Rapid Ordering in “Wet Brush” Block Copolymer/Homopolymer Ternary Blends Gregory S. Doerk* and Kevin G. Yager Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, New York 11973, United States S Supporting Information *
ABSTRACT: The ubiquitous presence of thermodynamically unfavored but kinetically trapped topological defects in nanopatterns formed via self-assembly of block copolymer thin films may prevent their use for many envisioned applications. Here, we demonstrate that lamellae patterns formed by symmetric polystyrene-block-poly(methyl methacrylate) diblock copolymers self-assemble and order extremely rapidly when the diblock copolymers are blended with low molecular weight homopolymers of the constituent blocks. Being in the “wet brush” regime, the homopolymers uniformly distribute within their respective self-assembled microdomains, preventing increases in domain widths. An order-of-magnitude increase in topological grain size in blends over the neat (unblended) diblock copolymer is achieved within minutes of thermal annealing as a result of the significantly higher power law exponent for ordering kinetics in the blends. Moreover, the blends are demonstrated to be capable of rapid and robust domain alignment within micrometer-scale trenches, in contrast to the corresponding neat diblock copolymer. These results can be attributed to the lowering of energy barriers associated with domain boundaries by bringing the system closer to an order−disorder transition through low molecular weight homopolymer blending. KEYWORDS: block copolymer, directed self-assembly, homopolymer, coarsening kinetics, ternary blends, pattern transfer
T
However, these methods frequently involve specialized equipment that may not be readily scaled up for high-volume manufacturing. They may also depend critically on the BCPs34 and substrates used.30,31 Relatively less attention has been paid to the impact of BCP thin-film composition on topological ordering of the domains. Arias-Zapata et al. demonstrated rapid directed self-assembly across 300 mm wafers by thermal annealing of cylinder-forming polystyrene-block-polydimethylsiloxane (PS-b-PDMS) blended with molecular plasticizers,35 but the evaporation of the plasticizer increases process complexity. Kim et al. showed that blending a small molecular weight “neutral” random copolymer of polystyrene (PS) and poly(methyl methacrylate) (PMMA) with lamellae forming PS-b-PMMA BCPs reduced the segregation strength. This resulted in a large increase in grain size for the smallest BCP studied, attributed to a “defect melting” mechanism.36 On the other hand, BCP/homopolymer ternary blends allow tuning of domain size and period37,38 and enable directed self-assembly into nonregular device-oriented geometries through homopolymer redistribution within the domains,39,40 but their impact on pattern ordering kinetics remain unclear.
hrough microphase separation of the constituent polymeric blocks, block copolymer (BCP) thin films self-assemble into highly uniform, periodic patterns potentially useful in the fabrication of microelectronic devices,1,2 patterned magnetic media,3,4 broadband antireflection surfaces,5 and superhydrophobic structures.6 In many cases, a substantial barrier to further development is the lack of long-range pattern order due to the presence of topological defects such as dislocations, disclinations, and grain boundaries.7 These defects are thermodynamically unstable,8,9 indicating that their ubiquitous presence is a product of slow defect annihilation kinetics. To generate globally ordered patterns via BCP self-assembly, researchers have developed a number of techniques to direct the organization of BCP domain patterns including shear,10 electric fields,11 magnetic fields,12 zone annealing,13,14 chemical patterns,15−18 and topographical templating.19−22 Nevertheless, defects persist even when using highly refined techniques for directed self-assembly,23 posing a severe impediment to the implementation of block copolymers in manufacturing processes with exacting defectivity constraints, as in the case of integrated circuits. To overcome this obstacle, researchers have explored alternatives to thermal annealing for accelerating BCP ordering that are complementary to directed self-assembly strategies. These include solvent vapor annealing,24−26 immersion annealing,27 rapid thermal processing,28 microwave annealing,29,30 and photothermal annealing.31−34 © 2017 American Chemical Society
Received: August 29, 2017 Accepted: December 1, 2017 Published: December 1, 2017 12326
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
www.acsnano.org
Cite This: ACS Nano 2017, 11, 12326−12336
Article
ACS Nano
Figure 1. Impact of composition on pattern order for blends with ∼3 kg/mol PS and PMMA homopolymer, after thermal annealing for 5 min at 250 °C. SEMs of lamellar patterns for (a) neat 36 kg/mol BCP and (b) the same 36 kg/mol BCP blended with 50% (w/w) hompolymer. (c) False color orientation maps for the 36 kg/mol and 74 kg/mol BCPs with varying homopolymer mass fractions. The scale bar represents 2 μm. (d) Plot of the correlation length, ξ (solid lines), and repeat period, LB (dashed lines) as a function of homopolymer mass fraction. Blue points represent data from the 36 kg/mol BCP, while red points represent data from the 74 kg/mol BCP. Lines are guides to the eye.
and PMMA domains can be readily oriented vertically on appropriately treated substrates, owing to the small surface tension difference between the two blocks at typical annealing temperatures (180−250 °C).42 To promote vertical domain orientation through the entire film thickness, Si substrates were first grafted with a PS-r-PMMA random copolymer “neutral brush” exhibiting a net surface energy intermediate between that of PS and PMMA. Diblock copolymers microphase separate to form a lamellar morphology when the volume fraction each block is close to 50%, with a repeat period (L0) that is a function of the total degree of polymerization (N) and hence the total molecular weight. The lamellar PS-b-PMMA BCPs employed here have number-average molecular weights (Mn) of 36 kg/mol (N = 353) and 74 kg/mol (N = 726), exhibiting repeat periods of ∼29 nm and ∼40 nm, respectively. Figure 1a shows a scanning electron micrograph (SEM) of a “fingerprint” pattern, characteristic of vertically oriented BCP
Here we investigate the pattern ordering kinetics of lamellaeforming PS-b-PMMA/PS/PMMA ternary blends. We find that blending small molecular weight PS and PMMA homopolymers that uniformly distribute within respective domains of significantly larger molecular weight (>10×) PS-b-PMMA BCPsknown as a “wet brush”41 system since the homopolymer “wets” the BCP chainssubstantially enhances pattern ordering kinetics, leading to an order-of-magnitude increase in the degree of lateral order during thermal annealing. Based on analyses of relative ordering kinetics scaling, we propose that the dramatic enhancement in these blends can also be ascribed to a reduction in the energy barriers for defect annihilation promoted by the addition of homopolymer chains.
