Article pubs.acs.org/Macromolecules
Control of Directed Self-Assembly in Block Polymers by Polymeric Topcoats Abelardo Ramírez-Hernández,†,§ Hyo Seon Suh,†,§ Paul F. Nealey,†,§ and Juan J. de Pablo*,†,§ †
Materials Science Division, Argonne National Laboratory, 9700 South Cass Avenue, Argonne, Illinois 60439, United States Institute for Molecular Engineering, The University of Chicago, Chicago, Illinois 60637, United States
§
ABSTRACT: The morphology of a block copolymer thin film is particularly sensitive to its boundary conditions. Lithographic applications of block polymers in the microelectronics and memory device industries require formation of morphologies with perpendicularly oriented domains. Current fabrication targets envisage the creation of dense arrays of structures with domain sizes in the sub-10 nm regime. Such length scales can be reached by resorting to block polymers with highly incompatible blocks (and a large Flory−Huggins parameter, χ). High χ values, however, generally lead to large differences in the surface energies of the corresponding blocks, thereby interfering with formation of the sought-after perpendicularly oriented domains. In this work, a coarse grain model is used to develop a topcoat strategy that enables control of the orientation of block copolymer domains in highly incompatible block polymer materials. A systematic study of a wide range of polymeric material combinations is presented, and the conditions leading to optimal assembly of perpendicular morphologies are clearly identified. We consider the effect of molecular weight, block polymer film thickness and architecture, and degree of incompatibility. Our results are summarized in the form of generic phase diagrams that should serve as a guide for experimental deployment of the topcoat strategy put forth in this work.
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INTRODUCTION Block copolymer thin films are of considerable interest for lithographic fabrication of dense arrays of nanostructures.1,2 In the bulk, the palette of structures offered by these materials includes morphologies whose symmetry and length scale depend on block copolymer composition, chain architecture, incompatibility between blocks, and molecular weight.3,4 In thin films, the morphology of block copolymers is also strongly influenced by the boundary conditions; a distinct morphology is the result of a subtle balance between interfacial and bulk contributions to the total free energy.5 Complex morphologies can be directed to form, for example, when there is a mismatch between the symmetry of the block copolymer morphology and that induced by the boundary conditions.6−9 By designing proper boundary conditions, block copolymers can also be directed to self-assemble into morphologies useful for lithographic patterning applications, which generally depend on the formation of perpendicularly oriented domains.2,10−12 A widely used approach to direct the self-assembly of block copolymer thin films relies on the use of chemically patterned surfaces.13,14 In directed self-assembly (or DSA), specific boundary conditions can be encoded into the assembly process in a systematic manner, thereby leading to engineered periodic structures that include spherical, cylindrical, lamellar,15−17 and device-oriented morphologies (e.g., bends, jogs, and Tjunctions).18−20 Recent efforts have sought to identify materials or materials combinations that enable patterning of features with characteristic dimensions smaller than 10 nm.11,21−25 This © 2014 American Chemical Society
target imposes severe constraints on the possible block polymer materials that could be used for DSA; in particular, the incompatibility interaction parameter, χ (Flory−Huggins parameter), must be large. A large χ parameter is generally accompanied by a large difference in the surface energy of the blocks.22 This leads to segregation of lower-surface-energy domains to the free interface, thereby disrupting the perpendicular orientation of ordered features. This problem can be overcome by placing chemically modified surfaces on top of the block polymer films; nonpreferential or weakly preferential surfaces have been shown to result in through-film perpendicular domains.26−36 One drawback of this approach, however, is that in order to access the perpendicular structures, one most separate that top surface from the film, which can result in severe damage to the underlying sample. An alternative approach is to use chemical treatments to create a nonpreferential upper boundary; recently Bates et al.21 reported topcoats consisting of novel tunable interfacial energy copolymers containing maleic anhydride that can be processed using aqueous solutions for ease of deposition and removal from the surface of non-base-sensitive organic solvent-soluble block copolymer films. The usefulness of their approach was demonstrated by assembling block copolymers on a homogeneous substrate producing perpendicular domains Received: February 24, 2014 Revised: April 22, 2014 Published: May 15, 2014 3520
