Hierarchical Structures in Thin Films of Miktoarm Star Polymers: Poly(n

Article pubs.acs.org/Macromolecules

Hierarchical Structures in Thin Films of Miktoarm Star Polymers: Poly(n‑hexyl isocyanate)(12K)−Poly(ε-caprolactone)1−3(5K) Young Yong Kim,† Sungmin Jung,† Changsub Kim,† Brian J. Ree,† Daisuke Kawato,‡ Naoki Nishikawa,‡ Daichi Suemasa,‡ Takuya Isono,‡ Toyoji Kakuchi,*,‡ Toshifumi Satoh,*,‡ and Moonhor Ree*,† †

Division of Advanced Materials Science, Department of Chemistry, Center for Electro-Photo Behaviors in Advanced Molecular Systems, Polymer Research Institute, Pohang Accelerator Laboratory, and BK School of Molecular Science, Pohang University of Science and Technology, Pohang 790-784, Republic of Korea ‡ Division of Biotechnology and Macromolecular Chemistry, Faculty of Engineering, Hokkaido University, Sapporo 060-8628, Japan S Supporting Information *

ABSTRACT: A series of miktoarm star polymers, [poly(nhexyl isocyanate)(12K)]−[poly(ε-caprolactone) 1−3 (5K)] (PHIC−PCL1−3) (composed of a rigid self-assembling PHIC arm and one to three flexible crystallizable PCL arms), were investigated to examine the polymers’ thermal properties and nanoscale thin film morphologies. The miktoarm polymers were stable up to 180 °C. The PHIC and PCL arm components underwent phase separation during the solution casting, drying, and post toluene-annealing processes, forming interesting but very complex thin film morphologies. The resulting thin film morphologies were examined in detail for the first time using synchrotron grazing incidence X-ray scattering (GIXS) measurements and quantitative data analysis. All of the miktoarm star polymer films formed vertically well-oriented lamellar structures, regardless of the number and length of PCL arms. These structures were quite different from the cylindrical structures commonly observed in conventional flexible diblock copolymer films having comparable volume fractions. The individual PHIC and PCL lamellar domains self-assembled to form their own respective morphological structures. The PHIC lamellae consisted of a mixture of horizontal and vertical multibilayer structure domains, as observed in the PHIC homopolymer film. The PCL lamellae formed fringed micelle-like crystals and/or highly imperfect folded crystals that differed significantly from the structures found in a PCL homopolymer film composed of typical folded lamellar crystals. These PCL crystals were formed with a mixture of vertical and horizontal orthorhombic lattices. Overall, the GIXS analysis revealed that the parameters that characterized the hierarchical structures in the thin films depended significantly on the number and length of the PCL arm and its crystallization characteristics as well as the chain rigidity and multibilayer structure formation characteristics of the PHIC arm.

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(PCPPO)13 (crystalline), or poly(ε-caprolactone) (PCL)10 (crystalline), whereas the coil arm component has been polystyrene (PS), polyisoprene (PI), or poly(ethylene oxide) (PEO). The morphologies and structures of the miktoarm star polymers in the bulk and solution states have been characterized using a variety of techniques.1−6,8−17 The morphologies and structures of the thin films, on the other hand, have not been extensively studied.7 In fact, functional miktoarm star polymers may constitute a novel class of advanced nanoscale thin film materials with potential applications in microelectronic, optical, and optoelectronic devices. The morphologies and structures of miktoarm star polymer thin films can differ significantly from the correspond-

INTRODUCTION Miktoarm star polymers (i.e., chemically asymmetric star polymers) have attracted significant attention over the past two decades due to their distinctive morphological structures, bulk properties, solution properties, and their potential utility in nanoscience applications.1−3 Much effort has been applied toward the development of high-performance miktoarm star polymers. As a result, several coil−coil miktoarm star polymers have been reported.1−7 A handful of rod−coil miktoarm star polymers have been synthesized8−12 in an effort to study the differences between their self-assembly behaviors and those observed in conventional coil−coil miktoarm star copolymers. The rod arm component in the rod−coil miktoarm systems studied thus far has been poly(glutamic acid) (PGA) and its derivative,8 poly(ε-tert-butyloxycarbonyl-L-lysine) (PBLL) and its derivatives,9 poly(n-hexylisocyanate) (PHIC),10,11 poly{2,5bis[(4-methoxyphenyl)oxycarbonyl]styrene} (PSPMPCS)12 (liquid crystal), poly(1-cyano-4-propyl-2,5-phenylene oxide) © 2014 American Chemical Society

Received: August 19, 2014 Revised: October 14, 2014 Published: October 21, 2014 7510

dx.doi.org/10.1021/ma501695b | Macromolecules 2014, 47, 7510−7524

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Scheme 1. Syntheses of PHIC−N3, PHIC−OH, PHIC−(OH)2, PHIC−(OH)3, PHIC−PCL, PHIC−PCL2, and PHIC−PCL3

Table 1. Fundamental Characteristics of the PHIC and PCL Homopolymers and Their Miktoarm Star Polymers Used in the Study PHIC−(OH)ma polymer PHIC PCL PHIC−PCL PHIC−PCL2 PHIC−PCL3

Mn,NMRb

11400 11300 11300

Mw/Mnc

Mn,clcdd

Mn,NMRb

Mw/Mnc

f PHICe (%)

Tcf (°C)

Tmg (°C)

ΔHfh (J/g)

Xci (%)

1.19 1.19 1.18

12000 5000 17000 17000 17000

10400 4600 16700 16400 16100

1.11 1.06 1.19 1.11 1.11

71 72 73

28.3 11.3 −1.1 −2.0

52.4 47.7 42.1 37.8

93.5 88.9 86.8 74.0

69.0 65.6 64.1 54.6

Macroinitiator (m = 1−3) used in the polymerization of ε-caprolactone. bNumber-average molecular weight determined in CDCl3 by 1H NMR spectroscopy analysis. cPolydispersity determined by size exclusion chromatography (SEC) in THF using polystyrene standards. dMn,clcd = ([M]0/ [I]0) × conversion × (molecular weight of monomer) + Mn,initiator; [M]0, initial concentration of monomer; [I]0, initial concentration of initiator. e Volume fraction of PHIC arm: f PHIC = (Mn,NMR,PHIC/dPHIC)/(Mn,NMR,PHIC/dPHIC + Mn,NMR,PCL/dPCL); dPHIC = 1.00 g/cm3 and dPCL = 1.15 g/cm3. f Crystallization temperature of PCL, which corresponds to the temperature at the peak maximum of the exothermic crystallization transition measured by DSC analysis with a rate of 10.0 °C/min. gMelting temperature of PCL crystals, which corresponds to the temperature at the peak maximum of the endothermic melting transition measured by DSC analysis with a rate of 10.0 °C/min. hHeat of fusion of PCL crystals, which was measured by DSC analysis with a rate of 10.0 °C/min. iOverall crystallinity of PCL homopolymer or PCL arms, which was estimated from the heat of fusion of PCL crystals with respect to that of perfect PCL crystals (ref 24). a

crystalline arm component, and PCL or poly(L-lactide) (PLLA) comprised the other crystalline component. These miktoarm star polymers were synthesized via the living ring-opening polymerization of L-lactide or ε-caprolactone, with either a hydroxyl end-functionalized PHIC (PHIC-OH or PHIC(OH)2 or PHIC-(OH)3) macroinitiator, respectively. These precursors were obtained via a click reaction of the azido endfunctionalized PHIC (PHIC-N3) and the ethynyl alcohol derivatives. The thermal and solution properties and the micelle formation characteristics were examined. The morphological

ing properties in the bulk state because their molecular configuration and structural ordering can be affected by the interface between the substrate and air or a vacuum. The morphologies and structures within the confined nanodomains of the thin films can differ significantly from those in the bulk state. Thus, it is important to investigate the morphology and structural ordering of the miktoarm star polymer thin films. We recently reported the precise synthesis of crystalline− crystalline miktoarm star polymers as a new class of miktoarm star polymers.18 In these syntheses, PHIC comprised one 7511


