Article Cite This: Macromolecules XXXX, XXX, XXX−XXX
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Morphological Evolution of Poly(solketal methacrylate)-blockpolystyrene Copolymers in Thin Films Duk Man Yu,† Darren M. Smith,‡ Hyeyoung Kim,† Jose Kenneth D. Mapas,‡ Javid Rzayev,*,‡ and Thomas P. Russell*,†,§,∥
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†
Department of Polymer Science and Engineering, University of Massachusetts Amherst, 120 Governors Drive, Amherst, Massachusetts 01003, United States ‡ Department of Chemistry, University at Buffalo, The State University of New York, Buffalo, New York 14260-3000, United States § Materials Science Division, Lawrence Berkeley National Laboratory, 1 Cyclotron Road, Berkeley, California 94720, United States ∥ Beijing Advanced Innovation Center for Soft Matter Science and Engineering, Beijing University of Chemical Technology, Beijing 100029, China S Supporting Information *
ABSTRACT: The morphological evolution of the lamellar microdomains in the thin films of symmetric poly(solketal methacrylate-b-styrene) (PSM-b-PS) copolymers that can be converted into poly(glycerol monomethacrylate-b-styrene) (PGM-b-PS) copolymers through acid hydrolysis reaction was investigated. This simple chemical transformation was performed in the solid state using trifluoroacetic acid vapor, markedly increasing the segmental interaction parameter (χ) from 0.035 to 0.438 at 25 °C and a 118 Å3 reference volume by changing the hydrophobic PSM block to the hydrophilic PGM block. To control the orientation of the lamellar microdomains using this responsive block copolymer (BCP), a protected random copolymer (PSM-r-PS) with 31 mol % SM, which was concurrently transformed into PGM-r-PS with the BCP, was used to tune the interfacial energies at the substrate. Atomic force microscopy and grazing-incidence small-angle X-ray scattering measurements as a function of exposure time to an acid vapor were performed to characterize the transition from the disordered state into the ordered state and to assess the orientation of the microdomains. As a result, a 9.4 nm full pitch lamellar microdomain morphology in the thin films was achieved after full conversion and thermal annealing, indicating that the modified substrate exhibited the surface neutrality toward the two blocks and successfully induced vertical orientation without any additional layers.
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is reduced to obtain smaller L0, χ must be increased. For a symmetric diblock copolymer, the critical value to form a lamellar microdomain morphology has been calculated to be 10.5.13−15 To achieve sub-10 nm periods using a high χ−low N system, various BCPs with strong immiscibility between the two blocks have been exploited, 16 such as siloxane-containing BCPs,11,17−21 tert-butyl-containing BCPs,10,22 cyclohexyl-containing BCPs,23 the fluorine-containing BCPs,24−27 dimethylazlactone-containing BCPs, 28 and hydroxyl-containing BCPs.29−31 However, perpendicular orientation of the microdomains to the film surface, which is required for most
INTRODUCTION As a bottom-up platform, the self-assembly of block copolymers (BCPs) has been attractive for many applications including storage devices, nanoporous membranes, and semiconductors, as highly ordered BCPs with cylindrical or lamellar microdomains have the potential to generate sub-10 nm periodic dot or line patterns by chemical differences between the two blocks.1−7 In the strong segregation limit, the periodicity of the microphase-separated domains (L0) is dictated as L0 ≈ χ1/6N2/3, where χ is the Flory−Huggins segmental interaction parameter and N is the total number of BCP segments. As shown by this relationship, N is a stronger variable than χ for L0, and therefore, decreasing N can allow further minimization of L0.8−12 However, the segregation strength parameter, χN, should be greater than a critical value for the microphase-separated morphology. Consequently, as N © XXXX American Chemical Society
Received: March 10, 2019 Revised: April 16, 2019
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DOI: 10.1021/acs.macromol.9b00488 Macromolecules XXXX, XXX, XXX−XXX
