Structure Formation in Two-Dimensionally Confined Diblock

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Structure Formation in Two-Dimensionally Confined Diblock Copolymer Films P. Mu¨ller-Buschbaum,*,† M. Wolkenhauer,‡ O. Wunnicke,§ M. Stamm,§ R. Cubitt,| and W. Petry† TU Mu¨ nchen, Physik-Department, LS E13, James-Franck-Str. 1, 85747 Garching, Germany, Max-Planck-Institut fu¨ r Polymerforschung, Ackermannweg 10, 55128 Mainz, Germany, Institut fu¨ r Polymerforschung Dresden e.V., Hohe Str. 6, 01069 Dresden, Germany, and ILL, BP 156, F-38042 Grenoble, France Received March 26, 2001. In Final Form: June 6, 2001 The surface topography and the chemical morphology of poly(styrene-block-paramethylstyrene) diblock copolymers confined within isolated drops are investigated. The drop geometry imposes a two-dimensional spatial restriction. With scanning force microscopy, the film thickness dependence of the mean drop distance as the basic topographical feature is determined. In addition to large drops, a second smaller droplet structure is observed among the drops. The comparison of grazing incidence small-angle scattering using X-rays and neutrons shows that inside the drops the diblock copolymer orients into perpendicular lamellae with respect to the substrate surface. Compared to the bulk, the lamellar spacing is stretched parallel to the substrate.

Introduction In the past, most experimental investigations of polymers considered chains in three dimensions. In terms of applications, this is the most natural situation; however, many recent high-tech applications demand two-dimensional polymer systems. Examples of substrate-supported polymer films are protective and lubricating coatings, adhesives, high-performance composite materials, and microelectronic encapsulants and dielectrics.1,2 In these thin films, the deviation from bulk behavior increases with decreasing film thickness. With decreasing film thickness, the spatial restriction gets a growing influence on the physical properties of the polymer films.3,4 Polymers confined in spaces smaller or comparable to their typical dimensions, like the radius of gyration of the unperturbed chain, dramatically differ in important physical properties from the equivalent bulk material. Due to the confinement key parameters, for example, the glass transition temperature,5-12 the thermal expansion13 or the chain movement14-17 are changed. Changes result from struc* Corresponding author. † TU Mu ¨ nchen. ‡ Max-Planck-Institut fu ¨ r Polymerforschung. § Institut fu ¨ r Polymerforschung Dresden e.V. | ILL. (1) Physics of Polymer Surfaces and Interfaces; Sanchez, I., Ed.; Butterworth-Heinemann: Stoneham, MA, 1992. (2) Polymer Surfaces and Interfaces II; Feast, W. J., Munrom, H. S., Richards, R. W., Eds.; John Wiley & Sons: Chichester, 1993. (3) Pai-Panandiker, R. S.; Dorgan, J. R.; Pakula, T. Macromolecules 1997, 30, 6348. (4) Binder, K. Annu. Rev. Phys. Chem. 1992, 43, 33. (5) Keddie, J. L.; Jones, R. A. L.; Croy, R. A. Faraday Discuss. 1994, 98, 219. (6) Forrest, J. A.; Dalnoki-Veress, K.; Dutcher, J. R. Phys. Rev. E 1997, 56, 5705. (7) Laschitsch, A.; Bouchard, C.; Habicht, J.; Schimmel, M.; Ru¨he, J.; Johannsmann, D. Macromolecules 1999, 32, 1244. (8) Torres, J. A.; Nealey, P. F.; de Pablo, J. J. Phys. Rev. Lett. 2000, 85, 3221. (9) Schwab, A. D.; Agra, D. M. G.; Kim, J. H.; Kumar, S.; Dhinojwala, A. Macromolecules 2000, 33, 4903. (10) Kim, J. H.; Jang, J.; Zin, W. C. Langmuir 2000, 16, 4064. (11) Fryer, D. S.; Nealey, P. F.; de Pablo, J. J. Macromolecules 2000, 33, 6439. (12) Forrest, J. A.; Mattsson, J. Phys. Rev. E 2000, 61, R53.

tural deviations in the conformation18-21 as well as in modified dynamics.22-25 Many theoretical and experimental investigations are limited to amorphous homopolymers which depict the most simple system for fundamental questions. In addition to the pure topological constraint, the interaction between polymer molecules and the boundary atoms gets a growing importance.26,27 Thus, turning from homopolymer systems toward more complex systems such as polymer blends or diblock copolymers increases the degree of complexity. In addition to dewetting, which is a common energetically driven structure formation process in the case of homopolymers,28-32 further (13) Wu, W.; van Zanten, J. H.; Orts, W. J. Macromolecules 1995, 28, 771. (14) Zheng, X.; Rafailovich, M. H.; Sokolov, J.; Strzhemechny, Y.; Schwarz, S. A.; Sauer, B. B.; Rubinstein, M. Phys. Rev. Lett. 1997, 79, 241. (15) Lin, E. K.; Kolb, R.; Satija, S. K.; Wu, W. Macromolecules 1999, 32, 3753. (16) Manias, E.; Chen, H.; Krishnamoorti, R.; Genzer, J.; Kramer, E. J.; Giannelis, E. P. Macromolecules 2000, 33, 7955. (17) Buenviaje, C.; Ge, S.; Rafailovich, M.; Sokolov, J.; Drake, J. M.; Overney, R. M. Langmuir 1999, 15, 6446. (18) Kuhlmann, T.; Kraus, J.; Mu¨ller-Buschbaum, P.; Schubert, D. W.; Stamm, M. J. Non-Cryst. Solids 1998, 235-237, 457. (19) Jones, R. L.; Kumar, S. K.; Ho, D. L.; Briber, R. M.; Russell, T. P. Nature 1999, 400, 146. (20) Kraus, J.; Mu¨ller-Buschbaum, P.; Kuhlmann, T.; Schubert, D. W.; Stamm, M. Europhys. Lett. 2000, 49, 210. (21) Brulet, A.; Boue, F.; Menelle, A.; Cotton, J. P. Macromolecules 2000, 33, 997. (22) Kremer, K.; Grest, G. S.; Carmesin, I. Phys. Rev. Lett. 1988, 61, 566. (23) Zheng, X.; Sauer, B. B.; van Alsten, J. G.; Schwarz, S. A.; Rafailovich, M. H.; Sokolov, J.; Rubinstein, M. Phys. Rev. Lett. 1995, 74, 407. (24) Anastasiadis, S. H.; Karatasos, K.; Vlachos, G.; Manias, E.; Giannelis, E. P. Phys. Rev. Lett. 2000, 84, 915. (25) Mu¨ller, M.; Wittmer, J. P.; Barrat, J. L. Europhys. Lett. 2000, 52, 404. (26) Dietrich, S. In Phase Transitions and Critical Phenomena; Domb, C., Lebowitz, J. L., Eds.; Academic: New York, 1988; Vol. 12. (27) Israelachvili, J. N. In Intermolecular and surface forces, 2nd ed.; Academic Press: London, 1991. (28) Brochard-Wyart, F.; Redon, C.; Sykes, C. C. R. Acad. Sci., Ser. II 1992, 19, 314. (29) Reiter, G. Phys. Rev. Lett. 1992, 68, 75. (30) Lambooy, P.; Phelan, K. C.; Haugg, O.; Krausch, G. Phys. Rev. Lett. 1996, 76, 1110.