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substrate is much stronger than the intensity of GIXS near the critical angle. Atomic force microscopy (AFM) measurements were performed to observe topography images of the miktoarm starpolymer thin films. The PHIC−PCLm star polymer films were imaged using a scanning probe microscope (model Multimode Nanoscope IIIa, Veeco, Santa Clara, CA) in tapping mode, which was equipped with a JV scanner; noncoated silicon etched probes (model LTESP, Veeco) were used in the measurements. Contact angle measurements of polymer solutions were conducted for Si substrates with a native oxide layer (Si/SiOx) by the sessile drop technique using a contact angle meter (KSV Instruments, Tokyo, Japan). The measurements were carried out at 25 °C using PHIC and PCL homopolymer solutions (0.5 wt %) in chloroform. The contact angle was measured 5 times per experiment, and the resulting contact angles were averaged out. Thermogravimetric analysis (TGA) and differential scanning calorimetry (DSC) measurements were carried out under a nitrogen atmosphere using a thermogravimeter (model TG/DTA 6200, Seiko Instruments, Tokyo, Japan) and a calorimeter (model DSC 6200, Seiko Instruments). A rate of 10.0 °C/min was employed for heating and cooling runs.

and structural details of miktoarm star polymers have not yet been examined. This study investigated the structural details of a series of miktoarm polymers, [poly(n-hexyl isocyanate)]−[poly(ε-caprolactone)]m (PHIC−PCLm: m = 1−3; Mn,NMR(PHIC) = 11 300−11 400, and Mn,NMR(PCLm) = 4800−5300, where Mn,NMR is the number-average molecular weight determined by nuclear magnetic resonance (NMR) spectroscopy) in nanoscale thin films using synchrotron grazing incidence X-ray scattering (GIXS) methods. Differential scanning calorimetry (DSC) and atomic force microscopy were employed as complementary techniques. Interestingly, toluene-annealed films of all the miktoarm polymers formed a well-defined vertical lamellar structure via phase separationa distinct feature compared to the structures observed in common, flexible diblock copolymers of comparable volume fractions. This unique vertical lamellar structure included the PHIC lamellae composed of a mixture of horizontal and vertical multibilayer structure domains, as observed in the PHIC homopolymer film. The PCL lamellae, however, formed fringed micelle-like crystals and/or highly imperfect folded crystals, which differed significantly from that of the PCL homopolymer film. The structural parameters were significanty influenced by the presence of the PHIC and PCL arms as well as the number and length of the PCL arms. Overall, the complex morphologies of the miktoarm star polymers in nanoscale thin films were attributed to the rigid chain properties and the multibilayer structure formation characteristics of the PHIC arm and in part by the morphology of the PCL arms rising from their crystallization.

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RESULTS AND DISCUSSION The PHIC−PCLm star polymers were found to undergo a twostep degradation process, regardless of the number of arms present, as shown in Figure 1a−c. The first step of the degradation process began around 180 °C (= Td1, the degradation temperature), and the second step of the degradation began around 290 °C (= Td2). The extent of degradation that occurred in the first step was less than that in the second step. Furthermore, the PHIC homopolymer was

EXPERIMENTAL SECTION

Materials and Thin Film Preparation. A series of PHIC−PCLm star polymers (m = 1−3) were prepared according to the synthetic methods reported previously,18 as shown in Scheme 1. In addition, PHIC and PCL homopolymers were synthesized. Details of the syntheses were given in the Supporting Information. The fundamental characteristics of the obtained polymer products are summarized in Table 1. Each polymer product was dissolved in chloroform and filtered using a disposable syringe equipped with polytetrafluoroethylene filter of pore size 0.2 μm, producing a 1.0 wt % solution. The obtained polymer solutions were spin-coated onto precleaned silicon substrates and dried in vacuum at room temperature for 1 day. The obtained polymer thin films were determined to have a thickness between 115 and 125 nm by using a spectroscopic ellipsometer (model M2000, J.A. Woollam, Lincoln, NE). A set of the film samples was thermally annealed in various conditions. Other sets of the film samples were put under various solvent-anneal treatments with toluene, tetrahydrofuran, chloroform, and carbon disulfide. Measurements. GIXS measurements were conducted at the 3C beamline19,20 of the Pohang Accelerator Laboratory (PAL), Pohang University of Science & Technology (POSTECH), Pohang, Korea. The samples were measured at a sample-to-detector distance (SDD) of 2922 and 928 mm for grazing incidence small-angle scattering (GISAXS) and 219 mm for wide grazing incidence wide-angle scattering (GIWAXS). Scattering data were typically collected for 30− 60 s using an X-ray radiation source of λ = 0.1180 nm (λ, wavelength) with a two-dimensional (2D) charge-coupled detector (CCD) (model Rayonix 2D MAR, Evanston, IL). The incidence angle αi of the X-ray beam was set in the range 0.140°−0.165°, which is between the critical angle of the polymer thin film and the silicon substrate (αc,f and αc,s). Scattering angles were corrected according to the positions of the Xray beams reflected from the silicon substrate with respect to precalibrated polystyrene-b-polyethylene-b-polybutadiene-b-polystyrene block copolymer or silver behenate powder (TCI, Tokyo, Japan). Aluminum foil pieces were applied as a semitransparent beam stop because the intensity of the specular reflection from the

Figure 1. TGA and DSC thermograms of the PHIC−PCLm (m = 1− 3) miktoarm star polymers and their homopolymers, which were measured with a rate of 10.0 °C/min in a nitrogen atmosphere: (a, f) PHIC−PCL; (b, g) PHIC−PCL2; (c, h) PHIC−PCL3; (d, i) PHIC; (e, j) PCL. 7512


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Table 2. Nanostructural Parameters of the Toluene-Annealed Thin Films of the PHIC and PCL Homopolymers and Their Miktoarm Star Polymers, Which Were Determined by GISAXS Analysis polymer thin film

nanostructure (observed)

La (nm)

S1b (nm)

S2 c (nm)

S3 d (nm)

PHIC PCL PHIC−PCL PHIC−PCL2 PHIC−CL3

not identified horizontal lamellae vertical lamellae vertical lamellae vertical lamellae

11.5 30.5 27.8 27.0

3.1 7.0 6.6 6.0

1.1 0.7 0.8 0.8

6.3 22.1 19.6 19.4

σ1e (nm)

σ2f (nm)

gg

ϕ̅ 1h (deg)

σφ1i (deg)

Osj

S k (nm)

0.8 0.3 0.5 0.5

0.9 0.5 0.5 0.5

0.280 0.135 0.175 0.215

0 90 90 90

7.9 8.0 8.9 11.1

0.972 −0.471 −0.465 −0.446

4.1−10.1 4.0−5.4 4.0−5.0

a

Long period of lamellar structure. bThickness of more dense (i.e., crystalline) layer (= Sc ) in the lamellar structure formed in PCL homopolymer films; thickness of more dense layer (i.e., PCL arm phase: SPCL ) in the lamellar structured PHIC−PCLm films. cThickness of interfacial layer (Si ) between the highly dense and less dense layers in lamellar structure. dThickness of less dense (i.e., amorphous) layer (= Sa ) in the lamellar structure formed in PCL homopolymer films; thickness of less dense layer (i.e., PHIC arm phase: SPHIC ) in the lamellar structured PHIC−PCLm films. e Standard deviation for the more dense layer in lamellar structure. fStandard deviation for the interfacial layer in lamellar structure. gParacrystal distortion factor along the direction parallel to the long period of lamellar structure. hMean value of the polar angle φ1 (i.e., orientation angle) between the orientation vector n1 (which is set along the direction parallel to the long period of lamellar structure) and the out-of-plane of the film. i Standard deviation for the orientation angle φ1 of lamellar structure. jSecond-order orientation factor. kMean interdistance of the crystals formed in the PCL arm phase layers along the out-of-plane of the film.

