nanostructures: surface-induced phase separation of polystyrene poly(vinyl alcohol ... · 2019. 7....

polymers

Article

Replication of Mesoscale Pore One-dimensionalNanostructures: Surface-induced Phase Separation ofPolystyrene/Poly(vinyl alcohol) (PS/PVA) Blends

Paritat Muanchan 1 , Takashi Kurose 2 and Hiroshi Ito 1,2,*1 Graduate School of Organic Materials Science, Yamagata University, 4-3-16 Jonan, Yonezawa,

Yamagata 992-8510, Japan; [email protected] Research Center for GREEN Materials and Advanced Processing (GMAP), Yamagata University, 4-3-16

Jonan, Yonezawa, Yamagata 992-8510, Japan; [email protected]* Correspondence: [email protected]

Received: 26 February 2019; Accepted: 7 June 2019; Published: 12 June 2019��

Abstract: Mesoscale pore one–dimensional (1D) nanostructures, or vertically aligned porousnanostructures (VAPNs), have attracted attention with their excellent hydrophobic properties,ultra−high surface area, and high friction coefficient, compared to conventional vertically alignednanostructures (VANs). In this study, we investigate the replication of VAPNs produced by thethermal nanoimprint process using anodic aluminum oxide (AAO2) templates (100 nm diameter).Polystyrene/poly(vinyl alcohol) (PS1/PVA) blends, prepared by the advanced melt–mixing processwith an ultra–high shear rate, are used to investigate the formation of porosity at the nanometer scale.The results reveal that domain size and mass ratios of PVA precursors in the PS matrix play a dominantrole in the interfacial interaction behavior between PS1–PVA–AAO2, on the obtained morphologiesof the imprinted nanostructures. With a PVA nanodomain precursor (PS1/PVA 90/10 wt%), theintegration of PVA nanodroplets on the AAO2 wall due to the hydrogen bonding that induces thephase separation between PS1–PVA results in the formation of VAPNs after removal of the PVAsegment. However, in the case of PVA microdomain precursors (PS1/PVA 70/30 wt%), the structuretransformation behavior of PS1 is induced by the Rayleigh instability between PVA encapsulatedaround the PS1 surfaces, resulting in the PS1 nanocolumns transforming into nanopeapods composedof nanorods and nanospheres.

Keywords: Porous Nanostructures; Phase Separation; Thermal Nanoimprint

1. Introduction

One-dimensional nanostructures inspired by the nature of gecko feet have attracted considerableattention because of their versatile applications such as self-cleaning surfaces, dry adhesives, and fornovel applications of thermally and electrically conductive materials [1–5]. Moreover, nanomaterialshave recently appeared based on structure-mediated surface functionalities produced by polymernanoengineering technology [6]. For example, the hydrophobic surface of the cicada wing is coveredwith natural vertically aligned nanostructures (VANs) which can penetrate and consequently killPseudomonas aeruginosa within several minutes of adhesion [7,8]. To enhance VANs bacteria-killingperformance, the characteristics and surface properties of VANs need improvement, including their (1)hydrophobic properties, (2) surface areas, and (3) bacteria adherence (ex. surface friction and surfacemodulus). For this reason, vertically aligned porous nanostructures (VAPNs) have attracted significantattention, with outstanding hydrophobic higher friction properties and ultra-high surface area, becauseof the nanoscale porosity and surface roughness. Particularly, VAPNs have been used as switchable

Polymers 2019, 11, 1039; doi:10.3390/polym11061039 www.mdpi.com/journal/polymers

http://www.mdpi.com/journal/polymershttp://www.mdpi.comhttps://orcid.org/0000-0001-7859-9146https://orcid.org/0000-0001-8432-8457http://dx.doi.org/10.3390/polym11061039http://www.mdpi.com/journal/polymershttps://www.mdpi.com/2073-4360/11/6/1039?type=check_update&version=2

Polymers 2019, 11, 1039 2 of 18

adhesive surfaces using humidity control [9,10]. Moreover, VAPNs with high value-added have beeneffectively applied in the fields of energy storage and nanosensing materials [5], and filters [9]. For thesereasons, VAPNs are regarded as a novel class of frontier nanomaterials in both industry and academia.

