is a low-cost, accessible lithographic technique with sub-10 nm patterning resolution. 1-5 There are two types of NIL processes, depending on the mechanism of the pattern replication: thermal embossing-nil (TE-NIL), and step-and-flash-nil (SF-NIL). TE-NIL utilizes the viscoelastic deformation of a polymer resist to faithfully replicate the features on a rigid mold pressed onto the polymer film. In contrast, the SF-NIL relies on the UVinduced crosslinking of reactive monomers, after they have filled the mold cavities, to replicate the mold features. TE-NIL offers unique capability and potential to directly pattern a range of functional materials. 3,6-12 Up to now, neat polymers have been widely used in NIL fabrication, where their viscoelastic properties dictate both the fidelity of the pattern replication and the stress state of the obtained structures However, neat polymers are intrinsically limited by their chemical and physical characteristics. Mixing or blending polymers has been a traditional processing strategy to improve the properties of the neat polymer materials in

23 2 bulk Polymer blending technology has been an attractive area in polymer science in the past several decades and its products are ubiquitous in our daily life. 20 Multifunctional polymer nanostructures with a combination of topographic features and diverse chemical functionalities are critical to many emerging technologies, and it remains unclear whether it is a viable approach to fabricate chemically heterogeneous nanostructures using NIL. In addition to the implication to the practical applications, the morphological evolution of multicomponent polymers under the NIL process is a fundamentally intriguing problem. At the length scale of the features commonly involved in NIL, surface and interfacial interactions between the constituent polymers and the confining environments become increasingly important In addition, the physical confinement of the cavity walls is stronger than that in the planar thin films, which will significantly influence the directionality and perhaps mobility of the phase separation and morphological evolution such as domain breakup or coarsening. 25,26 Under these motivations, this thesis work aims to understand the morphological evolutions of blends under the NIL conditions, in particular, the interplay between surface energy, surface topography and the processing parameters (pressure, temperature and time) of the NIL. This thesis is organized in the following way: This chapter, Chapter 1, gives a brief introduction to my thesis work and outlines other parts of the thesis.

24 3 Chapter 2 will briefly review the fundamental concepts and theories of polymer phase separation. Chapter 3 will describe the experimental setup, sample preparation and characterization methods used in this thesis work. Especially, the main processing method, NIL, will be introduced in this chapter. In Chapters 4, 5 and 6, we mainly investigated the phase evolution of a classical polymer blend pair of PS (polystyrene)/pmma (polymethyl methacrylate), whose morphological behaviors have been extensively studied in both bulk and thin films. 18,19,27-31 We first characterized the morphological evolutions in planar thin PS/PMMA films with varying compositions, on both preferential and neutral surfaces in Chapter 4. In addition, we also examined the effect of chemically patterned substrate on the morphology of as-cast PS/P2VP (poly2-vinylpyrrolidone) planar films. Chapter 5 includes the study of phase evolution of PS/PMMA under the confinement of two parallel rigid walls, which is realized by using a fresh piece of wafer as template during TE-NIL process. We particularly examined the influences of the imprinting temperature and film thickness on the phase morphology. We also detailed the interfacial morphology between PS and PMMA as well as polymers and confining environment for a slowly evolving system. In Chapter 6, we explored the coarsening process (or the intermediate to later stages of the phase separation) of PS/PMMA blends under two-dimensional physical confinement provided by NIL. The influences of the imprinting temperature, blend

25 4 composition, film thickness as well as substrate surface energy on the morphological evolutions were systematically investigated. In Chapter 7, we investigated the physical confinement effect on the phase evolution of initially mixed binary composite NOA65/5CB (4-cyano-4'- pentylbiphenyl) using SF- NIL. The pattern fabrication procedure was essentially a confined reaction-induced phase separation (RIPS) process that is a common way of demixing polymer-liquid crystal mixtures. 32 We discovered that under such confinement unique embedded liquid crystal layers can form at conditions that normally lead to polymer-dispersed liquid crystal (PDLC). Finally, I will summarize all the thesis work in Chapter 8.

26 5 CHAPTER II THEORY 2.1 Thermodynamics of polymer blends Flory-Huggins theory is a classical model for describing the mixing free energy of binary polymer mixture. 33,34 The mixing of polymers is considered from a thermodynamics perspective in the theory. 35 Specifically, the mixing process can be described in terms of Gibbs free energy change upon mixing a blend, where Hm is the enthalpy change and Sm is the entropy change on mixing. Flory- Huggins theory provides the Hm term and Sm term in the above equation, using a mean-field lattice model. There are several forms of the formula expression commonly used in literature and some terms are interchangeable. An equation with molar base is given in the following form: or in terms of degree of polymerization

27 6 where is volume fraction, is molar volume, is degree of polymerization of polymer component i, i = 1 or 2, is the interacting segment volume or reference volume, is Flory-Huggins interaction parameter, V is the total volume, is molar number of the reference unit, k is Boltzmann s constant and R is gas constant. 20 The phase diagram of binary polymer mixture can be estimated based on Flory-Huggins theory. 17 In a phase diagram, temperatures are plotted against one of the polymer component. If N1 is equal to N2, it is also common to plot N as a function of, where N = N1 = N2. A schematic phase diagram is shown in Figure 2.1. Figure 2.1 Schematic phase diagram for a binary polymer mixture. 20

28 7 The equilibrium condition for the binary system is where is the volume fraction or concentration of constituent polymer1. The curve representing the above equation is binodal curve or coexistence curve which divides the regions on phase diagram into two regions, one-phase region and two-phase region. Inside the two-phase region, it can be further divided into two sub-regions, metastable region and unstable region. The criterion deciding the two sub-regions is spinodal curve or stability curve, which corresponds to the equation The region between spinodal curve and binodal curve is in metastable state and the phase separation mechanism governing the phase evolution process in this region is nucleation and growth. The region inside the spinodal curve, where < 0, is the unstable region. Spinodal decomposition mechanism dominates the phase separation process in this region. Typical morphologies via nucleation and growth mechanism and spinodal decomposition mechanism are shown in Figure 2.2. Binodal curve and spinodal curve intersect at the critical point, which is related to the equation The critical point is typically associated with critical temperature and the corresponding critical volume fraction. In general, there are two types of critical

29 8 Figure 2.2 A schematic illustration of phase-separated structures initialized by nucleation and growth (left) and spinodal decomposition (right). temperatures, upper critical solution temperature (UCST) and lower critical solution temperature (LCST). UCST is the critical temperature above which two components can be mixed at any composition. The polymer system that exhibits UCST is called UCST system. Similar concept applies to LCST. Liquid-liquid and polymer-solvent mixtures usually are UCST systems. This is because one of the components in the system has relatively small molecular weight. As temperature increases, Sm increases correspondingly driving Gm to become smaller and eventually negative. However, this argument does not apply to polymer-polymer mixtures, because for high molecular weight materials, Sm is small. Other factors will be dominant and the situation could be complicated. In general, polymer blends show LCST, i.e., PS(polystyrene)/PVME(polyvinyl methyl ether), PMMA(polymethyl methacrylate)/san(styrene-acrylonitrile), PMMA(polymethyl methacrylate)/peo (polyethylene oxide), PMMA/PVC(polyvinyl chloride)

30 Temperature ( C) 9 binodal curve PS volume fraction (%) Figure 2.3 Phase diagram of PS190K/PMMA94K based on compressible regular solution model. In this thesis, I primarily examined PS/PMMA which is a UCST system. The main reason for selecting this polymer pair is that the bulk behavior of PS/PMMA is extensively studied in literature. It is easier to compare the confined cases in my studies with other references. Also it is much convenient to characterize polymer morphologies at room temperature by carrying out atomic force microscopy (AFM) measurement. A phase diagram of PS/PMMA for particular molecular weights used in this study, PS190K/94K, is shown in Figure 2.3. The phase diagram is calculated by a modified Flory-Huggins model, namely, compressible regular solution model. 40,41 Within the annealing temperature range in this study, i.e., 25 C C, PS/PMMA stays in two-phase region. What we observed is essentially the coarsening process or intermediate/late stage of the phase separation. PS/PMMA blends with other molecular weights used in this study also show similar phase diagram and annealing occurred within two-phase region. Since between PS and

31 10 PMMA is /T(K), which indicates T > 2, PS and PMMA are a pair of strongly segregating polymers. Accordingly, the interfacial regions between the separated PS/PMMA domains are very sharp, i.e., the interfacial width is typically a few segments (several nanometers). 19,42 Although it is possible to quench an LCST system rapidly to room temperature to freeze the morphologies developed at high temperature in some other experimental setups, the cooling process in our NIL system is slow and it is very likely that morphology evolution such as remixing occurs significantly during the cooling process and the true morphology developed at targeted phase separation temperature might not be accurately revealed. Another difficulty of using LCST system in this study is that most of the widely studied LCST systems have high critical temperatures and need long annealing time, i.e., several days, to develop morphologies in intermediate or late stage of the phase evolution. 37,43 The NIL system used in this study has a temperature limit of 250 C and a typical imprint cycle runs within one hour. Therefore, the confining method used in this study is not suitable for examining LCST system. I also used another polymer pair PS(polystyrene)/P2VP(poly2- vinylpyrrolidone) to examine the impact of substrate surface energy on the structure formation of spun-cast film. The ideal feature of the system is to be able to phase separate at room temperature, and this UCST system is desirable.

32 Wetting on surface Surface and interfacial interactions play important roles in phase separation process. It is very common for a polymer blend pair to have one of the two components preferentially wet the substrate. Thus, wetting phenomenon is crucial in determining the phase evolution process, especially for thin film cases. Wetting is closely related to surface and interfacial energies. The wettability of a liquid on a solid substrate is typically determined by the balance of surface and interfacial energies. Young s equation describes the equilibrium state of a liquid droplet on substrate, SG = SL + LG cos c where SG is the interfacial tension between solid substrate and ambient gas, SL is the interfacial tension between solid substrate and liquid phase, LG is the interfacial tension between liquid phase and ambient gas, and c is the liquid droplet contact angle on the substrate (Figure 2.4). 44,45 Figure 2.4 Contact angle and forces associated with a liquid droplet on solid substrate. Among all the parameters in Young s equation, contact angle is the easiest to be measured by experiments. It is therefore widely used as an indicator of the wetting tendency of a liquid on a solid substrate. c = 0 corresponds to the complete

33 12 wetting situation, where the liquid/solid interaction is very strong. When 0 < c < 90, the low contact angle indicates liquid would spread on the substrate upon contact. A wettable surface for water is typically termed as hydrophilic surface. When 90 c < 180, wetting is not favorable. The liquid/solid interaction is weak and the liquid tends to contract to minimize the overall surface and interfacial energies. A non-wettable surface for water is a hydrophobic surface. The parameter that quantifies the tendency of a liquid spreading on a given substrate is spreading coefficient S, which is defined as S = SG ( SL + LG ). For example, silicon oxide is a non-wettable surface for PS and SPS is less than A widely observed wetting phenomenon in polymer blend study is that PMMA has preferential wetting tendency on SiOx surface. The polymer pair that we mainly investigated in this study is PS/PMMA. PMMA wetting layer is thoroughly discussed in the following chapters. The formation of PMMA wetting layer dictates the coarsening process of phase evolution and therefore the final morphology. Furthermore, we manipulated the substrate surface energy to change the wettability of PMMA on the substrate by grafting a random copolymer layer on substrate surface. The treated substrate surface could eliminate the preferential wetting tendency of PMMA. It behaves as if it likes PMMA and PS equally and is neutral to both of the components in terms of wettability. In such case, the substrate is called a neutral surface. This type of random copolymer coating method changes the substrate surface energy physically as SiOx surface and copolymer layer are bonded by van der waals forces. Alternatively, substrate surface energy

34 13 could be altered chemically. For example, we added functional groups such as COOH or CH3 onto gold surface. The resulted gold substrate has high surface energy (hydrophilic) and low surface energy (hydrophobic), respectively. We next consider wetting in confined geometry. The simplest case is blend between two parallel walls. A free energy model predicting the final equilibrium configuration of two phases in confined geometry is proposed by Tanaka. 47 Based on the preferential wetting tendency of one phase, two possible structures could be resulted. The first one is layered structure or complete wetting configuration, and the second one is capsule structure of partial wetting configuration as shown in Figure 2.5. The free energy per unit area parallel to lateral direction for complete wetting is Fcw = 2 ( -wall + - ), and the corresponding free energy for partial wetting case is Fpw = 2 -wall + 2 -wall + - f(d), where, are volume fractions for phase and phase, -wall and -wall are interfacial tensions between phase and the wall, and phase and the wall, - is the interfacial tension between phase and phase, f(d) is the total area between phase and phase. The transition state could be reached when Fcw - Fpw = 2 - (1 - ( -wall - -wall)/ - ) = 0 When < - /( -wall - -wall), partial wetting is energetically preferred and capsule structure would be resulted. When > - /( -wall - -wall), the wettable component is dominant and wetting phenomenon occurs more significantly. As a result, layered structure forms. This model is successful in predicting the morphology of binary phases under weak confinement as surface energy and interfacial energies are keys

35 14 in determining the final equilibrium configuration. However, this model puts no restrictions on vertical direction and could be potentially insufficient to interpret phase behavior under strong confinement situation. Figure 2.5 Final possible equilibrium structure for partial wetting (a) and complete wetting (b) of wettable phase Cahn-Hilliard theory Cahn-Hilliard equation is a mathematical physics model initialized by Cahn and Hilliard for phase separation process in binary alloys. 48 It is not limited to alloys and is applicable to other phase separation processes related spinodal decomposition, such as binary polymer systems. 49 The equation in dimensionless form could be written as where c is the normalized concentration of the phase with c = ± 1 indicating pure component, D is diffusion coefficient with a unit of length 2 /time, is a constant to balance the units in the equation. 50 is the chemical potential. is a non-linear term, and is typically set as c 3 c.

36 15 Numerical methods used to solve Cahn-Hilliard equation are extensively studied The complexity of Cahn-Hilliard equation arises from its high order, fourth order, and the non-linear term (u). These issues are typically addressed by spectral method and finite difference method. Spectral method is a powerful solution technique for ordinary and partial differential equations. 54 The best known example is Fourier transform. One of the definitions of Fourier transform pair is given by 55 A significant feature of Fourier transform is that derivations could be calculated simply by multiplication operations under frequency domain Computationally, for a finite-sized domain, instead of expressing F(k) in terms of integral, it is typically expanded in a Fourier series

37 16 The following shows a simplified numerical solution to Cahn-Hilliard equation and its simulation results regarding the cases of our interests. In this study, we mainly concern phase separation with confinement such as between two parallel plates and within grating channels of nanoimprinting template. A relatively simple rectangular domain would be appropriate. The domain could be extended to a three dimensional lattice. However, due to the limitation of boundary condition in spectral method, which has to be periodic boundary condition, domain shape cannot be extended to arbitrary shape. Numerically is calculated by finite difference method. Similar to the continuous counterparty, there should exist a frequency independent parameter such that where is the distance between two neighboring elements, is Laplace operator to be determined as follows. The fast Fourier transform is two dimensional, which means

38 17 The rectangular domain is treated as an M N matrix. The concentration of one element in such matrix is denoted by c(m,n), where m and n are row index and column index, respectively. is calculated numerically as Therefore ( ) ( ) ( ) ( ) In practice, one wants to apply no-flux boundary condition to the domain of interest. Cosine discrete transform is more suitable as it satisfies Nuemann boundary condition automatically. According to Euler s formula, the Laplace operator derived above from fast Fourier transform should be the same for cosine discrete transform. With Laplace operator, Cahn-Hilliard equation then can be converted to Discretizing time, i.e., using semi-implicit Euler method, we have

39 18 where k indicates the k th step in the iteration. However, in real application, more advanced technique might be needed to discretize the time and space in order to achieve better stability. The dynamics of the process is therefore from where c k+1 could be obtained from inverse FFT or inverse cosine discrete transform correspondingly. Figure 2.6 shows a simulation result based on the implementation of the above method. A and B are the simulated results for phase separation in bulk state. The matrix size is set as Figure 2.6A is a typical morphology of spinodal decomposition at early stage of the phase separation. Phases with equilibrium concentrations have not completely formed at this stage as indicated by many transition regions. Figure 2.6B shows phase-separated morphology at an intermediate stage. Both the domain size and the correlation length increase as the coarsening process progresses. For confined phase separation, a simple idea is to reduce the size of one dimension in the simulation box. For example, the height is reduced to 1/32 of the width of the simulation region in Figure 2.6C. In the absence of preferential wetting of the boundary walls, the confined morphology at intermediate stage displays a type of alternating blocks structure. This structure is likely a metastable and further increasing the coarsening time in the simulation will not change it much. Albano et al., reported similar phase-separated structures

40 19 Figure 2.6 An example of simulation result based on Cahn-Hilliard equation. A, the morphology developed at an early stage of phase separation. B, the morphology after significant coarsening has occurred. C, the simulated morphologies within high-aspect ratio domains at different times. at different stages as we have shown in Figure 2.6C based on Monte Carlo simulation for Ising model. 56,57 They showed that there exists a stable correlation length at the thermal equilibrium stage which increases with diffusion coefficient and the interfacial tension between the two components. A similar 3D simulation result is shown in Figure 2.7. Figure 2.7 3D simulation results based on Cahn-Hilliard equation in bulk and confinement situations for two different blend compositions.

41 20 In practice, phase separation process is often complicated by the introduction of interfacial interactions between blend components and the confinement environment, particularly in a thin-film-like environment. Cahn-Hilliard model does not take wetting phenomenon into account and the above model assumes the confinement environment is neutral to both of the phases. Also spectral method has its limitations on boundary conditions. Therefore, although Cahn-Hilliard model describes well for phase separation in bulk, for confined cases, simulation results are rarely used for directly comparing with experimental reports. In reality, there are relatively few comprehensive theoretical investigations of phase separation in confined geometry due to the complexity risen from the interfacial interactions. 58 Some studies amended boundaries at surfaces or interfaces, i.e., augmenting a surface energy term in Cahn-Hilliard equation as shown by Kielhorn et al. 53 They introduced interaction strength to represent the interfacial interaction between one component of the blend and the substrate. When is 0, it indicates that the substrate surface is neutral to both polymer components. When is positive or negative, one component has preferential wetting tendency on the substrate surface. We applied the same method to demonstrate simulated phase-separated structures under asymmetric boundary condition and symmetric boundary condition with equal to 0.5. Figure 2.8 shows the simulation result with asymmetric boundary condition. It is somehow similar to the case of PS/PMMA film on silicon oxide surface. If the blend is in a bulk state, only the vicinity of the

42 21 Figure 2.8 Simulation results for phase separation with asymmetric boundary condition. The bottom of the simulated region has interaction strength = 0.5. A, simulated bulk-like phase-separated structure. B, simulated thin film cases with aspect ratio of 8 and different coarsening times. substrate surface will be affected (Figure 2.8A). When the blend is in thin film cases, the influence of the substrate surface energy is more obvious. Figure 2.8B shows the coarsening process in stripes ( ). As the phases evolve, the wetting layer grows, and more and more wettable phase migrates to the bottom layer. Eventually, bilayer-like structure is resulted. Note that the perfect bilayer equilibrium structure is not reached with the set of parameters used in Figure 2.8B due to stability issue (these parameter are good for illustrating the intermediate stages though), however it could be easily achieved by tuning the interaction strength, the aspect ratio or the diffusion coefficient. In practice, perfect bilayer structure might be never reached neither due to the extremely slow diffusion at the later stage. Also for PS/PMMA, PS tends to dewet on PMMA and to minimize the

43 22 total surface energy at the polymer-air interface, therefore surface-relief structure is typically observed. Furthermore, the substrate surface energy could also be chemically heterogeneous. Chemical pattern, enabled by lithographic techniques, is of particular interests in terms of its potential to guide the phase separation process. Figure 2.9 and Figure 2.10 demonstrate the chemical pattern impact on phase separation in simulated cases using the same method above. The chemical pattern periodicity is set to 32. Half of the periodicity has interaction strength = 0.5 and the other half has = Similar to preferential substrate, chemical pattern also only has influence at the vicinity of the substrate surface. Figure 2.9A shows little difference from the regular bulk phase separation regardless the imposed alternating surface energy at the bottom. As the film thickness decreases, the phase-separated structure shows strong guidance Figure 2.9 Simulation results for phase separation on chemical pattern, which has alternating = 0.5 and = -0.5 with periodicity of 32. A, simulated bulk-like phaseseparated structure. B, simulated thin film cases with different film thicknesses.

44 23 under the chemical pattern (Figure 2.9B). It is widely accepted that to realize a perfect pattern transfer from the chemical pattern to the blend film, the intrinsic correlation length in phase-separated structure needs to be commensurable with the chemical pattern periodicity within certain range of thicknesses (roughly half of the periodicity). 53,59 The simulation result in Figure 2.9 is consistent with literature. Figure 2.10 Simulated phase-separated morphologies (top view) with different interaction strength values. is 0, 0.05 and 0.5 in A, B and C, respectively. All simulated domains are and time t is Chemical pattern periodicity is 32. What we also found is that even when the intrinsic correlation length is much larger than the chemical pattern size, the chemical pattern effect still exists (Figure 2.10B) and it will result in a hierarchical structure revealing both the intrinsic correlated structure (Figure 2.10A) and the chemical pattern information. Figure 2.10B simulates that a relatively thick film (reflected in a reduced interaction strength = 0.05) with large molecular weights (reflected in long simulation time t = 2500, compared with t = 500 in Figure 2.9) phase separates on a small-sized chemical pattern. This type of structure is indeed observed in experiments in our study as will be shown in Chapter 5. In Figure 2.10C, the spinodal decomposition

45 24 structure is also arrested with the same interaction strength as shown in Figure 2.9, regardless of the larger intrinsic correlation length. This is caused by the small stepping in the simulation, which indicates small quenching depth and that the domain grows slowly. When the spinodal decomposition wave length is comparable with the chemical pattern size during the simulation, the structure is well guided and stable, and further iteration will not change it. For symmetric boundary condition case, stratified multi-layer (Figure 2.11A) formed near the boundary makes it hard to observe the gradual transition in confined cases (Figure 2.11B), which is very clear in asymmetric case as shown in Figure 2.8B. The stratified layer is also reported by other researchers simulation Figure 2.11 Simulation results for phase separation with symmetric boundary condition. Both top and bottom of the simulated region have interaction strength = 0.5. A, simulated bulk-like phase-separated structure. B, simulated thin film cases with different coarsening times.