Figure 2. GISAXS data of toluene-annealed PHIC and PCL homopolymer films (120 nm thick) deposited on silicon substrates, measured with an incidence angle αi of 0.140° at room temperature using an X-ray beam with a wavelength λ of 0.1180 nm; (a) 2D GISAXS pattern of PHIC film; (b) out-of-plane scattering profile extracted from the data in (a) along the αf direction at 2θf = 0.058° where αf and 2θf are the exit angles of the X-ray beam with respect to the film surface and to the plane of incidence, respectively; (c) in-plane scattering profile extracted from the data in (a) along the 2θf direction at αf = 0.172°; (d) 2D GISAXS pattern of PCL film; (e) out-of-plane scattering profile extracted from the data in (d) along the αf direction at 2θf = 0.058°; (f) in-plane scattering profile extracted from the data in (d) along the 2θf direction at αf = 0.172°; (g) 2D GISAXS image of PCL film reconstructed from the structural parameters in Table 2 using the GIXS formula derived for lamellar structure model; (h) lamellar structure model where n1 is the orientation vector of the structure model and φ1 is the polar angle between the n1 vector and the out-of-plane of the film. In (e) and (f), the black symbols are the measured data, and the red solid lines were obtained by fitting the data using the GIXS formula.

displayed a strong endothermic peak with a heat of fusion ΔHf of 93.5 J/g at 52.5 °C, which corresponded to the melting transition of the crystals (Tm,PCL). During the cooling run, the PCL homopolymer displayed a strong exothermic peak around 28.3 °C, which was attributed to the crystallization process (T c,PCL ) (Figure 1j). The DSC characteristics of the homopolymers suggested that the exothermic and endothermic transitions in the star polymers were attributed to the crystallization and crystal melting of the PCL arms. The PHIC−PCL polymer was characterized by Tc,PCL = 11.3 °C, Tm,PCL = 47.7 °C, and ΔHf,PCL = 88.9 J/g, all of which values were lower than those obtained from the PCL homopolymer, although the molecular weight of the PCL arm was nearly equal to that of the homopolymer (Table 2). The Tc,PCL, Tm,PCL, and ΔHf,PCL values were further reduced as the number of PCL

found to begin degradation at 180 °C, whereas the PCL homopolymer began degradation around 290 °C (Figure 1d,e). Considering these facts, the first step of degradation was attributed to the PHIC arm, and the second step of degradation was attributed to the PCL arms. The TGA results were complemented with a DSC analysis at a rate of 10.0 °C/min under a nitrogen atmosphere over the temperature range −50 to 170 °C. During the heating run, the PHIC−PCLm star polymers displayed a single endothermic peak over the temperature range 20−55 °C, whereas during the cooling runs, they exhibited a single exothermic peak over the temperature range −20 to 25 °C (Figure 1f−h). By contrast, the PHIC homopolymer did not display any discernible phase transitions during the heating and subsequent cooling runs (Figure 1i). During the heating run, the PCL homopolymer 7513


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horizontal lamellar structure and the out-of-plane direction of the film was 0°, with a standard deviation σφ1 of 7.9°. The positional distortion factor g of the lamellar structure was found to be 0.280. These results collectively indicated that a horizontal lamellar crystal structure developed in the tolueneannealed PCL thin film. The g and σφ1 values were not small, indicating the presence of a certain level of disorder in the horizontal lamellar crystal structure in the PCL thin film. The structural parameters were used to reconstruct the 2D GISAXS image. The reconstructed scattering image is displayed in Figure 2g and agreed well with the experimental data. The toluene-annealed PHIC−PCL films displayed a featured GISAXS pattern that differed significantly from that of the PCL homopolymer film. The representative GISAXS patterns are shown in Figure 3a,b. The in-plane and out-of-plane scattering

arms increased (Table 2). These reductions may have resulted in part from the lower molecular weight of the PCL arms as the arm number increased. The measured ΔHf,PCL values permitted estimation of the crystallinity Xc, found to be 68.8% for the PCL homopolymer, 65.4% for PHIC−PCL, 63.8% for PHIC− PCL2, and 54.4% for PHIC−PCL3, with respect to the heat of fusion (136 J/g) of a perfect PCL crystal.21 The DSC analysis results indicated that the PCL crystals formed in the PHIC− PCLm polymers were smaller than those formed in the PCL homopolymer. A series of GIXS studies were conducted on the thin films (115−125 nm thick) of the PHIC and PCL homopolymers and their miktoarm star polymers. The as-cast films of the miktoarm star polymers composed of highly disordered structures in random orientation. Thus, the as-cast films were subjected to different treatments, including thermal annealing under various conditions and solvent annealing in various solvents, to induce the formation of nanostructures. Toluene annealing at room temperature over 3 h prompted the polymer thin films to develop well-defined, highly ordered nanostructures, whereas the other annealing conditions produced highly disordered, randomly oriented structures. Figure 2a shows a representative 2D GISAXS pattern of the toluene-annealed PHIC homopolymer thin films. Out-of-plane and in-plane scattering profiles were extracted from the scattering pattern, as shown in Figure 2b,c. The GISAXS pattern was featureless, suggesting that in the toluene-annealed PHIC film there was developed either no nanostructure or a structure whose structural parameters could not be identified by this GISAXS analysis due to limits in the low and high scattering angle regions. Unlike the PHIC films, the toluene-annealed PCL homopolymer films produced a featured GISAXS pattern, as shown in Figure 2d. The out-of-plane and in-plane scattering profiles were extracted, as shown in Figure 2e,f. The out-ofplane scattering profile displayed two broad weak scattering peaks over the range 0.35°−0.85°. One peak was centered at αf = 0.48° and the other peak at αf = 0.62°. The breadth of the scattering peaks was attributed to the heavy overlap between the scattering features along the αf direction due to the reflected and transmitted X-ray beams, which is a feature of the GISAXS measurements gathered at the chosen grazing incidence angle αi (0.140°). Here, the peak at αf = 0.48° arose from the transmitted X-ray beam, whereas the peak at αf = 0.62° arose from the reflected X-ray beam. These peaks indicated a d-spacing of 11.5 nm. By contrast, the in-plane scattering profile was featureless. These scattering characteristics suggested that in the film the PCL polymer chains formed a horizontally oriented lamellar crystal structure. The out-ofplane and in-plane scattering profiles could be satisfactorily fitted using the GIXS formula22,23 derived for a lamellar structural model (Figure 2h). The GIXS formula derivation is given in the Supporting Information. The obtained structural parameters are listed in Table 2. The horizontal lamellar structure was found to have L = 11.5 nm. Each lamella was composed of three sublayers: a highly dense sublayer of 3.1 nm thick (= S1, which corresponds to the crystal layer thickness Sc ), an interfacial layer of 1.1 nm thick (= S2 = Si ), and a less dense layer 6.3 nm thick (= S3, which corresponds to the amorphous layer thickness Sa ). The lamellar structure was determined to have a second-order orientation factor Os,1 of 0.972. The mean polar angle ϕ̅ 1 between the orientation vector n1 of the

Figure 3. GISAXS data of toluene-annealed PHIC−PCL films (120 nm thick) deposited on silicon substrates, measured at room temperature using an X-ray beam (λ = 0.1180 nm): (a) 2D scattering pattern measured with αi = 0.165° at a sample-to-detector distance (SDD) of 928 mm; (b) 2D scattering pattern measured with αi = 0.140° at SDD = 2922 mm; (c) 2D scattering image reconstructed from the structural parameters in Table 2 using the GIXS formula derived for lamellar structure model; (d) in-plane scattering profile extracted from the data in (a) and (b) along the 2θf direction at αf = 0.160° where the black symbols are the measured data, and the red solid lines were obtained by fitting the data using the GIXS formula; (e) out-of-plane scattering profile extracted from the data in (a) and (b) along the αf direction at 2θf = 0.056°.

