Lamellar Orientation in Thin Films of Symmetric Semicrytalline

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J. Phys. Chem. B 2007, 111, 11921-11928

11921

Lamellar Orientation in Thin Films of Symmetric Semicrytalline Polystyrene-b-poly(ethylene-co-butene) Block Copolymers: Effects of Molar Mass, Temperature of Solvent Evaporation, and Annealing Guo-Dong Liang,† Jun-Ting Xu,*,†,‡ and Zhi-Qiang Fan†,‡ Key Laboratory of Macromolecular Synthesis and Functionalization, Department of Polymer Science & Engineering, Zhejiang UniVersity, Hangzhou 310027, P. R. China, and State Key Laboratory of Chemical Engineering, College of Materials and Chemical Engineering, Zhejiang UniVersity, Hangzhou 310027, P. R. China ReceiVed: June 7, 2007; In Final Form: July 25, 2007

Orientation of the lamellar microdomains in thin films of three symmetric polystyrene-b-poly(ethylene-cobutylene) block copolymers (S65E155, S156E358, and S199E452) on mica was investigated via atomic force microscopy (AFM), grazing incidence X-ray diffraction (GIXRD) and X-ray photoelectron spectroscopy (XPS). The results show that lamellar orientation in the SxEy block copolymers greatly depends on the molar mass of the block copolymers, the temperature of solvent evaporation, and annealing. The nascent thin film of the low molar mass block copolymer, S65E155, shows a multilayered structure parallel to the mica surface with the PS block at both polymer/mica and polymer/air interfaces, but the high molar mass block copolymers, S156E358 and S199E452, exhibit a structure with lamellar microdomains perpendicular to the mica surface. When the solvent is evaporated at a lower temperature, the crystallization rate is fast and a two-dimensional spherulite structure with the lamellar microdomains perpendicular to the mica surface is observed. Annealing of all the thin films with lamellar microdomains perpendicular to the mica surface leads to morphological transformation into a multilayered structure parallel to the mica surface. In all SxEy thin films on mica, the stems of PE crystals are always perpendicular to the interface between the lamellar PE and PS microdomains. A mechanism is proposed for the formation of different microdomain orientations in the thin films of semicrystalline block copolymers. When the thin film is prepared from a homogeneous solution, microdomains perpendicular to the substrate surface are formed rapidly for strongly segregated block copolymers or at a lower crystallization temperature and kinetically trapped by the strong segregation strength or solidification of crystallization, while for weakly segregated block copolymers or at slower crystallization rate, the orientation of the microdomains is dominated by surface selectivity.

Introduction It is well known that block copolymers can self-assemble into various supermolecular structures, including lamellae, gyroids, cylinders, and spheres, depending on the composition of the block copolymers. These nano-ordered structures have been considered as lithographic mask and template to fabricate new devices with novel properties.1-3 Orientation of the microphase-separated domains is one of the key steps in the fabrication process, especially for the anisotropic microdomains such as cylinder and lamellar, since application of the devices is determined by not only the shape but also the orientation of the microdomains. For example, cylinders aligned with axes perpendicular to the substrate can be used to produce film with nanochannels when the minor component is removed4,5 or nanopillars when the major component is removed,6 while the cylinders oriented uniformly parallel to the substrate may be used to fabricate nanowires.7 Lamellae aligned parallel to the substrate are candidates for photonic application when the long period is comparable to the wavelength of light,8 and the devices * To whom correspondence should be addressed. E-mail: [email protected]. Tel./Fax: +86-571-87952400. † Key Laboratory of Macromolecular Synthesis and Functionalization. ‡ State Key Laboratory of Chemical Engineering.

