Directed Assembly of Block Copolymers in Thin to Thick Films

May 13, 2013 - Daniel F. Sunday , Jiaxing Ren , Christopher D. Liman , Lance D. Williamson , Roel Gronheid , Paul F. Nealey , and R. Joseph Kline...
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Directed Assembly of Block Copolymers in Thin to Thick Films Adam M. Welander,† Gordon S. W. Craig,† Yasuhiko Tada,‡ Hiroshi Yoshida,‡ and Paul F. Nealey*,§ †

Department of Chemical and Biological Engineering, University of Wisconsin, Madison, Wisconsin 53706, United States Hitachi Research Laboratory, Hitachi Ltd., Hitachi City, Ibaraki 319-1292, Japan § Institute for Molecular Engineering, University of Chicago, Chicago, Illinois 60637, United States ‡

ABSTRACT: The extent to which lamellae-forming polystyrene-block-poly(methyl methacrylate) (PS-b-PMMA, bulk period L0 = 48 nm) can be directed to assemble on chemically nanopatterned striped surfaces (period LS), with few defects and with domains registered to and extending vertically away from the underlying pattern, was studied as a function of film thickness, commensurability between L0 and LS, and annealing temperature. Two regimes of behavior are identified: thin film ( 0.7 for the assembly it was classified as an example of Guided Assembly, and if S < 0.7 it was classified as unGuided Assembly. As shown in Figure 2, as |LS − L0| increased, the maximum thickness over which block copolymer domains could be directed to assemble decreased. For the thinnest films (t = 30 nm), the block copolymer was directed to assemble over the entire range of LS values, even when LS/L0 = 0.89. All of the

EXPERIMENTAL SECTION

Materials. Poly(styrene-b-methyl methacrylate) (PS-b-PMMA) block copolymer (Mn = 52-b-52 kg/mol, L0 = 48 nm) and hydroxyterminated PS brush, (PS-OH, Mn = 6 kg/mol) brush were purchased from Polymer Source, Inc. (Dorval, Quebec) and were used as received. Commercially available solvents (toluene and chlorobenzene) were used after filtering to remove large particles. Sample Preparation. The chemical patterns were prepared following the general technique used in previous work.2 PS-OH was covalently bonded to a silicon wafer, after spin-coating from a 1 wt % polymer solution in toluene, by baking the coated wafer at 190 °C for 18 h. PMMA photoresist was then applied to the PS brush and patterned using extreme ultraviolet interference lithography (EUVIL).17 Sections of the PS brush that were exposed after lithography and developing were chemically modified by a short oxygen plasma etch. The remaining photoresist was then stripped by sonication in B

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Figure 1. (a) Schematic representation of the EUV-IL process used to generate chemical patterns. Interference fringes were formed by coherent EUV light diffracted by a pair of gratings. The fringes were used to pattern photoresist on top of a layer of PS brush. The resist exposed by the interference fringes was developed, and the exposed PS brush was treated with oxygen plasma, resulting in a striped pattern with period LS in the PS brush that covers the width of the interference region. Beyond the interference region, the resist was completely exposed, and the PS brush was uniformly treated by the oxygen plasma. (b) Representative example of top-down SEM images of thin films of PS-b-PMMA that were spin-coated and annealed on a substrate treated with the process shown in (a). The substrate had three distinct regions: unexposed PS brush (pink border), oxygenplasma-treated PS brush (light blue border), and striped chemical pattern formed in the interference region (purple border).

Figure 2. Top-down SEM images of assembled PS-b-PMMA (bulk period L0 = 48 nm) on chemically patterned surfaces over a range of film thicknesses and pattern periods and annealed for 3 days at 190 °C. The order parameter S of each assembly is listed on its image. The border of each image denotes the classification of the system, with green, orange, and red representing Directed Assembly, Guided Assembly, and unGuided Assembly, respectively. Each image is 1 μm × 1 μm.