profiles, which were extracted at αf = 0.160° and 2θf = 0.213°, respectively, are shown in Figure 3d,e. The in-plane scattering profile clearly showed two peaks at 0.22° and 0.43° (Figure 3d). The relative scattering vector lengths from the specular reflection position were 1 and 2, respectively, indicating that these peaks arose from the same structural origin. The two peaks could not be distinguished in the out-of-plane scattering profile (Figure 3e). The scattering characteristics suggested the formation of a vertical lamellar structure in the PHIC−PCL film. Indeed, the two peaks could be assigned as the first- and second-order reflections of the vertical lamellar structure. The in-plane and out-of-plane scattering profiles could be satisfactorily fitted to the GIXS formula derived for a lamellar structure. The analysis results are summarized in Table 2. The lamellar structure was characterized by L = 30.5 nm, S1= 7.0 nm 7514


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(sublayer 1), S2 = 0.7 nm (sublayer 2), S3= 22.1 nm (sublayer-3), Os,1 = −0.471 (ϕ̅ 1 = 90° and σφ1 = 8.0°), and g = 0.135. The calculated structural parameters were used to reconstruct the 2D GISAXS image, as shown in Figure 3c. The reconstructed image agreed well with the experimental data. The GISAXS analysis results collectively indicated the development of a wellordered vertical lamellar structure in the toluene-annealed PHIC−PCL film. The dimensions of the lamellar structure were much larger than those of the lamellar structure formed in the PCL homopolymer film. This fact, together with the block copolymer chemical structure, suggested that a vertical lamellar structure developed in the PHIC−PCL film during the film formation process and subsequent toluene annealing step via the phase separation of the PHIC and PCL arms. The volume fractions of the PHIC and PCL arms suggested that the 7.0 nm thick sublayer 1 could be assigned as the PCL phase (S1 = SPCL ) and the 22.1 nm thick sublayer 3 could be assigned as the PHIC phase (S3 = SPHIC). The 0.7 nm thick sublayer 3 corresponded to the interfacial layer (S2 = Si ) between the PCL and PHIC phases. It should be noted that the interfacial layer present in the vertically oriented lamellar structured PHIC−PCL film was thinner than that formed in the horizontally oriented lamellar crystals in the PCL homopolymer film. As discussed above, the GISAXS analysis confirmed the formation of horizontally oriented lamellar crystals in the PCL homopolymer film; however, no nanostructures appeared to be present in the PHIC film. These observations raised questions about the morphological structure of the vertical PCL lamellae in the PHIC−PCL film. Additional structural information about the vertical PCL lamellar phase was sought by extracting an out-of-plane scattering profile along the αf direction at 2θf = 0.056° from the 2D scattering patterns shown in Figure 3a,b. As shown in Figure 3e, the out-of-plane scattering profile revealed a weak broad peak centered at αf = 0.63° (10.1 nm d-spacing) and another weak peak centered at αf = 1.67° (4.1 nm dspacing). The observation of these scattering peaks suggested that a heterogeneous morphology formed in the vertical PCL lamella, attributed to the crystallization of the PCL arms at a certain level. The crystals that formed were distributed over a mean distance of 4.1−10.1 nm over the vertical PCL lamella. The toluene-annealed PHIC−PCL2 and PHIC−PCL3 films also yielded featured GISAXS patterns similar to those obtained from the PHIC−PCL films (Figures 4a,b and 5a,b). The GISAXS patterns indicated the formation of vertically oriented lamellar structures in both toluene-annealed films. The in-plane and out-of-plane scattering profiles could be satisfactorily fitted using the GIXS formula derived for a lamellar structure, as shown in Figures 4d,e and 5d,e. The analysis results are summarized in Table 2. The determined structural parameters were used to reconstruct the 2D GISAXS images. The reconstructed scattering images are shown in Figures 4c and 5c, in good agreement with the experimental data. These GISAXS analysis results collectively indicated that a wellordered vertical lamellar structure had developed in the PHIC− PCL2 film as well as in the PHIC−PCL3 film through the film formation process and subsequent toluene annealing via the phase separation of the PHIC and PCL arms. Both the PHIC− PCL2 and PHIC−PCL3 films indicated the formation of a heterogeneous morphology in the vertical PCL lamellar phases due to the crystallization of the PCL arms, as evident in the outof-plane scattering profiles extracted along the αf direction at 2θf = 0.056° from the 2D scattering patterns (Figures 4e and

Figure 4. GISAXS data of toluene-annealed PHIC−PCL2 films (120 nm thick) deposited on silicon substrates, measured at room temperature using an X-ray beam (λ = 0.1180 nm): (a) 2D scattering pattern measured with αi = 0.165° at SDD = 928 mm; (b) 2D scattering pattern measured with αi = 0.140° at SDD = 2922 mm; (c) 2D scattering image reconstructed from the structural parameters in Table 2 using the GIXS formula derived for lamellar structure model; (d) in-plane scattering profile extracted from the data in (a) and (b) along the 2θf direction at αf = 0.160° where the black symbols are the measured data, and the red solid lines were obtained by fitting the data using the GIXS formula; (e) out-of-plane scattering profile extracted from the data in (a) and (b) along the αf direction at 2θf = 0.056°.

Figure 5. GISAXS data of toluene-annealed PHIC−PCL3 films (120 nm thick) deposited on silicon substrates, measured at room temperature using an X-ray beam (λ = 0.1180 nm): (a) 2D scattering pattern measured with αi = 0.165° at SDD = 928 mm; (b) 2D scattering pattern measured with αi = 0.140° at SDD = 2922 mm; (c) 2D scattering image reconstructed from the structural parameters in Table 2 using the GIXS formula derived for lamellar structure model; (d) in-plane scattering profile extracted from the data in (a) and (b) along the 2θf direction at αf = 0.160° where the black symbols are the measured data, and the red solid lines were obtained by fitting the data using the GIXS formula; (e) out-of-plane scattering profile extracted from the data in (a) and (b) along the αf direction at 2θf = 0.056°.

5e). The crystals that formed were distributed with a mean intercrystal distance of 4.0−5.4 nm over the vertical PCL 7515


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polymer backbones in the structure were present in a fully extended conformation and were packed laterally, whereas the n-hexyl bristles were fully extended and packed laterally with no interdigitation. The analysis determined a mean interpolymer distance of 0.67 nm (= d1) between the n-hexyl bristles of neighboring polymer chains, the repeat units of which were matched in position along the backbone. The n-hexyl bristles were characterized by an interpolymer distance of 0.50 nm (= d2) along the polymer backbone. The bristles between the nearest-neighboring polymer chains had a mean interpolymer distance of 0.45 nm (= d3). The n-hexyl bristles were characterized by a very small positional distortion factor (gdL = 0.056) along the out-of-plane direction of the multibilayer structure; however, relatively large positional distortion factors (gd1 = 0.140, gd2 = 0.120, and gd3 = 0.130) were measured along the lateral directions. The horizontally oriented multibilayer structure was characterized by Os,2,0 = 0.990 (ϕ̅ 2 = 0° and σφ2 = 4.66°; ϕ̅ 2 is the mean polar angle between the orientation vector n2 of the horizontal multibilayer structure and the outof-plane direction of the film), whereas the vertically oriented multibilayer structure was characterized by Os,2,90 = −0.437 (ϕ̅ 2 = 90° and σφ2 = 11.97°). The volume fraction ratio (ϕh,2/ϕv,2) of the horizontal and vertical multibilayer structures in the film was estimated to be 78/22. These results collectively indicated that the toluene-annealed PHIC film consisted of a horizontal molecular multibilayer structure as the major structural component and a vertical molecular multibilayer structure as the minor component. The polymer chains were fully extended in the 21 conformation, with no interdigitation between the nhexyl bristles of adjacent molecular layers. Figure 7b shows representative 2D GIWAXS patterns obtained from the toluene-annealed PCL homopolymer films. The out-of-plane and in-plane scattering profiles are displayed in Figure 7f,g. The PCL homopolymer films produced a featured scattering pattern. The PCL homopolymer molecule self-assembled very rapidly to form crystals with an orthorhombic lattice unit and a space group of P212121.24 These observations suggested that the arc scattering peaks at 7.6° (0.89 nm d-spacing) and 15.1° (4.47 nm d-spacing) along the αf direction at 2θf = 0° could be assigned as the {002} and {004} reflections, respectively. The scattering peaks at 15.3° (0.44 nm d-spacing) and 16.8° (4.40 nm d-spacing) along the 2θf direction at αf = 0.20° could be assigned as the {110} and {200} reflections, respectively. The other two scattering peaks could be assigned to be the {101} and {102} reflections, respectively, as shown in Figure 7b. The scattering pattern was successfully analyzed using the GIXS formula for an orthorhombic crystal lattice model (Figure 7f,g). The GIXS formula derivation is provided in the Supporting Information. The scattering data analysis results are listed in Table 3. The orthorhombic crystal lattices were characterized by the dimensional parameters, a = 0.80 nm, b = 0.53 nm, and c = 1.78 nm. The c value corresponded to the length of two repeat units plus one carbonyl carbon atom in a fully extended conformation. The crystal lattice was characterized by a very small positional distortion factor (g = 0.050), indicating that all orthorhombic crystals were stably positioned in the film. Interestingly, the orthorhombic crystals were present in two different orientations: the vertical and horizontal orientations. The vertically oriented crystals were characterized by Os,3,0 = 0.957 (ϕ̅ 3 = 0° and σφ3 = 9.79°). Here, ϕ̅ 3 is the mean polar angle between the orientation vector n3 of the orthorhombic