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processes such as phase separation33 and microphase separation34 become possible. In the case of diblock copolymer films, the thin film geometry favors the creation of lamellae in contrast to the bulk system. The lamellae in thin films are oriented parallel to the solid support.35-40 With decreasing film thickness, the confinement gives rise to an interplay between the intrinsic length scale of the bulk structure and the geometry of the film.41-45 Consequently, transitions between phases of identical symmetry but different orientation with respect to the confining walls become energetically possible, and lamellar domains reorient from a parallel to a perpendicular arrangement.46-50 With respect to applications, these arrays of nanoscopic structures with a controlled spatial orientation of the microdomains are desirable. The perpendicular arrangement especially attracts considerable attention due to its potential use for nanolithographic templates. In the case of homogeneous thin films, several different experimental approaches focusing on the perpendicular arrangement of the lamellae were reported.46-50 In general, polymer chains in a thin film geometry tend to orient parallel to the substrate surface. Chains in a small near-wall region have an elongated, flattened structure.51 The structural deviations from the bulk behavior become more pronounced with decreasing film thickness and temperature.52 However, a thin film introduces only a spatial restriction in one direction. One way to introduce a second spatial constraint is the replacement of a thin film by a drop structure. Experimentally, this is realized by preparing pancake structures on top of a solid support. Due to the asymmetries in the parameters defining the pancake shape (height , diameter), the imposed confinement is asymmetric as well. In addition to the inherent lengths introduced due to geometry, important lengths introduced by the diblock copolymer molecules are the radius of gyration Rg and the bulk lamellar spacing L0. In terms of length scales, the film thickness l and the drop diameter D are key (31) Xie, R.; Karim, A.; Douglas, J. F.; Han, C. C.; Weiss, R. A. Phys. Rev. Lett. 1998, 81, 1251. (32) Mu¨ller-Buschbaum, P.; Gutmann, J. S.; Stamm, M. Phys. Chem. Chem. Phys. 1999, 1, 3857. (33) Mu¨ller-Buschbaum, P.; Gutmann, J. S.; Cubitt, R.; Stamm, M. Colloid Polym. Sci. 1999, 277, 1193. (34) Hamley, I. W. The Physics of Block Copolymers; Oxford University Press: New York, 1998. (35) Anastasiadis, S. H.; Russell, T. P.; Satija, S. K.; Majkrzak, C. F. Phys. Rev. Lett. 1989, 62, 1852. (36) Russell, T. P.; Coulon, G.; Deline, V. R.; Miller, D. C. Macromolecules 1989, 22, 4600. (37) Stamm, M.; Go¨tzelmann, A.; Giessler, K.-H.; Rauch, F. Prog. Colloid Polym. Sci. 1993, 91, 101. (38) Mansky, P.; Russell, T. P.; Hawker, C. J.; Mays, J.; Cook, D. C.; Satija, S. K. Phys. Rev. Lett. 1997, 79, 237. (39) Torikai, N.; Noda, I.; Karim, A.; Satija, S. K.; Han, C. C.; Matsushita, Y.; Kawakatsu, T. Macromolecules 1997, 30, 2907. (40) Vignaud, G.; Gibaud, A.; Gru¨bel, G.; Joly, S.; Ausserre, D.; Legrand, J. F.; Gallot, Y. Physica B 1998, 248, 250. (41) Turner, M. S. Phys. Rev. Lett. 1992, 69, 1788. (42) Geisinger, T.; Mu¨ller, M.; Binder, K. J. Chem. Phys. 1999, 111, 5241. (43) Geisinger, T.; Mu¨ller, M.; Binder, K. J. Chem. Phys. 1999, 111, 5251. (44) Tang, W. H. Macromolecules 2000, 33, 1370. (45) Morkved, T. L.; Jaeger, H. M. Europhys. Lett. 1997, 40, 643. (46) Huang, E.; Russell, T. P.; Harrison, C.; Chaikin, P. M.; Register, R. A.; Hawker, C. J.; Mays, J. Macromolecules 1998, 31, 7641. (47) Fasolka, M. J.; Harris, D. J.; Mayes, A. M.; Yoon, M.; Mochrie, S. G. J. Phys. Rev. Lett. 1997, 79, 3018. (48) Fasolka, M. J.; Banerjee, P.; Mayes, A. M.; Pickett, G.; Balazs, A. C. Macromolecules 2000, 33, 5702. (49) Mansky, P.; Russell, T. P.; Hawker, C. J.; Pitsikalis, M.; Mays, J. Macromolecules 1997, 30, 6810. (50) Tang, W. H. Macromolecules 2000, 33, 1370. (51) Baschnagel, J.; Mischler, C.; Binder, K. J. Phys. IV France 2000, 10, 9. (52) Rivillon, S.; Auroy, P.; Deloche, B. Phys. Rev. Lett. 2000, 84, 499.