RESULTS AND DISCUSSION PS-b-PMMA is a convenient BCP material for patterning applications and for fundamental studies, since nanoscale PS 12327
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano
when the average homopolymer molecular weight is further increased to ∼12 kg/mol (11.5 kg/mol PMMA, 12.5 kg/mol PS).
lamellae that have no in-plane orientational preference, obtained by annealing a 36 kg/mol BCP film at 250 °C for 5 min. Homopolymers of PS (Mn = 3.5 kg/mol; NPS = 34) and PMMA (Mn = 3 kg/mol; NPMMA = 30) were blended in solution with the BCPs in equal proportions (by weight), so that the total homopolymer mass fraction equals ∼2 times the mass fraction of each individual homopolymer. The nearly equal molecular weights render the effective homopolymer degree of polymerization (NH = NPSNPMMA ) of 32 that is close to the degree of polymerization for either homopolymer. Self-assembly of the ternary blend (BCP with 50% homopolymer) after annealing at 250 °C for 5 min generates a line pattern with a significantly higher degree of long-range order (Figure 1b). False-color maps of the local line pattern orientation, generated through software-based image analysis (see Methods), provide a useful visual indication of the lamellar grain sizes. Figure 1c shows a series of such maps corresponding to blends with different homopolymer mass fractions for the 36 kg/mol (upper row) and 74 kg/mol (lower row) BCPs annealed at 250 °C for 5 min. It is evident that grain size increases with increasing homopolymer concentration for both cases, more dramatically in the case of the 36 kg/mol BCP. Though a substantial volume of homopolymer (>50%) can be blended in these ternary systems while maintaining the morphology of the neat BCP, there is a limit to the homopolymer fraction above which a loss of microdomain order occurs. At a 70% homopolymer mass fraction, the apparent grain sizes for 36 kg/mol BCP-homopolymer blend decreases to a level lower than the neat BCP. A higher resolution SEM (see Supporting Information Figure S1) indicates that while there is some texture present, the characteristic lamellae domains are no longer visible and the morphology appears to be disordered. The 74 kg/mol BCPhomopolymer blends generate a similar disordered or poorly ordered morphology at a 90% homopolymer mass fraction (see additional SEM in Figure S1). The grain size can also be estimated in a more quantitative manner based on the correlation length (ξ), which is a measure of the average spatial persistence of local pattern orientation in images of lamellae nanopatterns, calculated using software-based image analysis (see Methods). Figure 1d shows the trend in correlation length as a function of homopolymer mass fraction for the samples characterized in Figure 1c. Strikingly, ξ for the 36 kg/mol BCP 50% homopolymer blend is an order-ofmagnitude larger than that for the neat 36 kg/mol BCP, despite identical annealing conditions. Also plotted in Figure 1d is the repeat period of the blends (LB). Though ξ increases significantly with homopolymer addition, LB changes negligibly except in the case of the disordered 36 kg/mol blend. Previous studies have shown that homopolymer addition often swells BCP domains in both the bulk43,44 and thin films,37 though for very low NH values domain shrinkage can occur.38 This tuning of repeat spacing enabled by blending homopolymers with BCPs provides a convenient way to ensure commensurability between BCPs and templates used in directed self-assembly without resorting to cumbersome synthesis of new BCPs. When the average molecular weight of the homopolymers blended with the 36 kg/mol BCP is increased to 6 kg/mol, LB increases monotonically with increasing homopolymer mass fraction as a result of domain swelling (Figure 2a). The swelling is more substantial
Figure 2. Effect of homopolymer molecular weight on lamellar repeat spacing. (a) Repeat spacing as a function of homopolymer mass fraction for ∼3 kg/mol, ∼6 kg/mol, and ∼12 kg/mol homopolymer pairs blended with the 36 kg/mol BCP. Lines are guides to the eye. (b) Derived parameter βeff (defined in eq 2) as a function of the ratio of homopolymer to BCP degree of polymerization, α. Lines are linear fits according to the equations provided.
The swelling of BCP domains by their respective homopolymers has been described in terms of “wet brush” and “dry brush” systems38,41 based on the ratio of NH to N of the BCP (α = NH/N). For the dry brush case (larger α), the requirement to maximize conformational and translational entropy for both the homopolymer and BCP causes the homopolymer to localize to the domain interiors, in turn expanding the domain width. In the wet brush case, however, smaller α homopolymer chains distribute uniformly within their respective domains, where the loss in BCP conformational entropy by homopolymer solubilization is offset by the gain in translational entropy by the homopolymer. In this case, the homopolymer increases the interfacial area per block copolymer chain at the interdomain interface and may screen some of the interaction between the two polymer blocks there. Liu et al. developed an empirical model to describe the scaling of domain size with the volume fraction of homopolymer (ϕH): 12328
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano
Figure 3. Comparison of ordering kinetics between neat 36 kg/mol BCP and a 50% (w/w) blend of the same BCP with ∼3 kg/mol homopolymer. (a) False color orientation maps for films annealed at 210 °C. The scale bar represents 2 μm. (b) Plot of ξ as a function of annealing time at 220 °C. The equations represent fits to a power law relationship, ξ = Atν.
LB = (1 − ϕH)−β L0
PS-b-PMMA BCP blended with 1.5 kg/mol PS and PMMA homopolymers.38 The trends in domain scaling reported here disagree with those for lamellar PS-b-PMMA/PS/PMMA ternary blend thin films previously reported by Liu et al.37 However, the blend thin films described in that report featured comparatively larger homopolymers (α ≥ 0.2). To extend the analysis to α < 0.2, data were obtained from previous reports including lamellar ternary blends based on polystyrene-block-poly(2-vinylpyridine) (PS-b-P2VP/PS/P2VP)44 and polystyrene-block-polyisoprene (PS-b-PI/PS/PI)43 where self-assembly occurred in the bulk rather than in thin films. On the other hand, reports of lamellae forming binary homopolymer BCP blends have exhibited approximately constant or decreasing repeat spacings in cases where α < ∼0.12.45,46 Moreover, thin films of cylinder forming BCP-homopolymer ternary blends also exhibit smaller repeat spacings with respect to the neat BCP in the case of very small homopolymers (α < ∼0.02).38,47 Despite the variation of domain scaling behavior described in these reports, they follow a consistent explanation with regard to the effect of homopolymer blending on domain scaling: Relatively large molecular weight homopolymers localize to the center of domains to maximize configurational entropy (“dry brush” regime), causing significant domain swelling. As the homopolymer molecular weight is reduced, the number of chains is increased for the same homopolymer volume fraction, and the translational entropy grows relatively more important. This leads to a broadening of the homopolymer distribution within the domain, reducing swelling. Eventually, for α < ∼0.1 the homopolymers are small enough to distribute much more uniformly throughout the domain to maximize combined translational and configurational entropy, thereby “wetting” the
(1)
where the dependence of β on α was fit to a linear relationship.37 According to the data used for the linear fit, β approaches 0.477 in the limit α = 0, indicating that even the smallest homopolymers used for blending would swell domains, causing an increase in the repeat period, in contrast to the data reported here. To assess the relationship between homopolymer NH and domain swelling, an effective β (βeff) was extracted from the measured LB for each sample by rearranging eq 1: βeff =
ln(L0) − ln(LB) ln(1 − ϕH)
(2)