lamellar phases in both the PHIC−PCL2 and PHIC−PCL3 films. As discussed above, all of the miktoarm PHIC−PCLm star polymers examined in our study underwent favorably phase separation through the nanoscale film formation on silicon substrates and subsequent toluene annealing, inducing the selfassembly of well-defined vertical lamellar structures, regardless of the number and length of PCL arms present. The vertical lamellar structured morphologies were confirmed by AFM analysis, as shown in Figure 6.

Figure 6. AFM images (height image: a, c, and e; phase image: b, d, and f) in 1000 × 1000 nm2 of toluene-annealed PHIC−PCLm films deposited on silicon substrates: (a, b) PHIC−PCL; (c, d) PHIC− PCL2; (e, f) PHIC−PCL3.

The toluene-annealed polymer films were further examined by GIWAXS analysis. As shown in Figure 7a, the PHIC homopolymer films revealed a featured scattering pattern that differed significantly from the featureless GISAXS patterns discussed above. The films showed periodic arc peaks with regular spacings along the αf direction at 2θf = 0°: αf = 4.03°, 8.05°, and 11.80°. These arc peaks were also discernible along the 2θf direction at αf = 0°; however, their intensities were much weaker than those along the αf direction at 2θf = 0°. The appearance of these scattering peaks suggested that both the horizontally and vertically oriented lamellar structures were present in the film. The d-spacing of the first-order arc peak was determined to be 1.67 nm, larger than the diameter of the helical conformational PHIC polymer chain but close to twice the length (0.83 nm) of the fully extended n-hexyl bristle. The scattering pattern showed additional arc peaks at αf = 10.1° (0.67 nm d-spacing) and 2θf = 0°. Other weak arc peaks were discernible at 13.8° (0.50 nm d-spacing) and 14.8° (0.45 nm dspacing). These peaks may have resulted from the lateral packing of the polymer backbones and the n-hexyl bristles. Based on these results, the scattering pattern was analyzed in detail using the GIXS formula22,23 for a molecular multibilayer structure model (Figure 7f,g). The GIXS formula derivation is provided in the Supporting Information. The analysis results are summarized in Table 3. The molecular multibilayer structures were characterized as having a long period dL of 1.67 nm, in which the individual lamellae were composed of two sublayers, i.e., a more dense sublayer of 0.78 nm thick (= S1) and a less dense sublayer of 0.89 nm thick (= S2 ). The 7516


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Figure 7. GIWAXS data of toluene-annealed PHIC−PCLm and homopolymer films (120 nm thick) deposited on silicon substrates, measured with αi = 0.145° at room temperature using an X-ray beam (λ = 0.1180 nm); (a) PHIC; (b) PCL; (c) PHIC−PCL; (d) PHIC−PCL2; (e) PHIC−PCL3; (f) out-of-plane scattering profiles extracted from the data in (a−c) along the αf direction at 2θf = 0.0°; (g) in-plane scattering profiles extracted from the data in (a−c) along the 2θf direction at αf = 0.200°. In (f) and (g), the black symbols are the measured data, and the red solid lines were obtained by fitting the data using the GIXS formulas.

The nanoscale thin film morphology details determined above are schematically illustrated in Figures 8, 9, and 10. Overall, the quantitative GISAXS and GIWAXS analyses provided structural insights into the thin film morphologies that developed from the PHIC and PCL homopolymers and their miktoarm star polymers as follows. First, the PHIC homopolymer chains in the toluene-annealed thin films underwent self-assembly and favorably formed a wellordered multibilayer structure with no interdigitation. Both the backbone and the n-hexyl bristles were fully extended (Figure 8). The formation of this well-ordered structure was attributed to the favorable polar−polar interactions, between the carbonyl carbon units of the fully extended backbone and the tertiary amine units of the adjacent backbone, and to the nonpolar (i.e., van der Waals) interactions between the fully extended n-hexyl bristles along the backbone and those on the neighboring backbones. Interestingly, the multibilayer structure was found to form two types of orientational domain in the film: a horizontal multibilayer structure formed the major domain and a vertical multibilayer structure formed the minor domain. These orientational domains could not be detected in the GISAXS measurements, as discussed above. This failure to detect these domains may have resulted from either very low electron density contrast between the orientational domains or a small minor domain size that was beyond the detection limits of GISAXS measurements. The orientational domains observed here differed significantly from the horizontal multibilayer structures that uniformly formed in thin films of the PHIC homopolymer synthesized using a sodium benzanalide initiator.25 Thus, the benzyl ether end group of the PHIC

crystal lattice and the out-of-plane direction of the film, whereas the horizontally oriented crystals were characterized by Os,3,90 = −0.483 (ϕ̅ 3 = 90° and σφ3 = 6.11°). The volume fraction ratio (ϕv,3/ϕh,3) in the vertical and horizontal crystals in the PCL film was estimated to be 81/19. The PHIC−PCLm films yielded clearly featured scattering patterns, as shown in Figure 7c−e. The out-of-plane and inplane scattering profiles are shown in Figure 7f,g. All GIWAXS patterns consisted of scattering peaks from both the PHIC phases present in multibilayer structures and from the PCL phases present in orthorhombic crystal lattices. The scattering patterns were satisfactorily analyzed using the GIXS formula for the combination of a multibilayer structural model and an orthorhombic crystal lattice model (Figure 7f,g). The scattering data analysis results are listed in Table 3. These results indicated that the PHIC layer in the PHIC−PCLm films that formed a vertical lamellar structure was composed of both horizontal and vertical multibilayer structures, whereas the PCL layer consisted of both horizontal and vertical orthorhombic crystals, regardless of the PCL arm number. The multibilayer structures had structural characteristics identical to those observed in the PHIC homopolymer film; however, their orientation factors (Os,2,0 and Os,2,90) and volume fraction ratios (ϕh,2/ϕv,2) varied with the PCL arm number. Lower ϕh,2/ϕv,2 values resulted from higher PCL arm numbers. On the other hand, the orthorhombic crystals that formed in the PHIC− PCLm films had lattice characteristics identical to those observed in the PCL homopolymer film, but their orientation factors (Os,3,0 and Os,3,90) and volume fraction ratio (ϕv,3/ϕh,3) were influenced by the PCL arm number. 7517


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Table 3. Structural Parameters of the Toluene-Annealed Thin Films of the PHIC and PCL Homopolymers and Their Miktoarm Star Polymers, Which Were Determined by GIWAXS Analysis structural parameter

PHIC

PCL

PHIC− PCL

PHIC− PCL2

PHIC− PCL3

dLa (nm) S1b (nm) S2 c (nm) d1d (nm) d2e (nm) d3f (nm) gdLg (nm) gd1h (nm) gd2i (nm) gd3j (nm) ϕ̅ 2k (deg) σφ2,0l (deg) Os,2,0m ϕ̅ 2 (deg) σφ2,90 (deg)

