NANO LETTERS
Oriented Mesoporous Organosilicate Thin Films
2005 Vol. 5, No. 10 2014-2018
Erik M. Freer, Leslie E. Krupp, William D. Hinsberg, Philip M. Rice, James L. Hedrick, Jennifer N. Cha, Robert D. Miller, and Ho-Cheol Kim* IBM Research DiVision, Almaden Research Center, 650 Harry Road, San Jose, California 95120-6099 Received August 2, 2005
ABSTRACT Coassemblies of block copolymers and inorganic precursors offer a path to ordered inorganic nanostructures. In thin films, these materials combined with domain alignment provide highly robust nanoscopic templates. We report a simple path to control the morphology, scaling, and orientation of ordered mesopores in organosilicate thin films through the coassembly of a diblock copolymer, poly(styrene-b-ethylene oxide) (PS-b-PEO), and an oligomeric organosilicate precursor that is selectively miscible with PEO. Continuous films containing cylindrical or spherical pores are generated by varying the mixing composition of symmetric PS-b-PEO and an organosilicate precursor. Tuning interfacial energy at both air/film and film/substrate interfaces allows the control of cylindrical pore orientation normal to the supported film surfaces. Our method provides well-ordered mesoporous structures within organosilicate thin films that find broad applications as highly stable nanotemplates.
Because of the broad potential end uses of thin-film materials with engineered architecture, new tools to manipulate structure within these films are highly attractive. A bottomup approach to structured thin films is by the use of selfassembling molecules such as block copolymers that show unique microphase separated morphologies on nanoscopic length scales.1,2 The ordered periodic structures in organic block copolymer thin films have been used, for example, as etching masks for transferring patterns into solids to fabricate nanoscopic devices such as capacitors3 and memories.4 Alternatively, block copolymers can be used as structure directing agents in the synthesis of nanostructured inorganic materials. Since the pioneering work of Kresge et al.5 to create ordered mesoporous silica, numerous approaches have been developed using low or high molecular weight surfactants or amphiphilic block copolymers.6-9 The prevalent route to mesoporous silica using amphiphilic block copolymers or surfactants is through sol-gel chemistry. Efforts to generate continuous films with accessible porosity have been pursued as well;10,11 however, a simple method to prepare continuous inorganic films with controlled orientation of cylindrical pores through self-assembly has not been achieved. Here, we demonstrate the formation, scaling of domain size, and orientation control of cylindrical domains in porous organosilicate thin films from simple binary mixtures of an amphiphilic block copolymer and an oligomeric organosilicate precursor. * Corresponding author. Tel: 408-927-3725; fax: 408-927-3310; e-mail:
[email protected]. 10.1021/nl051517h CCC: $30.25 Published on Web 09/17/2005
© 2005 American Chemical Society
As shown schematically in Figure 1a, the porous structures in this study are generated through the coassembly of binary polymer mixtures of a block copolymer of polystyrene and poly(ethylene oxide) (PS-b-PEO) and an organosilicate precursor, silsesquioxane (SSQ), which is selectively miscible with PEO. The relative volume fraction of the PS and PEO + SSQ phases, which determines the morphology of the coassembled nanostructures, is controlled simply by varying the volume fraction of each block of PS-b-PEO and the mixing composition of PS-b-PEO and SSQ mixtures. Thin films (typically ∼300 nm thick) of the mixtures are exposed to solvent vapor for a given time and heated to form an organic-inorganic nanohybrid. Cross-linking of SSQ occurs above 150 °C. Because SSQ cross-linking preserves the structure, the morphology of the final porous film (created by further thermal treatment at 450 °C) is analogous to that of the nanohybrid.12 We focus on cylindrical and spherical morphologies generated using symmetric PS-b-PEO. Height-contrast AFM images of porous, SSQ thin films on the native oxide of silicon wafers are shown in Figure 1b and c. Long cylindrical pores oriented parallel to the surface are observed in Figure 1b, whereas hemispherical pores are shown in Figure 1c. We controlled the PS-b-PEO/ SSQ mixing ratios as 50/50 (wt/wt) and 30/70 for cylindrical and spherical pore morphologies, respectively. The internal microstructures of the porous films were evaluated by TEM with thin cross-sectional specimens prepared using a focused ion beam (FIB, FEI 830XL dual beam at 30 kV). As shown in the cross-sectional TEM micrographs in Figure 1d and e,
Figure 1. Block copolymer and organosilicate coassembled structures. (a) Schematic illustration of coassembled structures. The coassembly morphology is determined by the relative volume fractions of the PS and PEO + SSQ phases; heating to above 300 °C results in a porous, nanostructured SSQ. (b and c) 1 µm × 1 µm height contrasted AFM images (10-nm height scale) of the porous films containing a cylindrical (b) and a spherical morphology (c). The inset in b is the SAXS profile of the film. (d and e) TEM cross-sectional micrographs of films with a cylindrical (d) and a spherical (e) morphology (SSQ occupies darker regions). The film thicknesses are ∼285 nm. The molecular weight of PS-b-PEO used for both films is 19 000 g/mol. Samples for b-e were annealed at 20 °C in chloroform for 15 h.