For PS-b-PMMA, ϕH is nearly equivalent to the mass fraction of homopolymer. This βeff parameter can be qualitatively interpreted as the “responsiveness” of the system’s repeatspacing to the blending, that is, a small βeff value indicates a weakly dependent system, while a large value indicates a system that swells strongly as blending increases. This parameter is plotted vs α in Figure 2b using data obtained from SEMs of both BCPs blended with the 3 kg/mol, 6 kg/mol, and ∼12 kg/ mol homopolymer pairs. As expected, blending larger molecular weight chains (larger α) leads to a stronger swelling response. Data for the two BCPs appear to follow different scaling relationships with respect to homopolymer addition. Interestingly, a linear fit to the data for either BCP implies that for homopolymers smaller than ∼3 kg/mol, the domains of the blend will actually contract, a finding that is consistent with the smaller domains reported for thin films of a cylinder-forming 12329
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano BCP brush (“wet brush” regime). This reduces swelling further and also increases the area per BCP chain at the domain interface as they are spaced apart by homopolymer chains residing there. This in turn allows for the relaxation and concomitant contraction of the BCP chain of the opposing block, reducing the overall repeat spacing.48 Since small molecular weight homopolymer uniformly distributes within each block, the effect is symmetric in a ternary blend. The effects of thin-film confinement may also perturb the effect of homopolymer molecular weight on blend domain scaling in ways that are not fully understood. Nevertheless, the ∼3 kg/ mol homopolymer-BCP blends discussed here exists at the crossover between microdomain swelling and shrinking and so maintains a constant repeat spacing as homopolymer volume fraction is increased. The ability to maintain a consistent repeat spacing while achieving a substantially higher degree of ordering is especially advantageous as it could shorten annealing times and reduce defectivity in directed self-assembly using templates designed for the corresponding neat BCP. The significant increase in grain size for the blend films implies an increase in the kinetics of grain coarsening compared to the neat BCP. Time series studies at fixed annealing temperature confirm this is the case, as shown by representative orientation maps for thin films of neat 36 kg/mol BCP and the corresponding 50% homopolymer blend (3 kg/mol homopolymer) annealed at 210 °C for different time periods in Figure 3a. For neat BCP films, coarsening is exceedingly slow, with the grain size increasing by a barely perceptible amount during the first hour of annealing. On the other hand, for the blend films the progressive increase in grain size is clear across the entire time span. In both cases, no decomposition, dewetting, or loss of microphase order was observed, and the repeat spacing remained constant. For BCP thin films, grains coarsen according to a power law of the form ξ = Atν, where A is a prefactor encompassing the temperature dependence, t is the time, and ν may be referred to as the growth exponent.49,50 Figure 3b shows a log−log plot of ξ vs time for thin films of the neat 36 kg/mol BCP and the corresponding 50% homopolymer blend (3 kg/mol homopolymer) annealed at 220 °C. Each set of data was fit to this power law relationship. Though the prefactors are similar, the growth exponents are drastically different. The exponent for the neat BCP is ν ∼ 0.04. This very low growth exponent, indicative of extremely inhibited grain coarsening, has been observed previously for lamellae thin films49,51 and has been attributed to the effect of domain topology on defect diffusion and annihilation.49 In contrast, the blend films exhibit coarsening with a larger exponent of ν ∼ 0.29. Similar differences between the growth exponents for the 50% homopolymer blend (based on the 36 kg/mol BCP with the 3 kg/mol homopolymers) and the neat copolymers are observed across a broad temperature range (Figure 4a). Indeed, the blend films have an average ν of 0.22 ± 0.05, close to the growth exponents frequently observed in cylindrical morphology BCP coarsening.49,52 For longer annealing times, the grain size for the 50% homopolymer blend may plateau, as shown through analysis of a sample annealed under vacuum for more than 21 h at 210 °C (Figure S2); on the other hand, it should be noted that at 250 °C the blend films separated into an ordered and a disordered phase for longer annealing times (>300s), preventing robust correlation length analysis (see Figure S3 for an exemplary SEM image). This suggests that the highly loaded ternary blends studied here are close to a disordering transition.
Figure 4. Power law fit parameters for ordering of the neat 36 kg/ mol BCP (open red circles) and a 50% blend (w/w) of the same BCP with ∼3 kg/mol homopolymer (filled blue squares) for temperatures from 190 to 250 °C. The fit parameters are (a) the ordering exponent, ν, and (b) the prefactor, A.
Dynamic properties of polymer melts like viscosity and diffusivity are associated with the amount of available “free volume”, commonly correlated through the empirical logarithmic Williams−Landel−Ferry (WLF) equation,53−55 which can be referenced to the glass transition temperature (Tg). Small molecular weight homopolymer additives can act as plasticizers, reducing the glass transition temperature (Tg) and melt viscosity of a blend compared to the starting polymer. Differential scanning calorimetry performed on the 36 kg/mol BCP, the 3 kg/mol PMMA homopolymer, and the 3.5 kg/mol PS homopolymer yield glass transition temperatures (Tg) of 101 ± 1 (PS phase), 77 ± 1, and 74 ± 1 °C, respectively (see Figure S4). Using the Fox equation,56 the Tg for the blend can be estimated as 88 ± 2 °C. This ∼13 °C reduction in Tg is expected to play a role in enhancing chain and defect mobility in the ternary blends. The associated mobility enhancement should lead to a change of the power law prefactor A for ξ growth, especially since it has been suggested that this depends in part on the BCP diffusivity prior to microphase separation,49 where segregation strength plays a diminished role. However, we measured values of A for the blends that are roughly similar to or less than the values of A for the neat BCP in the range of 190−220 °C, though there is a substantial increase in the blend A value around 230 °C (Figure 4b). Considering the inherently pathway-dependent ordering of BCP domain patterns,33 more detailed investigations of initial ordering in thermally annealed BCPs generally would help disentangle the role of blending in early stage pattern coarsening. Annealing a BCP blend nominally similar to the neat BCP but with a reduced Tg would be equivalent to annealing the neat BCP at a temperature higher by the same magnitude as the 12330
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano Tg reduction; this temperature difference is covered within the experimental temperature range and no notable increase in the growth exponent is observed for the neat thin films, indicating that the reduction in Tg alone cannot adequately account for the increased growth exponent for the blend thin films. On the other hand, differences in growth exponents for thin-film BCP coarsening have been considered as manifestations for different mechanisms of defect movement and annihilation.49,57 We posit that the dramatically enhanced growth exponent can be attributed to the fact that blends are in close proximity to a disordering transition induced by the uniform distribution of homopolymer. Whether the addition of homopolymer promotes enhanced topological ordering depends critically on the molecular weight of the blended homopolymers. Consider the conceptual schematic isoplethic (constant total A/B polymer volume ratio) phase diagram for a “dry brush”, lamellar ternary diblock copolymer/homopolymer blend depicted in Figure 5a. The segregation strength for the BCP,
ymers to a single disordered homopolymer phase is located at χN = 2/α at ϕH = 1. Within mean-field theory, the macrophase, microphase, and disordered phase meet at an isotropic Lifshitz point, marked in Figure 5a by a black circle.59 However, it has been recognized experimentally that the Lifshitz point will occur at a ϕH value within a narrow channel between the macrophase separated region and the lamellar phase where the BCP molecules stabilize a microemulsion phase (μE),60 but not a lamellar domain structure. Starting with an ordered neat BCP, a distinct interface between A- and B-rich regions is preserved as homopolymer is added for all cases: the lamellar microdomains, the microemulsion phase, and the macrophase separated polymer domains. When χNH is 2. (b) “Wet brush” phase diagram in which χNH < 2.