1.67 0.78 0.89 0.67 0.50 0.45 0.056 0.140 0.120 0.130 0 4.66 0.990 90 11.97

1.67 0.78 0.89 0.67 0.50 0.45 0.056 0.140 0.120 0.130 0 7.74 0.973 90 18.98

1.67 0.78 0.89 0.67 0.50 0.45 0.056 0.140 0.120 0.130 0 7.45 0.975 90 12.34

1.67 0.78 0.89 0.67 0.50 0.45 0.056 0.140 0.120 0.130 0 9.79 0.957 90 6.15

Os,2,90 ϕh,2/ϕv,2n (v/v) ao (nm) bp (nm) cq (nm) gr ϕ̅ 3s (deg) σφ3,0t (deg)

−0.437 78/22

−0.352 91/9 0.80 0.53 1.78 0.070

−0.433 60/40 0.80 0.53 1.78 0.075 0 13.03

−0.483 49/51 0.80 0.53 1.78 0.075 0 24.19

0.80 0.53 1.78 0.050 0 9.79

Os,3,0u ϕ̅ 3 (deg) σφ3,90 (deg)

0.957 90 6.11

90 12.88

0.926 90 12.07

0.775 90 17.11

Os,3,90 ϕv,3/ϕh,3v (v/v)

−0.483 81/19

−0.428 0/100

−0.436 29/71

−0.377 58/42

a

Long period of the molecular multibilayer PHIC structure. Thickness of more dense layer in the molecular multibilayer PHIC structure. cThickness of less dense layer in the molecular multibilayer PHIC structure. dMean interdistance between the n-hexyl bristles of the neighboring polymer chains whose repeat units are matched in position along their backbones. eMean interdistance between the nearest n-hexyl bristles along the polymer backbone. fMean interdistance between the n-hexyl bristles of the nearest neighboring polymer chains. gParacrystal distortion factor along the direction parallel to the long period of molecular multibilayer structure. h Paracrystal distortion factor along n-hexyl bristles of the neighboring polymer chains whose repeat units are matched in position along their backbones. iParacrystal distortion factor along the nearest n-hexyl bristles along the polymer backbone. jParacrystal distortion factor along the n-hexyl bristles of the nearest-neighboring polymer chains. k Mean value of the polar angle φ2 (i.e., orientation angle) between the orientation vector n2 (which is set along the direction parallel to the long period of molecular multibilayer PHIC structure) and the out-ofplane of the film. lStandard deviation for the orientation angle φ2 of molecular multibilayer PHIC structure. mSecond-order orientation factor of the molecular multibilayer PHIC structure. nVolume fraction ratio of the horizontal and vertical multibilayer PHIC structures. oA unit cell dimension along the a-axis of orthorhombic PCL crystals. pA unit cell dimension along the b-axis of orthorhombic PCL crystals. qA unit cell dimension along the c-axis of orthorhombic PCL crystals. r Paracrystal distortion factor of the PCL crystal. sMean value of the polar angle φ3 (i.e., orientation angle) between the orientation vector n3 (which is set along the direction parallel to the long period of lamellar structure) and the out-of-plane of the film. tStandard deviation for the orientation angle φ3 of lamellar structure. uSecondorder orientation factor of lamellar PCL crystal structure. vVolume fraction ratio of the horizontal and vertical lamellar PCL crystal structures. b

smaller than that in films composed of poly(ethylene terephthalate) or its copolymers.27 In comparison, poorly developed microstructure phases (a minor component) may form fringed micelle-like crystals and/or highly imperfect folded crystals in which the ordered regions assume a horizontally oriented orthorhombic lattice. Third, the PHIC and PCL arms underwent phase separation in the PHIC−PCLm films to form a vertical lamellar structure in which the PHIC and PCL layers were alternatively stacked along the in-plane direction of the film. Surprisingly, these lamellar structures were quite different from the cylindrical structures observed in common flexible diblock copolymers with the same or similar volume fractions. Lamellar structure formation may be driven mainly by the rigid chain properties and multibilayer structure formation characteristics of the PHIC arm and in part by the crystallization of the PCL arms and the resulting domain morphology. Fourth, lamellar structured morphologies that assume a preferential orientation in the PHIC−PCLm films are very interesting. These morphologies and orientations may result from the selective and specific interactions between the immiscible arm components and the substrate surface during the film-casting process as well as between the arm components and the solvent molecules during the postsolvent annealing. Surface energy analysis revealed that the surface energy (46.9 mN/m = Es,Si) of the Si substrate with a native oxide layer (Si/ SiOx) was closer to that (40.0 mN/m = Es,PCL)28 of the PCL arm component than that (27.0 mN/m = Es,PHIC)29 of the

polymer used in our study, which originated from the benzyl alcohol initiation reaction during polymerization, may have played a role in forming of vertical multibilayer structure as a minor domain component, in addition to the horizontal multibilayer structure. Second, GISAXS measurements obtained from the PCL homopolymer films indicated the presence of a horizontal lamellar structure only. In the horizontal lamellar structure, the orthorhombic crystal lattice (i.e., the c-axis of the crystal lattice) was oriented along the out-of-plane direction of the lamella (i.e., the out-of-plane direction of the film) due to the folded lamellar crystal formation characteristics of the flexible PCL chains. Surprisingly, the GIWAXS analysis, however, confirmed that horizontal orthorhombic crystals were present as a minor structural component in the PCL films, in addition to the vertical orthorhombic crystals. These scattering results collectively indicated that the PCL films were morphologically heterogeneous and consisted of well-developed microstructure phases, poorly developed microstructure phases, and amorphous phases. The well-developed microstructure phases, as the major component, were composed of horizontal lamellar structures in which the lamellar crystals assumed a vertically oriented orthorhombic lattice. The lamellar crystals were 3.1 nm thick, which corresponded only to the length of four repeat units in the fully extended conformation. The amorphous layer was 6.3 nm thick, about twice the thickness of the crystal layer. The interfacial layer thickness was comparable to the interfacial layer thickness observed in polyethylene thin films26 but much 7518


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Figure 8. Schematic representations of molecular chain conformation and packing order in toluene-annealed PHIC thin films (A): (B) horizontal multibilayer structure where n2 is the orientation vector of the structure model and φ2 is the polar angle between the n2 vector and the out-of-plane of the film; (C) top and side views of molecular chain in 21 conformation and packing order with no interdigitation in the structure (B); (D) relative electron density profile along the out-of-plane of the multibilayer structure; (E) vertical multibilayer structure; (F) side view of molecular chain conformation and packing order with no interdigitation in the structure (E).

surface energy of 27.9 mN/m,30 was found to be another good solvent for both the PHIC and PCL arm components. During the toluene annealing step, the PHIC and PCL arm components anchored onto the substrate surface became mobilized, underwent phase separation, and formed lamellar domains. Nonpolar toluene had a higher affinity for the n-hexyl side groups of the PHIC arm component than for the amide backbone. The favorable interactions between the solvent molecules and the n-hexyl side groups introduced mobility into the flexible alkyl side groups. The mobilized alkyl side groups then increased the mobility of the amide chain backbone, albeit to a limited degree because the polymer backbone displayed a high chain rigidity and the nonpolar toluene did not provide good solvent characteristics. The limited mobility in the polymer backbone, in turn, limited the mobility of the nhexyl side groups. The overall mobility of the PHIC chains induced by the toluene vapor was relatively lower. Under the circumstances, the alkyl side groups appeared to form a smectic A-like ordering, rather than either a well-defined lattice ordering or a fully disordered state. This ordering further induced the polar amide backbones, which had a limited mobility, to pack closely together via their electron donor and acceptor interactions. These interactions collectively led the PHIC molecules to form a multibilayer structure with a 21 chain conformation. The multibilayer structure built up along the out-of-plane direction of the film to form vertical lamellar domains. As a result, vertical lamellar PCL domains also developed. Fifth, the vertical lamellar structure formed in the PHIC− PCLm films revealed very interesting, unique features. All three of the miktoarm star polymers had comparable volume