with a banded structure can be fabricated when the lamellar microdomains are perpendicular to the substrate. Lots of methods have been reported to control the orientation of the microdomains of block copolymers on a substrate. Surface property of the substrate is an important factor that determines orientation of the microdomains. Under asymmetric boundary conditions the microdomains are usually perpendicular to a neutral surface but parallel to a selective surface. Thus modification of the substrate surface is frequently used to control orientation of block copolymer microdomains.9-13 External fields such as electric field, magnetic field, and shearing are also used to orientate the microdomains.14-19 Solvent evaporation can induce orientation of microdomains as well.20-26 Alternatively, orientation of the microdomains may be directed by the roughness, pattern, or crystals on the substrate surface.27-35 Moreover, orientation of block copolymer microdomains is also affected by thermal annealing or commensurability between the film thickness and the long period.36-39 Recently thin film morphology of semicrystalline block copolymers has attracted increasing interests, including orientation of microdomains,33,34,40,41 registration of crystalline stems,32,42,43 nucleation mechanism and growth kinetics,44-46 and morphological reorganization.47-51 In the bulk of semicrystalline block copolymers, crystallization can compete with microphase

10.1021/jp074399q CCC: $37.00 © 2007 American Chemical Society Published on Web 10/11/2007

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TABLE 1: Molecular Characteristics of SxEy Block Copolymers SxEya

Mw/Mn

WEb

VEb

L0c (nm)

Tcd (°C)

TODTe (°C)

S65E155 S156E358 S199E452

1.12 1.09 1.08

0.45 0.44 0.43

0.51 0.50 0.49

21.3 43.3 50.6

91.5 94.3 89.1

199 272 285

a x and y are the polymerization degrees of styrene and ethylene units, respectively. b WE and VE are the weight and volume fractions of the PE block in SxEy block copolymers, respectively, and are calculated by WE ) 28y/(28y + 104x) and VE ) 1/[1 + (1/WE - 1)dE/ dS], where dE and dS are the densities of branching PE and PS at 120 °C, being 0.801 and 1.01 g/cm3, respectively.66 c L0 is the long period of SxEy block copolymers in bulk, L0, at room temperature. d Tc is the crystallization temperature of SxEy block copolymers at a cooling rate of 10 °C/min from 180 °C. e The order-disorder transition temperatures (TODT) of SxEy block copolymers were estimated based on mean-field theory67 and the method proposed by Han et al.68,69

separation. Both crystallization temperature (compared with glass transition temperature of the amorphous block) and segregation strength between the two unlike blocks have an effect on the final morphology of block copolymers.52-55 Therefore, it is expected that both crystallization and microphase separation may be used to tune the thin film morphology of block copolymers,56-62 which can be achieved by changing the crystallization temperature and segregation strength, respectively. It has been shown that crystallization temperature can affect the orientation of microdomains in thin film of polystyreneb-poly(L-lactide) block copolymer.62 It is also reported that amorphous block copolymers of higher molecular weight tend to adopt an orientation perpendicular to the substrate surface.63 A similar phenomenon was observed for crystalline/rubbery block copolymer in our previous work.64 However, for crystalline/glassy block copolymer the simultaneous effects of crystallization and microphase separation strength on thin film morphology have not been reported yet. In the present work, thin film morphology of three symmetric semicrystalline polystyrene-b-poly(ethylene-co-butene) block copolymers with different molecular weights was studied. We demonstrated here that the orientation of the lamellar microdomains could be tuned by a change of molar mass (segregation strength), temperature of solvent evaporation (crystallization rate), and annealing. Experimental Section Materials. Polystyrene-b-poly(ethylene-co-butene) block copolymers were synthesized by anionic living polymerization of styrene and butadiene monomers and subsequent selective hydrogenation of the polybutadiene block.65 The block copolymers are denoted as SxEy, in which the subscripts x and y refer to the polymerization degrees of styrene and ethylene units, respectively. The characteristic parameters of the SxEy block copolymers are summarized in Table 1. Preparation of Diblock Copolymer Thin Films. Block copolymer thin films were prepared by dip-coating an SxEy block copolymer/toluene solution (0.5 w/v %) onto mica at 90 °C. The mica surface is hydrophilic, while both polystyrene and poly(ethylene-co-butene) are hydrophobic, so mica provides a nearly neutral surface for both blocks. Another reason for using mica as substrate is that mica is transparent and smooth, which is suitable for both atomic force microscopy (AFM) and polarized optical microscopy experiments. The first several layers of the multiple-layered mica were removed to avoid dusts on the surface of mica prior to dip-coating. A mica substrate with size 10 × 10 mm2 was slowly dipped into a tank containing