χN, where χ is the Flory−Huggins interaction parameter, is on the ordinate, while ϕH increases from left to right on the abscissa. The blue arrow from left to right explicitly indicates the impact of blending in a progressively higher volume fraction of homopolymer. Accompanying the phase diagram are cartoons depicting the domain structure. Lamellar microdomains form in the neat BCP when χN is above the order− disorder transition (ODT), designated by χNODT. With increasing homopolymer volume fraction, ordered lamellar domains initially swell as the homopolymer is confined to the interior of the domains. On the homopolymer-rich end of the diagram (right side), the system macrophase separates into A and B homopolymer-rich phases if the homopolymer segregation strength, χNH, is >2.58 Consistent with the ordinate scaling, the transition from macrophase separated homopol12331
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano would be disordered rather than macrophase separated.58 The distribution of this homopolymer throughout the microdomains would promote mixing or disorder, reducing energetic barriers to chain diffusion and defect annihilation imposed by the domain structure.49,62 On the other hand, if χNH > 2, the homopolymer mixture would macrophase separate. In these cases, a ternary BCPhomopolymer blend is most appropriately described using the dry brush phase diagram, which implies a transition from microphase separated domains to a microemulsion phase, as the homopolymer fraction is increased, and finally to macrophase separated polymers. As a result, an interface between polymer regions is always preserved and there is little to no reduction in the energy barriers inhibiting chain redistribution. We tested this conjecture by examining the temporal ordering behavior of a blend of the 36 kg/mol BCP with a 50% weight fraction of 6 kg/mol PS and PMMA homopolymers at 220 °C, where χNH ≈ 2.1. From a power law fit to the extracted correlation lengths as a function of annealing time, the prefactor and ordering exponent are A = 148 ± 8 and ν = 0.015 ± 0.009, respectively (see Figure S5). The latter is comparable to the mean ordering exponent of ∼0.03 for the neat BCP and in stark contrast to the ordering exponent of ∼0.29 at 220 °C for a 50% blend of the same BCP with ∼3 kg/ mol PS and PMMA. Furthermore, a 70% blend of the 6 kg/mol PS and PMMA with this same BCP separates into several polymer macrophases (see Figure S6). No enhanced ordering is observed in a blend of the 36 kg/mol BCP with a 50% weight fraction of ∼12 kg/mol PS and PMMA homopolymers at 220 °C either (see Figure S7). These observations add substantial credibility to the hypothesis that a direct transition to a single disordered phase promoted by wet brush homopolymer addition is critical to the accelerated late-stage ordering in ternary BCP-homopolymer blends. The mechanism behind accelerated topological ordering discussed in this work is fundamentally different from a “defect melting” mechanism described previously by Kim et al.36 based on blending diblock copolymers with a random copolymer. In that work, the effective segregation strength, (χN)eff, of the blend decreases linearly with increasing volume fraction of the blended random copolymer, as evidenced by the accompanying decrease in lamellar repeat spacing. Assuming that the random copolymer screens the repulsive interaction between the two polymer blocks, its addition reduces (χN)eff primarily by reducing χ. Then the smallest BCP studied would approach the ODT ((χN)eff ≈ 10.5) at a volume fraction of random copolymer slightly greater than the value where the observed defect melting effect was strongest. However, the addition of random copolymer to lamellar BCPs with larger molecular weights did not contribute to any significant increase in order with increasing random copolymer volume fraction before macrophase separation, presumably because (χN)eff does not approach an ODT that would maximize the propensity for defect melting. It is worth noting that the random copolymer acts similarly to a neutral solvent, which dilutes the BCP and hence screens the repulsive interaction between the two blocks, effectively reducing χ. The resulting reduction in segregation strength decreases the BCP repeat spacing and enhances the ordering kinetics, but can also cause the polymer to approach or even cross over the ODT.25,65 On the other hand, the repeat spacing of BCPs blended with the ∼3 kg/mol homopolymers in this work remains effectively constant, and a significant ordering enhancement is observed
for the blends with the 74 kg/mol BCP at high homopolymer volume fractions as well, suggesting that the effective χ of the blends remains largely unchanged from that of the neat BCP. Instead the χN of the ODT itself increases with addition of more homopolymer, as depicted in Figure 5b. The practical implication is that blending in sufficiently low molecular weight homopolymer can significantly enhance ordering in larger molecular weight BCPs. Moreover, the combination of blending with thermal annealing offers a more deterministic and finely controlled alternative to solvent annealing to enhance ordering kinetics. The rapid ordering kinetics in the BCP-homopolymer blends provides a convenient yet powerful way to accelerate DSA processes for the formation of large area aligned nanopatterns. To demonstrate this, we compared the alignment capability of the neat 36 kg/mol BCP with its corresponding blend containing 50% (w/w) of the ∼3 kg/mol homopolymer using topographically confining grating templates prepared by photolithography of a negative tone photoresist, SU-8 (see Methods). The trenches in which the neat or blend polymer films were confined were larger than 2.5 μm widemore than 80 times the polymer repeat spacingpresenting a formidable challenge to large-scale pattern alignment, which propagates from the trench sidewalls. Figure 6a shows the neat BCP in a
Figure 6. Enhanced performance of wet brush BCP-homopolymer blends in directed self-assembly. (a) SEM of the neat 36 kg/mol BCP confined to a photoresist trench after annealing at 230 °C for 2 min. (b) SEM of the same 36 kg/mol BCP blended with 50% (w/ w) of ∼3 kg/mol homopolymers confined to a similar trench after annealing under the same conditions as in (a).
trench after annealing for 2 min at 230 °C. Lamellae are reasonably aligned with the grating direction within ∼500 nm from the sidewall, but defects emerge and the alignment degrades further into the trench. On the other hand, it is clear that unidirectional lamellae aligned with the sidewalls span the entire trench for the blend film annealed under the same conditions, as shown in Figure 6b. These results present unequivocal evidence of the benefit judicious polymer blending can provide for use in DSA. Moreover, the short annealing time of 2 min suggests that blending in low molecular weight homopolymers can help DSA meet the exacting time requirements required for practical application. Pattern transfer is an essential aspect for many envisioned applications for nanopatterning based on BCP self-assembly. However, this is very challenging using fully organic BCPs, like 12332
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano
section SEMs in Figure 7c,d, respectively. More details about the pattern transfer process are described in the Methods section. Interestingly, the pattern quality is improved upon transfer to the SiO2 layer, possibly due to erosion of the linebridging defects early in the etching process. It should be emphasized that this pattern transfer process has not been optimized for use with BCP-homopolymer blends; tailoring of the SIS and/or etch process could therefore improve the quality of transferred patterns.