PHIC arm component (Table 4). In our study, all PHIC− PCLm solutions used for the film formation processes were prepared in chloroform (which is a good solvent for both the PHIC and PCL arm components). The 0.5 wt % chloroform solutions of the PHIC and PCL homopolymers were prepared and used for contact angle measurements on the Si/SiOx substrates. The PHIC solution had a contact angle of 22.2° on the Si/SiOx substrate, relatively higher than that (13.5°) of the PCL solution (Table 5). These surface energy and contact angle results suggested the tendency of PCL arm component to anchor on or occupy (i.e., favor to interact with) the Si/SiOx substrate surface, unlike the PHIC arm component; however, the PCL arms were the minor component and, thus, could not cover the Si/SiOx substrate surface completely during the film formation process. The insufficient substrate surface coverage by the PCL arms gave the PHIC arm component (the major component) a chance to participate in the substrate surface coverage. Furthermore, the chloroform solvent used in the preparation of the polymer solutions had a surface energy of 26.7 mN/m (= Es,CHCl3),30 closer to that of the PHIC arm component than to the PCL arm component (Table 4). Thus, the chloroform solvent may promote the substrate surface coverage by the PHIC arm component during the film-casting process. The surface energy and contact angle measurements, as well as the volume fractions of the block components, suggested that both of the PHIC and PCL arm components could cover the Si/SiOx substrate surface during the filmcasting process. The substrate surface coverage by the PHIC and PCL arm components was likely to be proportional to their volume fractions in the miktoarm star polymer. Toluene, with a 7519


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Figure 9. Schematic representations of molecular chain conformation and packing order in toluene-annealed PCL thin films (A): (B) horizontal lamellar structure; (C) relative electron density profile along the out-of-plane of the lamellar structure; (D) molecular chain conformations in the amorphous layers and molecular chain order in the crystalline layers; (E) vertically oriented orthorhombic crystal lattice where n3 is the orientation vector of the structure model and φ3 is the polar angle between the n3 vector and the out-of-plane of the film; (F) fringed micelle-like crystals and/or highly imperfect folded crystals; (G) horizontally oriented orthorhombic crystal lattice.

The orientations of the orthorhombic crystals formed in the PCL lamella were significantly influenced by the number and length of the PCL arms (Table 3). The interfacial layer of the vertical lamellar structure in the PHIC−PCL film was only 0.7 nm thick. The interfacial layer thickness increased by 14% in both the PHIC−PCL2 and PHIC−PCL3 films. These interfacial layer thicknesses were smaller than those (ca. 1.1 nm) of the lamellar structures that formed in the thin films of conventional flexible diblock copolymers, for example poly(styrene-bisoprene) (PS-b-PI). They were also smaller than the lamellar crystal dimensions formed in the PCL homopolymer films and, furthermore, much smaller than the dimensions of the lamellar crystals formed by aromatic polyesters.27 Overall, the formation and ordering of the thin interfacial layers were remarkable, given the rigid PHIC arm component in the fully extended 21 conformation, the crystallization characteristics of the PCL arm component, and the spatially confined crowded geometry surrounding the arm jointer. The thin interfacial layers may have resulted from the self-assembly of the PHIC arm chains to form a well-ordered multibilayer structure that further built up a hierarchical structure, namely, vertical lamella. Sixth, the PHIC lamellae of the vertical lamellar structure formed in the PHIC−PCLm films were composed of horizontal and vertical multibilayer structure domains that resembled those observed in the PHIC homopolymer film; however, the orientation factor and volume fraction ratio varied, depending on the miktoarm star polymer system (Table 3). The orientation factor Os,2,0 of the horizontal multibilayer structure was 0.973 (ϕ̅ 2 = 0° and σφ2,0 = 7.74°) for the PHIC−PCL film,

fractions (71/29−73/27) of the PHIC and PCL arms (Table 1). Nevertheless, the orientation factor, positional distortion factor, and individual layer thicknesses of the vertical lamellar structures in the thin films varied, depending on the miktoarm star polymer system (Table 2). As the PCL arm number increased, the orientations of the vertical lamellar structures degraded slightly and the positional distortion increased somewhat. The vertical PHIC lamellae were thinner by 5.7% in the PHIC−PCL2 film and by 14.3% in the PHIC−PCL3 film with respect to the thickness of the PHIC−PCL film. In fact, the molecular weight Mn,NMR of the PHIC arm was reduced by only 0.9% in both the PHIC−PCL2 and PHIC−PCL3 films, relative to the molecular weight of the PHIC arm in the PHIC−PCL (Table 1). These Mn,NMR results suggested that the thickness values of the vertical PHIC lamellae may reflect the morphological changes that resulted from an increase in the PCL arm number. The morphology of the PHIC lamellae was influenced by the variations in the PCL arm number. The orientations of the multibilayer structure formed in the PHIC lamella were significantly influenced by the increase in the PCL arm number (Table 3). On the other hand, the vertical PCL lamellae were thinner by 11.3% in the PHIC−PCL2 film and by 12.2% in the PHIC−PCL3 film, relative to the thickness of the PHIC−PCL film (Table 2). Considering that the total Mn,NMR of PCL arms was reduced by 3.8% in the PHIC−PCL2 and by 9.4% in the PHIC−PCL3, with respect to the values of the PCL arm in PHIC−PCL, the thickness variations in the vertical PCL lamellae may have resulted in part from variations in the total molecular weight and further from the morphological changes due to variations in the number and length of the PCL arm. 7520


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Figure 10. Schematic representations of molecular chain conformations and packing orders in toluene-annealed PHIC−PCLm thin films: (A) vertical lamellar structure; (B) relative electron density profile along the out-of-plane of the lamellar structure; (C) horizontal multibilayer structure formed in the PHIC lamella where n2 is the orientation vector of the structure model and φ2 is the polar angle between the n2 vector and the out-of-plane of the film; (D) vertical multibilayer structure formed in the PHIC lamella; (E) top and side views of molecular chain in 21 conformation and packing order with no interdigitation in the structure (C); (F) relative electron density profile along the out-of-plane of the multibilayer structure; (G) side view of molecular chain conformation and packing order with no interdigitation in the structure (D); (H) fringed micelle-like crystals and/or highly imperfect folded crystals formed in the PCL lamella; (I) vertically oriented orthorhombic crystal lattice where n3 is the orientation vector of the structure model and φ3 is the polar angle between the n3 vector and the out-of-plane of the film; (J) horizontally oriented orthorhombic crystal lattice.

for the PHIC−PCL film, remarkably higher than that (78%) of the PHIC film; however the ϕh,2 value was again dramatically reduced to 60% in the PHIC−PCL2 film and 49% in the PHIC−PCL3 film. On the other hand, the orientations of the vertical multibilayer structure were significantly degraded in the PHIC−PCL film (Os,2,90 = −0.352; ϕ̅ 2 = 90° and σφ2,90 = 18.98°), and the volume fraction was dramatically reduced (only 9% ϕv,2) compared to the values (Os,2,90 = −0.437 (ϕ̅ 2 = 90° and σφ2,90 = 11.97°) and ϕv,2 = 22%) obtained in the PHIC film. The orientation of the vertical multibilayer structure improved as the volume fraction in the PHIC−PCL2 film increased (Os,2,90 = −0.433 (ϕ̅ 2 = 90° and σφ2,90 = 12.34°) and ϕv,2 = 40%). Surprisingly, the orientation improved significantly at a much higher volume fraction in the PHIC−PCL3 film (Os,2,90 = −0.483 (ϕ̅ 2 = 90° and σφ2,90 = 6.15°) and ϕv,2 = 51%). Overall, the orientations and volume fractions of the horizontal and vertical multibilayer structures formed in the vertical lamellar PHIC phases were significantly affected by variations in the number and length of PCL arms in the vertical lamellar PCL phase that interfaced with the lamellar PHIC phase. The orientations and volume fractions may be further influenced by the crystallization process and the resulting morphology of the PCL arms in the vertical lamellar PCL phase. Finally, the PCL phase domains in the vertical lamellar structures formed in the PHIC−PCLm films displayed very interesting morphologies that differed significantly from those observed in the PCL homopolymer film. The PCL phase

Table 4. Surface Energies of the Substrates, Polymers, and Solvents Used in This Study substance Si/SiOx substrate PHIC film PCL film a b

surface energy (mN/m)

substance

surface energy (mN/m)

46.9 (2.3)a

chloroform

26.7d

27.0b 40.0c

toluene

27.9d

The numbers in parentheses represent standard deviations. Reference 29. cReference 28. dReference 30.