the SxEy/toluene solution. After submersion for 5 min, the mica chip was withdrawn from the tank at a speed of about 10 cm/s. The solvent was allowed to evaporate at 90 °C for 5 min or at 50 °C for 10 min. Annealing of the SxEy thin films was conducted at 90 °C for various times under vacuum (10 Torr). AFM experiments were conducted immediately after evaporation of solvent or after annealing. The half-layered block copolymer thin film was prepared at a lower concentration (0.1 w/v %) and at a lower solvent evaporation temperature (20 °C). The film thickness was determined by comparing the height positions of the thin film and the naked mica surface in the large area of AFM images. Atomic Force Microscopy. The thin film morphology of SxEy block copolymers was investigated by tapping mode AFM. A commercial atomic force microscope (DI 3100 AFM, Digital Instruments) was used with a silicon microcantilever (spring constant 45 N/m and resonance frequency ∼300 kHz). The scan rate was ranged from 1 to 2.0 Hz to optimize the image quality. The set-point ratio, the ratio between the set-point amplitude and the free vibration amplitude (the lowest amplitude when tip and sample are not in contact), was chosen to be ∼0.6. The measurements were carried out in the repulsive regime of tapping mode AFM. For acquisition of the surface morphology and material distribution, both height and phase images were recorded. Since both difference between the material mechanical properties and height fluctuation can contribute to the phase contrast, the phase images were obtained under various conditions to make sure that phase contrast was caused by material mechanical properties. Parameters characterizing the thin films such as thickness of lamellar microdomains were obtained directly from cross-sectional profiles. To ensure the creditability of the data, cross-sectional profiles from different areas of height image were necessary. At least 10 values were obtained for every parameter. The average value and standard deviation of every parameter were presented. The height position of the substrate surface is referred to as zero in the height profiles, which was determined by comparison with the position of the naked substrate surface in the large area AFM images. Grazing Incidence X-ray Diffraction (GIXRD). GIXRD experiments were carried out on a high-resolution X-ray diffractometer (BEDE D1 system) with Cu KR radiation to investigate the orientation of PE crystal stems in block copolymer thin films. The diffracted beam is in the plane defined by incident beam and the surface normal. This geometry is sensitive to the structure parallel to the surface. Different grazing incidence angles were tested, and a better contrast of block copolymer signal over mica signal with small error induced by surface undulation at the same time was obtained at a grazing incidence angle of 0.6°. The XRD curves were scanned at room temperature in the 2θ range of 3.0-30°. X-ray Photoelectron Spectroscopy (XPS). The surface composition of the SxEy thin films was determined with PHI5300 X-ray photoelectron spectroscopy. The working vacuum of the analysis chamber was 3 × 10-7 Pa during the spectra measurement. The binding energies were calibrated based on the C-C bond with a binding energy of 285 eV to compensate for the sample charging. Transmission Electron Microscopy (TEM). The morphology of the SxEy thin films was observed with TEM (JEM-1230) using an acceleration voltage of 120 kV. The solution of S65E155 in toluene (90 °C) was spread on the water surface (50 °C), and the thin film formed was deposited on a copper grid for TEM observation without staining.

Lamellar Orientation of SxEy Block Copolymers

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Figure 1. AFM height image (a) and corresponding cross-sectional profile (b) of the nascent S65E155 thin film after evaporation of solvent at 90 °C. The image size is 10 × 10 µm.

Figure 2. XPS C 1s spectra of the nascent S65E155 thin film on mica at various incident angles of 90° (a); 60° (b); 30° (c) and spectra of half-layered S65E155 thin film on mica at an incident angle of 90° (d).