PS-b-PMMA, and may be even more difficult in blends with high homopolymer loadings (e.g., 50 wt % or more). To circumvent this issue, we used sequential infiltration synthesis66,67 (SIS) to convert the BCP patterns to a patterned inorganic hard mask. SIS is a process derived from atomic layer deposition where strong and selective interaction between a precursor molecule (often organometallic) and one polymer block causes the nucleation and subsequent growth of inorganic material primarily in that block only. After several cycles of precursor/water exposure, the polymer may be removed revealing an inorganic replica of the domain pattern. Here, we use trimethylaluminum (TMA), as its selective reaction with PMMA in PS-b-PMMA can be used to generate AlO x nanostructures via SIS that are excellent hard masks for pattern transfer into Si or SiO2.5,6 Figure 7a,b shows plan view and
CONCLUSIONS In conclusion, we have demonstrated that ternary blends of lamellae-forming PS-b-PMMA BCPs with low molecular weight PS and PMMA homopolymers exhibit dramatically enhanced ordering kinetics, resulting in an order-of-magnitude increase in topological grain size in practical time scales (minutes). This discovery can be attributed to the tendency of low molecular weight homopolymers that do not phase separate to distribute uniformly throughout BCP domains, lowering energy barriers to chain redistribution and defect annihilation, and eventually inducing disorder. This approach can be applied to BCPs with different molecular weights while minimally altering repeat spacing, making it easier to integrate it with existing process flows. As demonstrated here, these blends are capable of both robust DSA and pattern transfer, suggesting they could be combined with advanced directed self-assembly approaches68,69 to reduce defectivity, improve throughput, or lower required annealing temperatures. Moreover, we anticipate that the results discovered here are generalizable to other BCP/ homopolymer ternary blends beyond PS-b-PMMA/PS/ PMMA where a disordered homopolymer phase (χNH < 2) is accessible at annealing temperatures. METHODS Materials. The random copolymer neutral brush was obtained as a sample from Dow Chemical already dissolved in propylene glycol monomethyl ether acetate (PGMEA) and used as received. The lamellae-forming 36 kg/mol PS-b-PMMA (Mn = 18-b-18 kg mol−1; PDI = 1.07) and 74 kg/mol PS-b-PMMA (Mn = 37-b-37 kg mol−1; PDI = 1.08) BCPs, along with PS (Mn = 3.5 kg/mol, PDI = 1.05; Mn = 6 kg/mol, PDI = 1.10; Mn = 12.5 kg/mol, PDI = 1.04) and PMMA (Mn = 3 kg/mol, PDI = 1.14; Mn = 6 kg/mol, PDI = 1.10; Mn = 11.5 kg/mol, PDI = 1.22) homopolymers were obtained from Polymer Source and used as received. Homopolymers were blended in equal proportions to ensure that the total volume fraction of each polymer in the blend is nearly equivalent to the volume fraction of each block in the neat BCP. All neat and blended polymer solutions were prepared using PGMEA solvent at a concentration of 2% (w/w). Trimethyl aluminum (TMA) was purchased from Sigma-Aldrich. Sample Preparation. Substrate neutralization was achieved by spin coating a thin film of the neutral random copolymer brush onto Si substrates at 1500 rpm, baking on a hot plate at 250 °C for 5 min under nitrogen purging to graft the brush to the substrate, and removing ungrafted neutral brush by spin-rinsing in PGMEA at 3000 rpm. BCP blend and neat thin films were spin-cast at 3000 rpm to obtain 30−40 nm thicknesses (blend solutions spun slightly thinner films that the neat BCP solutions as a result of reduced viscosity). They were then baked on a hot plate under nitrogen purging at temperatures ranging from 190 to 250 °C for times ranging from 30 to 3600 s. To enhance domain contrast for imaging, samples were exposed to a brief O2 plasma etch (March Plasma CS1701F, 100W, 100 mTorr) for 6−7 s, where PMMA is preferentially etched with respect to PS. Image Analysis. Samples were imaged using a scanning electron microscope (Hitachi S-4800). Analysis was performed using ImageJ on SEM images 1280 × 960 pixels in size at a resolution of 0.1 pixels/nm.
Figure 7. Transfer of a self-assembled thin-film lamellae pattern from a 50% (w/w) blend of the 36 kg/mol BCP with ∼3 kg/mol homopolymers into SiO2. The sample was annealed at 200 °C for 20 min to promote ordering. (a) Plan view and (b) cross section SEMs of the AlOx line nanostructures after selective conversion of PMMA domains to AlOx and BCP ashing. (c) Plan view and (d) cross section SEMs after transfer of these AlOx line patterns into SiO2 by reactive ion etching.
cross-section SEMs, respectively, of the ∼20 nm tall AlOx lines on top of a ∼ 300 nm SiO2 film on Si. The pattern was formed by a 36 kg/mol BCP blend with 50% (w/w) of ∼3 kg/mol homopolymer annealed at 200 °C for 20 min. AlOx lines were obtained after 4 TMA/water cycles and subsequent polymer removal by O2 plasma etching. Though the presence of linebridging defects indicates some PMMA homopolymer is present within PS domains, the lamellae pattern is clearly defined. Subsequent reactive ion etching in a CHF3/Ar plasma transfers this pattern into the underlying SiO2 to generate lines with a total height of ∼60 nm, as shown by plan view and cross12333
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano ORCID
Lamellar repeat spacing was determined by analysis of the Fourier transform of SEMs and was found to be nearly identical for higher magnification images where the resolution was doubled to 0.2 pixels/ nm. Orientational correlation or grain size analysis was performed according to the approach established by Harrison et al.,52 where maps of in-plane orientation (θ(r⃗)) were computed from spatial derivatives using the OrientationJ plugin,70 in which r⃗ indicates the position. The local orientation map was converted to a map of the local orientational order parameter (ψ(r)⃗ = cos(2θ(r)⃗ )), from which the correlation function g(r) = ⟨ψ(0)ψ(r)⟩ was then calculated, where brackets denote the azimuthal average at a distance r. The correlation length, ξ, was obtained by fitting g(r) to an empirical exponential decay function, e−(r/ξ). Differential Scanning Calorimetry (DSC). DSC was conducted using a PerkinElmer Diamond DSC tool over a range of 35−150 °C. The values for Tg represent an average over 3 measurements. Directed Self-Assembly (DSA). Topographic grating pattern templates used to confine the BCP for DSA were fabricated by photolithography of a negative tone photoresist (MicroChem SU-8 2002, diluted with MicroChem SU-8 2000 thinner) directly upon the neutral brush grafted to Si, as described previously.71 A ∼700 nm photoresist layer was spin-cast directly upon the neutral brush grafted to Si, soft baked for 1 min at 95 °C, and subsequently patterned by exposure to I-line light through a photomask using a mask aligner (Karl Suss MA6). After exposure, the wafer was baked at 95 °C again for 90 s to cross-link the exposed photoresist. Unexposed photoresist was then removed by immersion in PGMEA for 1 min. Individual wafer pieces were then broken off, and BCP neat and blend thin films were cast onto the patterned wafer pieces at 800 rpm for 30 s. These films were baked on a hot plate under nitrogen purging at 230 °C for 2 min. Pattern Transfer. For demonstration of pattern transfer, PS-bPMMA thin films were spin coated at 3000 rpm onto Si wafer pieces with a ∼300 nm SiO2 film deposited by plasma enhanced chemical vapor deposition (Trion Orion III) and baked at 200 °C for 20 min. AlOx replicas of the PMMA domain patterns were generated by SIS with 4 cycles of exposure to TMA and water vapor (100s each) at 85 °C in a commercial atomic layer deposition tool (Cambridge Ultratech Savannah S100) with a base pressure of ∼0.5−1 Torr. The BCP template was subsequently removed by O2 plasma ashing (100 mTorr, 20W) for 300s. These AlOx nanostructures were then used as hard masks to etch into the underlying SiO2 by reactive ion etching in a March Plasma CS1701F tool using CHF3 and Ar gas in a ratio of 22:10 (total pressure of 100 mTorr) with 100W RF power.
Gregory S. Doerk: 0000-0002-2933-2047 Kevin G. Yager: 0000-0001-7745-2513 Notes
The authors declare no competing financial interest.