Mu¨ ller-Buschbaum et al.

parameters. Comparing these parameters, it is evident that the mean drop diameters have to be below the optical resolution limit. Thus, real-space pictures are accessible only with scanning force microscopy (SFM). Scattering experiments were performed to improve the statistical relevance of the local SFM information. Comparing X-ray and neutron experiments, that utilize the deuteration of one block, gains the chemical morphology in addition to the pure topological one.53 Because the amount of polymer is extremely small, the common transmission geometry is less advantageous and grazing incidence small-angle scattering (GISAS) has to be applied. GISAS overcomes problems related with the small sample volume and the heterogeneity of the samples due to the drop structure.54 The article is structured as follows: The introduction is followed by an experimental section describing the sample preparation and the techniques used. In the section GISAS, a short theoretical background of the scattering technique used is given. The sections on results and discussion are followed by a summary and an outlook. Experimental Section Sample Preparation. The used poly(styrene-block-paramethylstyrene) diblock copolymer, denoted P(Sd-b-pMS), was prepared anionically (Polymer Standard Service, Mainz). The nearly symmetric P(Sd-b-pMS) has a fully deuterated polystyrene block and a protonated polyparamethylstyrene block. The degree of polymerization of the PSd block compared to the total chain is fPSd ) NPSd/N ) 0.47. From the molecular weight Mw ) 230 000 g/mol (narrow molecular weight distribution of Mw/Mn ) 1.08) and the polymer-polymer interaction parameter of PSd and PpMS χ ) A + B/T with A ) -0.011 ( 0.002 and B ) 6.8 ( 1 K,55 a value of χN ∼ 24.0 is calculated. Thus, the investigated system belongs to the strong segregation regime.56 In a first preparation step, thin homogeneous films were prepared out of a toluene solution by spin coating (1950 rpm for 30 s). The film thickness was controlled from a variation of the polymer concentration of the solution used during the spin coating57 and checked by X-ray reflectivity measurements. To investigate confined thin films, film thicknesses between 1.5 and 35 nm were prepared. Due to the instability of free-standing films in this film thickness range, the homogeneous films were prepared on a solid support. Because of their small surface roughness, native oxide covered Si(100) surfaces (MEMC Electronic Materials Inc., Spartanburg) were used. The cleaning (immersing for 15 min at 80 °C in an acid bath consisting of 100 mL of 80% H2SO4, 35 mL of H2O2, and 15 mL of deionized water, rinsing in deionized water, and drying with compressed nitrogen) of these substrates in advance of the film preparation ensures reproducibility of the reported results. In the second preparation step, the initially homogeneous P(Sd-b-pMS) films were stored under toluene vapor for 12 h. After this exposure time, the samples were quenched to ambient air and examined. Several identical samples were prepared and investigated, exhibiting the presented results. Specular X-ray Scattering. With a laboratory X-ray source (Θ - Θ reflectometer Seifert XRD 3003TT), reflectivity measurements of the samples after the first preparation step were performed. A Ge(110) channel cut crystal is used to monochromatize the beam (λ ) 0.154 nm). The sample is placed on a specially designed vacuum chuck and is measured under air. The reflectivity data of the films measured right after preparation exhibit well-pronounced fringes due to the small surface roughness of typically 0.5 nm. From a fit, the film thicknesses l were (53) Mu¨ller-Buschbaum, P.; Gutmann, J. S.; Stamm, M.; Cubitt, R.; Cunis, S.; von Krosigk, G.; Gehrke, R.; Petry, W. Physica B 2000, 283, 53. (54) Mu¨ller-Buschbaum, P.; Casagrande, M.; Gutmann, J. S.; Kuhlmann, T.; Stamm, M.; Cunis, S.; von Krosigk, G.; Lode, U.; Gehrke, R. Europhys. Lett. 1998, 42, 517. (55) Schnell, R.; Stamm, M. Physica B 1997, 234, 247. (56) Helfhand, E.; Wassermann, Z. R. Macromolecules 1980, 13, 994. (57) Schubert, D. W. Polym. Bull. 1997, 38, 177.

Structure in Confined Diblock Copolymer Films

Figure 1. Schematic picture of the experimental setup used in the grazing incidence small-angle scattering experiments. The incident angle onto the sample surface is denoted Ri. A two-dimensional detector is used to measure one complete scattering array. The horizontal gray bar indicates one typical cut in the out-of-plane direction. This cut contains the smallangle scattering information and is used for further data evaluation. The two-dimensional detector collects the specularly reflected intensity (incident angle equals exit angle) as well as the diffusely scattered intensity (incident angle different from exit angle). obtained.58-60 The density resembles the mean density of both components PSd and PpMS. Due to the weak scattering contrast between PSd and PpMS, a resolution of the internal order is quite difficult.61,62 Due to the large surface roughness and the lateral heterogeneity of the polymer films after the second preparation step, reflectivity measurements on these samples yield no further useful information despite the fact that surface roughness is large. Grazing Incidence Small-Angle X-ray Scattering. GISAXS measurements were performed at the BW4 beamline at the synchrotron HASYLAB (DESY, Hamburg). In contrast to the common transmission geometry, we employed a reflection geometry. The sample was placed horizontally on a two-circle goniometer with a z-translation table.54 The beam quality was optimized by using a setup of high-quality entrance cross-slits and a completely evacuated pathway. The selected wavelength was λ ) 0.138 nm. Due to the large sample-detector distance of 12.12 m, a resolution better than 2.98 × 10-3 nm-1 was achieved in the out-of-plane direction. The scattered intensity was recorded with a two-dimensional detector which consisted of a 512 × 512 pixel array (see Figure 1). At one fixed incident angle Ri, the two-dimensional intensity distribution can be cut in several vertical and horizontal slices with respect to the sample surface.63-65 The horizontal slices contain the small-angle (58) Parrat, L. G. Phys. Rev. 1954, 55, 359. (59) Born, M.; Wolf, E. In Principles of Optics, 2nd ed.; Pergamon Press: Oxford, 1964. (60) James, R. W. In The Optical Principles of the Diffraction of X-rays; OxBow Press: Woodbridge, CT, 1962. (61) Stamm, M.; Schubert, D. W. Annu. Rev. Mater. Sci. 1995, 25, 325. (62) Tolan, M.; Press, W. Z. Kristallogr. 1998, 213, 319. (63) Salditt, T.; Metzger, T. H.; Peisl, J.; Reinecker, B.; Moske, M.; Samer, K. Europhys. Lett. 1995, 32, 331.