Results and Discussion Morphology of the Nascent Thin Films. Figure 1 shows the AFM height image and corresponding cross-sectional profile of the nascent (freshly dip-coated) S65E155 thin film on mica. The solvent is allowed to evaporate at 90 °C for 5 min. It can be seen that the nascent S65E155 thin film presents a terraced cross-sectional profile with discrete values for the film thickness. This is a characteristic of block copolymer thin film with lamellar microdomains parallel to the substrate surface. For the thin film of block copolymer with a lamellar structure, the thickness of the thin film is usually (n + 1/2)L0 (where n is an integral and L0 is the long period in the bulk) in the case of asymmetric wetting and nL0 in the case of symmetric wetting due to the presence of microphase separation. The crosssectional profile shows that the nascent S65E155 thin film is comprised of multiple polymer layers parallel to the mica surface with irregular holes. We note that the thickness of the nascent S65E155 thin film is integer times of the long period in the bulk (L0), showing that the S65E155 thin film on mica is symmetric and the same component is located at both polymer/mica and polymer/air interfaces. However, since both polyethylene and polystyrene are hydrophobic, the AFM image cannot provide the information of which block is located at the polymer/mica and polymer/air interfaces. Figure 2a-c shows the XPS C1s spectra of the nascent S65E155 thin film at various incident angles, normalized to the C-C bonding with a binding energy of 285 eV. A weak satellite peak located at about 292 eV is observed, which is attributed to PS block, arising from the shakeup transition (π f π*) accompanying core ionization.70 It is found

that the intensity of the satellite peak increases slightly with decreasing the incident angle, showing that the PS block is located at the polymer/air interface. However, the satellite peak located at 292 eV is very weak even in the neat PS homopolymer, and we cannot be completely sure that PS block is located at the polymer/air interface. Moreover, X-ray may penetrate the film and the signal at 292 eV may be from the PS block underneath the surface block. For the purpose of further confirmation, half-layered thin film of S65E155 was prepared, in which the thickness of the thin film is only half of the long period in bulk (see Figure S2 in the Supporting Information). The half-layered thin film is asymmetric, and different blocks are located at the polymer/mica and polymer/air interfaces. The XPS C1s spectrum of the half-layered thin film is shown in Figure 2d. No satellite peak appears at 292 eV. This shows that PE block is located at the polymer/air interface in the halflayered thin film. Based on the XPS results, we can draw the conclusion that PS block is located at both the polymer/mica and polymer/air interfaces in the symmetric thin film of S65E155 on mica. With increasing the chain length of SxEy block copolymers, the morphology of the thin film becomes quite different. Figure 3 shows the AFM image of the nascent S156E358 thin film on mica. It is found that the cross-sectional profile does not exhibit a terraced structure but a meandering texture. As a result, the lamellar microdomains in the nascent S156E358 thin film on mica are perpendicular to the substrate surface. The long period of the lamellar microdomains in the thin film is slightly larger than L0. Like S156E358, the nascent S199E452 thin film on mica also presents a lamellar structure perpendicular to the substrate surface (Figure S3 in the Supporting Information). However, the morphology of nascent S199E452 thin film is coarse, probably due to the stronger crystallizability of S199E452. As a result, we can conclude that the lamellar structure parallel to the substrate surface is preferential for the low molar mass block copolymer S65E155, but the lamellar structure perpendicular to the substrate surface is preferred for the higher molar mass bock copolymers (S156E358 and S199E452). Since similar results have been reported for the amorphous block copolymers63 and crystalline/rubbery block copolymers,64 the dependence of the orientation of the lamellar microdomains on molar mass (segregation strength) is common. Besides the orientation of the lamellar microdomains, orientation of the PE crystal stems should also be determined.71 Figure 4 shows the grazing incidence X-ray diffraction (GIXRD) curves of the nascent SxEy thin films. According to the geometry of GIXRD, only the reflections parallel to the substrate surface can be detected. It is found that no reflection is observed for

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Figure 3. AFM phase image (a) and corresponding cross-sectional profile (b) of the nascent S156E358 thin film after evaporation of solvent at 90 °C. The image size is 2 × 2 µm.

Figure 4. GIXRD curves of the nascent SxEy thin films on mica. The strong reflection peaks at 2θ ) 9.9°, 19.6°, and 29.1° are from mica.