ACKNOWLEDGMENTS This research took place at the Center for Functional Nanomaterials, which is a U.S. DOE Office of Science Facility, at Brookhaven National Laboratory under contract No. DESC0012704. REFERENCES (1) Black, C. T. Self-Aligned Self Assembly of Multi-Nanowire Silicon Field Effect Transistors. Appl. Phys. Lett. 2005, 87, 163116. (2) Tsai, H.; Pitera, J. W.; Miyazoe, H.; Bangsaruntip, S.; Engelmann, S. U.; Liu, C.-C.; Cheng, J. Y.; Bucchignano, J. J.; Klaus, D. P.; Joseph, E. A.; Sanders, D. P.; Colburn, M. E.; Guillorn, M. A. TwoDimensional Pattern Formation Using Grapho-Epitaxy of PS-BPMMA Block Copolymers for Advanced FinFET Device and Circuit Fabrication. ACS Nano 2014, 8, 5227−5232. (3) Hellwig, O.; Bosworth, J. K.; Dobisz, E.; Kercher, D.; Hauet, T.; Zeltzer, G.; Risner-Jamtgaard, J. D.; Yaney, D.; Ruiz, R. Bit Patterned Media Based on Block Copolymer Directed Assembly with Narrow Magnetic Switching Field Distribution. Appl. Phys. Lett. 2010, 96, 52511. (4) Albrecht, T. R.; Arora, H.; Ayanoor-Vitikkate, V.; Beaujour, J.-M.; Bedau, D.; Berman, D.; Bogdanov, A. L.; Chapuis, Y.-A.; Cushen, J.; Dobisz, E. E.; Doerk, G.; Gao, H.; Grobis, M.; Gurney, B.; Hanson, W.; Hellwig, O.; Hirano, T.; Pierre-Olivier, J.; Kercher, D.; Lille, J.; et al. Bit-Patterned Magnetic Recording: Theory, Media Fabrication, and Recording Performance. IEEE Trans. Magn. 2015, 51, 1−42. (5) Rahman, A.; Ashraf, A.; Xin, H.; Tong, X.; Sutter, P.; Eisaman, M. D.; Black, C. T. Sub-50-Nm Self-Assembled Nanotextures for Enhanced Broadband Antireflection in Silicon Solar Cells. Nat. Commun. 2015, 6, 5963. (6) Checco, A.; Rahman, A.; Black, C. T. Robust Superhydrophobicity in Large-Area Nanostructured Surfaces Defined by Block-Copolymer Self Assembly. Adv. Mater. 2014, 26, 886−891. (7) Darling, S. B. Directing the Self-Assembly of Block Copolymers. Prog. Polym. Sci. 2007, 32, 1152−1204. (8) Takahashi, H.; Laachi, N.; Delaney, K. T.; Hur, S.-M.; Weinheimer, C. J.; Shykind, D.; Fredrickson, G. H. Defectivity in Laterally Confined Lamella-Forming Diblock Copolymers: Thermodynamic and Kinetic Aspects. Macromolecules 2012, 45, 6253−6265. (9) Hur, S.-M.; Thapar, V.; Ramírez-Hernández, A.; Khaira, G.; Segal-Peretz, T.; Rincon-Delgadillo, P. A.; Li, W.; Müller, M.; Nealey, P. F.; de Pablo, J. J. Molecular Pathways for Defect Annihilation in Directed Self-Assembly. Proc. Natl. Acad. Sci. U. S. A. 2015, 112, 14144−14149. (10) Angelescu, D. E.; Waller, J. H.; Adamson, D. H.; Deshpande, P.; Chou, S. Y.; Register, R. A.; Chaikin, P. M. Macroscopic Orientation of Block Copolymer Cylinders in Single-Layer Films by Shearing. Adv. Mater. 2004, 16, 1736−1740. (11) Olszowka, V.; Hund, M.; Kuntermann, V.; Scherdel, S.; Tsarkova, L.; Boker, A. Electric Field Alignment of a Block Copolymer Nanopattern: Direct Observation of the Microscopic Mechanism. ACS Nano 2009, 3, 1091−1096. (12) Majewski, P. W.; Gopinadhan, M.; Jang, W.-S.; Lutkenhaus, J. L.; Osuji, C. O. Anisotropic Ionic Conductivity in Block Copolymer Membranes by Magnetic Field Alignment. J. Am. Chem. Soc. 2010, 132, 17516−17522. (13) Yager, K. G.; Fredin, N. J.; Zhang, X.; Berry, B. C.; Karim, A.; Jones, R. L. Evolution of Block-Copolymer Order through a Moving Thermal Zone. Soft Matter 2010, 6, 92−99. (14) Singh, G.; Yager, K. G.; Berry, B.; Kim, H.-C.; Karim, A. Dynamic Thermal Field-Induced Gradient Soft-Shear for Highly
ASSOCIATED CONTENT S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.7b06154. Additional SEMs of the BCP/homopolymer blends different homopolymer molecular weights and mass fractions, DSC traces of the 36 kg/mol BCP along with the ∼3 kg/mol PS and PMMA homopolymers in the vicinity of their Tg’s, plots of correlation length vs annealing time at 220 °C for a 50% (w/w) blend of the 36 kg/mol BCP with ∼6 kg/mol homopolymer as well as the neat 36 kg/mol BCP and its corresponding blend with 50% (w/w) of ∼3 kg/mol homopolymer at 210 °C, and a discussion of the estimation of the critical homopolymer segregation strength required for a direct transition from the lamellar phase to a disordered phase via homopolymer addition (PDF)
AUTHOR INFORMATION Corresponding Author
*E-mail:
[email protected]. 12334
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano Oriented Block Copolymer Thin Films. ACS Nano 2012, 6, 10335− 10342. (15) Cheng, J. Y.; Rettner, C. T.; Sanders, D. P.; Kim, H.-C.; Hinsberg, W. D. Dense Self-Assembly on Sparse Chemical Patterns: Rectifying and Multiplying Lithographic Patterns Using Block Copolymers. Adv. Mater. 2008, 20, 3155−3158. (16) Ruiz, R.; Kang, H.; Detcheverry, F. A.; Dobisz, E.; Kercher, D. S.; Albrecht, T. R.; de Pablo, J. J.; Nealey, P. F. Density Multiplication and Improved Lithography by Directed Block Copolymer Assembly. Science 2008, 321, 936−939. (17) Liu, C.; Ramírez-Hernández, A.; Han, E.; Craig, G. S. W.; Tada, Y.; Yoshida, H.; Kang, H.; Ji, S.; Gopalan, P.; de Pablo, J. J.; Nealey, P. F. Chemical Patterns for Directed Self-Assembly of Lamellae-Forming Block Copolymers with Density Multiplication of Features. Macromolecules 2013, 46, 1415−1424. (18) Doerk, G. S.; Liu, C.-C.; Cheng, J. Y.; Rettner, C. T.; Pitera, J. W.; Krupp, L. E.; Topuria, T.; Arellano, N.; Sanders, D. P. Pattern Placement Accuracy in Block Copolymer Directed Self-Assembly Based on Chemical Epitaxy. ACS Nano 2013, 7, 276−285. (19) Cheng, J. Y.; Mayes, A. M.; Ross, C. A. Nanostructure Engineering by Templated Self-Assembly of Block Copolymers. Nat. Mater. 