46 25 work. It is not real in practice, but it reflects the composition waves occurred at the boundary. 47 Regardless, the equilibrium state simulated is consistent with the stable complete wetting configuration discussed in previous section. The nonwettable phase forms continuous domain in Figure 2.11B. This is slightly different from Binder s model shown in Figure This is likely caused by large interaction strength (0.5) we used in our simulation. Figure 2.12 Schematic configurations of phase-separated films for different types of boundary conditions, as predicted by Binder s model. 58,60 A, equilibrium state under non-symmetric boundary conditions; B and C, equilibrium and non-equilibrium states under symmetric boundary conditions. B shows the situation when temperature T is lower than wetting transition temperature Tw, or wetting is less dominant. C shows the situation when T is larger than Tw and wetting of the wettable component plays an important role in the formation of phase separated structure. Liu provided a phenomenological wetting phase diagram (no spinodal decomposition considered) for confined binary liquid mixture in cylindrical pore based on Monte Carlo study (Figure 2.13). 61 The diagram is plotted base on

47 26 Figure 2.13 Phenomenological phase diagram for binary liquid mixture under the confinement of a cylindrical pore. The formation of the three configurations of plug, capsule and tube depends on the quenching depth t = ( )/ and, where is molecule length and is the pore diameter. 61 quenching depth t and the diameter of the cylindrical pore (presented as, where is molecule length). The diagram includes both of the wetting cases discussed above (Figures 2.11B and 2.12C) and accommodates the seemingly discrepancy. With a fixed, when the wetting tendency of one component is very strong, which means t is small since phase separation kinetics is relatively slow compared with fast wetting kinetics, the non-wettable phase has tube structure. This might be comparable to the tri-layer case in Figure 2.11B in the sense that the non-wettable phase is continuous, although the confinement is more severe in cylindrical pore. The continuous thread within strong confinement is thermally stable and can maintain the continuous state without breaking up into droplets. 62 When t increases, there might exist a state where the non-wettable phase is

48 27 capsules broken up from thread due to Rayleigh instability, similar to what was illustrated in Figure 2.12C. When t is sufficiently large, phase separation process dominates the wetting phenomenon. As a result, the confining environment behaves neutral-like and plugs structures (also see Figures 2.6C, 2.7 and 2.12B) are resulted.

49 28 CHAPTER III EXPERIMENTAL 3.1 Nanoimprint lithography (NIL) Nanoimprint lithography is an emerging nanopatterning technique invented by Prof. Stephen Chou and his students in ,63 It has obvious advantages of low cost, high throughput and high patterning resolution over traditional lithographic approaches. Because of its capability of patterning sub-50 nm features and rapid development since its invention, it is listed as one of the ten emerging technologies that will strongly influence the future. 3,64 International Technology Roadmap for Semiconductor (ITRS) has included NIL for the 32 and 22 nm nodes in future IC production. 65 There are many variations and implementations of NIL, among which two of them are most significant: thermal embossing NIL (TE-NIL) and step-andflash NIL (SF-NIL) Thermal embossing NIL (TE-NIL) TE-NIL is essentially a hot embossing process which had been practiced for many years even before Stephen Chou published the Science paper in The

50 29 principle of TE-NIL is very straightforward. Figure 3.1 shows a schematic process of the traditional TE-NIL. A layer of polymer resist is first coated on a substrate. The substrate is then heated from room temperature to an elevated temperature that is higher than the glass transition temperature of the polymer resist. When the targeted imprinting temperature is reached, a pressure is applied such that a mold with designed geometrical feature is pressed into the polymer resist for a certain amount of time. The polymer resist is mechanically deformed via viscoelastic deformation and filled into the cavities of the mold under the imprinting conditions. After the imprinting time runs up, without releasing the pressure, the whole system is cooled down to a temperature below the glass transition temperature of the polymer resist to freeze the replicated polymer features. Finally, the pressure is released and the mold is separated from the replica mechanically. Figure 3.1 Steps in a traditional TE-NIL process. Cast polymer film on a substrate, imprint at elevated temperature with pressure for a certain amount of time, and demold at temperature below the glass transition temperature of the polymer resist.

51 30 The mold used in TE-NIL is generally made of silicon with a native silicon oxide surface. The feature on the Si mold is fabricated by electron beam lithography or interference lithography. Often, there is a strong adhesion between the silicon oxide surface and polymer resist after the NIL process, and a self-assembled monolayer (SAM) with low surface energy is often deposited onto the mold surface for the ease of mold releasing. The polymer resist that can be used in TE-NIL includes a wide range of polymeric materials, not only thermal plastics, but also functional materials such as semiconducting polymers, block polymers, dielectric insulators. 15, Step-and-flash NIL (SF-NIL) Figure 3.2 Steps in SF-NIL process. Cast photoactive monomer resist onto a substrate, press and shine UV light to cross-link the resist, and separate the mold. SF-NIL, also called photo nanoimprint lithography, is another important type of NIL. The processing steps are similar to TE-NIL; however there are specific requirements for both the mold and the polymer resist. The polymer resist is typically a liquid monomer/crosslinker precursor which under ultra-violet (UV) light radiation will cross-link to form polymer networks. The UV light needs to pass through the imprinting mold, so the mold must be transparent materials such as glass. After UV curing, the liquid resist becomes solid and pattern transfer is completed. In comparison with the TE-NIL, SF-NIL requires only low pressure as

52 31 the viscosity of the monomer is dramatically lower than a polymer resist. However, the disadvantage of the SF-NIL lies in the sense that it cannot pattern functional materials that are not derived from photo-active monomers. Moreover, the crosslinking of the low-density monomer into polymers often induces appreciable volume shrinkage in the patterns and/or associated residual stress. 3.2 Substrate treatment As discussed in previous chapter, polymer/substrate interaction is essential in determining the phase-separated morphology. To study the influence of substrate surface energy on phase evolution, specific treatments of the substrate are needed. The substrate used in this study is standard silicon wafer which has a thin native silicon oxide layer of several nanometers thick. Such silicon oxide surface is hydrophilic and water contact angle is close to The following treatments were applied to generate a range of different surface energies Neutral surface The presence of the substrate often dominates the morphological evolution of the blend because the blend components are likely to have different wetting tendencies on the substrate This preferential wetting behavior can largely be suppressed by the addition of an organic or inorganic layer on the substrate to create a non-preferential surface. This type of non-preferential surface, often referred to as the neutral surface, has a surface energy lying between those of the blend components and is thus equally wettable by each component. The neutrality of such surfaces have been demonstrated in their ability to control the orientation of

53 32 block copolymer nanostructures in thin films, thereby enabling lamellar and cylindrical domains oriented perpendicular rather than parallel to the substrate For a blend film of PS/PMMA, a neutral-like surface can be created with a layer of random copolymer of styrene and methylmethacrylate, self-assembled monolayers, an organosilane with controlled surface coverage, or a mixture of PS and PMMA brushes. 78 Detailed procedure for preparing neutral surface to PS and PMMA is addressed in Chapter Fabrication of chemical pattern by NIL Figure 3.3 Schematic drawing of fabrication process of chemical pattern by NIL. Another widely used substrate surface treatment method is deposition of a metal layer on the substrate with subsequent SAM grafting. SAM layer with varying surface energy can be grafted on a substrate either homogeneously, or in a periodic fashion. A fabrication process of chemically heterogeneous substrate surface is shown in Figure 3.3. Two types of SAMs, a hydrophobic 1-octadecanethiol (ODT, Sigma-Aldrich) and a hydrophilic 1-mercaptoundecanoic acid (MUDA, Sigma- Aldrich), were grown onto gold surface in an alternating fashion. Specifically, a 50

54 33 nm-thick gold layer was first deposited onto the Si substrate by thermal evaporation. To enhance the adhesion between the gold layer and the silicon substrate, a 5 nm-thick chromium layer was deposited onto silicon wafer prior to the gold deposition. Subsequently, a PS grating pattern was fabricated by traditional TE-NIL on this gold-coated substrate. The patterned PS had a structure of a periodicity of 833 nm, line-space ratio of 1:1 and a line depth ~190 nm. Reactive-ion etching (RIE) was employed to etch away the PS residual layer such that the gold surface beneath the pattern trench was exposed to the air. The substrate was subsequently immersed in an ODT solution with a concentration of 0.05 wt. % for 3 min to deposit an ODT SAM onto the exposed gold surface. The excess ungrafted ODT molecules were washed away by ethanol rinsing and the first SAM deposition was completed. To deposit the MUDA, the remaining PS pattern was then removed by toluene dissolution. Surprisingly, the PS at the vicinity of the gold surface showed strong resistance to the toluene dissolution. In particular, after the toluene treatment, there typically existed a ~20-nm leftover PS pattern. The nature of this grafted or absorbed PS layer is currently unclear. Similarly, we observed reduction in PMMA solubility in toluene when we tried to use PMMA pattern instead of PS for the fabrication process. But this problem was solved by conducting a more drastic dissolution in a sonicator with toluene as medium for 30 min. After that a hydrophilic SAM was deposited in a similar way as the first ODT. The roughness of

55 34 the chemical pattern surface is around 2 nm which is close to the length of a single molecule of the SAM (Figure 3.4). Figure 3.4 AFM cross-section profile of the chemical pattern surface. 3.3 Characterization of polymer blend morphology In this study, we mainly characterize the surface and interface morphologies of polymer blend film by optical microscopy (OM) and atomic force microscopy (AFM) Optical microscopy (OM) Optical microscopy is a powerful tool to characterize the surface topographic features from several millimeters down to ~600 nm. It characterizes the film morphology, particularly in the lateral dimension, with sufficient field view to capture the morphological evolution over annealing. In a blend film, polymer phases typically contain topographic contrast reflecting the difference in domain heights. For example, when a PS/PMMA film was spun-cast from a toluene solution onto a silicon oxide substrate, isolated PMMA

56 35 domains tend to protrude from the continuous PS matrix. This is caused by the difference in the evaporation rate of toluene in PS and PMMA. However, under nanoimprinting condition, the intrinsic topographic difference might be eliminated completely due to the compression process. Fortunately, two constituent polymer phases in a blend film generally have different refractive indexes, i.e., for PS/PMMA, PS has a refractive index of while PMMA has a refractive index of The difference in refractive index then provides the optical contrast and thus reveals different phases even though the surface is completely featureless. Optical microscopy is thus convenient in such case provided that the features are within the resolution of the microscopy. Phase-separated structures, such as those evolved from spinodal decomposition, often show clear spatial correlation. One way of measuring correlation length directly is through X-ray or light scattering techniques, or fast Fourier transform processing on the OM images. By conducting FFT on micrograph, the constructive interference resulted from the structural correlation could be easily observed from resulting FFT image. Quantitatively, correlation length for such a structure can be calculated, provided that the length scale in the FFT image can be calibrated with known structural features such as periodic grating Atomic force microscopy (AFM) AFM, also called scanning force microscopy (SFM), is a high-resolution scanning probe microscopy with a resolution potentially less than one nanometer if ultrasharp tip is provided. It uses a cantilever with a sharp tip to scan the surface of

57 36 a sample. When the cantilever is brought to contact with the surface, forces between the cantilever and the sample cause a deflection of the cantilever, which is typically measured with a laser spot reflected from the back side of the cantilever. 79 The interactions as a function of the position of the cantilever on the sample surface are recorded in an image line by line during the scanning. Depending on the types of the mode and the cantilever used in the AFM operation, additional properties might be simultaneously obtained. 80 In polymer blend film study, tapping mode is the most commonly used approach. Besides the topographic feature, phase contrast could be also easily revealed by the difference of mechanical properties of polymers or by selectively removing one of the two components

58 37 CHAPTER IV PHASE EVOLUTION OF BINARY POLYMER THIN FILMS WITH FREE SURFACES 4.1 Introduction A polymer liquid in contact with other incompatible materials undergoes a variety of morphological changes as it proceeds towards its equilibrium state This evolution in morphology is manifested in both the dewetting of a thin polymer film from a non-wettable surface and the de-mixing of a polymer blend. 85,86 The latter system has been investigated for decades both because the transient morphologies are important for practical applications and the slow kinetics are convenient for fundamental investigations of the phase separation process. 86,89 Thin polymer blend films have been utilized to fabricate a variety of nano- to mesoscale structures that have a range of potential applications The ability to tune, with respect to size and structure, the phase-separated morphologies in blends is thus highly desirable, but it can be particularly challenging in thin films to achieve nanoscale structures

59 38 For a classical blend of PS/PMMA on a preferential surface such as silicon wafer with a native oxide layer, the morphological evolution of the blend film is dictated by the substrate wetting of PMMA. After longtime annealing, a so-called surface relief structure, featuring isolated PS droplets on a continuous PMMA wetting layer is observed for silicon wafer, glass substrate and mica. 30,50, Further, the morphologies of PS/PMMA blend films after longtime annealing were also reported on metal substrates (Au and Co), and even heterogeneous substrates. 103,104 Although the morphology of PS/PMMA films after annealing for long times (several hours to days at ~170 C) on neutral surface have been examined, the morphological paths remain unclear particularly at early and intermediate stages of the annealing. 78 In addition, the effect of the blend composition on the evolution of morphology on energetically different substrates is also unknown. Motivated by these questions, we characterized the evolution of morphology of thin PS/PMMA films with varying compositions, on both preferential and neutrallike surfaces. The morphological paths were found to be highly dependent on both the blend composition and the substrate energy, resulting in a rich set of nonequilibrium, micro- and nanoscale morphologies. Understanding and controlling these factors may therefore provide the opportunity to spontaneously create a range of spatially correlated micro- and nanostructures in thin films of polymer blends. Besides examining the effects of these chemically homogeneous substrates on phase evolution during annealing, at the end of this chapter, we also briefly

60 39 discussed the influence of chemically heterogeneous substrate or chemical pattern on the structure formation of polymer blend thin films during spin-coating process. 4.2 Samples and techniques Monodisperse PS (weight average molecular weight, Mw,PS = 48.1 kg/mol; polydispersity index, PDI = 1.01; Scientific Polymer Products Inc.) and PMMA (Mw,PMMA = 21.4 kg/mol; PDI = 1.07; Polymer Source Inc.) were used as received. The glass transition temperature (Tg) of PS and PMMA was determined to be, Tg,PS = 98 C, and Tg,PMMA = 125 C, from the second scan of differential scanning calorimetry (NETZSCH DSC 204 F1) with a scan rate of 20 C/min. Two types of surfaces on Si substrates were prepared for the study: a native SiOx layer and a random copolymer (RCP) layer. For the native SiOx surface, the Si substrates were treated in a piranha solution (a 70: 30 solution of concentrated sulfuric acid: hydrogen peroxide by volume) at 80 C for 30 min, and rinsed with DI water. Further, some of these piranha-cleaned substrates were coated with a layer of RCP (poly(s-r-mma-r-gma)) with a composition of styrene (S, 61 wt. %), methyl methacrylate (MMA, 35 wt. %), and glycidyl methyl methacrylate (GMA, 4 wt. %). 74 The RCP was synthesized by free-radical polymerization at 80 C for 72 h using an AIBN initiator and the desired ratio of monomers after purification on a basic alumina column and degassing in N2. The RCP product was recovered and purified by precipitation in methanol and subsequently dried under vacuum. The RCP was first spun-cast onto the silicon substrates from a dilute solution (0.03 wt. % in

61 40 toluene), and then crosslinked into a dense layer after 5 h annealing at 170 C in a vacuum oven. The uncrosslinked RCP chains were finally rinsed off with toluene. PS/PMMA blend films with thickness of ~50 nm were spun-cast from a toluene solution (1 wt. %) onto both surfaces described above. The as-cast films on both surfaces were then dried at 60 C for 2 h under vacuum to remove the residual toluene. The SiOx surface is preferentially wet by PMMA and not PS, while the RCP surface is either neutral to both polymers or slightly favorable to one of them depending on the blend composition. 27,73 For each surface, PS/PMMA blends with three compositions, 60/40 (the weight fractions of PS and PMMA are 0.6 and 0.4, respectively), 67/33, and 80/20, were examined. These blends are referred to as S60M40, S67M33, and S80M20 in the remaining text for convenience. All the blend films were annealed at 160 C on a homebuilt hot-stage (calibrated with 7-point standard materials from Aldrich) for different durations, and their surface morphologies were determined by atomic force microscopy (AFM, DI3100, Bruker) and optical microscopy (OM, Nikon LV150). To further distinguish the PS and PMMA phases, AFM measurements were carried out on the blend films after selectively removing the PS (or PMMA) by immersing the samples into cyclohexane (or glacial acetic acid) at ambient conditions for 2 h. 28,31 The low solubility of PMMA (or PS) in cyclohexane (or acetic acid) limits its dissolution, which keeps the PMMA domains intact. 28,31 However, slight swelling of the remaining structures during the selective dissolution is possible. 100

62 41 For examining the impact of chemical pattern on structure formation during spin-coating, PS (Mw,PS = 3 kg/mol; Scientific Polymer Products Inc.) and P2VP (Mw,P2VP = 4 kg/mol; Scientific Polymer Products Inc.) were used as received. PS/P2VP with weight ratio 1:1 was dissolved in THF with different concentrations. Then the PS/P2VP blend solution was spun-cast onto chemical pattern and the phase-separated morphologies were characterized by AFM. 4.3 Phase evolution of PS/PMMA films on SiOx surface Phase evolution of S60/M40 films Figure 4.1 shows the time-dependent morphological evolution of the S60M40 film on a SiOx surface. The as-cast film displayed dome-like domains with an average diameter of 200 nm and height of 10 nm across the film surface (Figure 4.1a). After selective removal of PS (inset of Figure 4.1a), PMMA cylinders with an average height of 36 nm were identified. In addition, AFM revealed nanoscale roughness (with an amplitude of a few nanometers and a lateral length scale of ~100 nm) at regions between the PMMA cylinders, indicating that there was a PMMA wetting layer because the SiOx surface would be significantly smoother. After selective dissolution of the PMMA domains, holes were left in the films, indicating that a majority of the PMMA cylinders were exposed to air, consistent with previous reports. 105,106 Note that all the as-cast film morphologies are kinetically trapped, non-equilibrium states that will evolve toward more stable states upon annealing. In general, the rate of the evolution depends on the film

63 42 thickness, composition, thermodynamic driving forces, and the viscosity of the constituents. Figure 4.1 Topographic AFM images (a-d) and optical micrograph (OM) (e) of S60M40 on native SiOx, after annealing at 160 C for different durations as labeled. The lower right insets in (a-e) are the topographic AFM images of the sample after selective removal of either PS (a-d) or PMMA (e). The lower left insets in (c-e) are the FFT images of the AFM image or OM (not to scale). (f) Schematic illustration of the early stages and evolution of the blend morphology. Upon annealing, the elevated PMMA domains quickly sank into the film leaving depressions at the film surface, and larger-scale height fluctuations across the film surface took place, forming valleys up to 28 nm deep (as marked in Figure 4.1b). This capillary fluctuation was highly correlated with the PMMA domain coalescence, with the coarsening occurring predominately at the valleys of the capillary wave (inset of Figure 4.1b). Such a phenomenon may be attributed to the

64 43 mass flow caused by surface fluctuations; the most significant lateral flow occurs at the valleys of the surface waves, which facilitates the coalescence of the PMMA domains. Such a phenomenon was also observed for the other films as discussed later. Figure 4.2 (a) Lateral domain correlation lengths, and (b) RMS surface roughness, of thin PS/PMMA films as a function of annealing duration, for a varying PS/PMMA composition. The amplitude of the capillary fluctuation increased over time, leading to a highly correlated ridge-like PS structure after 5 min (Figure 4.1c) that broke up into isolated PS droplets after 20 min (Figure 4.1d). Accordingly, the correlation length of the surface structures ( c) grew with the annealing time (ta, in seconds) as c [ m] ~ 0.24 ta 0.32, as shown in Figure 4.2. Such a growth rate is characteristic during the intermediate stage of phase separation and wetting of thin blend films. 72 During this period, the PMMA evolved into a continuous layer covering the entire SiOx surface (inset of Figure 4.1e). The RMS roughness of the film surface as a function of annealing time for all the systems is shown in Figure 4.2b. From 20 min up to 500 min, there was no evident change in the film morphology (and surface

65 44 roughness) featuring highly correlated PS droplets on top of the PMMA layer (FFT image shown in inset of Figure 4.1e). Such a phase inversion process shown in Figure 4.1, i.e., PS evolving from forming a continuous matrix into isolated domains, is consistent with the report of the relief structure of a PS/PMMA blend with 70 wt. % PS on a glass substrate. 101 This morphology is distinctive from the randomly-distributed polymer droplets that result from dewetting via random nucleation processes Phase evolution of S67/M33 films a) As-cast, 5 m 5 m b) 100 min, 280 m 200 m PS removed 2.5 m 2.5 m Figure 4.3 (a) Topographic AFM images of as-cast S67M33 on a native SiOx layer. The inset is the topographic AFM image of the as-cast film after the selective removal of PS. (b) Optical micrograph of the S67M33 after annealing at 160 C for 100 min. The inset is the FFT image of the optical image. Compared to the S60M40 film, the as-cast S67M33 film contained a smaller number of PMMA cylinders sitting on a thin PMMA wetting layer on the SiOx (Figure 4.3a). When annealed at the same temperature, the morphological evolution of the S67M33 was similar to that of S60M40: coalescence and downward migration of the PMMA domains, accompanied by the surface capillary wave, drove