A versatility technique for fabricating VAPNs is the templating method using nanoporous anodicaluminum oxide (AAO) assisted by solvent-wetting and thermal melt-wetting methods. The controlof surface-induced phase separation behaviors in the AAO pores is crucial for these techniques. Agreat deal of research has attempted to improve these methods to be faster, cheaper, more precise,less toxic, and more scalable, which are in conflict with the fundamentals of the solvent-wetting andthermal melt-wetting methods [11–14]. With the successful development of polymer engineeringtechnologies, the direct patterning method of thermal nanoimprint is an alternative method to replicatethe VAPNs because of its high replication capability with rapid speed, zero−solvent use, high precisionand reproducibility. For example, one-dimensional polymer nanofiber arrays with various lengths,diameters, aspect ratios, and patterns, were successfully produced by nanoimprint machine usingAAO templates [15]. However, the direct fabrication of VAPNs using the nanoimprint method withAAO templates has not been achieved yet. Therefore, it is a great challenge to develop the thermalnanoimprint process by controlling the phase separation of polymers in the AAO templates.

Study on the various polymer material behaviors at the nanoscale of AAO templates based on thesolvent-wetting and thermal melt-wetting methods have been successfully clarified. By using thesewetting methods, many publications have reported that polymer behaviors in AAO templates, consistof the flow ability and its gradients, phase separation behaviors, transformation behaviors, and so on.Atsushi Takahara et al. [16–18] studied flowing gradients by using various polymer blends in AAOtemplates. With the larger pore size of AAO, it promoted phase separation of polymer blends. In thecase of polymer blends with the same value of surface energy, polymer flow gradients were stronglydependent on the glass transition temperature (Tg) and polymer melt viscosity which are related tothe polymer chain mobility. At a constant melt-wetting temperature, the obtained nanostructuresshowed the higher mass ratios of the lower Tg segment due to the thermal-induced polymer chainmobility in the AAO templates. Jiun Tai Chen et al. [19] studied the coarsening mechanism ofpoly(methyl methacrylate) (PMMA) and tetrahydrofuran (THF) confined in AAO templates. In thatcase, the phase separation of PMMA and THF was induced by the difference in hydrogen bond affinitybetween PMMA-AAO and THF-AAO. With the stronger hydrogen bond of THF-AAO, the formationof THF droplets on the AAO pore wall was promoted by surface-induced phase separation, resultingin porous PMMA nanostructures after the removal of THF. The formation of porous nanotubes ofpolystyrene (PS) and fluorescent pyrene-ended PMMA (Py-PMMA) using the solvent-wetting methodwas also investigated by Jiun Tai Chen et al. [20]. Moreover, Thomas P. Russell et al. [21] havereported on the Rayleigh instability of PMMA/PS blends confined in AAOs using the solvent-wettingmethod assisted by thermal annealing. That method can produce PMMA/PS bilayer nanotubes. Thethermally-induced Rayleigh instability stimulated the transformation of nanostructures with polymernanotubes into nanorods and nanospheres. Various transformation and phase separation behaviors ofpolymer nanomaterials in AAO templates and other atmospheres using chemical solvent and thermalannealing methods have been investigated [22–33]. Furthermore, similar polymer phase separationsand transformation behaviors at the microscale have also been observed [34–38]. Because of the successof the AAO templating method, many publications have explored various AAO applications suchas thermal and physical properties of polymer nanostructures under nanoconfinement conditions,fabrication of higher−order nanostructures, polymerization and crystallization in AAO templates, andthe double infiltration nanoencapsulation process [39–49].

In this study, we mainly focus on the formation of VAPNs fabricated using thermal nanoimprintwith AAO templates. We investigate the phase separation and the transformation behaviors of polymerblends confined in AAOs. The material models used are polystyrene/poly(vinyl alcohol) (PS/PVA)blends. PS/PVA blends are prepared by the advanced melt-mixing process, using a high-shear machinewhich effectively produces the ultra-fine dispersion of PVA segments. Several surface properties

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including surface morphology, hydrophobicity, frictional properties, and the surface mechanicalproperties of the fabricated VAPNs have been experimentally evaluated. This novel technique has theadvantages of being a greener (non-chemical use), rapid process, of high−precision, high−profit, andgreat reproducibility.

2. Materials and Methods

2.1. Materials

Commercial grade PS of two types was used: PS1 (GPPS 679; PS Japan Corp., Tokyo, Japan, Tg =87 ◦C, melt flow rate (MFR) = 18.0 g/10 min at 200 ◦C) and PS2 (G210; Toyo Styrene Co., Ltd., Tokyo,Japan, Tg = 100 ◦C, MFR = 10.1 g/10 min at 200 ◦C). PVA (CP-1210; Kuraray Co., Ltd., Tokyo, Japan, Tg= 26 ◦C, melting temperature (Tm) = 161 ◦C, MFR = 4.4 g/10 min at 200 ◦C) was selectively used.