well-defined cylindrical and spherical pores are observed throughout the entire film thickness of the porous SSQ. The cylindrical pore morphology was confirmed with small-angle X-ray scattering (SAXS). A synchrotron X-ray source of 8.9 keV was used to obtain transmission SAXS profiles of the porous SSQ films on 80-µm thick, double-sided polished silicon wafers. As shown in the insets of Figure 1b, the SAXS profiles (PS-b-PEO molecular weight ) 8600 Dalton) show strong peaks originating from interpore scatterings. The relative positions of the peaks to the first-order peak (q/q*, q ) (4π/λ)sin θ, where λ is wavelength, θ is the angle between the scattered photon and transmitted beam, and q* is q at the first-order peak) correspond to 1, x3, x4, x7, and so forth, indicating hexagonally packed cylindrical pore structures. Figure 2 emphasizes the dynamics of domain orientation, as propagated from the two interfaces (air/film and film/ substrate) of the supported films. Cross-sectional TEM micrographs of the porous SSQ films, where cylindrical pores were generated using a symmetric PS-b-PEO (Mn,PS ) Mn,PEO ) 9500 Dalton) as a function of annealing time in chloroform Nano Lett., Vol. 5, No. 10, 2005
vapor. After annealing, samples were subsequently heated to generate porous structures. As shown in Figure 2a, annealing for 15 h results in 2-3 layers of cylindrical pores oriented parallel to the surface at both interfaces, which is attributed to the preferential interaction of the PEO + SSQ phase with both interfaces. Cylindrical domains oriented almost normal to the surface are observed in the middle of the film, suggesting the sample is not an equilibrium state. A plan-view TEM of the middle section of the film confirms a hexagonally packed cylindrical pore structure as shown in the inset of Figure 2a. The number of layers of parallel cylindrical pores at the two interfaces increases with increasing annealing time in chloroform vapor as shown in Figure 2b-d, corresponding to 67, 150, and 314 h, respectively. After 314 h annealing, the cylindrical pores orient parallel to the surface throughout the entire film thickness with exceptional longrange order. It is noted that the shape of pores is deformed because of the FIB process for TEM sample preparation.13 Scaling the dimension of block copolymer microdomains with molecular weight is well documented theoretically and experimentally.14-17 Here, we find that the cylindrical pores 2015
Figure 2. Cross-sectional TEM micrographs of the porous SSQ films containing cylindrical pores. The orientations of the pores are shown as a function of annealing time: (a) 15 h, (b) 64 h, (c) 157 h, and (d) 314 h in chloroform at 20 °C. The PS-b-PEO molecular weight is 19 000 g/mol. After 15 h, 2-3 layers of cylindrical pores orient parallel to the air/film and film/substrate interfaces. The inset in a is a plan-view TEM micrograph of the middle section of the film after etching top and bottom with an Ar ion mill, exposing perpendicular cylindrical pores of d ≈ 20 nm. With further annealing, the number of cylindrical pore layers oriented parallel increases (b-d).
Figure 3. Scaling behavior of the porous domain size with PS-b-PEO molecular weight. The block copolymers used are PS3b-PEO3, PS3.8-b-PEO4.8, PS9.5-b-PEO9.5, PS23-b-PEO20, and PS59b-PEO71 where the numerical subscript × 103 gives the block molecular weight in g/mol.
generated by this coassembly approach scales with the molecular weight of the PS-b-PEO block copolymers that coassemble with SSQ. Figure 3 shows a plot of the centerto-center distance, L0, of the cylindrical pores determined from AFM images as a function of PS-b-PEO molecular weight. The volume fraction of PS was kept at 0.22. The L0 value varies from 14 to 93 nm with varying the molecular weight of PS-b-PEO from 6000 Dalton to 130 000 Dalton. 2016
It follows power law behavior with molecular weight, and the scaling coefficient was determined as 0.64, which is in good agreement with the coefficient of 2/3 predicted from mean field theories in the strong segregation limit.14-16 The orientation of cylindrical pores normal to a surface is appealing especially when the substrate is directly accessible via a path from the film surface. Orientation control of cylindrical domains, however, is not trivial. Surfaces with controlled interfacial energy,18 strong external fields,19 and solvent evaporation20 have been used to achieve normal orientation of cylindrical domains of block copolymers. The number of organic block copolymers that show normal orientation of cylindrical domains in thin films is limited. For inorganic thin films, successful generation of supported, continuous films containing cylindrical mesopores oriented normal to the surface through self-assembly has not yet been reported. In this study, we adjust the interfacial energies at two interfaces by varying the substrate surface chemistry and the solvent vapor used for annealing. Vapor phase deposition of alkoxysilanes or deposition of a thin Au layer was used to modify the substrate surface. We find that substrate surface modification alone (either alkoxysilanes or Au) is insufficient to change the orientation of the cylindrical domains under chloroform vapor annealing. To adjust the interfacial energies at the air/film interface concomitantly, we used a mixed solvent vapor. Figure 4a and b shows