66 45 the PS into highly correlated droplets on top of a continuous PMMA layer (Figure 4.3b). Both the correlation length and average diameter of the PS droplets in the S67M33 film were larger than those in the S60M40 film (Figure 4.1e versus 4.3b). Quantitatively, the number of the PS droplets over a 280 μm 200 μm surface area for S60M40 and S67M33 were 7200 and 3800, with average diameters of 1.4 μm and 2.3 μm, respectively. The differences can be rationalized by the observation that the coarsening of PMMA domains occurred at the valleys of the capillary fluctuation, i.e., fewer PMMA domains correspond to a longer capillary wavelength Phase evolution of S80/M20 films Interestingly, morphological evolution in the S80M20 film was clearly different from that of the S60M40 and S67M33 blends. The as-cast S80M20 film showed a bilayer-like configuration (PS on PMMA) with a rough interface (Figure 4.4a). Upon annealing, the upper PS layer dewetted from the thin underlying PMMA layer via a conventional nucleation and growth (NG) mechanism (Figure 4.4b-d). After a relatively long incubation time (20 min), holes were randomly nucleated across the PS layer (Figure 4.4b), which was followed by the formation of PS stripes due to the impingement of holes (Figure 4.4c) and the final breakup of these stripes caused by the capillary instability (Figure 4.4d). 110 All these observed morphologies are similar to those characterized previously in planar PS films dewetting from a PMMA surface. 107,

67 46 a) As-cast, 5 m 5 m b) 20min, 400 m 400 m PS removed 2.5 m 2.5 m c) 100 min d) 1000 min Figure 4.4 (a) Topographic AFM image of as-cast S80M20 film on a SiOx surface. The inset in (a) is the topographic AFM image of the as-cast film after the selective removal of PS. (b-d) Optical micrographs of S80M20 after annealing at 160 C for different tas, showing: (b) random nucleation of holes in the top PS layer, (c) growth and impingement of the holes, and (d) formation of isolated PS droplets from the breakup of the impinged PS lines. Inset in (d) is the FFT of the optical image. The sizes of these optical images are 400 m 400 m. In contrast to the blends with lower concentrations of PS, the PS droplets in the annealed S80M20 film were not spatially correlated due to random nucleation dewetting (FFT image, inset of Figure 4.4d). Note that the thickness of a PS layer, if completely segregated on top of the PMMA, would range from 30 to 40 nm for the three PS/PMMA films. At this thickness range, NG is known to be the dominant mechanism for PS to dewet from the PMMA surface. 114 For the S60M40 and

68 47 S67M33 films, however, surface capillary fluctuations associated with coarsening of the PMMA domain occurred quickly before the bi-layer configurations necessary for the standard NG mechanism could develop. However, regardless of the spatial correlations, the surface relief structure, with isolated PS droplets on top of PMMA wetting layer, can be regarded as the final or later stage morphology on SiOx surface, which is stable over more extended annealing. 4.4 Influence of neutral surface on morphology evolution The poly(s-r-mma-r-gma) random copolymer-treated substrate, as described in the experimental section, may be regarded as a neutral surface to PSb-PMMA block copolymers or PS/PMMA blends with a specific composition. 115 The RCP surface treatment provides a surface energy that is intermediate to that of PS and PMMA, with the surface energy being tunable based on the ratio of styrene and methyl methacrylate units in the random copolymer. The degree of neutrality of the RCP-treated substrate also depends, however, on the composition of the overlying block copolymer or blend films. 73,74 Therefore the RCP-treated substrate used here, with 61 wt. % PS and 39 wt. % PMMA, might not be truly neutral to all of the PS/PMMA blends that are examined. Nevertheless the relative extent of the preferential interactions between the blend domains and the RCP-treated substrate will be greatly modulated in comparison to those interactions on the SiOx surfaces. The as-cast morphology of the S60M40 and S67M33 films on the RCP surface was surprisingly similar to those on the SiOx surface (Figure 4.5a versus Figure 4.1a and Figure 4.6a versus Figure 4.2a). All samples showed PMMA

69 48 wetting layers and cylindrical domains with areal densities that decreased with a decrease in PMMA concentration. The as-cast morphology of the S80M20 blend on the RCP surface (Figure 4.7a) largely resembled that on the SiOx surface (Figure 4.4a) in that both showed a bi-layer-like morphology, but a major difference was that the S80M20 blend on RCP surfaces exhibited a rougher interface between the PS and PMMA domains. We also examined PS/PMMA films 75 nm and 150 nm in thickness and with similar compositions but different molecular weights, which further confirmed that the morphology of as-cast PS/PMMA films was nearly identical on SiOx and RCP surfaces. Some studies in the literature have reported that the as-cast morphology of PS/PMMA films on SiOx and gold surfaces was identical, whereas other studies have shown different as-cast morphologies. 103,116,117 It is therefore suggested that the morphology of the as-cast blend film is relatively independent of the substrate surface energy and highly sensitive to the spin-coating and thin film processing conditions. 28, Phase evolution of S60/M40 films Upon annealing at 160 C, PMMA domains in the S60M40 blend film started to coalesce, while PS penetrated the PMMA wetting layer, as verified by AFM characterization after removing the PS domain (Figure 4.5b). From then on, both the PS and PMMA domains were continuous throughout the thickness of the film and contacted the RCP-treated substrate and the free surface (Figure 4.5b-e). The morphology showed high spatial correlation (FFT image, inset of Figure 4.5c), and the correlation lengths between PMMA phases ( c,pmma) increased from ~400

70 49 nm at 2 min to ~1.4 m at 100 min (Figure 4.5b-f), following c,pmma [ m] ~ ta 0.32, consistent with the characteristic growth rate of intermediate stage of the morphological evolution in thin blend films (Figure 4.2). 72 Clearly, the characteristic power-law growth behavior (with an exponent ~1/3) remains independent of substrate surface energy, despite the strongly different morphology involved. a) As-cast, 5 m 5 m b) 1 min, 5 m 5 m c) 2 min, 10 m 10 m PMMA PS removed PS removed PS removed bare surface 2.5 m 2.5 m 2.5 m 2.5 m 5 m 5 m d) 5 min, 10 m 10 m e) 20 min, 10 m 10 m f) 200 min, 10 m 10 m PS PS PS removed PS removed PMMA removed 5 m 5 m 5 m 5 m 5 m 5 m Figure 4.5 Topographic (a) and phase (b-f) AFM images of S60M40 on the random copolymer-treated substrate after annealing at 160 C for different tas as labeled. The lower right insets are the corresponding topographic AFM images after the selective removal of either PS (a-e) or PMMA (f). The lower left insets in (c-e) are the FFTs of the corresponding lower right insets (not to the scale). During annealing from 2 to 200 min, the areal coverage of the PS phase on the RCP-treated substrate remained around 50-55%, as estimated from the AFM images (over a 10 m 10 m area) after selective removal of PS. Since the

71 50 coverage was lower than its volume fraction (~60% as PS and PMMA have similar densities), the vertical thickness of PS domains was thicker than that of the PMMA domains. The fact that both PS and PMMA domains in the S60M40 film form stable contacts with the RCP surface during annealing confirms that the RCP is indeed effectively non-preferential to both polymers. This observation is in stark contrast to the morphological evolution of the S60M40 blend on a SiOx surface (Figure 4.1), where the preferential wetting of PMMA led to a completely different morphological path Phase evolution of S67/M33 films The PS phase in the S67M33 blend film also quickly penetrated the PMMA wetting layer to contact the RCP surface (inset of Figure 4.6b). At the film-air surface, large-scale fluctuations in thickness started to arise (Figures 4.6c and 4.6d). Both the PS/PMMA morphology at the RCP surface (insets of Figures 4.6c and 4.6d) and at the free surface (Figures 4.6c and 4.6d) were spatially correlated with distinctive length scales. From the FFT of the AFM images after removing the PS domain (insets of Figure 4.6d), two correlation lengths were observed: λc1 ~330 nm corresponding to the morphology at the substrate surface and λc2 ~ 1.8 μm for the morphology at the free surface. At the early stages of the evolution of the phaseseparated morphology, both structures have similar power-law dependences of λc1 [μm] ~ 0.056t 0.31 and λc2 [μm] ~ 0.27t 0.33 (Figure 4.2).

72 51 a) As cast, 5 m 5 m b) 1min, 5 m 5 m c) 2min, 10 m 10 m PS removed PS removed PS removed 2.5 m 2.5 m 2.5 m 2.5 m 5 m 5 m d) 5min, 10 m 10 m e) 20min, 10 m 10 m f) 200min, 10 m 10 m λ c1 PS removed PS removed PMMA removed λ c2 Bare surface 5 m 5 m 5 m 5 m 5 m 5 m Figure 4.6 Topographic AFM images of S67M33 on the random copolymer-treated substrate after annealing at 160 C for different tas as labeled. The lower right insets are the corresponding topographic AFM images after the selective removal of either PS (a-e) or PMMA (f). The lower left insets in (c and d) are the FFT images of the corresponding lower right insets (not to the scale). As shown in Figure 4.2, the correlation length of the S67M33 morphology at the RCP interface was similar to that of the S60M40 morphology at the RCP interface, while the correlation length at the free surface was comparable to that observed in the S60M40 blend on the SiOx substrates. From the inset of Figure 4.6d, it is clear that the PMMA domain coarsening occurred at the valleys of the capillary fluctuations. After 20 min, ridge-like PS domains developed while the PMMA evolved into a continuous domain (Figure 4.6e). Small but appreciable PMMA domains remained embedded in the PS phase under the ridges of the capillary

73 52 fluctuations. Finally, the PS ridges broke up into isolated droplets after 100 min due to the capillary instability (Figure 4.6f). In addition, holes were nucleated in the PMMA-rich phase in the S67M33 film, as highlighted in Figure 4.6f, suggesting that the RCP surface is not preferentially wet by PMMA at least for ultrathin PMMA layers (10-20 nm in Figure 4.6f) Phase evolution of S80/M20 films Figure 4.7 shows the morphological evolution of S80M20 on the RCP surface. The as-cast film displayed a corrugated PS/PMMA interface with a correlation length of λc = 210 nm (Figure 4.8a). According to AFM, the amplitude of the interfacial corrugation was as high as 20 nm, while the roughness at the free surface was less than 2 nm. After annealing for 1 min, isolated PMMA domains were observed with diameters ranging from several nm to ~100 nm (Figure 4.7b). The larger PMMA domains (~100 nm) were likely to evolve from the breakup of the PMMA-rich domain in the as-cast film, since their diameter was comparable with the width of the continuous PMMA phases in Figure 4.7a. In contrast, the smaller ones (10 nm or less) likely resulted from a secondary phase separation process via a nucleation and growth mechanism from the PS-rich phase, within which the effective PMMA concentration was significantly lower than 20% and thus lies in the metastable region of the PS/PMMA phase diagram. These PMMA domains became larger through coalescence, which led to depth-through PMMA cylinders with diameters of nm that were welldispersed within the PS matrix film (Figure 4.7c). This morphology enabled the

74 53 a) As-cast, 5 m 5 m b) 1 min, 5 m 5 m c) 5 min, 5 m 5 m PS removed PS removed PS removed 2.5 m 2.5 m 2.5 m 2.5 m 2.5 m 2.5 m d) 20 min, 10 m 10 m e) 100 min, 10 m 10 m f) 200 min, 15 m 15 m PMMA PMMA PMMA removed PS removed PS removed Bare surface Bare surface 5 m 5 m 5 m 5 m 7.5 m 7.5 m g 200 min h 1000 min Figure 4.7 Topographic (a-e) and phase (b-d and f) AFM images of S80M20 on the random copolymer-treated substrate after annealing at 160 C for different tas as labeled. The lower right insets are the topographic AFM images after the selective removal of either PS (a-c, e, and f) or PMMA (d). (g) and (h) are the optical images of the film after 200 min and 1000 min, correspondingly. formation of free-standing PMMA pillars (or depth-through holes) with a height (or depth) of 50 nm, as demonstrated in the inset of Figure 4.7c (or Figure 4.7d). Both

75 54 the height of the pillars and the depth of the holes were identical to the film thickness. Note that the topographic AFM images of the films after removing PS (insets of Figures 4.7b and 4.7c) showed larger PMMA domains compared to those in the phase image (Figures 4.7b and 4.7c), which is most likely due to convolution of the probe tip with the PMMA features. In comparison, the size of the holes left after PMMA removal in the inset of Figure 4.7d appeared similar to that of the phase image (Figure 4.7d). The depth-through PMMA cylinder morphology was established within 2 min and remained stable up to 10 min. Figure 4.8 Schematics of the later stage morphology after annealing on RCP surfaces for all three blend compositions. Capillary fluctuations started to evolve on the S80M20 surface after 20 min and led to a characteristic ridge and valley structure by 100 min (Figure 4.7e). Similar to that observed in Figure 4.1 and Figure 4.6, the PMMA domain coarsened mostly at the valleys of the capillary fluctuation (Figure 4.7e), while the rest of the PMMA cylinders remained within the PS-rich domain. Prior to the complete

76 55 breakup of the PS droplets, the correlation length of the surface structure was λc [μm] ~ 0.38ta In addition, holes within the PMMA-rich phase were observed and the bare RCP surface was exposed as marked in Figure 4.7e, similar to that observed in Figure 4.6f. This resulted in the formation of intriguing partial coreshell droplets, composed of a PS-rich inner core and PMMA-rich outer shell (Figure 4.7f), that were highly correlated in space (Figures 4.7h). Figure 4.8 schematically summarizes the later stage morphology of all three blend films after annealing on RCP surfaces. 4.5 Effect of chemically patterned substrate on structure formation of spun-cast blend films In pervious sections, we discussed the influence of chemically homogeneous substrate, either preferential or neutral, on the phase evolution of PS/PMMA films. From the effect of the neutral surface on morphology evolution of PS/PMMA films, we see that the substrate surface energy could radically change the pathway of phase separation. We also noticed that the starting points of the phase evolution, which are the structures formed immediately after spin-coating, are very similar on both SiOx surface and neutral surface. Therefore, homogenous surfaces, even with different surface energy, appears to be ineffective at controlling the as-cast film morphology. To be able to achieve a more diverse non-equilibrium film morphology during the annealing, it is beneficial to be able to manipulate the as-cast film morphology. To this end, chemical pattern has been shown to have stronger impact on the phase-separated morphology in blend films, as reviewed below. During the

77 56 spin-coating process, different from neutral surface, it can effectively guide the phase formation of polymer blend films by preferential interactions between the components and periodically patterned surface chemistry. Boltau et al., pioneered the studies of the guided phase separation on chemical patterns. 119 Specifically, the phase separation process of PS/P2VP and PS/PSBr blend films would result in isotropic, disordered morphology in bulk. They created chemically patterned substrate with alternating surface energy via microcontact printing technique, and demonstrated that the traditional phase separation process could be altered to form alternating stripes under the guidance of the chemical pattern. Kielhorn and Muthukumar based on Cahn-Hilliard-Cook model augmented by a surface energy term showed that spinodal decomposition could be arrested by delicately controlling the surface potential, the characteristic size of the pattern as well as the film thickness. They found that successful transferring a 1:1 line-space pattern from the substrate to the polymer film requires that the periodicity of the pattern λ needs to be commensurable with the inherent domain scale R. 53 Raczkowska et al., further investigated the morphologies of PS/P2VP blend spun-cast on chemically patterned substrate with different compositions and a wider range of R/λ ratios (0.2 R/λ < 1.8). 59 They found that compositional commensuration is also important in determining the morphologies. The conventional chemical patterns reported in the above studies have been normally fabricated through micro-contact printing process, which has a relatively low spatial resolution (i.e., 2 µm is the smallest reported so far). Several groups

78 57 used alternative fabrication methods and created chemical pattern with sub-micronor even nano-scale features. Mead and co-workers demonstrated the direct selfassemble of ordered PS/PAA blend on sub-micron-sized chemical pattern fabricated by electron-beam lithography (EBL). 120 Ginger et al., reported the guided phase separation of PS/PMMA on chemical patterns with resolution down to 150 nm by dip-pen nanolithography (DPN). 121 Stoykovich et al., using EUV interference lithography realized directed assembly of block copolymer blends on chemical nanopatterns of periodicity between 50 to 90 nm. 122 However, the above-mentioned lithography-based chemical patterns were mainly targeting successful transferring the geometric feature on the substrate to the polymer film and therefore the R/λ ratio was designedly selected to be around 1. In addition, these methods (E-beam, DPN, and EUV lithography) are intrinsically slow. Here we examine the structure formation of as-cast polymer blend on NILbased sub-micron-sized chemical pattern. We focus on the effect of R/λ ratio, or the concentration effect on the phase-separated, as-cast film morphologies. NIL-based chemical pattern offers a smaller λ than micro-contact-printing-based chemical pattern, so we would be able to examine the situations where the chemical pattern size is much smaller than the intrinsic correlation length of the polymer blend, i.e., R >> λ. The as-cast film from a 3 wt. % PS/P2VP (50:50 in weight percent) in toluene solution had sea-island structure with PS island domains dispersed in the P2VP

79 58 matrix. The isolated PS domains had concave shape with the inner region surrounded by high edges. The inner regions were even lower than the height level Figure 4.9 A comparison of structure formation of phase-separated PS/P2VP films spun-cast from 3 wt. % solution on chemical pattern with that on ODT surface. A and B are topographic AFM images of morphology formed on ODT surface with size 10 μm 10 μm and 50 μm 50 μm, respectively. C and D are corresponding topographic AFM images of morphologies on chemical pattern. of the P2VP matrix. These were believed to be caused by the substrate surface energy and is consistent with what Budkowski et al., had reported that dispersed domains tend to have concave surface on CH3-SAM while would be convex on COOH-SAM. 123 Such substrate-dependent domain surface shape is related to the solvent evaporation speed. The evaporation speed of toluene solvent is relatively

80 59 slower on CH3-SAM than on COOH-SAM due to the energetic favorability of its methyl groups with CH3-SAM. The vapor pressure increases during the spin-coating process. When it is less than the Laplace pressure of the toluene solvent trapped in PS domains at the vicinity of the surface, PS domain surface tends to be convex. If toluene evaporation is fast, i.e., on COOH-SAM, the convex shape could be frozen as PS solidifies. In contrast, on CH3-SAM, vapor pressure could increase to be larger than the Laplace pressure, which results in concave surface of PS domains. Figure 4.10 A comparison of structure formation of phase-separated PS/P2VP films spun-cast from 0.75 wt. % solution on chemical pattern with that on ODT surface. A and B are topographic AFM images of morphologies of as-cast film and cyclohexaneetched film formed on ODT surface, respectively. C and D are corresponding topographic AFM images of morphologies on chemical pattern. All AFM images are 10 μm 10 μm.

81 60 Although the chemical pattern feature was not clearly revealed at the film surface due to the mismatch of the inherent correlation length and the pattern periodicity, it did affect the structure formation of the blend film dramatically by inverting the relative height of the two components, transforming isolated domains to continuous domains. The height contrast between the edge and the middle part of the PS domains was eliminated on chemical pattern. The inner region was neither concave nor convex since the chemical pattern has a balanced energy of between those on CH3-SAM and COOH-SAM. The PS domains became absolutely higher than P2VP. Moreover, PS domains were no longer isolated and the overall morphology had a bi-continuous-like structure. The phase boundary was influenced even more obviously and became increasingly straight. Figures 4.9C and D show the guided and aligned bi-continuous hierarchical structure. When the solution concentration is reduced to 0.75 wt. %, the intrinsic domain size on ODT surface was reduced consequently. As a result, the film surface reflected the chemical pattern mostly. PS domains were elevated and elongated on ODT as shown in Figure 4.10C, which was confirmed by the selective dissolution experiment (Figure 4.10D). Different from the 3 wt. % case, PS domains appeared more isolated and embedded in continuous P2VP matrix, which covered larger surface area. This is due to the lower solubility of P2VP than PS in THF. During spin-coating process, P2VP solidified first and covered hydrophilic surface due to its strong affinity with MUDA-SAM. As the domain correlation length (~3 µm) was still much larger than the pattern periodicity (0.83 µm), perfect accommodation of P2VP

82 61 on MUDA surface was not possible. Some of the P2VP had to solidify on ODT surface. A more systematic comparison of the concentration-dependent film morphologies on the chemical patterns is shown in Figure When the solution concentration was high, i.e., 3 wt. % (Figure 4.11A), so were the correlation length and the domain size as well as the film thickness. The chemical pattern feature was merely reflected on the phase boundaries. When the concentration went down to 1.5 wt. %, the influence of the substrate surface energy became stronger and the two Figure 4.11 A transition from spinodal decomposition-based hierarchical structure to broken lines as the solution concentration decreases. A, B, C, and D are topographic AFM images corresponding to morphologies of as-cast films from solutions with concentration 3 wt. %, 1.5 wt. %, 0.75 wt. %, and wt. %, respectively, on chemical pattern. All AFM images are 10 μm 10 μm.