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morphologies were obtained with an L 0 of 18 nm corresponding to a line width of 9 nm after thermal treatment. This top coat process was further applied to a poly(5-vinyl-1,3benzidioxole-b-pentamethyldisilylstyrene) copolymer on guidelines patterned by nanoimprint lithography and finally achieved the aligned perpendicular lamellar pattern with a line width of 5 nm.33 Nealey and co-workers49 used initiated chemical vapor-phase deposition (iCVD) for a top coat on the poly(2-vinylpyridine-b-styrene-b-2-vinylpyridine) thin film. The iCVD method produced a solvent-free 7 nm-thick top coat with a cross-linked structure to tune the interface as a neutral layer. The authors further simplified the top coating process for commercial fabrication and showed the aligned perpendicular lamellar pattern with a line width as small as 9.3 nm after thermal annealing. Karim and co-workers50 exploited sharp thermal gradient cold zone annealing (CZA) as a different thermal annealing technique instead of a hot plate or a vacuum oven to induce a vertical orientation in PS-b-PMMA. They reported a one-step continuous process without any surface treatments or top coats on a quartz substrate, which has lower thermal conductivity than a silicon substrate, for roll-toroll manufacturing of large-scale thin films. The CZA process enabled the formation of the perpendicular lamellae in 100 nm-thick PS-b-PMMA films when a sharp thermal gradient (∼48 °C/mm) with an optimal sweep rate was applied, resulting in a transient vertical strain field. All of these approaches have potential for developing BCP thin-film patterning for nanoelectronic applications; however, further efforts are necessary to find the optimal strategies for various high-χ BCPs, enabling sub-10 nm pitch dimensions via thermal treatment. Recently, we reported a poly(solketal methacrylate-bstyrene) (PSM-b-PS) copolymer, as a new high-χ BCP, enabling the formation of 5.4 nm full pitch lamellar microdomains at low N (N = 16) in bulk.29,30 The hydrophobic PSM block in this BCP can be transformed into the hydrophilic self-associating poly(glycerol monomethacrylate) (PGM) block through acid-catalyzed hydrolysis, resulting in a massive increase in χ from 0.035 to 0.438. Thus, a highly ordered morphology was obtained from a phase-mixed PSM-b-PS copolymer through a chemical transformation without any additives. In addition, it should be noted that this simple, one-step chemical transformation was performed in a solid-state using an acid vapor without any additional preor post-treatment steps, which can facilitate potential scale-up and technological integration for the fabrication of nanopatterned thin films. In this study, morphological evolution of lamellar microdomains of symmetric PSM-b-PS copolymers in thin films is investigated as a function of exposure time to an acid vapor. A hydroxyl-containing random copolymer (PSM-r-PS) composed of SM (31 mol %), styrene (61 mol %), and 2hydroxyethyl methacrylate (8 mol %) was applied on the substrate to balance interactions at the polymer/substrate interface. Interestingly, as the SM segments in this random copolymer brush layer are also converted to the GM segments concurrently with the PSM-b-PS layer during acid-catalyzed hydrolysis, the interfacial energies between two layers can be effectively balanced after full conversion. Atomic force microscopy (AFM), scanning electron microscopy (SEM), and grazing-incidence small-angle X-ray scattering (GISAXS) measurements were used to observe not only the morphological change from the phase-mixed (disordered) to ordered
engineering applications, is an ongoing and unresolved challenge. Significant chemical differences between the two blocks engineered into the copolymer structure to achieve a high-χ system also result in a large difference between the surface energies of the two blocks, which leads to preferential surface interactions with the substrate.17,32−34 Solvent vapor annealing (SVA) is commonly used to achieve perpendicular microdomain orientation, as a saturated solvent atmosphere can plasticize the BCP thin films, increase chain mobility, mediate interfacial interactions, and minimize surface energy differences between two blocks.35−37 Typically, a good solvent for both blocks is used, and the SVA process can be performed at ambient temperatures.38,39 However, this process has drawbacks, including dewetting, deformation of the structure, strong dependence on the solubility of each block, solvent evaporation rate, relative humidity, and incompatibility with commercial nanomanufacturing processes.17,40−43 Therefore, inducing vertical orientations normal to the surface using highχ BCPs by thermal annealing process is desirable but challenging. When BCPs in thin films are thermally annealed above the glass-transition temperature (Tg) of both blocks without solvent or other additives, the resulting orientation is determined by preferential interactions at the polymer/ substrate interface and surface energies at the free surface of the film.1 To control these preferential interactions, the polymer/substrate interface can be modified by the attachment of random copolymers composed of the same repeating units as the BCP to control the orientation of the microdomains. Russell and co-workers44−46 first introduced