S65E155, but two peaks located at about 21.6° and 24.1°, associated with the (110) and (200) crystalline planes of polyethylene, respectively, are observed for high molar mass block copolymers, S156E358 and S199E451. The GIXRD result shows that the stems of PE crystals are perpendicular to the substrate surface for S65E155 but parallel to the substrate surface for S156E358 and S199E451. When combined with AFM results, we can see that the stems of PE crystals are always perpendicular to the interface of microphase-separated PE and PS microdomains, irrespective of the orientation of the lamellar microdomains. Effect of Annealing. The nascent thin film is usually metastable due to two-dimensional confinement and the interaction between the substrate and polymer. Annealing the nascent thin film will lead to morphological transformation. Annealing is conducted at 90 °C for 80 h in vacuum. Figure 5 shows the AFM height image and corresponding cross-sectional profile of the annealed S65E155 thin film on mica. It is found that annealing the S65E155 thin film on mica does not change the morphology of the nascent thin film, still exhibiting the terrace morphology. However, dewetting phenomenon is observed, and the surface of the thin film becomes smooth after annealing. Figure 6 shows the AFM images of the annealed thin film of S156E358 on mica. Again dewetting occurs. A striking difference between the nascent and the annealed S156E358 thin films is that the lamellar microdomains perpendicular to the mica surface in the nascent thin film disappear and the annealed S156E358 thin film exhibits a terraced morphology with approximately discrete values for the film thickness. The step change of the thickness of the thin film is similar to the long period of S156E358 in the

bulk. Therefore, we can draw the conclusion that for S156E358 annealing leads to morphological transformation from the lamellar structure perpendicular to the surface into the structure parallel to the mica surface. Similar morphological transformation is observed in the annealed S199E452 thin film (Figure S4 in Supporting Information). During morphological transformation of S156E358 and S199E452 thin films induced by annealing, the orientation of the PE crystal stems also changes from parallel to the mica surface into perpendicular to the mica surface, and no reflection peaks are observed in the GIXRD curves of the annealed thin films (see Figure S5 in the Supporting Information). Effect of Temperature of Solvent Evaporation. In order to investigate the effect of evaporation temperature on the thin film morphology, solvent is allowed to evaporate at 50 °C after dip-coating from solution at 90 °C. Figure 7 shows the AFM image of the nascent S65E155 thin film in which the solvent is evaporated at 50 °C. It is found that the thin film exhibits twodimensional spherulite morphology. A common center from which spherulites radiate is observed. Due to the presence of branches, we do not observe periodicity in the direction parallel to the surface nor a terraced structure in the cross-sectional profile. The S65E155/toluene solution (90 °C) was also spread on water at 50 °C to form a thin film for TEM observation. Since the mica surface is hydrophilic, it is expected that the thin film prepared in such a way has the same lamellar orientation as the S65E155 thin film on mica. As shown in Figure 8, spherulites with branching lamellae are observed. For the other two SxEy blocks, a similar morphology is formed at the lower temperature of solvent evaporation (Figures S6 and S7 in the Supporting Information). Comparing Figure 1 and Figure 7, one can see that the thin film morphology is strongly dependent on the temperature of solvent evaporation. For semicrystalline SxEy block copolymers, microphase separation and crystallization take place simultaneously during evaporation of the solvent. The temperature of solvent evaporation affects not only the evaporation rate of solvent (related to microphase separation) but also the crystallization rate of the crystallizable block. At higher temperature, the evaporation rate of solvent is fast but crystallization rate is slow. On the contrary, at lower temperature (50 °C), microphase separation takes place slowly, but the crystallization rate is fast and thus the crystallization process may dominate the thin film morphology, as revealed by the formation of spherulites. Evolution of the Two-Dimensional Spherulite Morphology. We have demonstrated that the S65E155 thin film on mica exhibited a multilayered structure parallel to the surface after

Lamellar Orientation of SxEy Block Copolymers

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Figure 5. AFM height image (a) and corresponding cross-sectional profile (b) of the annealed S65E155 thin film after evaporation of solvent at 90 °C. Annealing was performed at 90 °C for 80 h under vacuum. The image size is 10 × 10 µm.

Figure 6. AFM height image (a) and corresponding cross-sectional profile (b) of the annealed S156E358 thin film after evaporation of solvent at 90 °C. Annealing was performed at 90 °C for 80 h under vacuum. The image size is 10 × 10 µm.