nm), the assembled morphology at the top surface of the film resembled a fingerprint pattern of perpendicularly oriented domains and was therefore classified as Unguided Assembly. On chemical patterns with LS = 47.5 or 50 nm, which were approximately commensurate with L0, films exhibiting Directed Assembly were achieved with t up to 250 nm. When t = 280 nm, the films assembled on patterns with LS = 47.5 or 50 nm were classified as Guided Assembly. In addition to determining S of the films and classifying the quality of each assembly, we also used FFT of the SEM images to measure the characteristic period at the free surface of the film. All of the Directed Assembly samples had polymer domains with the same period as the underlying chemical pattern. Remarkably, even at a film thickness of 240 nm, the period of the block copolymer domains L = LS (47.5 and 50 nm). In contrast films classified as Guided Assembly had L ≠ LS at the free surface.

films with t = 30 nm were classified as examples of Directed Assembly. As t was increased, on the patterns with LS = 42.5 nm, which had the greatest incommensurability with L0, the assembled pattern changed from well-ordered lines (Directed Assembly, t = 30 nm) to lines with defects (Guided Assembly, t = 40−72 nm). The defects in the films with LS = 42.5 nm and t = 40−55 nm were primarily dislocation dipoles. This type of defect is consistent with those described previously by Kim et al. in 60 nm thick films of PS-b-PMMA (with L0 = 48 nm, as here) assembled on chemical patterns with LS = 45 nm.19 As shown in Figure 2, as t was increased to 70 nm on the substrates with LS = 42.5 nm, the nature of the defects changed such that the assembled domains were only partially aligned with the underlying chemical pattern. In thicker films (t ≥ 90 C

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We plotted the data as shown in Figure 3 to illustrate several important concepts. The 45° line is the boundary between guided and Unguided Assembly because when ΔFLS≠L0 > ΔFint, Unguided Assembly will occur. For a given ΔFLS≠L0, as ΔFint is increased from ΔFint < ΔFLS≠L0 to ΔFint = ΔFLS≠L0, the assembled block copolymer will exhibit Guided Assembly, with some defects still present. To achieve defect-free Directed Assembly, ΔFint must be further increased by an amount δ, which corresponds to the additional energy required to suppress low-energy defects. The dashed line in Figure 3 represents ΔFint = ΔFLS≠L0 + δ, the transition between Guided Assembly and Directed Assembly. Based on the data shown in Figure 3, δ would correspond to an additional ∼0.05 erg/cm2 beyond the line of equal energy. The solid and dashed lines show that for thin (∼30 nm) films very good assemblies were achieved with LS differing from L0 by up to 11%. This range of LS for which the same block copolymer is directed to assemble with a range of feature sizes exactly equal to LS/2, in other words, with feature sizes determined by the chemical pattern period, not the bulk lamellar period, mirrors the previous results reported by Edwards et al. with 40 nm thick films of lamellae-forming PS-bPMMA annealed for 70 h at 190 °C on their strongest interacting chemical patterns.2 As the film thickness t was increased, the range of LS over which Directed Assembly occurred decreased. The decrease in the range of LS with increasing t is interpreted as a consequence of averaging ΔFint over the entire volume of the film, such that in thicker films, the energy of the chemical pattern (the directing aspect of the system) on a per chain basis is less than that in thinner films on the same pattern. The kink in the dashed line at t = 72 nm in Figure 3 coupled with the continued formation of nearly perfect assemblies in films up to 240 nm thick suggests that in films thicker than a critical thickness tc (>72 nm in the system here) a different assembly mechanism occurs than in films with t < tc. When t < tc, the results here suggest that cooperative assembly of domains occurs throughout the entire film thickness. The cooperative nature of the assembly on 1:1 patterns was previously described by Edwards et al. under similar annealing conditions in experiments and molecular simulations.3 Edwards et al. showed that at early stages of assembly the lamallaeforming block copolymers first reorganized to wet the pattern stripes with the correct blocks near the substrate but formed structures with hexagonal symmetry farther away from the substrate. As the process continued, transient structures eventually evolved to yield perpendicularly oriented lamellae, aligned with and registered to the chemical patterns. Fingerprint morphologies at the top surface did not appear independently of the morphology driven by the chemical pattern, even at short times.3 To investigate the thick film regime of assembly in greater detail, we imaged the morphology present at the free surface of films as a function of time. Figure 4 shows the results for films with thickness 580 nm annealed on chemical patterns with LS ∼ L0 for times ranging from 1 min to 3 days. The annealing temperature was increased from 190 to 230 °C to improve the system dynamics. As shown in Figure 4, at short annealing times (less than 16 h) the top surface showed domains perpendicular to the surface that were organized into randomly oriented grains, and the grain size of domains at the surface increased with annealing time. After annealing for 16 h, the