2004, 3, 823−828. (20) Segalman, R. A.; Yokoyama, H.; Kramer, E. J. Graphoepitaxy of Spherical Domain Block Copolymer Films. Adv. Mater. 2001, 13, 1152−1155. (21) Bita, I.; Yang, J. K. W.; Jung, Y. S.; Ross, C. A.; Thomas, E. L.; Berggren, K. K. Graphoepitaxy of Self-Assembled Block Copolymers on Two-Dimensional Periodic Patterned Templates. Science 2008, 321, 939−943. (22) Jung, Y. S.; Chang, J. B.; Verploegen, E.; Berggren, K. K.; Ross, C. a. A Path to Ultranarrow Patterns Using Self-Assembled Lithography. Nano Lett. 2010, 10, 1000−1005. (23) Delgadillo, P. R.; Suri, M.; Durant, S.; Cross, A.; Nagaswami, V. R.; Van Den Heuvel, D.; Gronheid, R.; Nealey, P. Defect Source Analysis of Directed Self-Assembly Process. J. Micro/Nanolithogr., MEMS, MOEMS 2013, 12, 31112. (24) Sinturel, C.; Vayer, M.; Morris, M.; Hillmyer, M. A. Solvent Vapor Annealing of Block Polymer Thin Films. Macromolecules 2013, 46, 5399−5415. (25) Kim, E.; Ahn, H.; Park, S.; Lee, H.; Lee, M.; Lee, S.; Kim, T.; Kwak, E.-A.; Lee, J. H.; Lei, X.; Huh, J.; Bang, J.; Lee, B.; Ryu, D. Y. Directed Assembly of High Molecular Weight Block Copolymers: Highly Ordered Line Patterns of Perpendicularly Oriented Lamellae with Large Periods. ACS Nano 2013, 7, 1952−1960. (26) Kim, J. M.; Kim, Y.; Park, W. I.; Hur, Y. H.; Jeong, J. W.; Sim, D. M.; Baek, K. M.; Lee, J. H.; Kim, M.-J.; Jung, Y. S. Eliminating the Trade-Off between the Throughput and Pattern Quality of Sub-15 Nm Directed Self-Assembly via Warm Solvent Annealing. Adv. Funct. Mater. 2015, 25, 306−315. (27) Park, W. I.; Kim, J. M.; Jeong, J. W.; Jung, Y. S. Deep-Nanoscale Pattern Engineering by Immersion-Induced Self-Assembly. ACS Nano 2014, 8, 10009−10018. (28) Perego, M.; Ferrarese Lupi, F.; Ceresoli, M.; Giammaria, T. J.; Seguini, G.; Enrico, E.; Boarino, L.; Antonioli, D.; Gianotti, V.; Sparnacci, K.; Laus, M. Ordering Dynamics in Symmetric PS-BPMMA Diblock Copolymer Thin Films during Rapid Thermal Processing. J. Mater. Chem. C 2014, 2, 6655−6664. (29) Zhang, X.; Harris, K. D.; Wu, N. L. Y.; Murphy, J. N.; Buriak, J. M. Fast Assembly of Ordered Block Copolymer Nanostructures through Microwave Annealing. ACS Nano 2010, 4, 7021−7029. (30) Jin, C.; Murphy, J. N.; Harris, K. D.; Buriak, J. M. Deconvoluting the Mechanism of Microwave Annealing of Block Copolymer Thin Films. ACS Nano 2014, 8, 3979−3991. (31) Majewski, P. W.; Yager, K. G. Millisecond Ordering of Block Copolymer Films via Photothermal Gradients. ACS Nano 2015, 9, 3896−3906. (32) Majewski, P. W.; Yager, K. G. Block Copolymer Response to Photothermal Stress Fields. Macromolecules 2015, 48, 4591−4598.
(33) Majewski, P. W.; Yager, K. G. Latent Alignment in PathwayDependent Ordering of Block Copolymer Thin Films. Nano Lett. 2015, 15, 5221−5228. (34) Jin, H. M.; Park, D. Y.; Jeong, S.-J.; Lee, G. Y.; Kim, J. Y.; Mun, J. H.; Cha, S. K.; Lim, J.; Kim, J. S.; Kim, K. H.; Lee, J. K.; Kim, S. O. Flash Light Millisecond Self-Assembly of High χ Block Copolymers for Wafer-Scale Sub-10 Nm Nanopatterning. Adv. Mater. 2017, 29, 1700595. (35) Arias-Zapata, J.; Böhme, S.; Garnier, J.; Girardot, C.; Legrain, A.; Zelsmann, M. Ultrafast Assembly of PS-PDMS Block Copolymers on 300 Mm Wafers by Blending with Plasticizers. Adv. Funct. Mater. 2016, 26, 5690−5700. (36) Kim, B. H.; Park, S. J.; Jin, H. M.; Kim, J. Y.; Son, S.-W.; Kim, M.-H.; Koo, C. M.; Shin, J.; Kim, J. U.; Kim, S. O. Anomalous Rapid Defect Annihilation in Self-Assembled Nanopatterns by Defect Melting. Nano Lett. 2015, 15, 1190−1196. (37) Liu, G.; Stoykovich, M. P.; Ji, S.; Stuen, K. O.; Craig, G. S. W.; Nealey, P. F. Phase Behavior and Dimensional Scaling of Symmetric Block Copolymer−Homopolymer Ternary Blends in Thin Films. Macromolecules 2009, 42, 3063−3072. (38) Stuen, K. O.; Thomas, C. S.; Liu, G.; Ferrier, N.; Nealey, P. F. Dimensional Scaling of Cylinders in Thin Films of Block Copolymer− Homopolymer Ternary Blends. Macromolecules 2009, 42, 5139−5145. (39) Stoykovich, M. P.; Müller, M.; Kim, S. O.; Solak, H. H.; Edwards, E. W.; de Pablo, J. J.; Nealey, P. F. Directed Assembly of Block Copolymer Blends into Nonregular Device-Oriented Structures. Science 2005, 308, 1442−1446. (40) Stoykovich, M. P.; Kang, H.; Daoulas, K. C.; Liu, G.; Liu, C.; de Pablo, J. J.; Müller, M.; Nealey, P. F. Directed Self-Assembly of Block Copolymers for Nanolithography : Essential Integrated Circuit Geometries. ACS Nano 2007, 1, 168−175. (41) Dai, K. H.; Kramer, E. J.; Shull, K. R. Interfacial Segregation in Two-Phase Polymer Blends with Diblock Copolymer Additives: The Effect of Homopolymer Molecular Weight. Macromolecules 1992, 25, 220−225. (42) Black, C. T.; Ruiz, R.; Breyta, G.; Cheng, J. Y.; Colburn, M. E.; Guarini, K. W.; Kim, H.-C.; Zhang, Y. Polymer Self Assembly in Semiconductor Microelectronics. IBM J. Res. Dev. 2007, 51, 605−633. (43) Tanaka, H.; Hasegawa, H.; Hashimoto, T. Ordered Structure in Mixtures of a Block Copolymer and Homopolymers. 