Langmuir, Vol. 17, No. 18, 2001 5569 scattering information.32 For an improvement of the statistics, the intensity was integrated in the vertical direction by ∆qz ) (7.61 × 10-3 nm-1. The chosen incident angle Ri ) 0.534° is larger than the critical angles of the diblock copolymer as well as of the substrate material which enables an easy separation of the specularly and the diffusely scattered intensity. For further details concerning the beamline, see ref 66. Grazing Incidence Small-Angle Neutron Scattering. GISANS measurements were performed at the D22 beamline at the neutron reactor ILL (Grenoble). The used wavelengths were 0.6, 1.0, and 1.7 nm (wavelength selector, ∆λ/λ ) 10%). Details concerning the beamline are reported elsewhere.67 The basic concept of the experimental setup33 is equal to the one realized at the BW4 beamline described above. Extremely narrow crossslits with typical openings of mm, a large collimation distance, and evacuated flight paths were used. The scattered intensity at one fixed angle of incidence Ri is detected with a twodimensional detector. Due to the larger sample-detector distance of 17.66 m, a resolution between 4.45 × 10-3 and 1.91 × 10-3 nm-1 was achieved at a wavelength of λ ) 0.6 or 1.7 nm. Details in the setup like the vertically placed sample, the chosen incident angle, or the decreased pixel array (128 × 128) of the detector differ. Again, statistics are improved by integrating the intensity (∆qz ) (1.91 × 10-2 nm-1) and the small-angle scattering information is extracted from horizontal slices (with respect to the sample surface). A schematic picture of the experimental GISANS setup is shown in Figure 1. A typical slice used for the measurement of the GISAS data is depicted with the gray bar. Scanning Force Microscopy. Real-space pictures of the sample surface are obtained from SFM measurements using a PARK Autoprobe CP atomic force microscope. Micrographs were recorded at different sample positions using scan ranges from 80 µm × 80 µm down to 1 µm × 1 µm. Operating the SFM in noncontact mode minimizes a tip-induced sample degradation. The silicon gold coated conical cantilevers have resonant frequencies of about f ) 72.5 kHz and a spring constant of =2.1 N m-1. All measurements were performed in air at room temperature. From the raw data, the background due to the scanner tube movement is fully subtracted to determine the values of the root mean square (rms) roughness over the complete scan area. In addition to the rms roughness, which displays a statistical information perpendicular to the sample surface, a statistical information parallel to the surface is obtained from the power spectral density function (PSD).68,69 The PSD is calculated from the SFM height data by a 2D Fourier transformation and radial average of the isotropic Fourier space data. Due to the different scan ranges in real space, the PSDs cover different intervals in the reciprocal space. Thus, a combination of PSDs related to different scan ranges enlarges the covered interval in reciprocal space as compared to one individual PSD. In the following, the combined PSD data is called the master curve. The master curve is equivalent to a scattering signal and thus shows the existence of a most prominent in-plane length scale which might be present within the resolvable range. If a distinct peak is present in the master curve, the most prominent in-plane length Λ is extracted from its position. With the rms roughness and the master curve, the sample surface is statistically described.

Grazing Incidence Small-Angle Scattering In principle, grazing incidence small-angle scattering is an adaption of common small-angle scattering (SAXS) to surface and thin films by replacing the transmission by a reflection geometry.54,63,70 The sample surface is defined (64) Salditt, T.; Metzger, T. H.; Peisl, J.; Goerigk, G. J. Phys. D: Appl. Phys. 1995, 28, A236. (65) Salditt, T.; Metzger, T. H.; Brandt, Ch.; Klemradt, U.; Peisl, J. Phys. Rev. B 1995, 51, 5617. (66) Gehrke, R. Rev. Sci. Instrum. 1992, 63, 455. (67) In Guide to Neutron Research Facilities at the ILL; Bu¨ttner, H. G., Lelievre-Berna, E., Pinet, F., Eds.; ILL: Grenoble, France, 1997; p 32. (68) Gutmann, J. S.; Mu¨ller-Buschbaum, P.; Stamm, M. Faraday Discuss. 1999, 112, 285. (69) Mu¨ller-Buschbaum, P.; Gutmann, J. S.; Stamm, M. Macromolecules 2000, 33, 4886.

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as the (x,y) plane. The incidence beam is directed along the x-axis with an incident angle Ri. The (x,z) plane denotes the plane of incidence and reflection, and thus the condition for specular scattering is given by qx ) qy ) 0 and qz > 0, with the scattering vector b q ) (qx, qy, qz).62 In Figure 1, the specular condition (Ri ) Rf) is fulfilled at the position of the specular peak. Diffusely scattered intensity is observed for qx * 0 or qy * 0. Consequently, the 2D detector basically contains diffusely scattered information. GISAS in observed along the out-of-plane direction and satisfies the condition qy * 0. It is observable in horizontal cuts of the 2D intensity distribution as shown in Figure 1 (gray bar). A deviation from an ideal flat surface is requested to give rise to a diffusely scattered signal. A lateral disorder of the ideal flat surface introduced by the surface structure can be described in terms of a height-height correlation function

C(r b) ) 〈h(r b+b F ) h(F b)〉bF

(1)

which calculates the statistical average over all pairs of surface points separated by the distance b r.64 It gives rise to a GISAS signal. In the distorted-wave Born approximation (DWBA), the differential cross section is given by65,70

dσ Cπ2 ) 4 (1 - n2)2| Ti|2| Tf|2 F(q b) dΩ λ

(2)

where C is the illuminated surface area, Ti,f are the Fresnel transmission functions, and F(q b) is the diffuse scattering factor. The Fresnel transmission functions act only as overall scaling factors in the GISAS geometry, since Ri and Rf are fixed. If the incident or exit angle Ri,f is equal to the critical angle Rc(A) of the material A, the transmission functions have a maximum, which is called the Yoneda peak.71 Therefore, out-of-plane cuts at the positions Rf ) Rc(A) increase the scattering contribution of the material (A).72 For N identical and centrosymmetrical objects with a random orientation, the diffuse scattering factor can be approximated70

F(q b) ∼ NS(q b) P(q b)