this simple technique using hydroxyl-terminated random copolymers of styrene and methacrylate (PS-r-PMMA) grafted onto silicon substrates as a random copolymer brush with 5 nm thickness. They found that the interfacial energies of the PS-r-PMMA brush with the PS and poly(methyl methacrylate) (PMMA) homopolymers were balanced when the composition of styrene and methyl methacrylate in the random copolymer was adjusted to about 58:42, indicating that this random copolymer brush neutralized polymer/substrate interactions enabling control over the microdomain orientation. Perego and co-workers47 reported the effect of the total molecular weight and the grafting density of hydroxyl-terminated PS-rPMMA copolymers on microdomain orientations in the thin films. They showed that microdomains oriented parallel to the interface for brush thicknesses less than ∼5 nm but normal to the interface for thicker brush layers. This arises from the penetration of the diblock copolymer into the brush and whether the BCP segments can interact with the underlying substrate. 48 For recently reported high χ BCPs, the neutralization at the polymer/substrate interface was not enough to induce perpendicular orientation because of large surface energy differences between the two blocks, and so controlling both the polymer/substrate and polymer/air interfaces was required. Willson and co-workers17,19 demonstrated well-ordered lamellar morphologies oriented normal to the substrate using poly(styrene-b-trimethylsilylstyrene-b-styrene) (PS-b-PTMSS-b-PS) and poly(trimethylsilylstyrene-blactide) (PTMSS-b-PLA) copolymers through a top coat process. The materials designed as the top coats contained anhydride units and could be spin-coated using an aqueous solution to prevent damage to a block polymer layer during the coating process. By using such a process, the interfacial energies were balanced, and the perpendicular lamellar B
DOI: 10.1021/acs.macromol.9b00488 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules Table 1. Sample Codes and Characteristics of PSM-b-PS Copolymers sample code P(SM13-S14) P(SM5-S5) P(SM3-S4) P(SM3-S3)
total Mn (g/mol) 26 800 10 500 6700 5600
D̵ PSM‑PSa 1.22 1.08 1.13 1.08
PSM Mn (g/mol) 13 200 5400b 3000c 2600c
b
D̵ PSMa
NPSM
PS Mn (g/mol)
NPS
f PSMd
1.09 1.05 1.08 1.09
66 27 15 13
13 600 5100 3700 3000
131 49 36 29
0.47 0.49 0.42 0.44
a
Determined by size exclusion chromatography (SEC) in tetrahydrofuran using PS calibration. bDetermined by SEC with a light-scattering detector (T = 30 °C; λ = 630 nm) using a refractive index increment (dn/dc) of 0.067 for PSM. cCalculated from 1H NMR end group analysis. d Volume fraction of PSM. molecular weight was determined by 1H NMR.29 The information for the molecular characteristics and the sample codes of PSM-b-PS is summarized in Table 1. Trifluoroacetic acid (TFA), 1,4-dioxane, and toluene were purchased from Sigma-Aldrich Co. and used directly without further purification. Hydrochloric acid (HCl, 37%) from Fisher Scientific was also used as received. Acid-Catalyzed Hydrolysis of PSM-b-PS Copolymer in Thin Films. A silicon substrate with a native oxide layer (orientation of (100), International Wafer Service) (∼1 × 1 cm2 in size) was cleaned with a carbon dioxide snow jet followed by UV−ozone plasma (UVO cleaner model 342, Jelight Company Inc.). For a random copolymer brush layer, a hydroxyl-containing PSM-r-PS copolymer (Mn = 2400 g/mol, volume fraction of SM segments (f SM) = 0.44) dissolved into toluene (0.3 wt %) was first spin-coated onto the cleaned substrate. The low-molecular-weight random copolymer was used to prevent interpenetration of the random copolymer chain into the lowmolecular-weight BCP layer. Grafting PSM-r-PS to the substrate was achieved by thermal annealing at 150 °C for 12 h under vacuum. During this treatment, hydroxyl groups in the random copolymer bonded with the native oxide layer to anchor the polymer chains to the substrate. The thickness of grafted brush layer was measured to be 1.2 ± 0.1 nm by the ellipsometer (model LSE, Gaertner Scientific Corp.) after rinsing with toluene to remove the unreacted random copolymer chains from the substrate. Subsequently, PSM-b-PS thin films were deposited onto the modified substrate by spin-coating from PSM-b-PS solution dissolved in toluene, and the film thickness was adjusted by the concentration of the solution (0.3−1.5 wt %) as well as the spin-coating speed (2500−4500 rpm). To transform SM
state but also the lamellar pitch and orientation in the thin films. Using this simple surface modification, responsive PSMb-PS thin films displayed the vertically oriented 4.7 nm halfpitch microdomains after the solid-state transformation and thermal annealing without the need for any additional layers.