83 62 components were more easily to be guided. P2VP were more mobile and could occupy more hydrophilic substrate area. As a result, the large elevated PS domains appeared broken up. From A to B, the PS domain split into sub-domains with domain width identical to half of the periodicity of the chemical pattern, or the width of each SAM stripe. The splitting direction was perpendicular to the direction of chemical pattern line and less and less PS stayed on MUDA-SAM surface. Meanwhile the split PS domains grew longer along the SAM pattern line direction. In Figure 4.11B, individual broken-up PS domains were visible, but they were still interconnected, which is the intermediate stage between the two discussed in Figure 4.9 and Figure As the concentration decreased to 0.75 wt. %, the intrinsic PS domains were further split such that the broken-up was most completely and PS domains were almost isolated. When PS domains were completely separated along the direction perpendicular to the pattern lines, ideally they should cover all the ODT area. This might happen at a concentration slightly lower than 0.75 wt. % wt. % seemed to be lower than the constrained concentration that could lead to a perfect transferring the chemical pattern to blend films. The PS domains even broke up within pattern lines. This fell below the R/λ ~ 1 range and the morphology was consistent with the literature Conclusions To summarize, we systematically investigated the evolution of micro- and nanostructured morphologies in thin films of PS/PMMA blends as a function of the

84 63 blend composition and the surface energy of the substrate. For thin films of the three PS/PMMA blends on SiOx surfaces, PMMA evolved into a continuous layer that preferentially wet the SiOx surface, while PS droplets formed on top of the PMMA layer either driven by capillary fluctuations (S60M40 and S67M33) or a random nucleation process (S80M20). Non-preferential wetting was observed for the PS/PMMA blends on substrates treated with a random copolymer with a surface energy intermediate to that of PS and PMMA. Both PS and PMMA formed contacts with the substrate and highly correlated phase structures were obtained. Capillary fluctuations, correlated with the PMMA domain coalescence, also dominated the intermediate stage of the morphological evolution by forming PS-rich ridges and PMMA-rich valleys and thus the resulting isolated PS-rich structures were spatially correlated. It was observed that both the PS and PMMA domains formed direct contact with the RCP surface during the evolution of the morphology suggesting that the surface was indeed neutral-like to both polymers. For both the S67M33 and S80M20 polymer blends, holes in the PMMA phase were observed which further confirms that the RCP surface, unlike the SiOx surface, was not a uniformly wettable or preferential surface for the PMMA component. Because the thickness of the PMMA phase on the RCP surface decreased with its concentration in the blend, more dramatic dewetting of PMMA was observed for the S80M20 film than the S67M33 film. Similarly, no holes were observed in the PMMA layer for the S60M40 blend. In addition to the 50 nm films presented here, we have also examined PS/PMMA blend

85 64 films with larger molecular weight (slower kinetics) and in thicker films (75 nm and 150 nm). The morphological behavior of these other systems are consistent with the observations for the 50 nm films, in the sense that both the PS and PMMA domains contact the RCP surfaces in the annealed films. For the thicker films that were studied, no dewetting in the PMMA phases were observed. To the best of our knowledge, this is the first systematic experimental comparison on the phase evolutions of thin blend films on non-preferential and preferential surfaces. The unique non-equilibrium morphologies that are presented, ranging from the nanoscale bicontinuous phases and isolated PMMA cylinders with lateral domain correlations, offer a facile method to fabricate sub-micron structures over large areas. For completeness, we also studied the impact of chemically heterogonous substrate in additional to the chemically homogeneous silicon oxide or neutral surface. The NIL-based chemical pattern can provide smaller pattern periodicity than the widely used micro-contact-printing-based method, which renders the opportunities to examine the cases where intrinsic correlation length of the phaseseparated structure is significantly larger than the chemical pattern size. Unique hierarchical structure was obtained with large intrinsic correlation length and the morphological transition to perfect transfer case is revealed.

86 65 CHAPTER V PHASE EVOLUTION OF PLANAR THIN FILMS CONFINED BETWEEN TWO PARALLEL RIGID WALLS (1-D CONFINEMENT) 5.1 Introduction Mixing or blending polymers has been a traditional processing strategy to improve the properties of the neat polymer materials in bulk Processing polymers at nano-scale needs understanding and knowledge of properties of polymers that are different from their bulk. 124 At the length scale comparable to the size of molecules, surface and interfacial interactions between the constituent polymers and the confining environments could be dominant As we have seen in the previous chapter, for thin film with free surface, the interaction between substrate and polymer film has strong impact on the phase evolution. By tuning the substrate surface energy, we can manipulate the phase-separated structure in blend films. Similarly, the state of the other side of the film could also play an important role in the phase evolution. In the following two chapters, we impose confinement on the upper interface of the polymer film and investigate the impact of the physical confinement.

87 66 The simplest confinement on the phase separation of a blend film is to capping the film with two rigid substrates. Wendlandt et al., examined the composition depth profile of polymer blend films confined between two identical cross-linked polyolefin layers. They found that the interface of the two component polymer phases is perpendicular to the confining plane under symmetric boundary conditions, which is consistent with the results of Binder s simulation model. 60 Tanaka et al., reported the observation of interconnected structure following the formation of disk-like droplets in the PS/PVME films confined between two glass plates due to the wetting-induced long-range attractive interaction. 125 Dutcher et al., studied phase-separated morphology of PS/PMMA film capped with SiOx layer with varying thicknesses. Lateral structure was observed when the thickness of capping layer is smaller than a critical value, while lamellar structure was resulted when the mechanical confinement is stronger, i.e., the thickness of the SiOx capping layers is larger than the critical value. 126 In these studies, one side of the film is flexible and the confinement (rigidity of the substrate) is not strong enough to suppress the surface/interface roughening occurred during the domain-coarsening process. Bodensohn et al., confined 2,6-lutidine/water between two concentric cylinders. 127 However, the gap between the two concentric cylinders is large and comparable to the droplet diameter even at late stage. As during most of the coarsening process the droplet diameter is smaller than the gap distance, the confinement is not considered strong either.

88 67 Here we used nanoimprint lithography (NIL) to impose confinement environment and investigated the phase evolution of thermally annealed polymer blend film between two parallel plates. Using NIL provides strong confinement, including both rigidity of the substrates and/or external pressure, on both sides of the film and allows coarsening to occur only in a two-dimensional plane. The gap between the two plates depends on the thickness of the as-cast blend film. It could be as small as tens of nanometers. The confined droplets could result in thicknessdomain width ratio as small as 1%, which has never been reported in literature. 5.2 Samples and techniques Two pairs of PS/PMMA were prepared for examining the phase evolution under symmetric boundary conditions. The composition ratio (by mass) of PS and PMMA is 1:1. The first pair is PS48K (Mw = 48.1 kg/mol, PDI = 1.01; Scientific Polymer Products, Inc.)/PMMA15K (Mw = 15 kg/mol, polydisperse; Scientific Polymer Products, Inc.). This blend was mainly used to examine the morphology evolution during annealing under confinement condition since the low molecular weights would lead to high kinetics during coarsening. The glass transition temperatures (Tgs) of PS and PMMA were determined to be 98 C and 105 C respectively, from the second scan of differential scanning calorimetry (NETZSCH DSC 204 F1) with a scan rate of 20 C/min. PS/PMMA blend films were obtained by spin-coating toluene solutions onto silicon wafers at 2000 rpm for 1 minute. Film thickness was controlled by tuning the PS/PMMA concentration in the toluene solutions. Solutions with four different concentrations 1.5, 3, 5, and 10 wt. % were

89 68 used, and their corresponding thicknesses are reported below. The thickness of film spun-cast from 10 wt. % was measured to be around 600 nm, and thicker films are difficult to obtain directly from spin-coating due to the high viscosity of the blend solution. Instead, we prepared film with thickness around 1 m by sandwiching two pieces of spun-cast 600-nm-thick films with silicon substrates. All films were airdried for at least 72 hours prior to the NIL process. The symmetric boundary condition was realized by using silicon wafer as template during NIL process such that the polymer blend film was sandwiched between two pieces of silicon wafers. For the ~1-µm-thick film, the wafer template was spun-cast with film from 10 wt. % toluene solution as mentioned above. The NIL process was conducted as follows. A pressure of 4 MPa was applied to the system when temperature increased from room temperature to 110 C, which is slightly above the maximum Tg of PMMA and PS. Then the temperature was further increased until it reached the targeted imprinting temperature. The imprinting time was selected to be 30 minutes to allow the coarsening process to occur sufficiently at the imprinting temperature. Further increase of annealing time at such a high temperature under the imprinting pressure was prohibited by the NIL equipment. After the imprinting time was up, the whole system was cooled down to 40 C, which is below Tgs of both polymers. The adhesion between the two pieces of silicon wafer was very strong after the NIL process indicating good contact and strong confinement. They were mechanically separated by inserting a blade in between the two wafers.

90 69 The second pair is PS190K (Mw = 190 kg/mol, polydisperse; Scientific Polymer Products, Inc.)/PMMA94K (Mw = 94 kg/mol, polydisperse; Scientific Polymer Products, Inc.). This pair of PS and PMMA has lower kinetics during coarsening, compared with the first one. The purpose of studying this pair of PS/PMMA is to examine the molecular weight dependence of the phase evolution and the details of interface morphologies. The Tgs of PS and PMMA were determined to be Tg = 98 C, and Tg = 125 C respectively using the DSC measurements with the same procedure mentioned above. PS/PMMA films with thickness ~150 nm were prepared by spincoating 3 wt. % toluene solution at 2000 rpm for 1 minute. Pressure was applied at 130 C during the imprinting process and other imprinting conditions were the same as that described before. Selective dissolution of PS by cyclohexane was carried out at ambient conditions for 15 minutes to reveal the interface between PS and PMMA within the film. 128,129 Atomic force microscopy (AFM, DI3100, Bruker) and field emission scanning electron microscope (FE-SEM, JEOL JSM-7401F) were used to characterize the interface topography. 5.3 Confinement with symmetric boundary condition (Si wafers with native SiOx surfaces) Confined phase evolution with fixed film thickness We first carried out two control experiments for ~100-nm-thick films in order to examine the impact of imprinting pressure and confinement to be discussed later. The first one is simply to anneal the as-cast PS48K/PMMA15K film in air under

91 70 ambient pressure at 160 C for 30 minutes. Figure 5.1A shows the morphology of the as-cast film. The slightly elevated and isolated domains are PMMA. The resulted film morphology from annealing, as shown in Figure 5.1B, was isolated PS droplets with a correlation length of ~8.3 on top of a continuous PMMA wetting layer at the late stage of the phase evolution. This phase inversion phenomenon is consistent with phase-inversion process mentioned in the previous chapter and reported in literature. 30,50,98-100,130 Figure 5.1 Topographic AFM images of morphologies of PS48K/PMMA15K films with free surfaces. A, morphology of as-cast film. B, morphology of film annealed in air at 160 C for 30 min. C and D, morphologies of films after annealing in imprinter at 110 C for 2 min and then in air without external pressure at 160 C for 30 min. C is the morphology of the film left on the mold side and D is on the substrate side. All the AFM images have the same size of 50 m 50 m. The color (z-) scale is 40 nm in A and 100 nm in B, C, and D. Insets in B and D are FFT images obtained from corresponding optical microscope images.

92 71 The second experiment was to imprint the as-cast PS/PMMA film with a flat wafer at 110 C for 2 minutes to generate a good contact between the film and the template wafer. The film was indeed compressed at this processing condition as we verified that separating such two plates without further annealing would cause fracture of the film in between. The PS/PMMA film, sandwiched between two wafers, was then annealed under ambient pressure at 160 C for 30 minutes on a hotplate. After separating the two plates by inserting a blade between them, the PS/PMMA film was found to be left on both the substrate and the template with almost identical surface-relief morphology where the correlation length of the PS domain was ~4.17 (Figures 5.1C and 5.1D). The observations suggest that phase inversion occurred despite the existence of the upper confining plate, where the preferential wetting of PMMA onto both side of the silicon wafer dominates. The breakup of PS domains, as that observed in air (Fig. 5.1B), induces appreciable roughening at the PS/PMMA interface and corresponding pressure that caused the separation of the two wafers. Here we provided detailed morphologies for what Dutcher et al., had described as lamellar structure under similar weak confinement condition. 126 Theoretically, Tanaka predicted the equilibrium phase structure under such weak confinement (i.e., no strong confinement in vertical direction) of two parallel plates by integrating the wetting and interfacial free energies. 23,47 Specifically, there are two possible configurations under such confinement: layered structure (complete wetting scenario) and disc-like droplet structure (partial wetting

93 72 scenario). The transition between the two scenarios occurs at = - /( -wall - -wall), where is the volume fraction of non-wettable phase, - is the interfacial tension between phase and phase, -wall and -wall are the interfacial tension between phase and the wall, and phase and the wall, correspondingly. The PS/PMMA films studied here fall into the complete wetting scenario, as judged by the criteria PS < PMMA-PS/( PS-wall - PMMA-wall). Concretely, PS-wall - PMMA-wall = cos PMMA PMMA cos PS PS based on Young s equation, where PMMA are PS are contact angles of PMMA and PS on SiOx, PMMA and PS are surface tensions of PMMA and PS, respectively. At 160 C, PMMA-PS is 1.52 based on T ( C ) mn/m according to literature. 78,131 PS is reported to be ~7.5 and spreading coefficient is negative for PS on SiOx (-0.2 ~ mn/m in the temperature range of this study). 19,46,132 PMMA and PS are estimated to be mn/m and mn/m at 160 C for PMMA15K and PS48K, respectively, based on LeGrand and Gaines equation 19,133. PS-wall - PMMA-wall is larger than cos PS ( PMMA - PS) = 6.27 mn/m since PMMA has better wettability on SiOx than PS. Therefore PMMA-PS/( PS-wall - PMMA-wall) is smaller than 0.24 at 160 C. This value further decrease to less than 0.07 as the imprinting temperature increases to 230 C. Within the temperature range examined here, PMMA-PS/( PS-wall - PMMA-wall) is always smaller than the volume fraction of PS, ~ 0.5, indicating layered structure is energetically preferred. Thus, the identical surface-relief structure on both wafers we observed is essentially the complete wetting scenario where PMMA forms the two wetting

94 73 layers. However, the in-between PS layer is not a continuous phase as described in the theory, but rather forms isolated droplets with a clear spatial correlation. The breakup of the continuous PS phase is likely due to the Rayleigh instability similar to that in free-surfaced films, which induces significant roughening process at the PS/PMMA interfaces. The formation process of the symmetric surface-relief structure is shown in Figure 5.2. During the imprinting process, the initially elevated PMMA domains (Figure 5.2A) were flattened out (Figure 5.2B). Due to preferential wetting tendency of PMMA on SiOx, there formed a covering layer onto the upper plate. Upon further Figure 5.2 A schematic illustration of the formation process of the symmetric surface-relief structure. (A) as-cast morphology after spin-coating. (B) morphology formed after pressurization at 110 C for 2 min. Arrows indicate the direction of materials flow tendency upon further annealing. (C) thermally induced capillary wave occurred within the film during the annealing without external pressure. (D) final morphology of symmetric surface-relief structure, corresponding to the crosssectional profiles of Figures 5.1C and 5.1D.

95 74 annealing, PMMA domains in the middle of the film tend to migrate to either upper wetting layer or the bottom wetting layer. Therefore there exists a tendency for PMMA domains to shrink, which allows PS to expand laterally to form continuous middle layer (Figure 5.2C). Note that the transitional tri-layer shown in Figure 5.2C might not really exist. It is possible that a combined dewetting of PS on PMMA wetting layer and capillary fluctuations within the film occurred simultaneously as the migration of PMMA towards wetting layers. In our study, the elastic energy of the Si wafer (~500 m thick) is very large, and the roughening process occurred only within the films, causing the separation of the two plates. This is distinctively different from the morphological paths of PS/PMMA observed in the thin SiOx capping layers, where the roughening due to the morphological evolution induced the wrinkling (deformation) of the capping layers. 126 Figure 5.3 shows the phase evolution of PS48K/PMMA15K film with film thickness ~100 nm under a pressure of 4 MPa. AFM images were taken after the samples were selectively dissolved by cyclohexane and show the remaining PMMA morphologies within films. The elevated and isolated domains in AFM images are PMMA. The so-called surface relief structure was not observed in the film sandwiched between two wafers under the imprinting pressure. As shown in Fig. 5.3, PMMA remained as the dispersed phase throughout the whole annealing process for all the different temperatures examined. Strikingly, the external pressure applied on both sides of the wafers during the imprinting process apparently prevented the phase inversion process from occurring, as shown by

96 75 comparing Figures 5.1 and 5.3. The underlying mechanism is believed to be that the external pressure inhibits the growth and even thins the PMMA wetting layer as described in the following. Figure 5.3 Topographic AFM images for phase evolution of PS48K/PMMA15K films between two parallel plates. AFM images were taken at film surfaces revealed at the bottom interface after selective removal of PS. Imprinting temperatures are 160 C, 190 C, 210 C, and 230 C, respectively, as marked on each image. All the AFM images have the same size of 50 m 50 m and color (z-) scale of 200 nm. Although the driving force and the kinetics of the wetting layer growth is still an open question, one plausible explanation is based on the existence of pressure gradient between the inside and the boundary of the wettable domain. 125 It assumes that the wettable domain boundary towards the wall side has extremely small transverse curvature (1/ ) and the curvature difference 1/a leads to a pressure gradient of /a, where a is the radius of the domain and is surface tension (Figure

97 76 5.4A). The lateral growth of wetting domain and the subsequent wetting layer thickening are both based on the pressure gradient-driven mass flux. For PS/PMMA film with free surface, this mass flux leads to the migration of PMMA to the bottom of PS domains while lifting the overall height level of PS domains. In weakly confined cases, PMMA migrated towards both upper and lower Si wafers, squeezing PS in between. Figure 5.4 Schematic for the driven force of the wetting layer formation. A, wetting occurred in free surface. B, the prevention of wetting layer growth under confinement. In the presence of external pressure (i.e., 4MPa used here), the assumption of the extremely small curvature at the vicinity of the wall is invalidated. The pressure gradient that causes the mass flux might be not existent. More importantly, PS domain tends to generate pressure vertically due to Laplace

98 77 pressure. This Laplace pressure along with the external pressure would squeeze PMMA out from the region beneath PS and generate a large pressure in the PMMA wetting layer. The pressure in the PMMA wetting is roughly PMMA-PS/h, where h is PMMA domain height. As the PMMA domain height is smaller than the PMMA domain radius, even if there exists the pressure gradient PMMA-PS /a for the mass flux, the pressure inside the wetting layer PMMA-PS/h would be dominant. Figure 5.5 Schematic of thinning of PMMA layers. A, thinning of intrinsic PMMA wetting layer. B, thinning of artificial PMMA cushion layer. C, topographic AFM image with color (z-) scale of 200 nm corresponding to the coarsened morphology in B. We further verified our hypothesis that external pressure not only prevents the wetting layer growth but also causes the wetting thinning through the following experiment. We placed a PMMA cushion layer in contact with the PS/PMMA blend film, as schematically shown in Figure 5.5B. The PS/PMMA blend film is identical to the one described above, which is ~100 nm thick. This film is stacked onto a pure PMMA layer which is spun-casted from a 2 wt. % toluene solution to form a bi-layer. The annealing condition is the same as that in Figure 5.3D, 230 C for 30 min under a pressure of 4 MPa. Similar to the single layer case where PMMA in wetting layer is absorbed into the PMMA domains dispersed in PS matrix (as schematized in

99 Domain area (μm 2 ) Domain width and correlation length (µm) 78 Figure 5.5A), the PMMA layer is also absorbed into the isolated PMMA domains within the blend film. As a result, the PMMA domain height is ~170 nm which is much larger than the blend film thickness ~100 nm. In the following, we describe the detailed morphology evolution under the pressurized confinement environment. After being imprinted at 160 C for 30 minutes, the diameter of the PMMA domains grew to 3~10 times larger than that of the as-cast film, due to the coarsening and perhaps secondary phase separation. The height of PMMA domain, after removing the PS phase, was ~100 nm for all the imprinting temperatures studied, which roughly equals the as-cast film thickness. This suggests that during the pressurizations, no significant squeeze flow occurred at this thickness (film spun-cast from 3 wt. % solution), i.e., there was no appreciable PS/PMMA squeezed out of the edges of the two plates. The squeeze flow will become obvious when the film thickness increases to ~ 1 m as we will discuss later. Area Width Correlation length Temperature ( C) Figure 5.6 A plot of correlation length, PMMA domain area, and PMMA domain width as a function of imprinting temperature.

100 79 As the imprinting temperature increases, the correlation length increased from 1.79 µm at 160 C, to 5 µm at 190 C, and to 12.5 µm at 210 C as shown in Figure 5.6. At 230 C, anisotropic coarsening occurred dramatically and the 2-D correlation length is undefined and cannot be obtained. Instead, the average distance between domains along the direction perpendicular to the long axis s is estimated to be ~10 µm from the SEM image shown in Figure 5.7. Note that the alignment directions are not the same across the whole sample. Different regions have different alignment directions. Figure 5.7 is only a representative image of one region. This type of alignment is probably caused by the local pressure gradient. It is not possible for the two confining wafers to be perfectly parallel during the whole coarsening process. Figure 5.7 SEM image for film pressurized at 230 C for 30 minutes, which was taken from the same sample as Figure 5.3D. Dark domains in the image are PMMA.

101 80 Average PMMA domain area increases with the imprinting temperature as well (Figure 5.6), as the PMMA domains developed into highly anisotropic ribbonlike phases. To characterize the anisotropic domain shape, we analyzed the domain aspect ratio (AR), as well as domain width (Lmin). Based on the definition of AR and roundness Rd, we express Lmin as a function of domain area (S), Rd and AR, which can be directly obtained from ImageJ, Since both Lmax (length of major axis of domain) and Lmin grow as coarsening progresses, we used Lmin to characterize the kinetics of the coarsening process and AR to capture the additional anisotropic information of domains. Figure 5.8 shows the statistical distributions of Lmin of PMMA domains after 30 min at four different imprinting temperatures. All the distributions were Gaussian-type distribution. This is in contrast to Lmax or AR (Figure 5.9). The coarsening mechanism along Lmin direction for a domain seemed to be driven by minimizing the surface and interfacial energy via conventional collision-based coarsening. It is not possible, however, to quantitatively compare the kinetics of the growth of Lmin with unconfined counterpart, since in free-surfaced cases it had already been at the late stage when samples were annealed for 30 min at 160 C as indicated by well correlated isolated PS droplets. Under pressurized confinement condition, even at temperature as high as 230 C, the morphology was still not stable and domain coarsening kept evolving. From 160 C to 230 C, average

102 81 Figure 5.8 Histograms of PMMA domain width Lmin. Density defined as normalized frequency is used as y-axis for better visualization. A, B, C and D are for imprinting temperature 160 C, 190 C, 210 C, and 230 C, respectively. domain width Lmin increased from 0.74 μm to 6.64 μm (Figures 5.6 and 5.8). The maximum domain width at 230 C was more than 10 μm, which means a low thickness-domain width ratio (less than 1%) was achieved. The growing kinetics of Lmin is slower than that of the correlation length (Figure 5.6) and is also significantly slower than Lmax (reflected by the increase of AR). Notice that the standard deviation of Lmin maintained very large, around 30% of the mean value and the minimum value of Lmin is almost the same at different temperatures, i.e., in Figures 5.3A-C domains with size as small as the smallest ones in Figure 5.3D could be easily found. These broad distributions are another indication of the significant slower kinetics in confined case relative to free-surface case.