In Figure 3, we plotted the assembly classifications in terms of the difference in free energy per chain for a block copolymer

Figure 3. Assembly classification data plotted as a function of interfacial energy at the chemical pattern (ΔFint, eq 3) and energy required for the block copolymer to equilibrate with period LS different from its bulk period L0 (ΔFLS≠L0, eq 2). The solid line is where ΔFint equals ΔFLS≠L0 and represents the transition between Guided Assembly and Unguided Assembly. The dashed line represents the division between perfect Directed Assembly films and Guided Assembly films for the films with t ≤ 240 nm.

assembly with a period L = LS other than L0 (ΔFLS≠L0) versus the difference in energy per chain derived from interfacial interactions (ΔFint) for a film on a chemically patterned surface versus a neutral unpatterned surface. ΔFint was calculated using the interfacial energy acting at the film/substrate boundary, assuming perfectly assembled domains with L = LS on chemically patterned surfaces, and vertically oriented domains on neutral surfaces. To express ΔFint on a per chain basis, the difference in interfacial energy is averaged over the volume of the system or, in this case, the thickness of the film per unit area of the film. The value of ΔFint and ΔFLS≠L0 for each system was calculated using eqs 2 and 3:2 ΔFLS ≠ L0 kT

=

⎛ χ ⎞1/2 ⎛ 1 3 1⎞ (LS2 − L0 2) + 2aN ⎜ ⎟ ⎜ − ⎟ 2 ⎝ 6 ⎠ ⎝ LS L0 ⎠ 8a N (2)

where the constant a is the characteristic length, the parameter χ is the thermodynamic interaction parameter, and N is the degree of polymerization and ΔFint Mn = (γint ,p − γint ,n) kT ρNAtkT

(3)

where γint,n and γint,p are the interfacial energies of the block copolymer interacting with a neutral or chemically patterned surface, respectively, Mn is the molecular weight, and ρ is the density. All of the parameters in eqs 2 and 3 are known or can be calculated except for γint,p − γint,n (= Δγint). However, if we assume that ΔFint = ΔFLS≠L0 at the transition between Unguided Assembly and Guided Assembly, we can fit to the data shown if Figure 2 at the transition to calculate that Δγint = 0.15 ± 0.01 erg/cm2. D

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Figure 5. SEM image of a cross section of a 380 nm thick film of PS-bPMMA assembled on a chemical pattern. The PMMA domains have been removed, leaving a series of PS domains that are parallel to each other and perpendicular to the substrate. The collapse of the lamellae is an artifact of the sample preparation and imaging.