a height-contrasted AFM image and a FESEM micrograph of the porous SSQ film prepared on a (3-aminopropyl)triethoxysilane (APTES) treated silicon wafer and annealed in a vapor of a 50/50 (v/v) chloroform/octane liquid mixture. The images show hexagonally packed holes of ∼12 nm diameter, indicating cylindrical pores oriented normal to the surface. A fast Fourier transform of the AFM image shown in the inset of Figure 4a indicates hexagonally packed cylindrical pores with few defects. The internal structure of the sample is shown by a cross-sectional TEM micrograph in Figure 4c. It is clear that the cylindrical pores (the brighter areas) surrounded by cross-linked SSQ (the darker areas) are oriented normal to the surface. The normal orientation of cylindrical pores suggests that, on the APTES surface under a chloroform + octane vapor environment, the interfacial energy of PS and PEO + SSQ domains becomes nearly equivalent at both vapor/film and film/substrate interfaces. The neutrality of interfacial energy is likely due to absorption of octane in different concentrations into each domain. However, more investigation of the interfacial properties of swollen polymers (or polymer mixtures) is necessary to understand the physics of domain orientation of block copolymers under solvent vapor annealing. Similar to the APTES-coated surface, we find that the cylindrical domain orients normal to the substrate on 3-nm-thick Au-coated surfaces (Supporting Information, Figure S1), which opens the possibility of using this porous template for the Au-catalyzed high-temperature growth of inorganic nanomaterials. In summary, we demonstrate the formation of ordered mesoporous organosilicates thin films with controlled morphology, domain size, and orientation. The use of a mixed solvent vapor successfully controls the orientation of cylinNano Lett., Vol. 5, No. 10, 2005
Figure 4. Perpendicular orientation of cylindrical pores after annealing for 48 h in a chloroform + octane vapor at 20 °C. (a and b) 1 µm × 1 µm height contrasted AFM image (10-nm height scale) and FESEM micrograph of the surface. PS-b-PEO with a MW of 8600 g/mol was used. The inset in a is the Fourier transform of the AFM image. (c) TEM cross-sectional micrograph of the porous film shown in a and b. The darker and brighter regions correspond to SSQ and cylindrical pores, respectively. The film thickness is ∼250 nm.
drical pores in organosilicate thin films normal to the modified substrates. This approach provides a simple route to uniform, large area coatings of nanostructured organosilicates with controlled morphology that promises numerous end uses as a highly stable nanotemplate. Experimental Section. Diblock copolymers of polystyrene and poly(ethylene oxide) (PS-b-PEO) with molecular weights ranging between Mn ) 6000 and 130 000 g/mol were purchased from Polymer Source, Inc. and used as received. The organosilicate precursor, silsesquioxane (SSQ), is a copolymer of methyl trimethoxy silane and tetraethoxy silane with an approximate molecular weight of 2000 g/mol. Block copolymer solutions of 10 wt % in toluene were mixed with 10 wt % solutions of SSQ copolymers in propylene glycol propyl ether (PGPE). Thin films were prepared by spincasting the mixture solutions onto silicon wafers that were cleaned with a UV/ozone cleaner (UV-300H, SAMCO Inc.) at 100 °C. The surface energy of native oxide surfaces were modified with (3-aminopropyl)triethoxysilane (APTES) using a vapor phase silane deposition system (YES-1224, Yield Engineering Systems Inc.). Alternatively, 3-nm Au layers were deposited on clean silicon wafers. The spun-cast films were vacuum-dried at 20 °C for 24 h and then annealed under chloroform or chloroform + octane vapors at room temperature for varying amounts of time. After solvent vapor annealing, the wafers were heated to 450 °C (5 °C/minute) Nano Lett., Vol. 5, No. 10, 2005
and the temperature is held constant for 2 h under nitrogen to simultaneously cross-link SSQ and decompose PS-b-PEO. The surface topography is investigated using atomic force microscopy (AFM) in tapping mode using standard silicon cantilevers (Dimension 3100, Digital Instruments). Transmission electron microscopy (TEM) is used to characterize the cross-sectional thin-film structure. TEM samples are prepared with a focused ion beam (FIB) using a FEI 830XL dual beam at 30 kV. A Topcon 002B TEM operated at 180 kV was used to examine the FIB samples. Acknowledgment. E.F. would like to thank the National Science Foundation Center for Polymeric Interfaces and Macromolecular Assemblies for financial support (CPIMA grant NSF-DMR-0213618). We thank J. Frommer at IBM Almaden Research Center for valuable assistance with the AFM. Portions of this research were carried out at the Stanford Synchrotron Radiation Laboratory, a national user facility operated by Stanford University on behalf of the U.S. Department of Energy, Office of Basic Energy Sciences. The SSRL Structural Molecular Biology Program is supported by the Department of Energy, Office of Biological and Environmental Research, and by the National Institutes of Health, National Center for Research Resources, Biomedical Technology Program. We thank M. Toney, H. Tsuruta, and I. Smorsky for help with the SAXS experiments. 2017
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NL051517H
Nano Lett., Vol. 5, No. 10, 2005