103 82 Figure 5.9 shows the distributions of AR of the PMMA domains. In as-cast film, all the domains are cylinder-like and have AR equal to 1. In free-surfaced films on SiOx substrate, all the domains keep the apparent circular shape (no matter cylindrical or conical) during the coarsening process and therefore almost all the AR values accumulate at 1. Under pressurized confinement condition, at 160 C although most of the domains had AR closed to 1, there existed a few larger ones approaching 2.2. As the imprinting temperature increases, the overall AR distribution shifts to a relatively uniform distribution with maximum value much larger than 4. Figure 5.9D only shows a portion (up to 4) of the overall distribution for the purpose of direct comparison with the histograms of other imprinting temperatures. The broader distribution of AR at higher temperature indicates that the higher the imprinting temperature was, the more elongated domains were. Figures 5.8 and 5.9 together also suggest that the distribution of Lmax has a skewed Gaussian distribution with a fatter tail on the right side, which means as the temperature increases, Lmax increases more than Lmin and it causes AR to increase. The irregular shapes with AR much larger than 1 indicate that the shape relaxation driven by the uneven in-plane curvatures is slower than the domaincoarsening process under the strong confinement. The shape relaxation time, the time required to evolve into a circular shape, is proportional to the volume/surface ratio. 125 With a fixed volume, smaller interface tends to have longer relaxation time. Between the two plates, relaxation can only occur along in-plane directions. This means only a small portion of the domain boundary is possible to be the relaxation

104 83 frontier. This reduces the effective interface area. The relaxation frontier area fraction over the whole interface area is estimated to be 2 ah/(2 ah + 2 a 2 ) h/a, where h is the film thickness and a is the domain equivalent diameter. At high temperature such as 230 C, this number could be as low as 2% (~100 nm / ), which means only 2% of the total interface is free to relax. The anisotropic domain shape therefore becomes very stable and to relax to disc shapes takes an extremely long time. Figure 5.9 Histograms of PMMA domain aspect ratio AR. Density defined as normalized frequency is used as y-axis for better visualization. A, B, C and D are for imprinting temperature 160 C, 190 C, 210 C, and 230 C, respectively.

105 Thickness effect on phase evolution Figure 5.10 Phase evolution of PS48K/PMMA15K films with different film thicknesses between two parallel plates. A-E, topographic AFM images of film surfaces revealed at the bottom interface after selective removal of PS. Film thicknesses are ~50 nm, ~100 nm, ~200 nm, ~600 nm and ~960 nm, respectively. F is the plot showing the relationship between film thickness and blend solution concentration.

106 85 Figures 5.10 and Figure 5.11 show AFM and SEM images of remaining PMMA after PS was selectively dissolved by cyclohexane, for films with different thicknesses imprinted under the same condition (4 MPa at 230 C for 30 minutes). Here the film thickness is the thickness (excluding the several nanometer thick PMMA wetting layers on both sides) after imprinting rather than the as-cast one. The four as-cast films were obtained from PS/PMMA solutions with concentrations of 1.5, 3, 5, and 10 wt. %, respectively. The imprinted film thickness, estimated from the PMMA domain height after selectively removing PS, increased with the solution concentration as plotted in Figure 5.10F. Figure 5.11 SEM images for the corresponding AFM images in Figure A, B, C and D correspond to A, C, D, E in Figure 5.10, respectively.

107 86 Specifically, the imprinted film thicknesses are ~50 nm, ~100 nm, ~200 nm, and ~600 nm with the increase of solution concentrations. We also used two 10 wt. % films and stacked them together face to face to achieve a thicker film, whose PMMA morphology after the annealing is shown in Figure 5.10E. The film thickness is estimated to be ~960 nm, which is thinner than twice of the ~600 nm film possibly due to the squeeze flow during the imprinting process suggesting 25% compressive strain. When the film thickness is ~50 nm, the domain width is smaller than that of ~100-nm film and AR is larger. As the film thickness increases, the domain becomes increasingly more circular. Migler et al., studied the capillary instability of polymer threads under the confinement of two parallel plates. 62 They found that the thread state is very stable when the ratio between the gap distance and the thread diameter, h/d0, is sufficiently low, i.e., below 1.3. The reason could be twofold. First, theoretically the thermal fluctuation decays significant within such parallel plate confinement. 134 Second, the reduction in thread radius caused by the fluctuation decreases in confined case and the threads are less likely to break up while maintaining the current shape. The flat ribbon and disc structures we obtained during confined coarsening process are similar to the beginning stage of Migler s study. For 3 wt. % film, h is ~100 nm and the initial domain width could be as large as ~400 nm. Even before the coarsening progresses, the h/d0 or h/lmin ratio is less than ~1.3. After imprinting at 230 C for 30 minutes, the domain width is ~4 μm and the h/lmin ratio

108 87 is ~0.025, much lower than the critical value ~1.3. The wavenumber at that state is zero and the fluctuation is zero too, both of which mean no breaking up tendency and the state is stable. For thicker film such as ~960-nm film, the as-cast domain diameter (still ~400 nm, despite the fact that the film thickness is larger) is much smaller than the gap between two plates. So thicker films undergo certain stages where h/d0 is larger than ~1.3. Thermal fluctuation starts to play such that wavenumber is larger than 0, preventing domains from forming elongated structures. Therefore, reasonably circular domains are resulted. Similar to the domain width distribution for ~100-nm film at different temperatures, the width distribution for different thicknesses maintained Gaussian distribution. The general trend as the film thickness increases is that the mean value increases, which is consistent with typical coarsening process. ~600-nm film seems to be a stable state two-dimensionally since further increasing the film thickness to ~960 nm does not increase the width, or diameter. However, it does coarsen more in ~960-nm film when we consider the volume instead of the crosssection area. To illustrate how the domain deviates from isotropic shape, we plot Lmin against domain area as shown in Figure This type of scatter plot is unique to anisotropic coarsening. For isotropic coarsening, i.e., in bulk or PS/PMMA film on silicon oxide surface (in plane), Lmin is the domain diameter which always equals to regardless of the film thickness. For anisotropic coarsening, there exists an deviation from the ideal isotropic state, which is reflected by the curve y =.

109 88 The deviation decreases as the film thickness increases. Figure 5.12 also implicitly indicates that AR approaches to 1 as film thickness increases. Figure 5.12 Scatter plots of Lmin as a function of domain area for films with different thicknesses. A, B, and C are scatter plots for films with thickness ~50 nm, ~100nm, and ~600 nm, respectively. D is the normalized deviation of scatter points from ideal isotropic coarsening curve y = Molecular weight effect on phase evolution As we discussed in Figure 5.3, for PS48K/PMMA15K, the PMMA domain height which is roughly the gap distance maintains ~100 nm regardless of the imprinting temperature. The morphology evolution is mainly two dimensional and underwent significant change. The structures obtained after imprinting process are

110 89 believed to be at intermediate stage. We used PS190K/PMMA94K to examine the situation of a somewhat earlier stage and to reveal the gradual change in the interface regions of interest, as this polymer pair has lower chain mobility (or higher viscosity) and therefore slow kinetics during coarsening process. Figure 5.13 Morphology of confined PS190K/PMMA94K film between two flat plates imprinted at 180 C. A, B and C are AFM topographic images of interface between polymer film and template wafer. A is the morphology of film surface. B is the corresponding leftover on the template wafer. C is the film morphology after selective removal of PS. D, E and F are AFM topographic images of interface between polymer film and substrate, corresponding to A, B and C, respectively. All AFM images are 10 m 10 m with color (z-) scale 200 nm. When PS190K/PMMA94K film was imprinted with flat Si wafer at 180 C, the resulted morphology shows symmetry to some extend despite the initial asymmetric bi-layer-like structure after spin-coating. Figures 5.13A-C show morphologies of the interface between polymer film and template wafer and Figures 5.13D-F show morphologies on the other side. Qualitatively, the two interfaces show very similar

111 90 structure. Figure 5.13A shows the surface topography of film supported by the substrate after demolding. It comprises three different types of regions (PMMA, PS and interface between them) reflecting three typical fracture mechanisms. The PMMA cylinders, actually PMMA disks considering the small gap between the two plates, fractured through brittle fracture mechanism and resulted in rather flat surfaces. The matrix of the film is made of PS and the surface roughness of PS region is relatively large, compared with that in PMMA region. Both PS and PMMA are expected to undergo brittle fracture at room temperature, as the percent elongations of PS and PMMA are very close to each other, and , respectively. 135 The slight ductile fracture of PS is probably caused by the introduction of PMMA from the wetting layer on the template wafer. The ductile fracture phenomenon is more obvious at the interface between PS matrix and PMMA disks across the film. The edge surrounding each PMMA disk domain is as high as ~100 nm, or the percent of elongation could be as large as ~100. Considering the width of the edge, the aspect ratio could be on the order of ten. Jeong et al., have reported the fabrication of high-aspect-ratio nanohairs (>20) through nanodrawing, a similar process as used in this study. Based on the estimation of adhesion energy ratio, the ratio of the production of work of adhesion and surface area at the polymer/substrate interface to that at the mold/polymer interface, they found that with this ratio closed to 1, PS or PMMA film undergoes delamination, fracture and elongation concurrently on the same sample when demolding is executed within the temperature range of C. 136 In our

112 91 symmetric boundary condition case, the adhesion energy ratio is exactly 1 and the demolding temperature is room temperature. Our observations of high-aspect-ratio nanostructure, delamination at the transition region, and that the fracture occurred between at polymer/plate interface are all mostly consistent with literature. Figure 5.13B presents the surface morphology on template wafer corresponding to the film shown in Figure 5.13A. The ultra-small surface roughness with Ra ~2nm indicates that the fracture of the film occurs mostly at the top of the film (Figures 5.13A-C) or at the bottom of the film (Figures 5.13D-F). The materials left on the template wafer is believe to be PMMA grafting layer due to the interaction between PMMA and hydroxyl groups on SiOx substrate, which maintained after selective dissolution by cyclohexane. Materials distribution of the film was verified as shown in Figure 5.13C. Disk-like PMMA domains were left after removal of PS, which have similar diameter and correlation length as the as-cast film. PMMA remained isolated domains as in the as-cast film. The kinetics of coarsening process dramatically slows down. The slowed-down kinetics is caused by the suppression of the wetting layer growth and therefore the translational motion. Hydrodynamic transport is less dominant than collision-based diffusion. Furthermore, the friction between the polymer film and the plate due to the existence of the pressure applied during the imprint process also inhibits the motion of coarsening. The slowed-down kinetics of confined polymer films could be also predicted by simulation based on Cahn-Hilliard

113 92 equation, for example, in Puri s work, the domain size and the correlation length stabilize much earlier for the partial wetting case than complete wetting case. 58,137 Figures 5.13 D-F capture the interface morphologies of between polymer film and the substrate. They appear similar to their counterparties on the other side (Figures 5.13 A-C). The similarity in morphology is resulted from the symmetric boundary condition applied during the confined coarsening process. PMMA domain sizes are close to each other on both sides (Figures 5.13A and 5.13D). The domain heights in Figures 5.13C and 5.13E are almost the same ~130 nm, indicating that the holes are penetrating the film mostly from top to bottom. Notice that the domain heights are smaller than the film thickness ~150 nm. This suggests that there exists a wetting layer with thickness of several tens of nanometers. This thick wetting layer is found to be existent at the bottom of the film. The stark contrast between Figure 5.13E and Figure 5.13B shows that there are more PMMA at the bottom interface. This is the combined consequence of the structure formation after spinning-coating and the slow kinetics. After spinning-coating, the structure is highly asymmetric with much more PMMA aggregated onto the substrate as a wetting layer due to the preferential wetting tendency onto the substrate, meanwhile on the surface side the PS matrix is exposed to the air and contributes to the major component at the upper interface. During the phase evolution, the morphology is transforming towards a more symmetric state. However, as the kinetics is slow, after imprinting at 180 C for 30 minutes, the slight asymmetry still exists. The big PMMA bumps in Figure 5.13E justify this slight asymmetry.

114 93 Figure 5.14 Morphology of confined PS190K/PMMA94K film between two flat plates imprinted at 210 C. A, B and C are AFM topographic images of interface between polymer film and template wafer. A is the morphology of film surface. B is the corresponding leftover on the template wafer. C is the film morphology after selective remove of PS. D, E and F are AFM topographic images of interface between polymer film and substrate, corresponding A, B and C, respectively. All AFM images are 10 m 10 m with color (z-) scale 200 nm. When the confined coarsening occurs at 210 C, the resulted morphology is very similar to that obtained at 180 C, as shown in Figure The negligible change in morphology from 180 C to 210 C further reveals that the kinetics is slowed down under confinement conditions, which otherwise would be huge resulting well developed relief structure at 210 C. The disk-like PMMA domain size as well as the domain circularity both increases a bit at this relatively high temperature. More interestingly, the torus-shaped edge contouring the PMMA domain boundary becomes sharper. The elongation percentage is larger than the

115 C counterpart and the aspect ratio could be as high as 50. The surface roughness Ra in Figure 5.14B is estimated to be ~2 nm, which is very close to the 180 C case. This means the polymer/upper plate interface is stable and the mass transfer to the wetting layer vertically is negligible. The surface roughness in Figure 5.14E dropped to 4.16 nm at 210 C from 14.9 nm at 180 C, and the big bumps shown in Figure 5.13E disappear. The domain heights in Figures 5.14C and 5.14F increase to ~140 nm, which is very close to the film thickness. All of these indicate the thinning of PMMA wetting layer at the bottom of the film, consistent with earlier observations in the smaller molecular weight system. 5.4 Effect of substrate surface energy We also examined the impact of substrate surface energy on the confined phase-separated morphology, particularly under similar imprinting pressure. Silicon wafer is treated by coating random copolymer layer onto it to form neutral surface as discussed in chapter 4. The results turned out to be that for both the ascast films and the confined films, the film morphologies are essentially the same as those obtained with silicon oxide surfaces. Figure 5.15 shows two examples of morphologies of ~100-nm PS48K/PMMA15K films confined between two plates with neutral surfaces at 160 C and 230 C for 30 min. The reason could be found in role of the wetting layer during the coarsening of the confined film. With silicon oxide surfaces, external pressure inhibits the growth of the wetting layer. Upon annealing PMMA has a tendency to dewet on the neutral surface. However, dewetting of PMMA on some region of free neutral

116 95 surface will cause PMMA thickness increase of neighboring regions, i.e., the gain in interfacial energy is dissipated in the rim of the hole. 138 Under confinement, such energy dissipation becomes difficult since vertical thickness growth is prevented. Therefore the external pressure not only inhibits the wetting layer from growing but also from dewetting. Figure 5.15 Morphologies of ~100-nm PS48K/PMMA15K films confined between two plates with neutral surfaces at 160 C and 230 C for 30 min. (A) topographic AFM image for film pressurized at 160 C. The color (z-) scale of 200 nm. (B) SEM image for film pressurized at 230 C. A and B are very similar to their SiOx counterparts Figure 5.3A and Figure 5.7, respectively. 5.5 Conclusions In this chapter, we examined the phase evolution of PS/PMMA with strong physical confinement under symmetric boundary conditions by using NIL. The confinement generates domains with thickness-domain width ratio as small as 1%, which is much lower than what had been reported in literature. Phase inversion that is commonly observed in free-surfaced case is avoided due to the prevention of

117 96 preferential wetting tendency of PMMA on silicon oxide substrates and the suppression of surface roughening. PMMA domains maintain dispersed in PS matrix throughout the confined coarsening process. The domain shape depends on the film thickness. When film thickness is less than 200 nm, the elongated domain caused by the slow relaxation has less tendency to break up due to the low thermal fluctuation. Dramatically slowed-down kinetics is observed in PS/PMMA films with moderate molecular weights. Phases behave as if they were frozen even at relatively high annealing temperature. We also revealed detailed interfacial morphologies within the film and at the boundaries.

118 97 CHAPTER VI COARSENING PROCESS OF POLYMER BLENDS UNDER TE-NIL CONDITIONS (2-D CONFINEMENT) 6.1 Introduction NIL offers unique capability and potential to directly pattern a range of functional materials. 3,6-12 Up to now, neat polymers have been widely used in NIL fabrication, where their viscoelastic properties dictate both the fidelity of the pattern replication and the stress state of the obtained structures However, neat polymers are intrinsically limited by their chemical and physical characteristics. Multifunctional polymer nanostructures with a combination of topographic features and diverse chemical functionalities are critical to many emerging technologies. Mixing or blending polymers has been a traditional processing strategy to improve the properties of the neat polymer materials in bulk. 17,19,139 Currently, it remains unclear whether it is a viable approach to fabricate chemically heterogeneous nanostructures, i.e., demixed polymer blend, using NIL.

119 98 Besides the implication to the practical applications, the morphological evolution of multicomponent polymers under the NIL process is a fundamentally intriguing problem. At the length scale relevant to the NIL, surface and interfacial interactions between the constituent polymers and the confining environments may be dominant In addition, the physical confinement of the cavity walls is stronger than that in the planar thin films, which will significantly influence the phase separation and morphological evolutions such as domain breakup or coarsening. 25,26 Under these motivations, this chapter presents the fabrication of polymer blend patterns, focusing on the morphological evolutions of the blend under systematically varying NIL conditions, blend compositions, and film thickness. Compared with the phase evolutions of blend films within two planar rigid substrates discussed in the previous chapter, the superstrate used here is topographically patterned Si mold. The results clearly demonstrate the significant influence of preferential wetting and physical confinement on the morphological evolutions of the confined polymer blend, which leads to a range of unique topographically uniform and chemically heterogeneous patterns. 6.2 Samples and techniques PS and PMMA with corresponding molecular weight of 190,000 g/mol and 94,000 g/mol were purchased from Scientific Polymers and used as received. The glass transition temperature (Tg) for PS and PMMA were determined to be 100 C ± 2 and 126 C ± 2 C, respectively, by differential scanning calorimetry (DSC) with a scanning rate of 10 C/min. PS/PMMA films on silicon wafers (with a native oxide

120 99 layer) were prepared by spin-coating, at 2000 rpm for 1 min, from their toluene solutions. From a predetermined spin-coating curve, the PS/PMMA film obtained was ~150 nm. The films were subsequently annealed at 50 C under vacuum for 2 h to remove the residual toluene. Three different compositions of the PS/PMMA films, 30/70, 50/50, and 70/30 in weight ratio, were prepared. Figure 6.1 A schematic of the geometry of the nanoimprinting mold used in Chapter 6, which has a periodicity of 834 nm, a line-to-space ratio of 1:1, and a cavity depth of ~195 nm. NIL processes on the PS/PMMA films were carried out on a nanoimprinter (Eitre 3, Obducat, Inc.). The mold applied is a parallel line-and-space grating consisting of SiOx with a periodicity of 834 nm, a line-to-space ratio of 1:1, and a cavity depth of ~195 nm. A schematic of the mold is shown in Figure 6.1. The mold was treated with piranha solution prior to use. To facilitate the mold separation after imprinting, a low surface energy self-assembly monolayer of CF3(CF2)5(CH2)2SiCl3 (tridecafluoro-1,1,2,2-tetrahydrooctyltrichlorosilane, Aldrich, Inc.) was deposited onto the mold surface through a vapor deposition process. For each PS/PMMA composition, NIL was carried out at three temperatures, 150, 180, and 210 C, under a pressure of 4 MPa for 30 min. The mold was separated from the PS/PMMA replica after the system was cooled down to 50 C. To distinguish different components within the films and patterns, cyclohexane (CLE) and acetic

121 100 acid (AA) were used to selectively dissolve PS and PMMA, respectively. AFM (DI3100, Vecco) and a high resolution optical microscope (Nikon LV 150) were used to examine the morphologies of the films or patterns obtained. 6.3 Structure formation of PS(190K)/PMMA(94K) films For all the PS/PMMA films, phase separations were observed after the spincoating, as expected from the phase diagram of PS/PMMA with the given molecular weights, and as demonstrated in Chapter We note that all the temperatures that the films and patterns were exposed to in this study are within the two-phase region of the PS/PMMA phase diagram. The phase separation causes the film surfaces to be rather rough, and the exact film thicknesses were not determined but estimations from the solvent treated patterns show that the thickness is close to the targeted value. Figure 6.2 presents optical micrographs and topographic AFM images of the as-cast PS/PMMA films with three designated compositions: S30/M70, S50/M50 and S70/M30, revealing two-phase morphologies for all the films. Each of the two phases was identified by AFM measurements on the films after selective dissolution of either PMMA by acetic acid (AA) or PS by cyclohexane (CLE). Islands or cylindrical domains were observed in the as-cast films. The relative heights of these islands and cylindrical domains increased after the PS dissolution (Figures 6.2 A3, B3, and C3), while holes were found after corresponding PMMA removal (Figures 6.2 A4, B4, and C4). This clearly shows that PS forms the continual phase for all the films even in the 30/70 composition where PS is the minor phase. It is known that bulk

122 101 morphology for a pair of immiscible polymers is determined by both the composition and the viscosity ratio between the two components. 18 The major component of the blend is likely to form the continuous phase. According to the minimum energy dissipation mechanism, the lower (higher) viscosity component tends to form the Figure 6.2 Optical micrographs (OM) and topographic AFM images of PS/PMMA blend films with film thickness around 150 nm and three different compositions. The first, second and third rows show morphologies of S30/M70, S50/M50 and S70/M30, respectively, as marked on the left. The images in the first column are optical micrographs of as-cast films. The second, third and fourth columns are topographic AFM images corresponding to the morphologies of as-cast films, films after PS was selectively dissolved by cyclohexane (CLE), and films after PMMA was selectively removed by acetic acid (AA), respectively. All AFM images are 10 µm 10 µm in size with color scale 400.