Figure 7. AFM height image of the nascent S65E155 thin film after evaporation of solvent at 50 °C. The image size is 10 × 10 µm.

long-time annealing at 90 °C. As a result, the two-dimensional spherulite morphology with lamellar microdomains perpendicular to the surface may not be stable for S65E155 thin film on mica. Figure 9 shows the morphological change of the S65E155 thin film on mica, which was prepared by evaporation of solvent at 50 °C for 10 min, during annealing at 90 °C at different times. One can see that the two-dimensional spherulite structure remains after annealing for 5 h. Annealing for a longer time (50 h) leads to disappearance of the two-dimensional spherulite morphology and the formation of the multilayered lamellar structures parallel to the mica surface. However, stripe-like lamellar texture can still be seen in the upper polymer layer,

Figure 8. TEM micrograph of S65E155 thin film spread on water at 50 °C.

indicating that the lamellar microdomains are perpendicular to the substrate surface in the upper polymer layer (Figure 9c). During annealing, the holes develop as well, especially in the upper polymer layer, showing a dewetting behavior. After annealing for 80 h, the lamellar microdomains perpendicular to the substrate surface disappear completely, and a terrace structure with holes perforating the thin film is formed (Figure 9e). The cross-sectional profile shows that the thicknesses of the polymer layers are integer multiples of the long period in bulk, indicating that the lamellar microdomains become parallel to the mica surface. A similar morphological transformation is observed for S156E358 and S199E452 thin films upon annealing (Figures S8 and S9 in the Supporting Information). The parallel orientation for the lamellar microdomains in the annealed thin

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Figure 9. AFM height images and corresponding cross-sectional profiles of S65E155 thin film at various annealing times after evaporation of solvent at 50 °C: (a, b) 5 h; (c, d) 50 h, and (e, f) 80 h. The image size is 10 × 10 µm.

film shows that the mica surface is selective to the S block, and such a morphological transformation is driven by the minimization of surface free energy. Since the annealing temperature (90 °C) is higher than the glass transition temperature of the S block, the block copolymers are allowed to adjust their conformation to some extent. Figure 10 shows the GIXRD profiles of the S65E155 thin film prepared by evaporation of solvent at 50 °C at various annealing times. It is found that two weak peaks located at 21.6° and 24.1°, associated with the (110) and (200) crystalline planes of polyethylene, respectively, are observed for the nascent S65E155 thin film, showing that crystal stems of PE are parallel to the substrate surface. This indicates that the lamellar microdomains are perpendicular to the mica surface, since our previous result shows that the stems of the PE crystals are perpendicular to the

interface of the PS and PE microdomains. Annealing the S65E155 thin film leads to the decrease in the intensity of (110) and (200) crystalline planes, until complete disappearance after annealing for 80 h. This clearly shows that the orientation of the PE stems changes with the orientation of the lamellar microdomains. Mechanism for Orientation of Microdomains. The thin film morphology of different SxEy block copolymers shows that the shorter SxEy block copolymers with weak segregation strength form a lamellar structure parallel to the mica surface, while the longer SxEy block copolymers with strong segregation strength tend to form a lamellar structure perpendicular to the mica surface. A similar finding has been also reported for the thin films of amorphous block copolymers.63 On the other hand, a lower temperature of solvent evaporation, thus a faster crystallization rate, leads to the formation of spherulites with lamellar

Lamellar Orientation of SxEy Block Copolymers

Figure 10. GIXRD curves of the S65E155 thin film at various annealing times after evaporation of solvent at 50 °C.