(3)

to depend on the form factor of the individual objects P(q b) and to depend on the structure factor S(q b). A mathematical description of the form factor depends on the type of object like sphere, cylinder, or slab, while the structure factor directly yields the most prominent in-plane length ξ. Thus, directly from the measured intensity in the out-of-plane direction (GISAS signal) the surface structure can be determined. The given approximation assumes one “effective surface” which can be described by one “effective height-height correlation function”, which implies the existence of a dominant surface morphology. Results and Discussion In the present investigation, we address the next step from a spatial constraint in one direction, namely, the thin film geometry, toward a spatial constraint in two dimensions. The corresponding geometry is a pancake or drop structure. The diblock copolymer is confined into pancakes or drops, and thus the polymer molecules are placed on a solid support that fills one half space and are (70) Naudon, A.; Babonneau, D.; Thiaudiere, D.; Lequien, S. Physica B 2000, 283, 69. (71) Yoneda, Y. Phys. Rev. 1963, 131, 2010. (72) Mu¨ller-Buschbaum, P.; Stamm, M. Physica B 1998, 248, 229.

surrounded by air in the second half space. With some limits, this geometry can be regarded as a reverted microporous system. Basic differences are the nonhomogeneous surrounding which gives rise to a directiondependent excess free energy density. After the first preparation step, the spin coating of the polymer solution, smooth thin P(Sd-b-pMS) films with small surface roughnesses result. The X-ray reflectivity data exhibit well-pronounced fringes as shown in Figure 2. The prepared film thicknesses are l ) 1.5, 3.0, 4.5, 6.0, 10.0, and 35.0 nm. Compared to the common molecular lengths of the polymer, given by the radius of gyration Rg ) 13.6 nm73,74 of the unperturbed chain and the bulk lamellar spacing L0 ) 45.0 nm,75,76 the films are very thin and can be regarded as confined4-6 except the thickest one with l ) 35.0 nm. Only the acid cleaning and a very diluted solution enable the preparation of continuous and smooth films in this regime. However, films are confined only with respect to one spatial direction introduced by the limited film thickness. In the perpendicular direction (parallel to the substrate), the samples are extremely large with respect to molecular dimensions. One way to impose a second spatial constraint parallel to the sample surface is offered by a pancake or drop geometry. Recently, the storage under toluene vapor was reported as one possible experimental route to prepare small polymeric drops in the case of homopolymer32 and polymer blend samples.33 Thus in the second preparation step, the initially continuous P(Sd-b-pMS) films are stored under toluene vapor. After a storage time of 12 h, the samples were quenched to ambient air. Independent of the actual preparation within every repetition, we detect a morphology with the same statistics. Thus, while each individual picture shows a different arrangement of the individual surface features, the statistical features of the film morphology remain the same. In the following, the morphological information is divided into a topological and a chemical one. A. Topological Information. Whereas optical methods do not show any surface morphology, with SFM the topology signal is visualized. Figure 3 exhibits typical examples of the observed surface structures. Despite the largest film thickness l ) 35 nm (Figure 3f) for all confined thin films isolated drops evolved after the second preparation step. All presented data have the same scan size of 10 × 10 µm2 to visualize the increasing drop size with increasing initial film thickness l. The height scaling depicted by the color coding is chosen individually to enhance the topological features, because the mean drop height increases with increasing l. The contact line of the drops deviates from the ideal circular shape with increasing l. Thus, the drops are no longer spherical caps, for example, for l ) 10 nm (Figure 3e). In addition to the dominant large drops, a second smaller droplet structure is present on the surface. Large drops and small droplets are placed on top of the silicon substrate. In contrast to the destabilized thin films, the thickest examined film remained stable. The surface exhibits just a statistical roughness and no sign of a destabilization. The peak-tovalley surface roughness is comparable to the values measured on top of the drops. The inset in Figure 2 displays the topology as measured on top of a large drop at a sample with an initial film thickness of l ) 10.0 nm. The mean (73) Jung, W. G.; Fischer, E. W. Makromol. Chem., Macromol. Symp. 1988, 16, 281. (74) Bartels, V. T.; Abetz, V.; Mortensen, K.; Stamm, M. Europhys. Lett. 1994, 27, 371. (75) Giessler, K. H.; Rauch, F.; Stamm, M. Europhys. Lett. 1994, 27, 605. (76) Giessler, K.-H.; Endisch, D.; Rauch, F.; Stamm, M. Fresenius’ J. Anal. Chem. 1993, 346, 151.


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Figure 2. Typical X-ray reflectivity curve (crosses) as measured right after preparation. The solid line is a model fit based on a matrix formalism. It yields the originally prepared film thickness, e.g., 3.0 nm, for the presented data. The inset shows a SFM picture as measured on top of the large drops, which were installed after the toluene vapor treatment (here, the l ) 10.0 nm sample). The height variation of the topology signal is shown by the color coding.

Figure 3. SFM pictures (scan range 10 × 10 µm2) of the topology signal as observed after storage under toluene vapor for 12 h. The shown samples differ in the originally prepared continuous film thickness: (a) 1.5, (b) 3.0, (c) 4.5, (d) 6.0, (e) 10.0, and (f) 35.0 nm. The height scaling is different for each pattern to depict the in-plane structure more clearly.

shape of the drop was subtracted from the data to enable an observation of the quite small remaining microroughness. Thus, in the SFM picture shown in Figure 2 (inset) within the presented area, the surface appears to be flat. The microroughness appears to be statistical and shows no sign of order like perpendicular domains. One cannot, on the other hand, expect to see differences between PS and PpMS in this mode. To obtain a statistical description

beyond the pure pictorial one, the rms surface roughness σrms and the PSD were calculated from the SFM data. The rms surface roughness contains the height information perpendicular to the substrate surface, and the PSD displays the in-plane information parallel to the substrate surface. A.1. Perpendicular Direction. Due to the strong heterogeneity of the created surface topography, calculated rms surface roughness values depend on the scan range of the SFM experiment. On very small scan ranges less than 1 µm2, the measured surface is either on top of a drop or between the drops and thus statistically significant for these two types of surface areas. Intermediate scan ranges yield quite unreliable σrms values due to the strong surface position dependence, whereas on large scan areas above 100 µm2, again statistically significant data are calculated due to the large number of drops inside these areas. Consequently, in Figure 4 two different types of measured rms surface roughness are plotted as a function of the initially prepared film thickness l. The microroughness calculated from very small scan areas is shown with crosses, and the mean surface roughness calculated from the large scan areas is shown with circles. The microroughness is small as compared to the mean surface roughness. As a function of initially prepared film thickness, the microroughness is constant for all samples within the experimental error, while the mean surface roughness increases with increasing l with the exception of the nonconfined continuous film (l ) 35.0 nm). The dotted line in Figure 4 visualizes the regime of confined film thickness with l < Rg. This increase in σrms is in good agreement with the observations of the X-ray reflectivity experiments which exhibit a large surface roughness from the strong damping of the fringes.58-60 As seen from the bearing analysis from SFM data, the density profile is smeared out completely due to the drop structure. The solid line in Figure 4 shows a fit to the data following

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Figure 4. Root mean square surface roughness σrms plotted as a function of the initially prepared film thickness l. Values calculated from small scan areas are plotted with crosses; ones from large scan areas, with circles. The symbol size pictures the error bar of the experimental data. The solid line is a model fit as explained in the text.