123 102 continuous (dispersed) phase, as frequently witnessed by melt processing of immiscible polymer blends. 18 The interplay between these two factors determines the final morphology of the blend. In the present case, PMMA is the more viscous component and tends to form the dispersed phase according to the minimum energy dissipation mechanism. The observation that the major component PMMA forms the dispersed phase in S30/M70 films suggests that the viscosity ratio is probably the dominant factor. In addition, in blend thin films, the morphology is also strongly affected by the surface and interfacial energies of the components. In this regard, PS has slightly lower surface tension than that of PMMA, which could also favor higher surface coverage of PS. 140 Further, studies have shown that autophobic dewetting of PMMA due to the favorable interaction of PMMA with the surface hydroxyl groups at the substrate surface could also lead to similar morphologies. 141,142 Regardless, PS forms continuous phase and PMMA domains are elevated over the PS domains during spin-coating. All the morphologies are consistent with literature. 98 Figure 6.2 also shows that the size of the PMMA domains progressively increases with its concentration. For the S30/M70 films, the spatial distribution of the PMMA islands clearly indicates a bi-continuous-like morphology resulting from spinodal decomposition during the spin-coating process. 17 The wavelength (λf) or the correlation length, from the FFT (inset) of Figure 6.2A1, was determined to be μm. For S50/M50 and S70/M30 films, the PMMA rich phases have evolved into cylindrical shape. For the S50/M50 films, the diameter (dpmma) of PMMA domains

124 103 was estimated, from the holes left after the AA dissolution, as up to 0.6 μm. Similarly, dpmma was estimated to be ~0.35 μm for the S70/M30 films. Consistent with the studies by Morin et al., smaller domains of PS (or PMMA) within the PMMA (PS) rich domains are clearly observed in the AFM images in Figure Patterned morphology of 150-nm PS/PMMA films For each composition, three temperatures (Timps) were applied, 150 C, 180 C, and 210 C, to imprint the PS/PMMA films. Because the mobility of PS and PMMA strongly depends on temperature, the selected Timps will provide a kinetic pathway of the morphological evolution of PS/PMMA during NIL process. Specifically, the steady state viscosities of PS are estimated to be Pa s, Pa s, and Pa s at 150 C, 180 C, and 210 C, respectively. 143 In comparison, the viscosities of PMMA are roughly Pa s, Pa s, and Pa s at 150 C, 180 C, and 210 C, respectively. 140,144 In contrast to the dramatic change in chain mobility, the surface tensions (γs) of both PS and PMMA only change slightly: γpss are 31.3, 29.2, and 27 mn/m, while γpmmas are 30.8, 28.5, and 26.3 mn/m at 150 C, 180 C, and 210 C, respectively. 19 The following subsections discuss the morphologies of the patterned PS/PMMA films with corresponding compositions of S30/M70, S50/M50, and S70/M30, respectively. Determined by the mold geometry (refer to Figure 6.1), the minimum thickness of the resist film to completely fill the cavities is ~100 nm. Below this thickness, incomplete filling of the cavity will occur, which results in significant broken polymer lines during imprinting. This is due to the nucleated dewetting of

125 104 the PS/PMMA within the cavities, which has been observed even in the neat polymer resist Thicker films such as 150 nm will provide enough polymeric materials to completely fill the cavities and leave no space for nucleated dewetting. This was indeed the case for all our 150 nm and 300 nm blend films. High fidelity of pattern replications was achieved across the whole field of the patterns (2.5 cm 2.5 cm) at all three Timps, as confirmed by AFM mapping. The as-imprinted PS/PMMA lines have heights of ~195 nm, precisely matching the depths of the mold cavities. In addition, the patterned lines display sharp corners, which is consistent with complete cavity filling. Such good pattern replication is determined by the viscoelastic properties of the PS/PMMA, i.e., both components are capable of large deformation under the imprinting conditions. In the following, detailed morphologies of PS/PMMA within the uniform patterns are described Morphology of patterned S30/M70 films In order to distinguish the features that are resulted from confinement effect, S30/M70 films were annealed in air as controls. Figure 6.3 shows the morphologies after the films were annealed for 30 minutes at three different temperatures, which are the same as the three imprinting temperatures. After the film was annealed at 150 C, the surface roughness dropped to less than 5 nm, which was caused by surface leveling driven by the surface tension of the PS and PMMA. The isolated PMMA domains are surrounded by valleys with depth ~20 nm, and the height level of these domains is almost the same as that of the continuous PS matrix. The correlation length 3.3 µm, estimated from the FFT of the image, is almost the same

126 105 as that of as-cast film. This unnoticeable change in correlation length indicates that the kinetics of both PS and PMMA at this temperature so low that domain coarsening process is insignificant. Nor did the roughness at interface (Ra, 28.5 nm before annealing and 29 nm after annealing, obtained after selective dissolution of the films) change much, suggesting again the kinetics was slow and the morphology change happened only at the surface. Figure 6.3 Optical micrographs (OMs) and topographic AFM images of 150-nm S30/M70 blend films annealed at three different temperatures as marked on the left. The images in the first column are optical micrographs of surface morphologies. The second, third columns are topographic AFM images corresponding to surface morphologies, and interfacial morphologies after PS was selectively dissolved by cyclohexane (CLE). All AFM images are 10 µm 10 µm in size.

127 106 In comparison, the morphology would evolve into ridge-valley structure at higher temperature. When film was annealed at 180 C, the ridge had already started to break up into isolated spheres but the breakup was not complete, which is more obvious in AFM image B2 than optical micrograph B1 owing to the better resolution of AFM. More PMMA migrated into bottom layer resulting in enlarged interface coverage of PMMA as shown in B3 compared with A3. At the same time PS was driven upwards contributing the domain growing in the coarsening process. Holes at the interface correspond to the spheres at the surface. Irregular holes at the interface match the incomplete breakup at the surface, both of which indicate that the film morphology at this stage is not thermally stable. When film was annealed at 210 C, PS droplets became more isolated. Domain size increased while the domain density decreased, as shown in Figure 6.3 C2 and C3, compared with B2 and B3, respectively. Besides, interface was smoother and holes were more rounded, all of which minimized the surface and interfacial energy and led to a more thermally stable state. For the imprinted S30/M70 patterns (Figure 6.4), the optical micrographs (A1, B1, and C1) clearly reveal the two distinct phases, in contrast to the topographic AFM images (A2, B2, and C2). The dark and bright regions in A1, B1, and C1 correspond to the PS and PMMA phases, which were further verified by the AFM measurements on the patterns after selective dissolution. Similar to the ascast film (Figure 6.2 A1) and annealed free-surface planar film (Figure 6.3 A1), the PS/PMMA patterns created at Timp = 150 C (A1) display an inverted morphology

128 107 across the patterned lines: correlated PMMA-rich domains dispersed in the PS-rich matrix despite that PMMA is the major component. This unique structure is manifested in the FFT of the optical micrograph (inset of A1): the scattering ring corresponds to the spatial correlations of the PMMA domains, and the two symmetric spots are the first order diffraction peaks (corresponding to a pitch of ~834 nm) of the periodic PS/PMMA pattern. The λf of the PMMA domains is determined to be μm, which is slightly larger than those obtained both from the as-cast film and annealed film (~3.3 µm). The PMMA phase boundaries remain smooth and continuous across different pattern lines (A1), which is consistent with the AFM images on the solvent treated patterns (A3 and A4). Apparently, the flow of PS/PMMA during cavity filling does not generate appreciable domain breakups; neither does it involve considerable mass flow in the directions parallel to the lines. Moreover, no significant morphological evolutions were evident after the cavity filling, which is probably due to the slow mobility of both polymers at 150 C. At Timp = 180 C and 210 C, high fidelity of pattern replications was also achieved, as demonstrated by representative AFM images (B2 and C2). The asimprinted patterns (B1 and C1) show that significant morphological evolutions occurred during the NIL process at both 180 C and 210 C. At 180 C, PMMA domains are still highly correlated across different lines, but the phase boundaries become noticeably rough (B1). This trend becomes more evident for PS/PMMA patterns created at 210 C (C1). From the FFT image (inset of C1), the correlation

129 108 for PMMA domains along the pattern line direction is completely lost, while noticeable correlations remained at the direction perpendicular to the lines. Such anisotropic disintegration of the PMMA domains indicates that the mass transport that contributed to the domain breakup mostly occurred along the line direction due to the confinement of cavity walls. Clearly, both PS and PMMA become increasingly mobile with the increase of Timp, which enables the domain breakup during the imprint process. Figure 6.4 Images of S30/M70 patterns imprinted at three different temperatures, as marked on the left. From left to right, different columns correspond to optical micrographs of films after imprinting, corresponding topographic AFM images of the as-imprinted films, after the removal of PS with CLE, and after the removal of PMMA with AA. All AFM images are 10 µm 10 µm. Insets are the FFT images of the optical micrographs.

130 109 Compared with annealed planar films, at 180 C, ridge-valley structure did not form and the initially formed spinodal decomposition structure was kept (B1) even at such a high temperature. The confinement not only slowed down the kinetics of coarsening process, but also changed the pathway of the coarsening process by suppressing the vertical growth of PS domains. Therefore, domain coarsening (especially PS, since majority of it was within the pattern lines, not in the residual layer) can happen only along pattern line direction. The slight morphology change around phase boundary was the result of this confined movement. At 210 C, mass transport took place significantly along the pattern line direction so that the original PMMA spatial correlation completely disappeared. No isolated spheres formed such as those in the planar film, since there was not even the pathway to the ridge-valley structure. The remaining PS patterns after PMMA removal (A4, B4, and C4) are ~240 nm, which is ~45 nm taller than the as-imprinted pattern (A2, B2, and C2). In contrast, after the CLE dissolution, the maximum height of the remaining PMMA pattern was ~195 nm (A3, B3, and C3). The discrepancies indicate that the residual layers in the patterned S30/M70 were mostly PMMA. From the mass conservation consideration, a 45 nm residual layer thickness implies an initial film thickness of ~145 nm, which is close to the targeted 150 nm. This suggests that PMMA may prefer the extrusion of the mold (SiOx surface) and the Si substrate, 148,149 the two surfaces confining the residual layer, which will drive the PS out of the residual layer or isolate them into smaller domains. Upon AA treatment, the residual layer

131 110 is thus completely removed. Note that such preferential segregation of PMMA will effectively cause the increase of PS composition within the patterned lines. However, a quantitative assessment is difficult from the AFM images shown in Figure 6.4. Additionally, one would also expect that the PMMA will segregate underneath the patterned lines, which could lead to the complete dissolution/removal of the patterns after acetic acid treatment. Indeed, further dissolution of the remaining pattern in A4 reveals a layer of PMMA with thickness around nm underneath the PS in the patterned lines (image not shown). The fact that these PMMA did not dissolve in acetic acid within the short time is likely attributed to both the protection from the PS layer and the strong interactions between the PMMA and the substrate. Interestingly, the AFM images of the AA dissolved pattern (A4, B4, and C4) revealed that some thin polymer lines (marked by the arrows) left within the PMMA rich domains. These lines were nm tall and nm wide. These remaining lines were identified to be undisclosed PMMA, through the following two experiments. First, these patterns were completely removed by using more rigorous dissolutions with AA at 50 C. Second, we determined the Tg of these small lines to be close to the bulk value (126 C) of the PMMA, by using the newly developed Nanothermal analysis (Nano-TA). 150 Note that the dissolution of PMMA thin films supported on substrate is a complex phenomenon. The origin of the reduced solubility of PMMA in the 30/70 patterns remains unclear. Xue et al., have shown that favorable interactions between PMMA and hydroxyl groups on the substrate

132 111 surface can be strong enough to prevent PMMA dissolution. 142 Similar reduction in the solubility of high molecular weight PS in toluene after being imprinted at high temperatures was also observed recently Morphology of patterned S50/M50 films Figure 6.5 Optical micrographs (OMs) and topographic AFM images of 150-nm S50/M50 blend films annealed at three different temperatures as marked on the left. Listed in the first column are the optical micrographs of surface morphologies (A1, B1, and C1). The dark and bright domains are the PS and PMMA, respectively in all three OM images. The second and third columns are topographic AFM images corresponding to surface morphologies, and interface morphologies after PS was selectively dissolved by cyclohexane (CLE). All AFM images are 10 µm 10 µm in size.

133 112 Phase evolution for blend film at this composition also followed typical coarsening process as shown in Figure 6.5. After surface leveling, ridge-valley structure formed then broke up into droplets due to Rayleigh-like instability. The difference from S30/M70 films is that the as-cast film contained PMMA cylinders instead of PMMA island domains due to the reduced PMMA fraction. At 180 C, some of the PMMA cylinders merged and formed larger concave domains while other PMMA cylinders maintained their original shape mostly. As these domains grew large enough, the inner part would be relatively flat. These concave domains were surrounded by sharp edges and the edges behaved as growing fronts. As phase evolved, the concave domains spread out and finally became the matrix. Similar to S30/M70 films, PMMA interface coverage increased as temperature increased. Similar morphological analysis on the patterned S50/M50 films is shown in Figure 6.6. Successful pattern replications were also achieved at all three Timps (A2, B2, and C2), forming topographically uniform patterns onto phase separated PS/PMMA (A1, B1, and C1). The solvent dissolution experiments revealed that the residual layers (~35 nm) of the S50/M50 at all three Timps mostly consisted of PMMA, similar to the S30/M70 samples. Within the patterned lines, the morphological evolutions progressively increase with the increase of Timp. Specifically, PMMA domains are mostly isolated within the lines (A3 and A4) at Timp = 150 C, while growing large enough to bridge across the lines and interconnect with the residual layer at Timp = 180 C (B3 and B4). Such morphological change is more evident at Timp = 210 C (C3 and C4).

134 113 Correspondingly, the average height of the remaining PMMA pattern, in regions where the PS was removed, decreased from ~60-80 nm (A3) and nm (B3), to nm (C3) as Timp increased from 150 C and 180 C to 210 C. Figure 6.6 Images of the S50/M50 patterns imprinted at three different temperatures, as marked on the left. From left to right, different columns corresponding to optical micrographs of films after imprinting, corresponding topographic AFM images of the as-imprinted film, after the removal of PS with CLE, and after the removal of PMMA with AA. All AFM images are 10 µm 10 µm. Insets are the FFT images of the optical images. The dark and bright regions in the optical images correspond to the PS and PMMA phases, respectively. These observations suggest that the PMMA domains within the lines gradually coalesce with the residual layer and drive the original continuous PS

135 114 matrix into encapsulated blocks within the line. The surface coverage of the PS along the lines decreases simultaneously. Such a phase inversion in the patterned PS/PMMA blends is driven by the stronger segregation tendency of PMMA into the residual layer and coarsening within the cavity walls. Intriguingly, the encapsulated block morphology formed via anisotropic growth of the PMMA domain in the S50/M50 pattern (C1, C3, and C4) is similar to that developed via the anisotropic breakup of the PMMA domain in the S30/M70 pattern (C1, C3, and C4). Such similarity maybe the result of the complex interfacial interactions among PS, PMMA, and the cavity walls Morphology of patterned S70/M30 films Excellent pattern replications were also achieved for the S70/M30 at all Timps (Figure 6.7 A2, B2, and C2). The optical micrographs of the as-imprinted structures cannot distinguish the two-phase morphology and only line-space structure was captured. This is because the domain size was beyond the resolution of microscope at 150 C (A1) and well-encapsulated continuous tube-like structures formed at 180 C and 210 C (B1 and C1). At 150 C, the lateral feature size of the intrinsic phase-separated structure was almost the same as that of as-cast film due to the low kinetics. After NIL at this temperature, these features almost maintained as the cases of other compositions, and were therefore beyond the resolution of microscopy. As at this stage, the final morphology was the result of the superposition of the intrinsic phase structure and the template geometry, line structure could be identified in micrograph (A1). With

136 115 the selective dissolution experiments, dispersed PMMA domains were found distributed randomly across the patterned lines and the residual layers (A4). The depths of holes left in the residual layer after AA-treatment were found to be around 50 nm. This is consistent with the residual layer thickness obtained from other imprinting temperatures to be discussed later (B4 and C4). Some of these domains were deformed by the squeeze flow during cavity filling (A3). A3 and A4 justified once more that the PS/PMMA morphology is almost identical to that in the as-cast films, which is consistent with the S30/M70 and S50/M50 films patterned with Timp = 150 C. In particular, because the PMMA domain ( µm, Figure 6.2 C3 and C4) is smaller than the widths of the lines and residual layer (~0.42 µm), they are completely isolated. Significant coarsening of the PMMA domains occurred with increase of Timp, within both the patterned lines and the residual layers. Because of the confinement effect, the regular long-range PS coarsening was inhibited. The concentration of PMMA is so low that there is no real PMMA-rich region (although at Timp = 150 C and 180 C, after CLE-treatment, sharp corners remained, the major component of the pattern line was PS as characterized by B4 and C4) throughout the whole process. Instead of developing blocks leaving space for PMMA to form PMMA-rich region, PS shaped into continuous tube within the cavity encapsulated by PMMA. At Timp = 180 C, after the PMMA removal, the remaining PS lines become noticeably thinner, indicating that the top of the lines are covered by a thin layer of PMMA (Figure 6.7 D). This is consistent with the observations that most of the line

137 116 surfaces are intact after the CLE dissolution (B3). In addition, the remaining PS lines are significantly smoother than that in A4. Moreover, PMMA domains start to coalesce within the residual layer, forming larger holes after the AA dissolutions (B4). Such morphological evolutions become more dramatic in the patterns created at 210 C. After the PMMA removal (C4), the remaining PS pattern has uniform height ~245 nm, also suggesting that almost all residual layers (~50 nm) were consisted of PMMA. In addition, the remaining PS lines grow thicker than that observed at B4, suggesting that more PS moved into the patterned lines due to volume conservation. This is also consistent with the change of PMMA in the residual layer. As more and more PMMA migrated into residual layer, more and more PS would move into the cavities contributing the thicker lines at 210 C. The residual layer beneath PS domains was estimated to be ~40 nm based on the volume fraction of both polymer components. PS formed tube-shaped structure covered by PMMA. These PS tubes and the covering PMMA layer are uniformly-formed so that under either solvent treatment, the integrity of the tube was preserved. As temperature went up from 150 C to 210 C, less and less PS was removed (A3-C3). More and more PS was protected by gradually formed PMMA covering layer due to the preferential wetting of PMMA on SAM-treated mold surface. At 210 C almost no lines were dissolved by CLE (C3), indicating that the entire surfaces of the lines are covered by PMMA. However, the thickness of this PMMA layer seems to be thinner than that observed in the patterns with Timp = 180 C, which means the morphology of encapsulated PS

138 117 Figure 6.7 Images of the S70/M30 patterns imprinted at three different temperatures, as marked on the left. From left to right, different columns corresponding to optical images of films after imprinting, corresponding topographic AFM images of the as-imprinted film, after the removal of PS with CLE, and after the removal of PMMA with AA. All AFM images are 10 µm 10 µm. Insets are the FFT images of the optical images. D, The cross-sectional profiles of B2, B4, and C4, as marked in the AFM images.

139 118 dominated the overall morphology and affected the thickness of the PMMA covering layer. The observation that the PMMA surface coverage in the S70/M30 system is the most dramatic among all three compositions is rather surprising because PMMA concentration is the lowest in this system. Such a counterintuitive phenomenon could be driven by multiple factors. First, the size of the PMMA domains in the S70/M30 films/patterns is the smallest, creating the largest interfacial tension that favors domain coarsening. Second, the average spacing between the PMMA domains is only ~350 nm, which is significantly smaller than the other two compositions. This leads to the smaller diffusion length required for PMMA to bridge and cover the surfaces both in the residual layer and within the lines. In addition, PS is the less viscous component, and the higher composition of PS is likely to make the matrix more mobile, thus easier for PMMA to diffuse within the matrix Substrate surface energy effect on morphology Figure 6.8 shows the surface and interface morphologies of as-cast films, as well as films annealed at 150 C, 180 C and 210 C on neutral surface, compared with that on SiOx surface. Similar to PS48K/PMMA21K, the as-cast films had similar morphologies on SiOx surface and on neutral surface (A1 and A2). However, rather than the phase evolution pathway change, large molecular weight PS/PMMA underwent a phase evolution with more obvious slowed-down kinetics. On SiOx

140 119 surface, a typical coarsening process goes through a series of stages of surface leveling, formation of ridge-valley structure, breakup of ridge-valley structure into droplet, and stabilization of the droplet growth. Characteristics of each step can be reflected on both surface and interface morphology. For 150-nm S50/M50 films on SiOx surface, morphologies at 150 C, 180 C and 210 C (C2-C4 and D2-D4) were representative for the three stages mentioned above. Compared with films on SiOx surface, the interface and surface morphology changes were less dramatic on neutral surface. At 150 C, some of the PMMA cylinders started to coalesce while other PMMA cylinders maintained their original shapes and sizes. This coalescence was different from that on SiOx surface and there was no vertical movement during the merging. At 180 C, the coalescence continued and there formed spinodal decomposition-like PMMA domains. At 210 C, the spinodal decomposition-like PMMA domains grew larger while other PMMA cylinders maintained mostly as before. The PMMA area fraction at the interface increased from ~34% for the as-cast film to ~56% after the film was annealed at 210 C. The increase of PMMA area fraction at the interface indicates that there were less and less PS at the bottom part of the film, meanwhile more and more PS moved to the topper layer of the film contributing more roughness at the surface and long-range order fluctuation, which can be justified from either surface morphology (A3 and A4) or surface roughness. However the surface roughness increase was slow. After surface leveling, the surface roughness dropped from ~15 nm to less than 5 nm. Further annealing did

141 120 not increase the surface roughness much, at 180 C it was ~9 nm and at 210 C it was only ~10 nm, even less than the roughness in the as-cast film. Figure 6.8 Phase evolution of 150-nm free-surfaced planar S50/M50 films annealed for 30 min at different temperatures on neutral surface. The first column is AFM topographic images for as-cast films. The second, third and fourth columns are AFM topographic images for films annealed at 150 C, 180 C and 210 C, respectively. The first row reveals the surface morphology while the second row shows the interfacial morphology after selective removal of PS by cyclohexane. The third and fourth rows are for planar films on SiOx. All AFM images are 10 µm 10 µm in size.