= 620 nm, it was still possible to see differences in the top surface of the film due to the chemistry of the underlying substrate. SEM images of the top surface of the PS-b-PMMA film over the PS brush showed perpendicular domains, while SEM images over the PMMA wetting region showed no contrast, implying that only one block of the block copolymer was present at the air interface. Remarkably, the films assembled over the chemical pattern (Figure 6a, part C) could still be classified as Directed Assembly over the patterned area, and there was a clear demarcation between fingerprint and Directed Assembly at the boundary of the patterned area. The results shown in Figure 6 represent almost a 3-fold increase in the thickness at which Directed Assembly could be achieved, compared to annealing at 190 °C. The lamellae in the assembled 620 nm thick film had an aspect ratio greater than 25. Further increasing the film thickness to 785 nm allowed us to test the limits of thick film Directed Assembly for the specific annealing conditions of 230 °C for 3 days. The top surface of the 785 nm film showed no influence of the underlying surface chemistry, and it was impossible to distinguish between the three different substrate sections, unlike the films shown in Figures 2 and 6a. We hypothesize that in the assembly mechanism that occurs in films thicker than tc the block copolymer morphology nucleates at both the air−polymer surface as well as the interface with the chemical pattern. In this hypothesized assembly mechanism, grains form and coarsen at the free surface, while at the chemical pattern interface there is just one grain over the entire pattern. At the free surface, the grains adopt a random orientation and grow laterally and downward. Nucleation at the interface with the chemical pattern is not random, and the grain propagates upward. With sufficient annealing and block copolymer mobility (enhanced by increasing the annealing temperature), the grains from both interfaces eventually come together in the middle of the film, as shown in Figure 5, and the grains nucleated at the free surface may align with the underlying, commensurate chemical pattern. The thermodynamic driving force that aligns the domains initiated at the free surface with the domains initiated at the chemical pattern interface is the elimination of defects. Thus, in films with t > tc, the ability to achieve Directed Assembly will depend in large part on the kinetics of the system. The idea that a different assembly mechanism could prevail in thicker films matches previous work that showed that as the film thickness increased, the formation of the block copolymer morphology, studied in terms of the kinetics of growth of grains within the film, can no longer be described adequately by either two-dimensional (2D), thin film results or unconstrained threedimensional (3D) bulk results. For 2D films of lamellar block copolymer films, Ruiz et al. reported that the two-dimensional

Figure 4. SEM images of 580 nm thick PS-b-PMMA films annealed at 230 °C for various times. Samples annealed for less than 16 h showed no indications of surface patterning, only changes in the grain size of domains at the air interface. After 16 h of annealing, traces of the effect of the underlying chemical pattern appeared, but the assembly still showed many defects. After 3 days of annealing, the domains assembled with a high degree of order. Each image shows a surface area of 1.5 μm × 1.5 μm.

domains at the surface suddenly began to organize with respect to the underlying chemical pattern, more than a halfmicrometer below the surface. After 3 days of annealing, the lamellae at the surface of the film were perfectly registered with the underlying chemical pattern, and the films could be classified as examples of Directed Assembly. The ability to assemble through 580 nm thick films at 230 °C compared to 240 nm at 190 °C is attributed to the 45-fold faster dynamics of the system at the higher temperature.20 One concern with top-down SEMs is whether the structures seen at the free surface of the film continue through the thickness of the film. In previous work on the assembly of cylinder-forming PS-b-PMMA on nonpreferential surfaces, Han et al. showed that cylindrical domains perpendicular to the substrate could persist across submicrometer thick films, but well-defined grain boundaries between domains nucleated at the surface and domains nucleated at the substrate were sometimes present.21 For self-assembly of lamellae-forming PSb-PMMA on a strongly PS-preferential substrate, Xu et al. observed that the lamellar domains were parallel to the substrate at the chemical pattern and at the free surface but that a mixed domain orientation appeared in the center of the film.22 In this work, to determine if the domains in the assembled thick films that we classified as Directed Assembly were perpendicular to the substrate and traversed the thickness of the film, we used an SEM to image a cross-section of a 380 nm thick assembled film, as shown in Figure 5. The PMMA domains were purposely removed or degraded in an e-beam to improve image contrast. The remaining PS domains were parallel to each other across the thickness of the film. The collapse of some of the PS domains was a result of the PMMA removal for imaging. As shown in the work associated with Figure 4, by simply increasing the annealing temperature, we can significantly increase the block copolymer film thickness over which the chemical surface pattern can direct the assembly. To test the extent of t over which we could achieve Directed Assembly, we repeated the substrate configuration shown in Figure 2, but with a 620 nm thick PS-b-PMMA film annealed at 230 °C. At t E