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EXPERIMENTAL SECTION
Materials. Symmetric PSM-b-PS copolymers (Mn = 26 800, 10 500, 6700, and 5600 g/mol) were synthesized by a sequential
Scheme 1. Solid-State Chemical Transformation of PSM-bPS to PGM-b-PS by TFA-Vapor-Catalyzed Hydrolysis Reactiona
a
Acetone is produced as a byproduct of the hydrolysis during the reaction.
reversible addition−fragmentation chain-transfer polymerization of SM and styrene. In our previous work, detailed synthetic procedures for the monomers and polymers were reported, and the polymer
Figure 1. Morphological evolution of P(SM13-S14) (N = 197) in the thin films as a function of exposure time ranging from 0 to 80 min to TFA vapor: (a,f) 0, (b,g) 20, (c,h) 40, (d,i) 60, and (e,j) 80 min. The samples were thermally annealed at 150 °C for 12 h and then measured by (a−e) AFM height and (f−j) height profiles (red line). The color contrast is shown in the AFM height images; dark area is thinner (holes) and bright area is thicker (islands). The scale bars represent 250 nm. C
DOI: 10.1021/acs.macromol.9b00488 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules segments to GM segments in the solid state (Scheme 1), PSM-b-PS thin films were exposed to TFA vapor in a sealed glass jar (volume of the jar = 46.5 cm3) with 150 μL of 10 M TFA at room temperature. The TFA exposure time was varied from 0 to 80 min. When the thin films were taken out from the glass jar, the residual TFA vapor and acetone, a byproduct of the transformation, were removed by purging with dry nitrogen gas for several minutes, and then the converted thin films were placed in a vacuum oven for thermal annealing at 150 °C for 12 h prior to morphological characterization studies. Small-Angle X-ray Scattering Analysis. The morphologies of PSM-b-PS samples in bulk were investigated by SAXS before and after conversion. A detailed acid hydrolysis procedure for the bulk samples was reported in our previous work.30 For SAXS measurements, the sample powder was compression-molded into a small steel washer (0.5 mm in thickness) at 150 °C for 10 min and then Kapton films (0.06 mm in thickness) were attached on both sides of the steel washer to prevent leaking. Subsequently, thermal annealing for all of the samples was conducted at 140 °C for 24 h under vacuum to equilibrate their morphologies. Using a Ganesha SAXS-LAB instrument with Cu Kα radiation (λ = 0.1542 nm) and an incident beam diameter of ∼0.3 mm under vacuum, SAXS patterns were obtained as a function of the scattering vector (q = (4π/λ)sin θ) from 0.05 to 3.0 nm−1 using a two-dimensional (2D) detector (Pilatus 300K), where λ is the X-ray wavelength and 2θ indicates the scattering angle. The sample-to-detector distance was calibrated with a silver behenate standard. GISAXS measurements for thin films were performed using the same instrument at room temperature under vacuum, and the incident angle (αi) was fixed at 0.18°, which is larger than the critical angle of the polymer film (0.16°). The SAXS and GISAXS data were collected for 10 min. Absolute intensities were determined by standardless calibration using a PIN diode and a Pilatus 300K detector to measure the intensity of the incident beam and transmission for the sample and the scattered intensity, respectively.29,51 Surface Characterization. The surface morphologies of PSM-bPS thin films were investigated by AFM (Dimension 3100, Digital Instrument) as a function of exposure time to acid vapor. The AFM tip, which has a pyramidal shape with a radius of curvature of 6 nm at the tip apex, was operated in the tapping mode. Reactive ion etching (RIE, STS Vision 320 Mark II RIE System, RF power: 100 W, pressure: 10 mTorr) with O2/Ar (3/1 in volume ratio) plasma was used to generate the contrast prior to SEM measurement. The etching thicknesses of PS, PSM, and PGM homopolymers in the thin films as a function of etching time ranging from 5 to 20 s were obtained to adjust the etching time depending on the film thickness. SEM experiments were performed using an FEI Magellan XHR-400 field emission-SEM operated at 1 kV acceleration voltage and 50 pA beam current after an RIE process.