Table 5. Contact Angles of the PHIC and PCL Homopolymer Solutions (0.5 wt %) in Chloroform onto Si Substrates with an Oxide Layer homopolymer solution (0.5 wt %)

contact angle (deg) on Si/SiOx substrate

PHIC PCL

22.2 (1.7)a 13.5 (1.2)

a

The numbers in parentheses represent standard deviations.

0.975 (ϕ̅ 2 = 0° and σφ2,0 = 7.45°) for the PHIC−PCL2 film, and 0.957 (ϕ̅ 2 = 0° and σφ2,0 = 9.79) for the PHIC−PCL3 film. Overall, the orientation of the horizontal multibilayer structure degraded to some extent upon incorporation of the PCL arms relative to the structure (Os,2,0 = 0.990; ϕ̅ 2 = 0° and σφ2,0 = 4.66°) formed in the PHIC homopolymer film. The volume fraction ϕh,2 of the horizontal multibilayer structure was 91% 7521


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domains in the PHIC−PCLm films did not yield scattering signals indicative of a lamellar structure (Figures 3b, 4b, and 5b). By contrast, the PCL homopolymer film exhibited scattering signals that were typical of a lamellar structure (Figure 2d). Instead, the PCL phase domains in the PHIC− PCLm films revealed very weak broad scattering peaks only along the in-plane direction of the vertical lamellar PCL phase (Figures 3e, 4e, and 5e). These peaks had d-spacings of 4.0− 10.1 nm (Table 2), which were much smaller than or comparable to the long period L (11.5 nm) of the lamellar crystals formed in the PCL homopolymer film. Furthermore, the PCL phase domains consisted of orthorhombic crystals and amorphous phases, as observed in the PCL homopolymer film. The orthorhombic crystals were significantly influenced in orientation and population by the number and length of PCL arms (Tables 1 and 3). The overall crystallinity Xc of the PCL phase decreased as the PCL arm number increased and as the PCL arm length decreased (Table 1). Surprisingly, only the horizontal orthorhombic PCL crystals formed in the PHIC− PCL film, which differed significantly from the PCL homopolymer film, which consisted of orientationally mixed orthorhombic crystals; nevertheless, the orientational ordering was not better than that of the horizontal orthorhombic crystals in the PCL homopolymer film (Table 3). The PCL lamellar phases in the other miktoarm polymer films formed orientationally mixed orthorhombic crystals, namely, horizontal and vertical crystals, as observed in the PCL homopolymer film. Their orientations and volume fractions varied dramatically in the PHIC−PCL2 and PHIC−PCL3 films with respect to those of the PCL homopolymer film. Overall, the PCL phase domains of the vertical lamellar-structured PHIC−PCLm films were formed from fringed micelle-like crystals and/or imperfectly folded crystals, rather than from typical lamellar crystals. The crystals had an orthorhombic lattice and were limited in size as a result of the geometrical confinement due to the width (6.0−7.0 nm, depending on the arm number and length) of the vertical PCL lamellar domains. The orthorhombic crystals may have developed an intercrystal distance of 4.0−10.1 nm along the in- plane direction of the vertical PCL lamella. Moreover, the orthorhombic crystals formed with a horizontal orientation or a mixture of horizontal and vertical orientations. The degree of orientation and volume fraction of the horizontal and vertical orthorhombic crystals that formed were significantly influenced by the number and length of PCL arm as well as by the structural ordering in the PHIC lamellae that interfaced with the PCL lamellar phase. The formation of horizontal crystals was more favorable in the PCL phase comprising fewer and longer arms. In particular, the PCL phase of the PHIC−PCL3 polymer film, which was formed from polymers having the shortest and most numerous arms, formed almost equal volume fractions of horizontal and vertical crystals.

respect to the perfect PCL crystals. The PCL homopolymer displayed Xc = 68.8%. The miktoarm polymers displayed very interesting and very complex thin film morphologies. They successfully underwent phase separation during the solution casting, drying, and toluene annealing processes, revealing vertical lamellar structure formation in an orientationally controlled manner. Additional crystal formation was observed in the phase-separated lamellar domains. Surprisingly, the vertical lamellar structures were quite different from the cylinder structures observed in common flexible diblock copolymers of comparable volume fractions. Lamellar structure formation may have been driven by the rigid chain properties and the multibilayer structure formation characteristics of the PHIC arm and in part by the crystallization and resulting morphology of the PCL arms. Furthermore, the vertical orientations of the lamellar structures may have resulted from the competitive anchoring of the immiscible arm components onto the substrate surface during the film-casting process and their phase growth via phase separation during the postsolvent annealing process. At the same time, the arm components in the phase-separated domains underwent self-assembly (ordering). The PHIC arm phase formed a mixture of horizontal and vertical multibilayer structure domains, as observed in the PHIC homopolymer film, whereas the PCL arm phase formed fringed micelle-like crystals and/or highly imperfect folded crystals. The PCL arm phase structure differed significantly from that of the PCL homopolymer film, which formed typical folded lamellar crystals. The formation of an imperfect structure was attributed to the geometrical confinement caused by the narrow vertical PCL lamella and the spatially confined crowded arm jointer. Geometrical confinement further limited the size of the imperfect structure. The crystals assumed an orthorhombic lattice. The orthorhombic crystals formed with a horizontal orientation or a mixture of horizontal and vertical orientations. Overall, the parameters describing the hierarchical structures that formed in the thin films were significantly influenced by the presence of the PHIC and PCL arms as well as by the number and length of the PCL arms. In summary, synchrotron GIXS measurements and quantitative data analysis revealed detailed information about the hierarchical structures that developed in the nanoscale thin films formed by the PHIC− PCLm star polymers.

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ASSOCIATED CONTENT

* Supporting Information S

Synthesis and GIXS data analysis. This material is available free of charge via the Internet at http://pubs.acs.org.

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CONCLUSIONS The thermal properties and nanoscale thin film morphologies of PHIC−PCLm star polymers (m = 1−3; Mn,NMR(PHIC) = 11 300−11 400 and Mn,NMR(PCLm) = 4800−5300) and their homopolymers were investigated using thermal analysis and synchrotron GIXS analysis. The miktoarm polymers were stable up to 180 °C. By comparison, the PHIC and PCL homopolymers were stable up to 180 and 290 °C, respectively. Among the samples cooled at a rate of 10.0 °C/min, the crystallinities Xc of the PCL phases were 65.4% for PHIC− PCL, 63.8% for PHIC−PCL2, and 54.4% for PHIC−PCL3, with

AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected], Tel +82-54-279-2120, Fax +82-54279-3399 (M.R.). *E-mail [email protected], Tel +81-11-7066606, Fax +81-11-706-6603 (T.S.). *E-mail [email protected], Tel +81-11-7066602, Fax +81-11-706-6602 (T.K.). Notes

The authors declare no competing financial interest. 7522


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ACKNOWLEDGMENTS This study was supported by the National Research Foundation (NRF) of Korea (Doyak Program 2011-0028678 and Center for Electro-Photo Behaviors in Advanced Molecular Systems (2010-0001784)) and the Ministry of Science, ICT & Future Planning (MSIP) and the Ministry of Education (BK21 Plus Program and Global Excel Program). This work was also supported by the Japan Society for the Promotion of Science (JSPS) Grant-in-Aid for Scientific Research (B): No. 25288093. The synchrotron X-ray scattering measurements at the Pohang Accelerator Laboratory were supported by MSIP, POSTECH Foundation, and POSCO Company.

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Hierarchical Structures in Thin Films of Miktoarm Star Polymers: Poly(n

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