142 121 In contrast, on SiOx surface, after the drop to less 5 nm after surface leveling, the surface roughness increased drastically due to the formation of ridge-valley structure and the coarsening of the broken-up isolated droplets, more than 50 nm at 180 C and close to 80 nm at 210 C. A comparison of film surface/interface roughness on neutral surface with on SiOx surface is shown in Figure 6.9. Figure 6.9 Surface and interface roughnesses of S50/M50 planar films annealed at different temperatures on SiOx surface and neutral surface. A 150-nm S50/M50 film on neutral-surfaced substrate was nano-imprinted at 180 C with exactly the same processing parameters as those on silicon wafer. The differences observed between the two-types of substrates are similar to that discussed in Chapter 4, despite the differences in the kinetics due to the difference in the molecular weight of the PS. Despite the dramatic difference in the free planar films between on SiOx surface and on neutral surface, it is surprising to see that the nano-imprinted films have very similar morphology on different substrate with different surface energy.

143 122 This is illustrated in Figure As discussed before, on SiOx surface, successful pattern replication was achieved (A2), PS formed encapsulated block domains within channels (B2 and C2), PMMA formed wetting layer at the bottom of the film and a covering layer on the top of the pattern. All of these could apply to the patterned films on neutral surface (A1-A3). In other words, qualitatively the morphologies are essentially the same. Three reasons might be the explanations for the above observation. First, the morphologies of as-cast films were similar, which means the starting points of the phase evolution under 2D-confinement in both cases are similar. Second and more importantly, the nanoimprint process is the key factor that determines the final morphology. It prevents the vertical growth of PS domains during coarsening process, and therefore changed the phase evolution pathway by avoiding the formation of ridge-valley structure and isolated PS droplets. It guides the domain coarsening along the pattern line direction, which resulted in encapsulated elongated domains for lower PS concentration (i.e., S30/M70, S50/M50) and continuous PS tube-like lines for higher concentration (i.e., S70/M30) in channels. The surface energy of the template treated with SAM also played an important role in the determination of the final morphology, especially the top part of the patterned film which could subsequently influence the material distribution beneath it. All of these are the same for a certain films on these two different substrates. Third, even if the PMMA wetting layer was eliminated for ~60-nm PS48K/PMMA21K films, there could be still a PMMA wetting layer in a thicker film

144 123 or PS/PMMA with larger molecular weight such as 150-nm S50/M50 as thicker bottom layer or slow kinetics caused by polymer long chains would make it hard for the dewetting to happen. Figure 6.10 Morphology of 150-nm S50/M50 films nano-imprinted at 180 C for 30 min on neutral surface. A1, B1 and C1 are topographic AFM images for asimprinted film, film etched by cyclohexane and film etched by acetic acid, respectively, as marked on top of the figure. A2, B2 and C2 are taken from Figure 6.6 and are included for comparison. All AFM images are 10 µm 10 µm in size. 6.5 Film thickness dependence of morphology The films examined in previous sections are thick enough for complete filling during the NIL process. However, the morphology evolution actually occurred not only within the channels of the imprinting template, but also at the regions beneath the template, i.e., within the residual layer. In this section, we briefly investigated the film thickness dependence of morphology, i.e., the situations when there are not

145 124 enough polymeric materials to fill into the channels of the template and when the residual layer has thickness comparable to the pattern height. Figure 6.11 shows the morphology of patterned ~75-nm film (S50/M50, at 180 C for 30 minutes) on substrate with neutral surface. The template was treated with the low surface energy SAM as usual. The morphology obtained with SiOx substrate is essentially the same. Figure 6.11 Morphology of as-imprinted ~75-nm S50/M50 film (at 180 C for 30 minutes) on neutral surface and size distribution of elevated domains. A, Topographic AFM image of pattern after NIL; B, Gaussian-like apparent area distribution of elevated domains. The incomplete filling caused dewetting of polymers on the surface of the template. The resulted broken-line structure is shown in Figure 6.11A. Some regions of the template were filled completely while other regions were filled only slightly at the entrance of the cavities. The corresponding pattern height was ~195 nm at the fully filled region indicating locally successful pattern replication. In the incomplete filling region, there existed a ~10-nm tall pattern with dual-peak. This observation is very similar to the profile of the dual-peak deformation mode

146 125 appeared near the entrance of the cavities, which is attributed to the stronger shear thinning of polymer resist around template corners. 83 The size or the length of elevated domains roughly follows a Gaussian distribution as shown in Figure 6.11B. This broken-line structure is not unique to polymer blend, as pure PS under similar condition could also form this type of structure. 145 Clearly, the dewetting tendency in the incompletely filled cavity is dominant, regardless of whether the resist is pure polymer or blends. Figure 6.12 Topographic AFM images of surface and interface morphologies of patterned ~75-nm S50/M50 film (210 C for 30 minutes). A, as-imprinted morphology; B, morphology of PMMA after PS was selectively dissolved by CLE; C, morphology of PS after PMMA was selectively removed. All AFM images are 10 μm 10 μm. Within the blocks of the lines, the spatial distribution of the PS/PMMA component is very similar to that of ~150-nm films discussed above. Specifically, PS blocks are encapsulated by PMMA (Figure 6.12). Therefore the dewetting of PS inside the channels is likely dominant during the confined coarsening process and the broken lines were imposed on the coarsened morphology.

147 126 Figure 6.13 Topographic AFM images and cross-sectional profiles of interfacial morphologies of patterned ~300-nm S50/M50 films after PS was selectively dissolved by CLE. A, B and C are topographic AFM images of films imprinted for 30 minutes at 150 C, 180 C and 210 C, respectively. All AFM images are 10 μm 10 μm. D is the cross-sectional profiles corresponding to the marks drawn on A, B and C. For films with thickness of ~300 nm, the general observations in ~150-nm films still apply. The major difference arises from the thicker residual layer of ~300- nm films, which contains both PS and PMMA instead of PMMA mostly. We take S50/M50 films to show the phase evolution as an example. The as-cast morphology is similar to that of ~150-nm films which is elevated PMMA domains dispersed in

148 127 PS matrix. The diameter of PMMA domains in ~300-nm films is ~1.3 μm, larger than ~0.6 μm in ~150-nm films. After imprinting, the morphologies were more complicated and the interface boundary was less clear than the thinner films we have seen before. At Timp= 150 C, the morphology of as-cast film was preserved mostly. The circular or spherical PMMA domains could be observed from Figure 6.13A. As the domains had diameter larger than the pattern periodicity of the imprinting mold, they were deformed such that some parts of the domain flew into the cavities of the mold while other parts were squeezed into the residual layer. The pattern trench regions in Figure 6.13A are rather rough. The variation in the roughness was caused by the thick residual layer of ~300-nm films. For the 150-nm films, the ~50-nm residual layer contains the PMMA wetting layer and small amount of PS. During the imprinting process, PS in the residual layer tends to be absorbed into PS-rich domains in the cavities (details shown in next section), while PMMA flew in replacing PS in the residual layer quickly. ~300-nm films have ~200- nm residual layer. The thick residual layer unavoidably includes a significant part of the PS matrix of the as-cast morphology. Also, the residual layer accommodates 2/3 of the overall blend materials. However, the PMMA volume fraction is only 50%. It is not possible for the residual layer to contain PMMA purely. As Timp was increased to 180 C, coarsening occurred both within pattern lines and residual layer. As show in Figure 6.13B, PMMA domains with full pattern height grew longer. Meanwhile, the PMMA initially belonged to elevated domains but squeezed into residual layer also coarsened along the pattern line direction. At Timp = 210 C,

149 128 the coarsening tendency observed at Timp= 180 C continued and encapsulated PS domains developed clearly. As shown in the cross-sectional profile (blue curve) in Figure 6.13D, encapsulated PS domains have height larger than the pattern height. This means PS in the residual layer could coalesce with the encapsulated PS-rich domains formed within the pattern line. Note that due to the preferential wetting of PMMA on the mold surface, at the entrances of the cavities PS might be partially protected by PMMA. This would cause selective dissolution with CLE to be less effective in removing PS within the residual layer, i.e., it is very likely that beneath the blue curve in Figure 6.13D, there exists significant amount of PS. This could be indeed confirmed by selective dissolution of PMMA to reveal PS as shown in Figure Figure 6.14 Topographic AFM images and cross-sectional profiles of interfacial morphologies of patterned ~300-nm S50/M50 films after PMMA was selectively dissolved by AA. A is the topographic AFM image of films imprinted for 30 minutes at 210 C. B is the cross-sectional profiles corresponding to the marks drawn on A. The coarsening process occurred within ~300-nm film suggests that there is no significant migration of PMMA from the pattern line region into the residual

150 129 layer. PMMA coarsens both within pattern line region and residual layer without substantial mass flow vertically. This is consistent with the conclusion that the vertical growth/coarsening of PMMA wetting layer is less dominant at the presence of external pressure. 6.6 A comparison between phase evolutions in thin films under 1-D confinement and 2-D confinement In this section, we demonstrate the similarity and differences between phase evolutions in thin films under our 1-D confinement and 2-D confinement through 150-nm thick S50/M50 films. Figure 6.15 A comparison of morphologies between pressurized ~150-nm S50/M50 film and patterned S50/M50 film. A and B are extracted from Figures 5.13 and 5.14, and they are topographic AFM images for films pressurized for 30 minutes at 180 C and 210 C, respectively. C and D are extracted from Figure 6.6, and they are topographic AFM images for films imprinted for 30 minutes at 180 C and 210 C, respectively. All AFM images are 10 μm 10 μm.

151 130 As we discussed in Chapter 5, under pressurization condition, the morphology of ~150-nm S50/M50 film behaved as if it was frozen (Figures 6.15A and 6.15B). Due to the suppressing of wetting layer growth, the coarsening is essentially a collision-based diffusion process, and the kinetics of phase evolution is extremely slow. In contrast, morphology of the patterned film evolved dramatically from 180 C to 210 C. Coarsening occurred significantly with in pattern lines. This is probably because the imprinting process disturbed the initial meta-stable state obtained during spin-coating, and broke the circular/spherical PMMA domains into small interconnected pieces with irregular shapes, which were highly energetically unstable. The coarsening occurred within the pattern lines in turn affected the residual layer and increase the overall mobility of the system. Figure 6.16 Schematic of major flow directions during coarsening under imprinting condition. For ~150-nm films, it turns out to be that the residual layer is made of PMMA mostly. The PMMA wetting layer indeed grew under nanoimprinting condition regardless the external pressure. This is likely due to the combined effect of PMMA preferential wetting tendency on the mold and the Ostwald ripening of PS. A schematic showing the flows that would cause PMMA to migrate to the bottom layer is illustrated in Figure The residual layer is ~50 nm thick at the

152 131 bottom of which is the PMMA wetting layer. As discussed in Chapter 5, the PMMA wetting layer after spin-coating could be of several tens of nanometers (no direct measurement, but ~10 nm after thinning at 180 C). The PS domain or layer beneath the trench of the mold is therefore at most ~40 nm. This thickness would be reduced when considering PMMA also wets the mold trench surfaces. And this wetting has tendency to drive PS out of the residual layer. Further, when PS in the residual layer connects to a PS-rich domain (~190 nm in height and ~417 nm in width) in the cavity, Ostwald ripening would drive the PS from the residual layer into the cavity to merge with larger PS-rich domain. In fact, such PS-rich domain always exists when we consider two dimensionally. If PS is relatively isolated in the residual layer, the dramatic coarsening in the channel could cause potential lateral flow in the residual layer to mobilize it. Finally, dewetting of PS on PMMA and/or trench surfaces within residual layer might also play a role in driven PS out of residual layer. In sum, the force that prevents the PMMA layer from growing due to the Laplace pressure from neighboring PS domains in the planer pressurization situation in Chapter 5 does not exist in the patterning condition for 150-nm films, because such PS domains are highly unstable and tend to flow to the channels/cavities of the mold.

153 Conclusions Figure 6.17 A transition of encapsulated PS domains from blocks to continuous threads as PS concentration increases from 30% to 70%. AMF images are extracted from figures shown before and are morphologies of AA-etched films imprinted at 210 C. All AFM images are 10 μm 10 μm. To conclude, we demonstrate the successful fabrication of topographically uniform patterns over phase separated polymer blend films, focusing on the morphological evolutions of the blend during the patterning process. The results demonstrate that for patterns created at relatively low temperature such as 150 C, the PS/PMMA morphology within the patterns is similar to that in the as-cast films. With increase of imprinting temperature, the morphological evolutions become more dramatic. The morphological evolutions at higher Timp (180 C and 210 C) vary with the compositions. In the S30/M70 patterns, the large PMMA domains start to break anisotropically within the cavities during imprinting. This leads to randomly distributed blocks or plugs of PS (or PMMA) domains within the lines. For the S50/M50 patterns, the anisotropic coarsening and PMMA preferential wetting during imprinting cause a phase inversion: the originally dispersed PMMA domains evolve into the continuous phase across the residual layer and patterned lines which drives the PS into isolated blocks within the patterned lines. For the

154 133 S70/M30, PMMA migrates to cover the surface of the patterned lines in addition to occupying the residual layer. This also leads to encapsulated long PS threads within the lines, which is different from the block morphologies observed in the other two compositions. Figure 6.17 highlights the transition from PS blocks to PS threads as PS concentration increases. These results demonstrate that the complex morphological evolutions under the nanoimprint conditions are due to the interplay of domain coarsening and preferential wetting of different components. By controlling the processing parameters and surface energy of the substrates and superstrates, designing different polymer blends, and adjusting their compositions, one can achieve uniform topographic surface patterns with chemically highly heterogeneous multicomponent polymers.

155 134 CHAPTER VII CONFINED REACTION-INDUCED PHASE SEPARATION UNDER SF-NIL CONDITIONS 7.1 Introduction In this last chapter, we investigated the physical confinement effect on the phase evolution of initially mixed binary composite NOA65/5CB (4-cyano-4'- pentylbiphenyl) using step-and-flash NIL (SF-NIL). The pattern fabrication procedure was essentially a confined reaction-induced phase separation process (RIPS) that is a common way of demixing polymer-liquid crystal mixtures. 32 In bulk films, a typical binary composite resulted from RIPS is the so-called polymerdispersed liquid crystal (PDLC), consisting of spherical liquid crystal droplets uniformly distributed in the polymer matrix The birefringence of liquid crystal offers PDLC devices the capability of displaying electro-optic responses, and plenty of studies have investigated the responses of various types of PDLC cells. 154, The most well-known types of RIPS methods include thermal-induced phase separation (TIPS) and polymerization-induced phase separation (PIPS). While TIPS

156 135 incurs phase separation by quenching an upper critical solution temperature (UCST) system, normally a thermoplastic polymer and liquid crystal, below its critical temperature; for PIPS, the UV irradiation-triggered polymerization of monomer causes liquid crystal to separate from the cross-linked network. For PIPS, Vorflusev and Kumar found that if the UV light had an intensity gradient, the phase separation pathway would be changed and the resulted morphology would be a stratified composite film (PSCOF) which consists of adjacent layers of liquid crystal and polymer. 162 Such a PSCOF-based device, if incorporated ferroelectric liquid crystal, can have switching time 100 times faster at low fields than conventional surface-stabilized devices. To date, the effect of geometry confinement on the RIPS is still unclear. PDLC and PSCOF cells typically have thickness on the order of several microns and confinement is almost nonexistent Practically, the photo-induced RIPS is completely compatible with the pattern replication process in SF-NIL. Here we study the influence of physical confinement on RIPS, and discover that unique embedded liquid crystal layers can form at conditions normally lead to PDLC. Section 7.3 illustrates the morphology arrangement of NOA65/5CB after confined RISP and Section 7.4 discusses the electro-optic responses of a liquid crystal cell formed by above-mentioned confined RISP. 7.2 Samples and techniques NOA65 (Norland Inc.) / 5CB (Sigma-Aldrich Inc.) mixtures with different compositions were prepared according to literature Confined reaction-induced

157 136 phase separation was realized by using SF-NIL with the NOA/5CB mixture. In a typical SF-NIL process, liquid resist (photoactive monomer) was squeezed into the cavities of imprinting template via combined pressure and capillary effect and then was cured into a rigid replica with UV irradiation. For SF-NIL, either the mold or the substrate has to be transparent for the UV radiation. In this study, transparent substrate was used, and the schematic diagram of the whole fabrication process is illustrated in Figure 7.1. Figure 7.1 A schematic illustration of confined reaction-induced phase separation process. PAA pattern was first made by TE-NIL. Then while NOA/5CB mixture was confined within PAA trenches mostly, SF-NIL was carried out to induce the phase separation of NOA and 5CB. Finally, after selectively dissolving PAA and 5CB, the NOA-5CB interface was revealed and characterized. A polyacrylic acid (PAA) sacrificial pattern was first prepared by conventional thermal embossing NIL. 169 The PAA solution (50 wt. % in H2O, Mw = 2 kg / mol, Sigma-Aldrich Inc.) was used as received without any further dilution. This highconcentration PAA solution created a thick PAA resist film after spin-coating and a PAA pattern with thick residual layer subsequently, which is beneficial to the dissolution process described later. The as-imprinted PAA pattern was used as a substrate to carry out SF-NIL on a nano-imprinter (Eitre 3, Obducat Inc.). The hydrophilic PAA is immiscible with the precursor solutions used. A drop of NOA/5CB mixture liquid was put onto the PAA substrate with a glass slide covering

158 137 on top of it. NIL was conducted at 40 C with a pressure of 4 MPa. The imprinting temperature 40 C was selected to be above the nematic-isotropic temperature (TNI) of 5CB, 35 C, in order to help liquid crystal align along the pattern line direction. 170 After the pressure had been applied for 1 minute, the precursors confined between the PAA pattern and the glass slides was under UV exposure for 10 minutes. The UV light induced the crosslinking of NOA and hence the two components phaseseparated. After the system was cooled down to the room temperature, the sandwiched NOA/5CB was put into methanol to dissolve PAA and 5CB to reveal the interface between NOA65 and 5CB. 171 AFM (Dimension 3100, Bruker) was then used to characterize the topographic profile of the interface. 7.3 Morphology characterization Figure 7.2 Phase-separated morphologies of NOA/5CB mixtures confined by 833 nm-width pattern. A, B, and C are topographic AFM images of NOA-5CB interfacial morphologies after 5CB was selectively removed for NOA87/5CB13, NOA50/5CB50 and NOA30/5CB70 films, respectively. All AFM images are 5-µm wide. D is interfacial cross-section profiles for different compositions.

159 138 Figure 7.2 depicts the topographic AFM images of the phase-separated NOA/5CB with three different compositions: NOA87/5CB13, NOA50/5CB50 and NOA30/5C70. Pure NOA pattern had a periodicity of 833 nm, a line-space ratio of 1:1, and a depth of ~ 185 nm indicating high fidelity of pattern replication (see cross-section profile in Figure 7.2D, AFM image not shown). However, the crosssectional profile revealed that the height of the replicated pure NOA was ~5 nm smaller than the depth of the original PAA pattern or the depth of the master template possibly due to the shrinkage of the NOA after crosslinking. When the 5CB concentration was 50%, the remaining NOA pattern height after selective dissolution of 5CB decreased and line width at the top part of the pattern was slightly narrower than that of pure NOA pattern (Figures 7.2B and 7.2D). The reduction of the NOA pattern height (to ~120 nm) was resulted from the washed-away 5CB, and the corresponding region was occupied by 5CB after phase separation but before dissolution. Figure 7.3 Schematic cross-sectional profile of the PSCOF device structure. 4, cross-linked NOA; 5, liquid crystal; 6, UV source. 162 The phase-separated bilayer-like structure along with the smooth interface between NOA65 and 5CB suggest that the phase separation process was similar to

160 139 what happens during the formation of phase-separated composite films (PSCOF). 162, In a PSCOF, NOA crosslinks first at the area where is closest to the UV light (Figure 7.3). At that region UV intensity is the strongest. Then the interface between cross-linked NOA and uncross-linked NOA behaves as a growing front. It grows towards a direction away from the UV light source as more and more NOA monomers are gradually cross-linked. Meanwhile, the growing front keeps pushing uncross-linked NOA monomers and liquid crystal molecules away towards the bottom of the composite film until NOA is completely cured. The key condition that triggers the formation of PSCOF is the UV irradiation intensity. As reported by Qian et. al., there exists a transition from PDLC structure to PSCOF structure with UV intensity decreasing from mw/cm 2 to mw/cm 2, other conditions being the same. 172 Under the UV irradiation, two competing mechanisms that determine the final morphology compete with each other. One is the abovementioned vertical diffusion-based gradual crosslinking process of NOA. The other one is the conventional reaction-induced phase separation occurred almost homogeneously across the whole composite film. When UV irradiation intensity is strong, the fast kinetics of phase separation makes the phase separation process dominant. The crosslinking of NOA at different places of the film occurs spontaneously resulting in uniformly dispersed liquid crystal domains in the cross-linked NOA matrix. When the UV irradiation intensity is not strong enough to lead the phase separation process ahead of the gradual crosslinking of NOA, liquid crystal molecules diffuse towards the opposite side of

161 140 the UV source. Typical UV light intensity that used to fabricate PDLC cells is on the order of several tens of milli-watts per centimeter squared. 163,165 However, the UV light intensity at the sample stage of the nano-imprinter used in this study is 60 mw/cm 2, a typical intensity used for PDLC cell fabrication with small LC domains. Normally, under such a strong UV intensity and the concentration of 50% liquid crystal, PDLC-like structure is expected to form in micron-sized films. The current study probes how such UV induced phase separation occurs under the 2D confinement. 163,165,170 As shown in Fig. 7.2, the phase-separated NOA50/5CB50 film clearly showed that no liquid LC droplets were found, which implies that the conventional phase separation did not happen. The inhibition of the traditional phase separation was caused by the confinement effect. As the conventional reaction-induced phase separation occurs through a nucleation and growth mechanism. 158 The intrinsic randomly distributed nuclei could be disturbed by the existence of the physical confinement. On the other hand, the PAA might be a preferential surface for the 5CB, which leads to the fast segregation of 5CB on the PAA surface prior to or during the NOA crossliking. The width of the remaining NOA pattern line was narrower than the pure NOA pattern, with narrowness extending from the top part of NOA pattern to the bottom. Notice the bottom part of the NOA pattern was slightly wider than the top part. These might be caused by the complex convolution of the AFM tip geometry and the sample surface topography. But more likely, it was due to the uneven intensity distribution of UV light in these 417 nm-wide slots. It is widely reported