σrms ∼ ln(l). A logarithmic increase of the surface roughness is observed in capillary wave models. However, it should be underlined that the drop morphology is far away from an homogeneous film surface. Thus, the slow increase in σrms might be attributed to creation of pancakelike drops. With increasing initial film thickness, the drop diameter increases stronger than the drop height (compare for example parts c and e of Figure 2). A possible origin can be the interaction between the polymer and Si surface. The contact angle increases with increasing difference in the surface energies between polymer and Si. A small contact angle yields a buildup of pancakelike drops rather than spherical caplike ones. Thus, the mean surface roughness is dominated by the shape of the evolved surface structure. In contrast, the microroughness is independent of the surface structure. The dashed line exhibits the constant mean microroughness on top of these pancakes as well as on top of the continuous film. On top of a drop or on a continuous film, the microroughness is not affected by the presence of a drop structure at all. As expected, microroughness does not depend on the film thickness rather than on polymer specific parameters such as surface tension. A.2. In-Plane Direction. With respect to the in-plane information, two different length scales have to be addressed separately as well. From the calculated PSD data combined into one master curve, a statistical lateral information is gained. In Figure 5, an example of a resulting master curve (crosses) is plotted as a function of in-plane wave vector qy. The most prominent in-plane length gives rise to the peak at position qI marked with I. The related length ξI ) 2π/qI corresponds to the mean distance between the large drops. This most prominent in-plane length ξI is calculated for all initially prepared film thicknesses. Data ξI extracted from the master curves are plotted in Figure 6 (crosses). In general, Figure 6 presents lateral lengths Λ as a function of the initially homogeneous film thickness l in a linear plot. The lines are fits to the data. The mean distance between the large drops increases with increasing film thickness following ξI ∼ l (solid line in Figure 6). This observed film thickness dependence is comparable to the observed increase of lateral structures as measured in the case of ultrathin blend films of dPS/PpMS after toluene vapor storage.78 During the second preparation step, the toluene vapor

Figure 5. Comparison between the master curve calculated from the SFM data (crosses) and the grazing small angle scattering data measured with X-rays (triangles) and neutrons (circles). The corresponding initial film thickness for the presented example is l ) 6 nm. Typical length scales are indicated by the arrows and labeled I, II, and III as explained in the text. The dashed lines depict the resolution limit which is different for the three experimental techniques. In this double logarithmic presentation, all curves are shifted along the y-axis for clarity.

Figure 6. Lateral lengths Λ as a function of the initially homogeneous film thickness l. The largest in-plane distances ξI, visible only in the master curves calculated from the SFM data, are plotted with crosses. The additionally observed most prominent in-plane lengths ξII are shown with circles (GISANS), triangles (GISAXS), and filled circles (SFM). The solid line is a linear fit to the data; the dashed line is a fit assuming a constant.

wets and swells the initially homogeneous P(Sd-b-pMS) film.33 Both parts of the block copolymer PS and PpMS are comparably strongly swollen due to the chemical similarity of PS and PpMS. The incorporation of solvent (77) Sadiq, A.; Binder, K. J. Stat. Phys. 1984, 35, 517. (78) Mu¨ller-Buschbaum, P.; Stamm, M. Colloid Polym. Sci. 2001, 279, 376.


molecules leads to a plastification of the glassy film due to an effective reduction of the glass transition temperature.7 A highly concentrated polymer-toluene solution layer is created on top of the substrate. Thus, a two-phase system of the solvent toluene and the diblock copolymer P(Sd-b-pMS) sets in. The viscosity is drastically decreased when compared to the value of the bulk block copolymer, and therefore kinetics is accelerated. The system destabilizes into isolated drops. The thin film geometry imposes a spatial constraint in the direction perpendicular to the substrate. Thus, the growth of the domains is first influenced by this constraint. Consequently, the domain size scales with this characteristic length introduced by the film thickness (ξI ∼ l).77 During the quench to ambient air conditions, the solvent molecules are rapidly removed from the system. This immobilizes the polymer molecules. Regions which contain no polymer molecules appear as a bare substrate area after the quench. The surface is covered by drops surrounded by air as observed with SFM (see Figure 3). To improve the statistical relevance of the results from the SFM measurements, in addition grazing incidence small-angle scattering measurements were performed. With this type of measurement, lateral in-plane length scales are probed from the molecular scale-up to the mesoscopic regime. Experiments were performed with X-rays as well as with neutrons. The GISAXS (triangles) and the GISANS (circles) data are shown in Figure 5 together with the master curve (crosses) calculated from the SFM data. Because the mean distance between the drops is large, it cannot be resolved within the scattering experiments. The dashed lines in Figure 5 indicate the resolution limits of the individual techniques. The largest resolvable in-plane length decreases from the SFM technique to GISAXS and GISANS. Thus, the scattering methods are blind with respect to the information on large length scales which is however easy to determine from the SFM master curve. On the other hand, this invisibility of the large drops helps to obtain additional information about the second in-plane feature, namely, the small droplets. While the structure factor of the droplets is not detected, the form factor gives rise to a signal. In the master curve from the SFM data, it is visible only as a shoulderlike peak at a position qII marked with II. Thus, the master curve shows only weak evidence for the presence of a second length scale ξII which can be attributed to the small droplets. In contrast, in the scattering data the intensity decay (bending point) is more pronounced. Nevertheless, a marked peak is not present in this range. The replacement of a peak by a bending point results from deviations of a perfect monodispersity of the droplet diameters. This imperfection of the mean droplet diameter yields a smearing of the equivalent form factor in the scattering data. The effect is less pronounced in the SFM results because SFM measurements average only over a much smaller surface area, which increases the apparent monodispersity. Nevertheless, both scattering techniques depict the presence of this second topographical surface feature as a statistically relevant one. In Figure 6, the film thickness dependence of this droplet structure is compared with the behavior of the large drops. Results of the GISAXS data are plotted with triangles; the GISANS data, with circles; and the SFM data, with filled circles. The dashed line is the fit to the data using a constant value Λ ) 580 nm. Thus, the second structure is not dependent on the initially prepared film thickness. This can be explained by an initially continuous and ultrathin metastable polymer film which destabilizes in a second