Figure 11. Schematic pathways for formation of the different orientations of the lamellar microdomains in thin films of SxEy block copolymers.

microdomains perpendicular to the mica surface. However, the lamellar structure perpendicular to the mica surface, which is formed in the higher molecular mass SxEy block copolymers or at a lower temperature of solvent evaporation, is not stable, and it transforms into the structure parallel to the mica surface. On the basis of these findings, we proposed two pathways for formation of different orientations of lamellar structure in SxEy thin films, as shown in Figure 11. At initial state, the S block and E block are miscible, due to the presence of solvent, and both blocks contact the substrate surface, which is favorable for formation of the lamellar structure perpendicular to the mica surface. Such a structure is formed more rapidly at stronger segregation strength (such as S156E358 and S199E452) or at a faster crystallization rate (at lower temperature of solvent evaporation, for example, 50 °C). What is most important is that the strong segregation strength and/or solidification of crystallization prevent the conformational adjustment of both blocks in the thin films and that the lamellar structure perpendicular to the mica surface is kinetically trapped. In contrast, when microphase separation (such as S65E155) and crystallization takes place slowly (at lower temperature of solvent evaporation, for example, 90 °C), the surface free energy dominates the orientation of the lamellar structure and the block copolymers have enough time to adjust their conformations. Although both the S and E blocks are hydrophobic, the surface free energies of these two blocks cannot be identical, and mica prefers one of them more or less. Therefore, the lamellar structure parallel to the mica surface is formed. It should be noted that the initial state of the block copolymers plays a very important role in determination of the orientations of the lamellar microdomains. As reported in our previous work, a poly(oxyethylene)-b-poly(oxybutylene) block copolymer (E224B113) forms a lamellar

J. Phys. Chem. B, Vol. 111, No. 41, 2007 11927 structure perpendicular to the silicon surface when the thin film is prepared by spin-coating a E224B113/CHCl3 solution.64 However, when this thin film is melted at 70 °C (below the orderdisorder transition temperature) and recrystallizes at 35 °C, the lamellar structure parallel to the silicon surface is formed (Figure S10 in the Supporting Information). In the latter case, E224B113 crystallizes from a microphase-separated melt without the presence of solvent. Reiter and He have reported that a faster crystallization rate (at a lower crystallization temperature) may lead to formation of the lamellar structure perpendicular to the substrate surface,41,62 but in their work the semicrystalline block copolymers are already microphase-separated melt with a lamellar structure parallel to the substrate surface before crystallization. As a result, in the present work the crystallization rate affects the orientation of the lamellar microdomains in thin films of semicrystalline block copolymers via a different mechanism from that proposed by Reiter et al.41 Our findings provide two useful methods for controlling the perpendicular orientation of the lamellar microdomains in thin films of block copolymers when block copolymers are used as template for fabrication of nanodevices. If the nanodevice requires a larger size of lamellar domain, a longer block copolymer (thus larger microdomain size) with stronger segregation strength should be chosen. However, if the nanodevice requires a smaller size of lamellar domain, a shorter but crystalline block copolymer is a good candidate, and the thin film should be prepared at a lower crystallization temperature (faster crystallization rate). Conclusions The above results show that the orientation of the lamellar microdomains in SxEy block copolymers thin film on mica greatly depends on the molar mass of the block copolymers (segregation strength), temperature of solvent evaporation, and annealing. The nascent thin film of low molar mass block copolymer, S65E155, shows a multilayered structure parallel to the mica surface, but the high molar mass block copolymers, S156E358 and S199E452, exhibit a lamellar structure perpendicular to the mica surface. For high molar mass block copolymer, S156E358 and S199E452, annealing the nascent SxEy thin films results in morphological transformation into a multilayered structure parallel to the mica surface. For the thin film of S65E155 on mica, when the solvent is evaporated at 50 °C, a twodimensional spherulite structure with the lamellar microdomains perpendicular to the mica surface is observed. Annealing the two-dimensional spherulite structures leads to the formation of the multilayered structure parallel to the mica surface. In all SxEy thin films on mica, the stems of PE crystals are always perpendicular to the interface between the lamellar PE and PS microdomains. It is proposed that the perpendicular orientation of the microdomains results from kinetic trapping by the strong segregation strength or solidification of crystallization, while the parallel orientation of the microdomains is due to surface selectivity. Acknowledgment. This work was supported by the National Natural Science Foundation of China (20374046 and 20674073) and the New Century Supporting Program for the Talents by the Chinese Ministry of Education. Supporting Information Available: SAXS profiles of SxEy block copolymer in the bulk, AFM images of half-layered

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