162 141 that the pattern created by photo-lithography has similar profiles just because of the slightly non-uniform light intensity distribution For a negative photoresist, at the center part of such slot, the light intensity is stronger than those on two sides. Accordingly, the pattern height at the center is larger than two sides and the pattern line width goes wider moving away from the light. The nonuniformity was found to be caused by Fresnel diffraction. 178 Because the halfperiodicity of the pattern is very close to the wavelength, this effect became more obvious in this study. Another reason that could explain the shape of the crosssectional profile lies in the diffusion pathway of liquid crystal molecules. At the region closed to the PAA wall, the wall restricted the molecule movement directed towards it, and liquid crystal molecules were likely to collide with the wall (or even preferentially wet the PAA wall). As a result, the probability of collision among liquid crystal molecules increased, reducing the diffusion speed and increasing the density of the molecules. The accumulated liquid crystal molecules not only blocked NOA monomers from flowing towards the UV light source, they also reduced the UV light intensity of the region where they resided and thus further inhibited the crosslinking process of NOA. In contrast, near the center region of the slot, liquid crystal molecules had more freedom. They could diffuse towards any direction and more NOA monomers tended to flow to that region where eventually formed crosslinked NOA with larger height. Further the existence of the confining wall generated a friction between the NOA monomer/liquid crystal molecule and the

163 142 wall. It lowered the mobility of liquid at the vicinity of the wall and slowed-down the diffusion rate, which led to a smaller amount of cross-linked NOA. Liquid crystal occupies 51.2% of the pattern cross-section area, which is similar to the weight percentage of the 5CB in the mixture, but by coincidence. Note that the density of liquid crystal 1.58 g/cm 3 is larger than NOA 0.83 g/cm 3, so if assume that all the 5CB is accumulated at the bottom of the PAA trenches, there should be extra NOA cross-linked beyond the pattern region. There must be a covering layer next to the glass slide during the phase separation to accommodate these extra cured NOA because of the mass conservation. This covering layer could be revealed as a residual layer in the cross-section profile plot, with a thickness ~35 nm. When there was a smaller amount of 5CB in the mixture, i.e., 13%, the remaining NOA pattern height (~168 nm) was taller than that of NOA50/5CB50. Line width at the top part of the pattern was only slightly narrower than that of pure NOA pattern (Figures 7.2A and 7.2D). Cross-sectional profile shows that the area fraction of liquid crystal was 12.7%. Based on mass conservation, the thickness of the covering layer could be estimated as ~50 nm. This thickness is larger than that of NOA50/5CB50. The increase of this layer thickness was a result of decreasing the liquid crystal fraction. As liquid crystal has a lower viscosity than NOA monomer, NOA87/5CB13 mixture was estimated to have a higher viscosity. The increased viscosity provided less mobility for the squeeze flow occurred beneath

164 143 the imprinting template during NIL process. Less resist was squeezed out from the patterning region and a thicker residual layer of NOA was resulted. When the 5CB was the major component, i.e., 70%, the NOA-5CB interface was no longer as smooth as that observed in lower concentration systems. The remaining NOA surface, after removing 5CB, appeared rough. The amount of 5CB was so large that it even covered the surface of the plateau of the PAA substrate. As revealed in the cross-section profile plot, the straight line part indicating the PAA- NOA boundary did not exist at this composition. Different from the cases with other compositions, isolated 5CB individual layer in each PAA channel was connected by the 5CB layers on top of the PAA plateaus and formed continuous layer beneath cross-linked NOA. Figure 7.4 Topographical AFM images of surface morphology (A) for pure NOA pattern and NOA-5CB interfacial morphologies (B and C) after selective removal of 5CB for NOA85/5CB15 and NOA50/5CB50 films. All films were confined with 150 nm-width pattern and all AFM images are of size 1 µm wide.

165 144 Figure 7.4 showed the results of similar confined reaction-induced phase separation process for the NOA/5CB system, but with much smaller pattern size. The PAA template had a periodicity of 150 nm, line-space ratio of 1:1 and a depth of only 45 nm. For such a small pattern size, the sacrificial PAA pattern did not show the square-wave-like cross-section: neither the trench part of the pattern was flat nor did sharp corners appear on the top of the pattern. The patterned pure NOA pattern showed good replication of the PAA pattern despite a slightly smaller height due to shrinkage. As the 5CB concentration increased, both the height and the width of the NOA pattern revealed after the dissolution of 5CB decreased, consistent with the results obtained with 833 nm-width PAA pattern. Clearly, the conventional PDLC-like phase separation was suppressed in both cases examined. PSCOF-like structure formed due to the loss of competitiveness of the driving force of the PDLC-like phase separation. Regardless of the confinement size, similar morphology was observed: a top layer of cured NOA mimicking the substrate geometry, and a liquid crystal layer sandwiched between NOA layer and the substrate. The similarity provides a way of fabricating embedded crystal layer. The line-space template geometry used in this study could be extended to be any type of topography, such as a designed pattern. The confined phase separation process would deposit a liquid crystal layer immediately onto the top of the trench of a pattern with flexible thickness by varying the mixture composition.

166 Electro-optic response of ultra-thin confined phase-separated composite film (C- PSCOF) A schematic diagram of the fabrication process of ultra-thin C-PSCOF cell is illustrated in Figure 7.5. A pure NOA pattern was made first with the same process mentioned above, but instead of on a regular glass slide substrate, it was fabricated on an ITO glass. Subsequently, the pure NOA pattern was used as a substrate, similar to the PAA pattern shown in Fig. 7.1, for the deposition of the NOA/5CB precursor. Another piece of ITO glass was used as the capping layer on top of the NOA/5CB that was finally UV-crosslinked within the trenches of the pure NOA pattern. Since NOA is a good adhesive for bonding glass, after the second SF-NIL, the two pieces of ITO glass were well-bonded and the cell formed. Figure 7.5 A schematic illustration of fabricaiton of ultra-thin C-PSCOF (confined phase-separated composite film) cell. With a PAA pattern, pure NOA pattern was replicated by the same method as in Figure 7.1. Then NOA pattern was used as substrate to conduct SF-NIL to induce the confined phase seperation of NOA65 and 5CB forming C-PSCOF cell. We pick NOA50/5CB50 C-PSCOF cell as a representative cell to show the electro-optic responses. The input signal was AC 10 V with frequency 100 Hz. Barannik et al. reported that if there are ions, even of a tracer amount, i.e., %

167 146 presented in the liquid crystals, an internal electrical field due to the diffusion of the ions would rise in opposition to the externally applied electrical field (Figure 7.6B). 179 Unfortunately, the 5CB used in this study contains ~ 2% impurity which is most likely to include a significant amount of dissolved ions. Thus, the ion effect was expected in the electro-optic response of the cell. Figure 7.6 Ion effect in PDLC cell. A, Schematic of a bilayer of ion charges that produce an internal electric field opposite to the external field; B, Ion concentration dependence of ion effect on dynamics of the optical response of PDLC film. 179 Separation of the cation and anion layers at the boundaries between liquid crystal and NOA is expected when the external electric field is applied (Figure 7.6A). The rates of both the ion layer formation (essentially a charging effect) and the randomization (discharging) are determined by the diffusion rate of ions in the confined liquid crystal domains. These rates and the associated electro-optic responses were found to be temperature-dependent and concentration-dependent. 179 When the temperature is low such that the diffusion rate of the ions is close to zero, the PDLC cell has very similar behavior to the PDLC cell with pure liquid crystal. In contrast, when the diffusion rate is very high, in the voltage-on state, the

168 147 internal electric field will quickly cancel out the external electric filed, making the transmittance gradually decay to zero after it reaches the peak value, while in the voltage-off state, the internal electric field comes to play solely and generates similar response as the external field does in the voltage-on state (Figure 7.6B). The experimental setup used to characterize the electro-optic response of the cell is shown in Figure 7.7. A laser beam was focused on a photodiode passing through the C-PSCOF cell. Meanwhile, electrical signal generated by a function generator was applied to the C-PSCOF cell. The detected optic signal was recorded Figure 7.7 Experimental setup for the measurement of electro-optic response of C- PSCOF device. by an oscilloscope, which synchronized the applied electrical signal. Figure 7.8A captures one period of the input and output signals. Notice that the zero point in the time axis is not the exact starting point of the whole response; it is in an intermediate stage and when input signal jumped. It is well-known that liquid crystal molecules tend to align along the physically guided direction. 162,170 In the current cell, the liquid crystal molecules aligned along the grating line direction due to the physical confinement when the external electrical field is off. The orientation of the cell was such that the pattern grating direction was perpendicular to the

169 148 incident plane (the plane parallel both to the polarization direction of the incident light and the propagation direction of the incident light) as shown in Figure 7.8B. When no AC voltage was applied, the alternating 5CB and NOA blocks along the polarization direction of the incident light have alternating refractive index with 5CB block 1.7 and NOA block This causes strong diffraction of the light inside the cell and therefore the transmittance is reduced (the total incident energy is distributed among different diffraction orders, and transmittance is the 0 th order diffraction). We define the state of liquid crystal molecules at this time initial state and will refer to this term in the following discussion. When the cell is in the voltage-on state, the liquid crystal molecules will tend to align along the electrical field direction, which is, if ideally, parallel to the propagation direction of the incident light and perpendicular to the pattern grating lines. The refractive indices of the liquid crystal block and the NOA block are now almost identical, both being around The transmittance would be correspondingly higher, compared with that in the initial state. The corresponding state of liquid crystal molecules is defined as 90-degree state for convenience. When AC voltage is applied, liquid crystal molecules in this C-PSCOF cell aligned between these two ideal states. Raw data were measured in voltage and reported as obtained, but could be converted into and explained in terms of transmittance. We use them interchangeably. When liquid crystal molecules were aligning towards 90-degree state (or initial state), the diffraction became weaker (or stronger) and the transmittance increased (or decreased).

170 149 Figure 7.8 A typical electro-optic response of C-PSCOF cell corresponding to an applied electric field (AC voltage of 10 V, 100 Hz) along with schematics of alignment directions of the liquid crystal molecules at different stages (A), and a schematic illustration of the ideal alignment states of liquid crystal molecules in voltage-on state and voltage-off state (B). We take the electro-optic response in Figure 7.8A as an example to describe molecular mechanism that give rise to the time-dependent transmittance through

171 150 the cell. First, ions started to diffuse when voltage was applied with an external field (E0), with cations and anions moving towards opposite directions, which induced an internal electrical field (Ei). As the separation of the charges further increase, the strength of Ei increased gradually, which caused the strength of the effective electric field (E0 - Ei) to decrease correspondingly. The total electrical field reached an asymptotic values once all the ions were completely separated. At the same time, within the cell, liquid crystal molecules started to align with the effective electric field direction that gradually reduced. The orientation time of the LC is expected to be much faster than that of the diffusion of the ions, because it only requires the 5CB molecules to rotate locally. In comparison, the ions need to diffuse across a gap of ~100 nm. Therefore, the transmittance increases upon the external electrical field due to the quick orientation of 5CB molecules, and reached it maximum value for this given voltage and frequency (state 1 in Figure 7.8A, we call it max-aligned state#1). Note that the driving forces (E0 - Ei) decreases as the ions diffusion continuous. Once the maximum transmittance is reached, the liquid crystal molecules started to align back due to the lack of strong effective electrical field to maintain the current state (state 2 in Figure 7.8A), driven by the elastic energy of the LC molecules associated with the different aligned states. The internal electric field kept increasing, which effectively act as a slow switching off process, even the external electrical field is still on. Immediately before the voltage turned into the off-state, the liquid crystal molecules were not in the initial stage (Figure 7.8A).

172 151 This could be due to the upper limit of the intensity of the internal electric field, or because the AC frequency was so high that the decay time was not long enough for liquid crystal molecules to align back. In the presented case, it was more likely that the latter was true as will be shown later that this state would vary with frequency. Also the average external voltage acted on the cell was larger than the average internal voltage during the voltage-on stage. Correspondingly, the energy stored in liquid crystal molecules contributing towards the 90-degree state was larger than that towards the initial state. The overall energy guided the liquid crystal molecules to align along the external electric field direction to a certain extent, and a residual aligning was resulted (state 3 in Figure 7.8A, residual aligning state#1). After all, imagine if there were no ion effect, the liquid crystal molecules could have stayed in the 90-degree state. When the cell was in the voltage-off state, internal electric field generated during the previous half-period came to play solely and it itself was the effective electric field. The internal electric field drove the liquid crystal molecules to align towards its direction, which was opposite to the external electric field direction (state 4 in Figure 7.8A). The intensity of the internal electric field kept decreasing as ions diffused back to a randomly distributed state. At some point, the aligning reached its limit when correspondingly the output signal was at its peak value (state 5 in Figure 7.8A, max-aligned state#2). Then the electric field was not strong enough to maintain this aligned state. The liquid crystal molecules started to align back towards the initial state (state 6 in Figure 7.8A). However, the initial state

173 152 was not reached at the end of voltage-off state. A residual aligning remained again (state 7 in Figure 7.8A, residual aligning state#2). The next-cycle voltage-on state started with this residual aligning state#2 (from state 7 to state 8 in Figure 7.8A) and the process discussed above rolled on and on. Interestingly, the output signals in the voltage-on state and the voltage-off state were almost the same. This was the result of the balancing of the two residual aligning states at the end of voltage-on state and voltage-off state. Again notice that the signals presented in Figure 7.8A was not taken from the first cycle, but one of the cycles after the cell had responded for some time. It was taken from a stabilized state. When the first 10-V square wave input was applied, the starting point was the initial state. Compared with the residual aligning state#1, there would result in a larger residual aligning at the end of the first voltage-on state. This in turn would make internal electric filed appear smaller in the following voltage-off state. So did the residual aligning at the end of the voltage-off state. This residual aligning was rolled into the second cycle and made the voltage-on state output signal weaker than the first one, i.e., the residual aligning at the end of voltage-on state was larger than that in the first cycle. These tendencies continued cycles after cycles: the residual aligning at the end of voltage-on state kept increasing while the residual aligning at the end of voltage-off state kept decreasing, until both of them stabilized, or even converge to each other resulting in two duplicate output signals within one single input cycle as in the case presented. Notice that the final signal

174 153 captured in the voltage-off state is not necessarily the same as that in the voltageon state if the effect of the internal electric field was not strong enough. The maximum output values appeared when the aligning angle was maximized or the liquid crystal molecules were at max-aligned states. The minimum values were observed at the end of voltage-on state or voltage-off state where the residual aligning states were defined. Figure 7.9A shows the output signals with different input voltages. When the input voltage was relatively small, e.g., far below 6 V, the aligning degree was negligible, and the transmittance change was not obvious enough to be observed. The output signal captured was simply a flat line. As the input voltage increased, the overall output signal plot was lifted up and the minimum was raised to a certain base level. The larger the input voltage was, the more the residual aligning, and the higher the base level. Figure 7.9B shows the positive correlation between the maximum/minimum output voltage and the input voltage. Within each cycle, minimum voltage (it did not matter whether it is the one at the end of the voltage-on state or the one at the end of the voltage-off state, actually these two local minimum values converged to each other for all the input voltages) increased with the input voltage. The growth followed an exponential-like increasing process within the input range. The convergence of the two minimum values implied that the input voltages were not strong enough to make the liquid crystal molecules reach the 90-degree state. Otherwise the decay time (from max-aligned state to residual aligning state within each voltage-on/off state) within voltage-on state would be shorter than that in

175 154 Figure 7.9 Voltage dependence of electro-optic responses of C-PSCOF cell. A, electro-optic responses of C-PSCOF cell to different applied voltages ranging from 6V to 10V. B, plot of relationship between maximum/minimum output value and the input voltage. voltage-off state, which was not true. Similar to the minimum value, the maximum output signal also increased as the input value increased. The stronger the input was, the more the liquid crystal molecules would align and therefore the larger value of transmittance. The amplitude of the output also increased with the input voltage, which means the voltage measured at max-aligned state has larger variation than that measured at residual aligning state, and hence the

176 155 transmittance is more sensitive to the voltage change when the state was more closed to 90-degree state. Figure 7.10 Frequency dependence of electro-optic responses of C-PSCOF cell. A, a comparison of electro-optic responses under 200-Hz 10-V applied voltage with 100- Hz 10-V applied voltage. B, plot of relationship between maximum/minimum output value and input frequency. The base level of output signal is also AC frequency dependent. Figure 7.10A demonstrates a direct comparison between the output signals obtained under 100- Hz input and 200-Hz input. This frequency dependence was caused by the difference in the decay time of detected transmittance (output voltage). The larger the frequency was, the shorter the decay time, the more the max-aligned state would be maintained and the more residual aligning would be resulted. For the case

177 156 of 200-Hz input, after the max-aligned state was reached, there was less time for internal field to cancel out the external field. The overall ion effect within voltage-on state was smaller, which resulted in larger residual aligning and in turn lifted the base level of the output signal. A boarder view of the frequency dependence is illustrated in Figure 7.10B. As frequency increased, the base level as well as the peak value both increased. Notice that the amplitude of the output signal increased first then decreased. When the frequency was 120 Hz, the amplitude of the output signal was the largest. This is the balanced result between the extent of the residual aligning and the sensitivity of transmittance to the aligning angle change. When the frequency was higher than 120 Hz, the residual aligning ruled. The increase of frequency would shorten the decay time and reduced the difference between the maximum and minimum output signals. In contrast, when the frequency was below 120 Hz, the decay at the end of each voltage-on/off state varied less with frequency change. Thus, the reduction of the amplitude caused by the increase of the frequency could be easily compensated. The compensator was believed to be the sensitivity of transmittance to the aligning angle change, which implies that at higher base level (but still below that corresponding to ~120 Hz) where liquid crystal molecules could be aligned more parallel to the propagation direction of the incident light, transmittance tended to vary more. This was similar to what had been concluded at the end of the voltage dependence discussion, the transmittance was more sensitive to changes when the state was more closed to 90-degree state.

178 Conclusions To conclude this section, confined reaction-induced phase separation was realized by nano-imprinting initially mixed NOA65/5CB syrup. Both sub-micron-sized and nano-scale-sized templates were used to examine the size effect of confinement on phase-separated morphology. The mixture was confined within the trenches of the NIL templates mostly while the phase separation was occurring. Although there existed a covering/residual layer with a thickness of several tens of nano-meters depending on the composition or the viscosity of the NOA65/5CB mixture, the phase boundary was found to be within the pattern region for most of the cases. Embedded liquid crystal layer, rather than isolated domains were obtained and sandwiched between the cross-linked NOA and the NIL templates. Even though the intensity of the UV light in SF-NIL was strong enough to create PDLC-type morphology, it did not happen under the confinement because the conventional phase separation which was supposed to rise through nucleation and growth mechanism was inhibited. The top-down gradual crosslinking process of NOA became dominant during the overall phase separation process and resulted in PSCOF-type bilayer-like structure in the NOA/5CB film. The confining walls also influenced the morphology in a short range. Along the center line of the trench of the polymer substrate, the thickness of the liquid crystal layer was smallest, increasing towards the wall directions. This confined reaction-induced phase separation procedure has a potential to fabricate arbitrary shape of liquid crystal layer with tunable thickness. Further, this procedure was applied to fabricate ultra-thin C-PSCOF cell. The internal electric field caused by the ions within liquid crystal domains complicated the measured transmittance. Specifically, the residual aligning accumulated at the end of each voltage on/off state was found to be the key to explain the base level as well as the overall shape of the output

179 158 signal. The seemingly weird doubled frequency of the output signal found roots in it. The external voltage dependency and frequency dependency also correlated with the residual aligning. When the applied voltage was increased (decreased), so did the residual aligning. As a result, both the base level and peak value raise (fell). Similar situation also applied to frequency. But different from the voltage dependence, where the sensitivity of transmittance to the cell (designated by the amplitude of the output signal) increased monotonously, there existed certain frequency (~120 Hz) to which the transmittance was most sensitive.

180 159 CHAPTER VIII SUMMARY In this study, we examined how the surface tension and interfacial interaction between constituent polymers of polymer blend and confinement environment influence the phase evolution, and revealed unique micro or nanostructures by demixing of polymers under imprinting process. Figure 8.1 summarizes the main work of this thesis. The classical polymer blend pair PS/PMMA was investigated extensively in thin film cases with different conditions. We started from the phase evolution of PS/PMMA film on SiOx surface, which is well known to form surface-relief structure due to the dictation of the preferential substrate-wetting of PMMA. We systematically examined the kinetics for the first time in literature and found that the resulted PS relief structures on the PMMA wetting layer varied with the blend composition, transitioning from capillaryfluctuation-mediated breakup to random nucleation with the increase of PS concentration. By tuning the substrate surface energy neutral to both PS and

181 160 PMMA, the formation of the conventional PMMA wetting layer was prevented upon annealing and both PS and PMMA domains could directly contact with the substrate. Compared with chemically homogeneous substrate such as SiOx or neutral surface, chemical pattern has stronger impact on the structure formation of thin film on it. It could alter the intrinsic phase-separated structure dramatically and guide the phase separation according to the geometry of the pattern during spin-coating process. With NIL-based chemical pattern fabrication approach which provides smaller pattern periodicity than the traditional micro-contact-printingbased method, we were able to obtain unique hierarchical structures and revealed the morphological transition to perfect transfer case. Figure 8.1 Schematic summary of the thesis including thin films on chemically homogenous substrates (SiOx surface, neutral surface), on chemical pattern, between two parallel rigid substrates, and under traditional TE-NIL/SF-NIL conditions.