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step. A comparable behavior was observed in homopolymer films of polystyrene quite recently.79 In coexistence with the large drops, a continuous ultrathin film is present during the storage under toluene vapor. The coexistence of macroscopic large drops and an homogeneous ultrathin film in the dewetting of diblock and triblock copolymer films was attributed to an autophobic dewetting by Hamley et al.80 An ultrathin film might be stabilized due to shortrange interactions on top of the substrate. The thickness of this film is determined by the interaction of the polymer solution with the substrate. Because both the solution concentration and the substrate material are not changed as a function of initial film thickness, they are constant for all investigated samples. During the quench, the increase in surface tension, due to the removal of solvent molecules, reduces the stabilizing influence of the shortrange interaction. The previously constant ultrathin film decays into droplets of constant diameter via a second destabilization. The broadened size distribution might be attributed to hydrodynamic effects. B. Chemical Morphology. Besides the pure topological information in thin diblock copolymer films, the internal morphology introduced by a microphase separation process is important. Because both blocks PSd and PpMS differ only by one methyl group, despite the deuteration, their chemical and mechanical properties are very similar. Therefore, surface characterization methods such as friction and stiffness measurements do not yield enough contrast to distinguish between the components and a selective dissolution of one component is not possible. A chemical morphology is gained from a comparison of X-ray and neutron scattering data. With X-rays, both polymers are hardly distinguishable (δ(dPS)/δ(PpMS) ) 0.9), while with neutrons the scattering of both differs markedly (δ(dPS)/δ(PpMS) ) 4.1). Thus, comparing GISAXS and GISANS data helps to separate a topological from a chemical morphology. A signal present in the neutron data but absent in the X-ray data has to result from a chemical morphology. B.1. Perpendicular Direction. The perpendicular direction is investigated with vertical cuts from the 2D detector area. Like in the common reflectivity curve, the presence of a lamellar stack oriented parallel to the substrate would give rise to a Bragg peak. As the example shown in Figure 1 pictures, we do not detect any sign of a Bragg peak in the neutron data. In the X-ray data, however, no peak is expected due to the missing contrast. Thus, in the samples no structure with a repeat unit parallel to the substrate surface is present. B.2. In-Plane Direction. In GISAS experiments, only in-plane lengths are probed because the wave vector component perpendicular to the substrate surface is constant. Thus, GISAS is insensitive toward orientations in the perpendicular direction but directly sensitive to structures which are parallel to the substrate surface. As visible by the example depicted in Figure 5, the GISANS data (circles) exhibit a well-pronounced peak at position qIII which is not present in the GISAXS data. Thus, qIII resembles an internal structure which is only visible due to the enhanced contrast between PSd and PpMS when probed by neutrons. The related real-space length is ξIII ) 72 nm. While the peak at qIII is quite well established in the case of an initial film thickness of l ) 6 nm with decreasing film thickness the peak at qIII becomes less pronounced as shown in Figure 7. In the scattering data (79) Wunnicke, O.; Mu¨ller-Buschbaum, P.; Wolkenhauer, M.; Mahltig, B.; Lorenz-Haas, C.; Stamm, M. J. Colloid Interface Sci., submitted. (80) Hamley, I. W.; Hiscutt, E. L.; Yang, Y. W.; Booth, C. J. Colloid Interface Sci. 1999, 209, 255.

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Figure 7. Double logarithmic plot of GISANS data (measured at a wavelength of 0.6 nm) from samples stored for 12 h under toluene atmosphere. From the bottom to the top, the initially prepared film thickness increases (l ) 1.5, 3.0, 4.5, 6.0, 10.0, and 35.0 nm). The arrow marks a position related to a peak which is originated from an internal structure ξIII as explained in the text. The dashed line indicates the resolution limit. All curves are shifted along the y-axis for clarity.

of the thinnest films examined, the peak at qIII is only hardly resolved from the background (see two curves from the bottom in Figure 7). This results from the extremely small scattering volume of these drops which consists of a few molecules only and the limited available neutron flux. The influence of the neutron flux is demonstrated by a variation of the wavelength used for the GISANS measurements. In Figure 8, measurements with the initial film thickness of l ) 3 nm using wavelengths of 0.6, 1.0, and 1.7 nm are compared. By increasing the wavelength, the neutron flux is decreased due to the characteristics of the beamline whereas the resolution and the coherently illuminated surface area are increased. How drastically the peak intensity depends on the neutron flux is visible best by a comparison of the peak at the position marked with II. While a peak at qII is well visible at wavelengths of 0.6 and 1.0 nm, it is extremely weak at 1.7 nm. As a consequence, the peak marked with III is only weakly present at a wavelength of 0.6 nm and vanishes in background and noise at 1.0 and 1.7 nm. The typical accumulation times were above 6 h, and thus an improvement of the peak statistics by an increase of the accumulation time is unlikely. This elucidates the limits of the GISANS method although the experiments were performed at a high flux beamline. With increasing drop diameter, the scattering volume is increased as well, and the peak resulting from the internal order gets a higher intensity. The increase in scattering volume increases the number of internal repeat units inside each drop and thus gives rise to an improved signal. The presence of the peak at qIII is a clear sign of a perpendicular oriented lamellar, but unfortunately deviations in the mean peak position as a function of film thickness are difficult to detect, as Figure 7 shows. Within the experimental error, a value of ξIII ) 72 ( 5 nm is obtained. We observe no film thickness dependence or


Figure 8. Double logarithmic plot of GISANS data of a sample with an initial film thickness of l ) 3 nm measured at different wavelengths of λ ) 0.6 nm (triangles), 1.0 nm (circles), and 1.7 nm (filled circles). The experimental resolution is indicated by the dashed lines. Typical peaks are present at positions marked with II and III. All curves are shifted again for clarity.