2.2. Preparation of AAOs

High-purity aluminum (Al) sheets (99.99% pure) 5.0 cm × 5.0 cm of 0.5 mm thickness weredegreased in acetone, followed by electropolishing in a mixed solution of HClO4/H2O/EtOH(10.0/7.0/83.0 wt%) at a constant temperature of 5 ◦C, with 20 V applied for 10 min. The polishedaluminum sheets were anodized using a two-step anodization process: The first anodization was donewith the conditions shown in Table 1, then the porous alumina was dissolved in a solution containing6.0 wt% of H3PO4 and 1.8 wt% of H2CrO4 at 60 ◦C for 3 h. The second anodization was appliedagain with the anodizing conditions shown in Table 1. The remaining aluminum film was removedby dissolving in copper chloride solution. Pore expansion was done by immersion in 8.5 wt% of theH3PO4 solution at 40 ◦C. For this study, AAO templates with different two pores size were prepared:AAO1 and AAO2.

Table 1. Anodizing conditions. Anodic aluminum oxide (AAO).

AnodizationAAO1 AAO2

1st 2nd 1st 2nd

Electrical Current (A) 12.1 12.1 12.1 12.1Applied Voltage (V) 40 40 60 60

Oxalic Acid (M) 0.2 0.2 0.3 0.3Time (h) 17 24 17 48

Temperature (◦C) 10 10 5 5

2.3. Preparation of Polymer Blends

The PS1/PVA pellets with mass ratios of 90/10 and 70/30 wt% were introduced into a high-shearmachine (NHSS2-28; Niigata Machine Techno Co., Ltd., Niigata, Japan) for the blending process witha rotation speed of 500 rpm at 180 ◦C for 10 s. The obtained PS1/PVA blends were pelletized with acrushing machine.

2.4. Preparation of Polymer Films

The polymer pellets (neat PS1, neat PS2, and PS1/PVA blends) were melted in a template witha thickness of 500 µm at 200 ◦C for 5 min, followed by pressing 10 MPa in a hot press machine(IMC-11FA; Imoto Machinery Co. Ltd., Tokyo, Japan). The molten film was cooled in a cold pressmachine (IMC-181B; Imoto Machinery Co., Ltd., Tokyo, Japan) at 5 MPa for 3 min.

2.5. Fabrication of Nanostructures by Thermal Nanoimprint

In this experiment, one–dimensional nanostructures were fabricated with a thermal nanoimprintmachine (Izumi Tech., Miyagi, Japan) using AAO templates. The experimental procedure is shownin Figure 1. The polymer films were placed onto the AAO template and inserted into the thermal

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nanoimprint machine. The imprinting conditions are presented in Table 2. After imprinting, the AAOwas removed by dissolution in 4.0 M of NaOH solution for 24 h, followed by water washing at 60 ◦Cand then the drying process.

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imprinting, the AAO was removed by dissolution in 4.0 M of NaOH solution for 24 h, followed by 133 water washing at 60 °C and then the drying process. 134

135

Figure 1. Fabrication of one-dimensional nanostructures by thermal nanoimprint. 136

Table 2. Imprinting conditions. 137

Materials Temperature (oC) Pressure (MPa) Press Time (min) Obtained Nanostructures

PS1 160−220 5 30 Nanopillars

PS2 160−220 5 30 Nanopillars

PS1/PVA (90/10 wt%)

120−160 5 30 Nanopillars

PS1/PVA (90/10 wt%)

180 5 15–30

60–120 Porous Nanopillars

Nanofragments

PS1/PVA (70/30 wt%)