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Figure 2. GISAXS measurements of P(SM13-S14) thin films as a function of exposure time to TFA vapor. The GISAXS data were collected at room temperature for 10 min after thermal annealing at 150 °C for 12 h: (a) 2D GISAXS pattern taken after exposure of 80 min and (b) in-plane scattering profiles corresponding to a horizontal cut at qz = 0.270 nm−1, which were obtained from the 2D GISAXS patterns at different exposure times. The intensity profiles are vertically shifted for clarity, and the arrows indicate the primary peak (q*) of the lamellar morphology. The inset graph shows the inverse of a maximum intensity (1000/I(q*)) determined from the SAXS profiles.
by 1H NMR spectra in DMSO-d6 (Figure S2) displaying the ratio of the signal integral areas for the hydroxy protons in the GM segment at 4.55−4.96 ppm and the methyl hydrogens of the backbone at 0.65−1.05 ppm.52 After conversion, the primary scattering peak of P(SM13-S14) shifted from q* = 0.379 nm−1 to q* = 0.229 nm−1, resulting in an increase in L0 from 16.6 to 27.4 nm arising from the enhanced stretching of the copolymer chains at the interface, and multiple higherorder reflections were also observed at scattering vector ratios of 1:2:3:4 ≈ relative to q*, corresponding to an ordered lamellar morphology. For the disordered P(SM5-S5), P(SM3S4), and P(SM3-S3) samples, acid-catalyzed hydrolysis of the SM blocks induced microphase separation and resulted in ordered lamellar microdomains with an L0 of 13.7, 11.7, and 10.2 nm with corresponding q* = 0.460, 0.535, and 0.615 nm−1, respectively. For the thin films of PSM-b-PS copolymers, a random copolymer brush layer, composed of a hydroxyl-containing PSM-r-PS copolymer, was spin-coated on the silicon substrate to control the polymer/substrate interfacial interactions. This anchored random copolymer can effectively balance interfacial interactions by the simultaneous conversion of SM units to GM units with the PSM-b-PS layer. To obtain the surface and interfacial energies of PSM, PGM, and PS with PSM-r-PS and poly(glycerol monomethacrylate-b-styrene) (PGM-r-PS), the
RESULTS AND DISCUSSION
The SAXS intensity profiles of a series of PSM-b-PS copolymers in bulk [Mn = 26 800 g/mol (N = 197), 10 500 g/mol (N = 76), 6700 g/mol (N = 51), and 5600 g/mol (N = 42)] were measured at room temperature after thermal annealing at 140 °C for 24 h, as shown in Figure S1. P(SM13-S14), the highest molecular weight copolymer in this study, shows only one primary scattering peak at q* = 0.379 nm−1 with L0 = 2π/q* of 16.6 nm. For the three lower molecular weight copolymers, P(SM5-S5), P(SM3-S4), and P(SM3-S3), no scattering peaks were observed, indicating a phase-mixed morphology because of the small χ for PSM-b-PS (0.035 at 25 °C). Because acid-catalyzed hydrolysis of the hydrophobic PSM to the hydrophilic PGM markedly increased the χ between two blocks of the copolymer (0.438 at 25 °C), the formation of ordered morphology from a disordered state was found even for the very low molecular weight copolymers.30 This chemical transformation was corroborated D
DOI: 10.1021/acs.macromol.9b00488 Macromolecules XXXX, XXX, XXX−XXX
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Figure 3. Morphological evolution of P(SM5-S5) (N = 76) in the thin films as a function of exposure time ranging from 0 to 80 min to TFA vapor: (a,f) 0, (b,g) 20, (c,h) 40, (d,i) 60, and (e,j) 80 min. The samples were thermally annealed at 150 °C for 12 h and then measured by (a−e) AFM height and (f−j) height profiles (red line). The color contrast is shown in the AFM height images; dark area is thinner (holes) and bright area is thicker (islands). The scale bars represent 100 nm.