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PS-b-PMMA over the same temperature range.20 For example, χ of PS-b-PMMA varies by only 4% over the range of annealing temperatures used in this study, such that the value of χN remains relatively constant in the intermediate segregation limit.25 In the case of γ, Wu showed that the γ values of PS and PMMA are nearly equal from 100 to 200 °C. Mansky et al. found that at 170 °C Δγ for the blocks of PS-b-PMMA was less than 1%27 and subsequently determined that Δγ of PS and PMMA decreased with increasing temperature up to 250 °C.28 Given that on the top surface of the 785 nm thick film annealed at 230 °C we observed fingerprint patterns, which are a sign of self-assembly on a nonpreferential (Δγ = 0) surface, we can conclude that Δγ at 230 °C was ∼0, and therefore Δγ could only have varied by at most ∼1%. The small variation in χ and γ over the temperature range in our study reinforces our hypothesis as it relates to the importance of dynamics in the assembly of thick films. The importance of dynamics leads us to conclude that the image in Figure 6b showed a nonequilibrium structure in the 785 nm thick film and that with continued annealing, or annealing at higher temperature, Directed Assembly may be achieved in the film. However, the interplay between thermodynamics and kinetics, and the nucleation and growth of grains, as well as the elimination of defects and trapped defect sites may be more complicated.

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CONCLUSION In summary, we have demonstrated that two regimes of assembly behavior exist for the assembly of block copolymer films: thin film and thick film. In our work here the critical thickness that separated these regimes was ∼72 nm. In the thin film regime, the assembly process is driven from the chemical pattern−block copolymer interface upward through the film, and the system follows the phenomenological models used in previous studies of Directed Assembly of block copolymers. In the thick film regime, nucleation and growth of grains at the free surface of the film becomes decoupled from Directed Assembly near the chemically patterned surface. In the thick film regime, polymer dynamics has a great effect on the ability to achieve high quality assemblies. When annealing at 190 °C for 3 days, the thickness limit for achieving nearly defect-free assemblies with our system was 240 nm. In films thicker than 240 nm, annealing at higher temperatures accelerated the assembly dynamics, such that domains could nucleate at the free surface of the film, grow perpendicularly into the film, and merge and align with domains that nucleated at the interface with the chemical pattern. For example, by annealing at 230 °C it is possible to direct the assembly of block copolymer domains in films up to 620 nm thick, yielding domains oriented perpendicular to the substrate and having very few defects in their assembly. The results of the accelerated assembly process were well-ordered, high aspect ratio domains, which could prove useful in a number of applications, including catalysis, separation membranes, and solar cells.

Figure 6. (a) Low-magnification image of a 620 nm thick film of PS-bPMMA annealed at 230 °C on a substrate coated with a chemically patterned area (indicated by the purple box), an oxygen plasma modified brush (light blue box), and an unmodified PS brush (pink box). Three high-magnification images, from approximate areas denoted by the letters A, B, and C, show details of the top surface over the PS brush, at the boundary of PS-brush and the chemical pattern, and directly over the chemical pattern, respectively. (b) Lowmagnification image of 785 nm thick film of PS-b-PMMA over a chemically patterned substrate. The top surface shows perpendicular block copolymer domains with no correlation with the underlying chemical pattern.

correlation length (ξ), associated with grain size of assembled domains observed at the free surface of the film, can be described by the power law ξ(t) ∼ tϕ, where ϕ (growth exponent) is 0.11 ± 0.01 for a film thickness of one-half the bulk lamellar period L0, while film thicknesses of L0 exhibited nearly constant ξ.23 Garetz et al. measured 3D correlation lengths, κ and ω, in bulk lamellar block copolymers. From their bulk results ϕ can be estimated to be between 0.066 and 0.131,24 but bulk studies do not include the effect of the additional nucleation and growth sites provided by the two surfaces of a thick film. In consideration of our hypothesis that improved assemblies in thick films at higher annealing temperatures stems primarily from improved dynamics, we evaluated other factors, such as χ and γ, that impact block copolymer assembly and are usually temperature dependent. In the case of PS-b-PMMA, the change of χ and γ with temperature is very small when compared to the 2−3 order of magnitude increase seen in the assembly speed of

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest. F

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ACKNOWLEDGMENTS This work was supported by the Semiconductor Research Corporation (SRC) and the NSF through the Nanoscale Science and Engineering Center (NSEC, DMR 0832760). This work made use of the facilities and staff at the UW center for Nanotechnology, the Synchrotron Radiation Center at UW Madison, and the Swiss Light Source at the Paul Scherrer Institute.

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REFERENCES

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