Figure 9. Schematic cross sections of the morphology inside the drops created by the toluene storage. The top double arrow indicates the bulk lamellar spacing L0. The bottom double arrow depicts experimentally observed stretching of the lamellae with a perpendicular arrangement. The deuterated block PSd is denoted with A, and the protonated block PpMS is denoted with B. The gray scaling visualizes the scattering contrast as present for neutrons.

dependence from the drop diameter. As compared to the value of the bulk lamellar spacing L0 ) 45 nm, this would impose a stretching parallel to the substrate surface. In Figure 9, the arrangement of the diblock copolymers inside a drop is schematically shown by a cross section. A comparison between the two double arrows indicating L0 and ξIII in Figure 9 visualizes this stretching. Possible origins are the interaction between the substrate atoms and the blocks, the confinement imposed by the limited volume supplied by a drop geometry, and the used preparation procedure. The confinement of the diblock copolymer molecules into drops was realized by the toluene vapor storage. The internal order (perpendicular lamellae) installs during the plastification of the film. The related swelling due to the incorporation of solvent molecules might originate the observed increase by a factor of ξIII/L0 ) 1.6. On the other hand, a stretching of the parallel component of the chain dimension was reported from confined homopolymer samples.18-21 For PSd, a stretching by a factor of 1.5 was reached in the limit of a film thickness comparable to the radius of gyration of the unperturbed chain.20 In the examined film thickness regime, the confinement gives rise to an interplay between the intrinsic length scale of the bulk structure L0 and the geometry of


the film.41-45 Because the dimensions of the prepared drops differ markedly between the in-plane (diameter) and the perpendicular (height) direction, the imposed spatial constraint in the perpendicular direction is significantly stronger when compared to the confinement in the inplane direction. With respect to the nearest, confining walls, transitions between phases of identical symmetry but different orientation become energetically possible and lamellar domains can reorient from a parallel to a perpendicular arrangement.46-50 In addition, the polymer molecules have to fill the geometrical space available by the drop volume which can be the origin of the stretching parallel to the substrate surface. Another reason is maybe the anisotropic origin of the film: chains are swollen in solution and dry fast. Unfortunately, it is not possible to determine the origin of the lamellar stretching without doubt. The plastification during toluene vapor storage is a preparation step in the presented investigation. In general, a destabilization of thin diblock copolymer films is only rarely reported in the literature and thus common methods such as annealing are quite ineffective. Frequently, a lamellar orientation parallel to the substrate stabilizes diblock copolymer films. Nevertheless, with respect to applications these arrays of nanoscopic structures with a controlled spatial orientation of the microdomains are desirable. Especially the perpendicular arrangement of the lamellae was reported only for homogeneous thin copolymer films.46-50 This thin film geometry is addressed in the presented investigation as well. In the case of the thickest film, which remained homogeneous after toluene vapor storage, the GISANS data show only a quite broad peak with a weak intensity. However, the presence of this peak again is a fingerprint of at least a partial perpendicular arrangement of the created lamellae inside the film. The lamellar spacing seems to be stretched further, but the width of the peak and its weak intensity make a quantitative comparison difficult. Despite the influence of specific interactions in the thin film geometry an elongated and flattened structure of the polymer chains is expected theoretically51 and observed experimentally for homopolymers.52 Perpendicularly arranged lamellae were not detected in the system P(Sd-b-pMS) in previous investigations.75,76 This might by originated by the differences in the applied preparation technique. However, perpendicularly arranged lamellae cannot be detected directly by the recently applied nuclear reaction analysis technique. In the case of other sample systems such as P(Sb-MMA), the investigated perpendicular lamellae exhibited a lamellar spacing comparable to the bulk value L0.47,48 A doubling of the periodicity in the triblock copolymer P(S-b-B-b-MMA) from bulk to perpendicular arrangement in the thin film geometry was attributed to a surface reconstruction effect.81 Despite the vertical shift of alternating lamellae, the pure in-plane distance is again well given by the bulk value.81 (81) Stocker, W.; Beckmann, J.; Stadler, R.; Rabe, J. P. Macromolecules 1996, 29, 7502.

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Due to the presence of the small droplets among the drops, it is evident that the diblock copolymer molecules can be confined into even smaller volumes. This enhances the confinement in the in-plane direction. With the described technique, the small droplets cannot be prepared without the presence of the large drops. On the other hand, this would yield a further reduced polymer volume and thus would impose severe problems even with the GISANS technique. Consequently, the internal arrangement of the diblock copolymers inside the small droplets is still an open question which goes beyond the scope of the present investigation. Summary and Outlook We showed that after a toluene vapor storage within the limited range of strongly confined initially homogeneous diblock copolymer films a destabilization into isolated drops is achieved. The dependence of the mean drop distance from the film thickness is explainable within a model based on the spatial constraint induced by the initially prepared film thickness. The geometry of the drops introduces a second spatial restriction, while the thin film geometry which frequently is investigated in the case of confinement offers only a one-dimensional confinement. The available space for polymer molecules is limited in the perpendicular direction (with respect to the substrate surface) by the drop height and in the parallel direction by the drop diameter. Due to the large differences between both related length scales, the realized confinement is not symmetric. From this asymmetry, an arrangement into perpendicularly aligned lamellae inside the drops results. The observed stretching of the lamellar spacing when compared to the bulk is compatible with chain stretching observed in strongly confined homopolymers. A stretching due to swelling during the plastification cannot be excluded due to the sample preparation. Different preparation routes without vapor storage or a tuning of the solvent used might address this question in more detail. However, the nanostructured drops containing perpendicular lamellae might be important with respect to applications such as nanotemplates. In summary, the drop geometry offers interesting possibilities for the investigation of other polymer systems with respect to spatial confinement in two dimensions. With respect to more immiscible diblock copolymers such as P(S-bMMA), it is questionable whether the perpendicular alignment is characteristic for the drop geometry. Of course, a further increase of the in-plane confinement due to smaller droplets will be interesting as well. Acknowledgment. We thank S. Cunis and G. von Krosigk for their help at the BW4 beamline at the HASYLAB. Additionally, we owe many thanks to R. Gehrke for his general support of the experiment at HASYLAB. This work was supported by the BMBF (Fo¨rderkennzeichen 03DUOTU1/4). O.W. was supported by the DFG Schwerpunktprogramm “Benetzung und Strukturbildung an Grenzfla¨chen” (Sta 324/8-2). LA010448T

Structure Formation in Two-Dimensionally Confined Diblock

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