contact angles of each thin film were measured using water and methylene iodide, and the surface energies were calculated using the method described by Wu.53,54 For PSM (43.7 mJ/ m2), a value similar to that of PS (40.7 mJ/m2) was found; however, upon hydrolysis of PSM to more hydrophilic PGM, a large difference between PGM (52.8 mJ/m2) and PS surface energies was observed. The random copolymer provided a sufficient pathway for mitigating this large difference of the interfacial energies between PGM and PS blocks in the BCP and led to an orientation of the lamellar microdomains normal to the substrate surface. The small difference of the interfacial energies between PGM and PS on the random copolymer brush layer was found to be Δγ = 3.3 mJ/m2 at room temperature and can be diminished further at temperatures above Tg. All of the data for the contact angles and the surface and interfacial energies of polymers are summarized in Tables S1 and S2. The morphological evolution of the lamellar microdomains in the thin films as a function of exposure time to acid vapor was investigated by AFM. Since there is an ∼18% decrease in the volume of the PSM block when SM units are fully converted to GM, the symmetric PSM-b-PS copolymers show 8−9% decrease in the total volume. For the thin films, this volume reduction is evidenced in a thickness direction because of lateral constraints on the film. In consideration of the decrease in thickness by hydrolysis, the thickness of the fully converted thin film was fixed at 0.75L0 based on the period of PGM-b-PS copolymers, as the free-energy model of thin films introduced by Walton et al.55 predicts that a perpendicular orientation is energetically and highly favored at 0.75L0 when the effect of the polymer/substrate interface on the BCP orientations is minimized, such as a neutral interface.17,56 Consequently, the initial film thicknesses of P(SM13-S14) and P(SM3-S3) were set at 22.5 and 8.4 nm, corresponding to 0.75L0 for the respective PGM-b-PS copolymers. Figure 1 shows the AFM height images (a−e) and the height profiles (red line, f−j) for the surface morphological changes of
Figure 4. GISAXS measurements of P(SM5-S5) thin films as a function of exposure time to TFA vapor. The GISAXS data were collected at room temperature for 10 min after thermal annealing at 150 °C for 12 h; (a) 2D GISAXS patterns taken after exposure of 80 min and (b) in-plane scattering profiles corresponding to a horizontal cut at qz = 0.270 nm−1, which were obtained from the 2D GISAXS patterns at different exposure times. The intensity profiles are vertically shifted for clarity, and the arrows indicate the q* of the lamellar morphology. The inset graph shows the 1000/I(q*) determined from the SAXS profiles.
E
DOI: 10.1021/acs.macromol.9b00488 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules
Figure 5. Lamellar morphology of P(SM3-S4) (N = 51) and P(SM3-S3) (N = 42) in the thin films after exposure of 80 min to TFA vapor. The sample was thermally annealed at 150 °C for 12 h and then measured by (a,b) AFM height and (c−e) GISAXS. The in-plane scattering profile corresponding to a horizontal cut at qz = 0.270 nm−1 was obtained from the 2D GISAXS pattern, and the arrows indicate the q*y of the lamellar morphology. The scale bars represent 100 nm.
film.41,57 The 2D GISAXS patterns for the P(SM13-S14) thin films at each exposure time are shown in Figure S3, where qy indicates the in-plane scattering vector and qz is the out-ofplane scattering vector. For the pristine P(SM13-S14) thin film before conversion, no in-plane scattering peak was seen from vertically oriented lamellae; however, the primary scattering peak along the in-plane direction was generated after TFAvapor-catalyzed hydrolysis for 20 min, and then the intensity of the scattering peak increased with longer hydrolysis times, indicating stronger microphase separation of the copolymer. Figure 2 shows the 2D GISAXS pattern of P(SM13-S14) after TFA exposure for 80 min and the in-plane intensity profiles from the 2D GISAXS patterns for different exposure times. The in-plane primary scattering reflection arising from a perpendicular lamellar morphology was found at q*y = 0.241 nm−1 with an L0 of 26.1 nm, which was similar to the value in bulk (27.4 nm). In addition, the decrease in the inverse of a maximum intensity (1000/I(qy*)), quantified from the in-plane intensity profiles, with increasing exposure time to TFA vapor was attributed to improved vertical orientations, in agreement with the AFM results. Figure 3 shows morphological evolution of thermally annealed P(SM5-S5) thin films, measured by the AFM height images (a−e) and the height profiles (red line, f−j), as a function of exposure time to TFA vapor. For the pristine copolymer thin film, no orientation was observed with a flat surface consistent with a phase-mixed morphology as observed by the SAXS data in bulk. At 20 min of the exposure time, the hole-island structure with a 5.2 nm depth was seen because of the microphase separation arising from stronger segmental interactions; however, this early separation did not arise from a lamellar-ordered structure as, in our previous study,30 we confirmed that P(SM5-S5) showed only single primary peak
P(SM13-S14) in the thin films with 0−80 min exposure time after thermal annealing at 150 °C for 12 h. The pristine copolymer thin film features the hole-island structure with an 11.7 nm depth originating from a primary scattering peak with no higher-order reflections; however, after 20 min of exposure to TFA vapor, the perpendicular lamellar morphology with a short length scale was partially observed by the conversion of SM to GM, and the height difference was reduced to 5.7 nm. In contrast to the chemical transformation in solution, SM segments do not convert randomly to GM in the solid state, thus making the surface morphologies nonuniform.29 With increasing time for acid-catalyzed hydrolysis, longer lamellar microdomains were developed with a smoother surface because of stronger microphase separation. When the P(SM13-S14) thin film was exposed to TFA vapor for 80 min, a flat surface was found with