1. Solubilization of Low Molecular Weight Homopolymers. Macromolecules 1991, 24, 240−251. (44) Torikai, N.; Takabayashi, N.; Noda, I.; Koizumi, S.; Morii, Y.; Matsushita, Y. Lamellar Domain Spacings of Diblock Copolymer/ Homopolymer Blends and Conformations of Block Chains in Their Microdomains. Macromolecules 1997, 30, 5698−5703. (45) Winey, K. I.; Thomas, E. L.; Fetters, L. J. Swelling of Lamellar Diblock Copolymer by Homopolymer: Influences of Homopolymer Concentration and Molecular Weight. Macromolecules 1991, 24, 6182−6188. (46) Quan, X.; Gancarz, I.; Koberstein, J. T.; Wignall, G. D. Effect of Homopolymer Molecular Weight on the Morphology of Block Copolymer/homopolymer Blends. Macromolecules 1987, 20, 1431− 1434. (47) Mykhaylyk, T. A.; Mykhaylyk, O. O.; Collins, S.; Hamley, I. W. Ordered Structures and Phase Transitions in Mixtures of a Polystyrene/Polyisoprene Block Copolymer with the Corresponding Homopolymers in Thin Films and in Bulk. Macromolecules 2004, 37, 3369−3377. (48) Matsen, M. W. Phase Behavior of Block Copolymer/ Homopolymer Blends. Macromolecules 1995, 28, 5765−5773. (49) Ruiz, R.; Bosworth, J. K.; Black, C. T. Effect of Structural Anisotropy on the Coarsening Kinetics of Diblock Copolymer Striped Patterns. Phys. Rev. B: Condens. Matter Mater. Phys. 2008, 77, 54204. (50) Majewski, P. W.; Yager, K. G. Rapid Ordering of Block Copolymer Thin Films. J. Phys.: Condens. Matter 2016, 28, 403002. (51) Kim, B. H.; Lee, H. M.; Lee, J.-H.; Son, S.-W.; Jeong, S.-J.; Lee, S.; Lee, D. Il; Kwak, S. U.; Jeong, H.; Shin, H.; Yoon, J.-B.; Lavrentovich, O. D.; Kim, S. O. Spontaneous Lamellar Alignment in 12335
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
Article
ACS Nano Thickness-Modulated Block Copolymer Films. Adv. Funct. Mater. 2009, 19, 2584−2591. (52) Harrison, C.; Adamson, D. H.; Cheng, Z.; Sebastian, J. M.; Sethuraman, S.; Huse, D. A.; Register, R. A.; Chaikin, P. M. Mechanisms of Ordering in Striped Patterns. Science 2000, 290, 1558−1560. (53) Milhaupt, J. M.; Lodge, T. P.; Smith, S. D.; Hamersky, M. W. Composition and Temperature Dependence of Monomer Friction in Polystyrene/Poly(methyl Methacrylate) Matrices. Macromolecules 2001, 34, 5561−5570. (54) Tong, Q.; Sibener, S. J. Visualization of Individual Defect Mobility and Annihilation within Cylinder-Forming Diblock Copolymer Thin Films on Nanopatterned Substrates. Macromolecules 2013, 46, 8538−8544. (55) Williams, M. L.; Landel, R. F.; Ferry, J. D. The Temperature Dependence of Relaxation Mechanisms in Amorphous Polymers and Other Glass-Forming Liquids. J. Am. Chem. Soc. 1955, 77, 3701−3707. (56) Fox, T. G. Influence of Diluent and of Copolymer Composition on the Glass Temperature of a Polymer System. Bull. Am. Phys. Soc. 1952, 1, 123. (57) Vega, D. A.; Harrison, C. K.; Angelescu, D. E.; Trawick, M. L.; Huse, D. A.; Chaikin, P. M.; Register, R. A. Ordering Mechanisms in Two-Dimensional Sphere-Forming Block Copolymers. Phys. Rev. E 2005, 71, 61803. (58) Scott, R. L. The Thermodynamics of High Polymer Solutions. V. Phase Equilibria in the Ternary System: Polymer 1Polymer 2 Solvent. J. Chem. Phys. 1949, 17, 279−284. (59) Broseta, D.; Fredrickson, G. H. Phase Equilibria in Copolymer/ homopolymer Ternary Blends: Molecular Weight Effects. J. Chem. Phys. 1990, 93, 2927−2938. (60) Bates, F. S.; Maurer, W. W.; Lipic, P. M.; Hillmyer, M. A.; Almdal, K.; Mortensen, K.; Fredrickson, G. H.; Lodge, T. P. Polymeric Bicontinuous Microemulsions. Phys. Rev. Lett. 1997, 79, 849−852. (61) Broseta, D.; Fredrickson, G. H.; Helfand, E.; Leibler, L. Molecular Weight and Polydispersity Effects at Polymer-Polymer Interfaces. Macromolecules 1990, 23, 132−139. (62) Dalvi, M. C.; Lodge, T. P. Diffusion in Block Copolymer Melts: The Disordered Region and the Vicinity of the Order-Disorder Transition. Macromolecules 1994, 27, 3487−3492. (63) Ryu, H. J.; Sun, J.; Avgeropoulos, A.; Bockstaller, M. R. Retardation of Grain Growth and Grain Boundary Pinning in Athermal Block Copolymer Blend Systems. Macromolecules 2014, 47, 1419−1427. (64) Russell, T. P.; Hjelm, R. P.; Seeger, P. A. Temperature Dependence of the Interaction Parameter of Polystyrene and Poly(methyl Methacrylate). Macromolecules 1990, 23, 890−893. (65) Xiong, S.; Wan, L.; Ishida, Y.; Chapuis, Y.-A.; Craig, G. S. W.; Ruiz, R.; Nealey, P. F. Directed Self-Assembly of Triblock Copolymer on Chemical Patterns for Sub-10-Nm Nanofabrication via Solvent Annealing. ACS Nano 2016, 10, 7855−7865. (66) Kamcev, J.; Germack, D. S.; Nykypanchuk, D.; Grubbs, R. B.; Nam, C.-Y.; Black, C. T. Chemically Enhancing Block Copolymers for Block-Selective Synthesis of Self-Assembled Metal Oxide Nanostructures. ACS Nano 2013, 7, 339−346. (67) Peng, Q.; Tseng, Y.-C.; Darling, S. B.; Elam, J. W. Nanoscopic Patterned Materials with Tunable Dimensions via Atomic Layer Deposition on Block Copolymers. Adv. Mater. 2010, 22, 5129−5133. (68) Doerk, G. S.; Cheng, J. Y.; Singh, G.; Rettner, C. T.; Pitera, J. W.; Balakrishnan, S.; Arellano, N.; Sanders, D. P. Enabling Complex Nanoscale Pattern Customization Using Directed Self-Assembly. Nat. Commun. 2014, 5, 5805. (69) Stein, A.; Wright, G.; Yager, K. G.; Doerk, G. S.; Black, C. T. Selective Directed Self-Assembly of Coexisting Morphologies Using Block Copolymer Blends. Nat. Commun. 2016, 7, 12366. (70) Rezakhaniha, R.; Agianniotis, a.; Schrauwen, J. T. C.; Griffa, a.; Sage, D.; Bouten, C. V. C.; van de Vosse, F. N.; Unser, M.; Stergiopulos, N. Experimental Investigation of Collagen Waviness and Orientation in the Arterial Adventitia Using Confocal Laser Scanning Microscopy. Biomech. Model. Mechanobiol. 2012, 11, 461−473.
(71) Jeong, S.-J.; Moon, H.-S.; Kim, B. H.; Kim, J. Y.; Yu, J.; Lee, S.; Lee, M. G.; Choi, H.; Kim, S. O. Ultralarge-Area Block Copolymer Lithography Enabled by Disposable Photoresist Prepatterning. ACS Nano 2010, 4, 5181−5186.
12336
DOI: 10.1021/acsnano.7b06154 ACS Nano 2017, 11, 12326−12336