dx.doi.org/10.1021/ma500411q | Macromolecules 2014, 47, 3520−3527
Macromolecules
Article
without long-range order. In different work, Kim et al.24,25 also used chemical treatments to induce blocks to exhibit similar surface energies at the relevant annealing temperatures, thereby forming through-film perpendicular domains. Previously, we have experimentally demonstrated DSA of through film perpendicular structures in block copolymer films using topcoats.37,38 Block copolymers (PS-b-PMMA and PS-bP2VP) were directed to assemble on lithographically defined chemically nanopatterned surfaces with topcoats consisting of cross-linked polymer mats.37,38 The mats were composed of homopolymers corresponding to one or the other blocks of the copolymer and were thus preferential wetting toward one of the blocks or were random copolymers with composition to illicit nonpreferential wetting toward the copolymer. Cross-sectional SEM analysis revealed perpendicular through film structures with nonpreferential wetting topcoats, whereas mixed perpendicular and parallel oriented structures were observed for preferential topcoats.38 In addition, we showed that a topcoat of consisting of PMMA when used as a topcoat with PS-b-P2VP could induce DSA of perpendicular structure as the difference in interfacial energy between PMMA and PS and PMMA and P2VP was sufficiently small. Interestingly, the PMMA topcoat could be patterned in an additional lithographic step to define the regions of assembled block copolymer to be available for pattern transfer.38 In this work we explore the use of polymeric topcoats, to be placed on top of block polymer films, which could lead to formation of perpendicular domains. These topcoats could then be selectively removed by relying on solvents or exposure to radiation, thereby providing a simple and useful strategy for lithographic patterning with high-χ materials. By relying on theory and simulations, we are able to go beyond past studies and carry out a systematic study of the materials characteristics that influence the directed self-assembly of block polymers under topcoats. We are able to delineate the conditions leading to perpendicular assembly of block copolymer domains, and we examine the effects of molecular weight, film thickness, architecture, and degree of incompatibility.
/nb = ρ0 kBT
⎤ κ + (ϕA + ϕB + ϕC − 1)2 ⎥ ⎦ 2
Λμ(x , y) −z 2 /2ξ 2 0), and the B-stripe prefers block B (Λb > 0). The top hard surface is neutral for all polymers. The exploration of the equilibrium morphologies is performed by Monte Carlo simulations. The local densities are defined on a lattice with a spacing ΔL and computed from the beads’ positions by a particle-to-mesh interpolation. The configurations are sampled according to Metropolis criteria, Pacc = min[1, exp(−βΔ/ )], where Δ/ is the energy difference between and original and a trial configuration. It includes both intra- and intermolecular contributions as well as any change of the one-body potential energy induced by the pattern. The trial moves considered here include reptation-like displacements, local displacements of the beads, and translations of an entire chain. The parameters that define the melt have the following values: κNAB = 100, χABNAB = 73 and 5̅ = ρ0Re3/NAB = 112. Block polymer chains are discretized into NAB = 32 beads; their Re is set as the reference length scale. The bond length is fixed
MODEL The model adopted here is based on the standard Hamiltonian used in self-consistent-field theory to examine the morphology of block polymers.39−41 Such a model has been shown to provide a quantitative description of the experimental morphologies observed in block polymer DSA.42,43 Polymers are represented by flexible linear chains described by the discretized Gaussian chain model, where the position of the tth bead in the ith chain is denoted by ri(t). The system is composed by n = nAB + nC macromolecules, where nAB = diblock copolymer chains and nC = homopolymer chains in a topcoat layer. Each AB diblock copolymer consists of NAB beads. Each homopolymer C consists of NC beads. Intramolecular interactions acting along the chains are given by
∑∑
(2)
where ϕμ(r) is the local dimensionless density of beads of type μ (= A, B, C) and ρ0 is the bead number density. The repulsion between unlike monomers is represented by the first three terms in eq 2, where the Flory−Huggins parameters χμν describe the strength of the corresponding interactions. The last term in eq 2 is the Helfand’s quadratic approximation,45 which assigns a finite compressibility to the melt. The parameter κ is related to the inverse isothermal compressibility.46 To mimic the polymer thin films, polymer chains are confined by two impenetrable hard surfaces located at the z = 0 and z = Lz planes. The pattern on the bottom surface is represented by a one-body potential acting on each bead and depends explicitly on the position and type of the bead. Thus, a potential of the form = = ∑j