120−180 5 15−120 Nanopeapods

2.6. Characterization 138 Surface morphologies of AAO templates and polymer nanostructures were observed using 139

scanning electron microscopy (SEM, JSM-6510; Jeol Ltd., Tokyo, Japan and TM-1000; Hitachi High-140 Technologies Corp., Tokyo, Japan) and field emission scanning electron microscopy (FE-SEM, SU-141 8000; Hitachi High-Technologies Corp., Tokyo, Japan). The surface chemical property of the polymer 142 blends films were analyzed by attenuated total reflection (ATR, AIM-8800 Infrared Microscope; 143 Shimadzu Corp., Kyoto, Japan). The wavenumber range of 3,200–3,600 cm-1 was selected for 144 investigation due to differences in the chemical structures between PS and PVA (the presence of a 145 hydroxyl group). Thermal transition behavior of polymers were evaluated using modulated 146 differential scanning calorimetry (M-DSC, DSC Q200; TA Instruments Inc., Delaware, USA). A 147 contact angle meter (DM 500; Kyowa Interface Science Co., Ltd., Saitama, Japan) was used to measure 148 the surface interactions between molten polymers and AAO surfaces, and between water droplets 149 and polymer surfaces. The friction coefficient of polymer surfaces at the macroscale was evaluated 150 by a friction and wear tester (EFM-III-F; Orientec Co., Ltd., Tokyo, Japan) using stainless steel 151 materials. The friction test conditions were evaluated using an applied force of between 21–40 N with 152 the rotational speed of 10 rpm for 30 s. All of the tests were performed at ambient conditions (25 °C, 153 relative humidity 35%). The surface mechanical properties at the nanoscale were evaluated by 154 nanoindenter (G200; Agilent Technologies Inc.). 155

Figure 1. Fabrication of one-dimensional nanostructures by thermal nanoimprint.

Table 2. Imprinting conditions.

Materials Temperature (◦C) Pressure (MPa) Press Time (min) ObtainedNanostructures

PS1 160−220 5 30 NanopillarsPS2 160−220 5 30 Nanopillars

PS1/PVA(90/10 wt%)

120−160 5 30 Nanopillars

PS1/PVA(90/10 wt%)

180 5 15–3060–120Porous Nanopillars

NanofragmentsPS1/PVA

(70/30 wt%)120−180 5 15−120 Nanopeapods

2.6. Characterization

Surface morphologies of AAO templates and polymer nanostructures were observed usingscanning electron microscopy (SEM, JSM-6510; Jeol Ltd., Tokyo, Japan and TM-1000; HitachiHigh-Technologies Corp., Tokyo, Japan) and field emission scanning electron microscopy (FE-SEM,SU-8000; Hitachi High-Technologies Corp., Tokyo, Japan). The surface chemical property of thepolymer blends films were analyzed by attenuated total reflection (ATR, AIM-8800 Infrared Microscope;Shimadzu Corp., Kyoto, Japan). The wavenumber range of 3200–3600 cm−1 was selected forinvestigation due to differences in the chemical structures between PS and PVA (the presence ofa hydroxyl group). Thermal transition behavior of polymers were evaluated using modulateddifferential scanning calorimetry (M-DSC, DSC Q200; TA Instruments Inc., New Castle, DE, USA). Acontact angle meter (DM 500; Kyowa Interface Science Co., Ltd., Saitama, Japan) was used to measurethe surface interactions between molten polymers and AAO surfaces, and between water droplets andpolymer surfaces. The friction coefficient of polymer surfaces at the macroscale was evaluated by afriction and wear tester (EFM-III-F; Orientec Co., Ltd., Tokyo, Japan) using stainless steel materials. Thefriction test conditions were evaluated using an applied force of between 21–40 N with the rotationalspeed of 10 rpm for 30 s. All of the tests were performed at ambient conditions (25 ◦C, relative humidity35%). The surface mechanical properties at the nanoscale were evaluated by nanoindenter (G200;Agilent Technologies Inc., Santa Clara, CA, USA).

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3. Results and Discussion

3.1. Characterization of AAOs

Figure 2 shows SEM images of the surface morphologies of fabricated AAO templates. TheAAO templates consist of dense arrays of cylindrical nanopores displaying abundant honeycomb–likeporous cavities caused by the formation of highly ordered hexagonal closely packed morphology,formed by the two–step anodization process. The narrow pore size of the AAO templates can befabricated under our experimental conditions. Figure 2a shows the prepared AAO1 template with a50 nm diameter pore, 25 nm inter–pore distance, and 120 µm depth. For the AAO2, Figure 2b shows atemplate with 100 nm pore size, 50 nm inter–pore distance, and 130 µm depth. In this experiment, itis noted that increased pore size is mainly influenced by increased applied electrical potential. Theslight increase in pore depth occurs because of the increased anodizing time and concentration of theelectrolytic solutions. In this study, the AAO2 templates were selectively used to examine the phaseseparation behaviors and structural transformation behaviors of polymer nanostructures. Figure 2cshows a cross–sectional side view of AAO2 nanocavities of 100 nm pore diameter. The SEM micrographdepicted in Figure 2c confirms that an AAO2 having a smooth surface can be prepared in this study.The pore wall of the AAO2 shows an almost perfectly smooth surface and corresponds to the expecteddimensions and invariance of the pore diameter through the length of AAO2 templates.


3. Results and Discussion 156

3.1. Characterization of AAOs 157 Figure 2 shows SEM images of the surface morphologies of fabricated AAO templates. The AAO 158

templates consist of dense arrays of cylindrical nanopores displaying abundant honeycomb–like 159 porous cavities caused by the formation of highly ordered hexagonal closely packed morphology, 160 formed by the two–step anodization process. The narrow pore size of the AAO templates can be 161 fabricated under our experimental conditions. Figure 2a shows the prepared AAO1 template with a 162 50 nm diameter pore, 25 nm inter–pore distance, and 120 µm depth. For the AAO2, Figure 2b shows 163 a template with 100 nm pore size, 50 nm inter–pore distance, and 130 µm depth. In this experiment, 164 it is noted that increased pore size is mainly influenced by increased applied electrical potential. The 165 slight increase in pore depth occurs because of the increased anodizing time and concentration of the 166 electrolytic solutions. In this study, the AAO2 templates were selectively used to examine the phase 167 separation behaviors and structural transformation behaviors of polymer nanostructures. Figure 2c 168 shows a cross–sectional side view of AAO2 nanocavities of 100 nm pore diameter. The SEM 169 micrograph depicted in Figure 2c confirms that an AAO2 having a smooth surface can be prepared 170 in this study. The pore wall of the AAO2 shows an almost perfectly smooth surface and corresponds 171 to the expected dimensions and invariance of the pore diameter through the length of AAO2 172 templates. 173

174 Figure 2. Scanning electron microscopy (SEM) images of surface morphologies of (a) AAO1 175 (b) AAO2 and (c) cross−sectional AAO2. 176

3.2. Fabrication of One–dimensional Nanostructures by Thermal Nanoimprint 177 In this part, the flowability and confinement effects of PS1 and PS2 at the nanoscale will be 178

clarified. Figure 3 shows representative SEM images (side view) of PS1–VANs and PS2–VANs 179 obtained by thermal nanoimprint using the AAO1 template. For the each of the SEM images, PS1–180 VANs and PS2–VANs were produced using the imprinting conditions of 180 °C, 5 MPa, for 30 min. 181 The crucial point for performing the thermal nanoimprint is that the imprinting temperatures must 182 be above Tg or Tm of the polymers which is sufficient for chain mobility and adequate for flowability 183 at the nanoscale in AAO cavities. In Figure 3, PS1–VANs and PS2–VANs with lengths of 70 and 40 µm 184 respectively can be produced. The obtained VANs, having bundle–like structure, exhibit orderly 185 array structures but the tendency to agglomerate. The appearance of agglomerated nanostructures is 186 explainable by surface adhesion induced between individual nanostructures. Generally, the quality 187 of agglomerated nanostructures was enhanced by using lower aspect ratios and higher interpore 188 distances in AAO templates. 189

As Figure 3, the difference in the length of VANs due to the different flowability of the PS 190 precursors in the range of 160–220 °C is shown, with fixed imprinting pressure and imprinting time 191 of 5.0 MPa and 30 min respectively. The longer length of PS1–VANs is compared with PS2–VANs, 192 due to the higher MFR and lower Tg of PS1. The lower value of Tg indicated higher chain mobility of 193 polymer due to larger free–volume of the polymer chain packing. The PS1 showed the higher 194 flowability within AAO1 templates as compared with PS2. The flowability of the polymers was 195 enhanced by increased imprinting temperature due to the reduction of polymer viscosities, and 196

Figure 2. Scanning electron microscopy (SEM) images of surface morphologies of (a) AAO1 (b) AAO2

and (c) cross−sectional AAO2.

3.2. Fabrication of One–dimensional Nanostructures by Thermal Nanoimprint

In this part, the flowability and confinement effects of PS1 and PS2 at the nanoscale will be clarified.Figure 3 shows representative SEM images (side view) of PS1–VANs and PS2–VANs obtained bythermal nanoimprint using the AAO1 template. For the each of the SEM images, PS1–VANs andPS2–VANs were produced using the imprinting conditions of 180 ◦C, 5 MPa, for 30 min. The crucialpoint for performing the thermal nanoimprint is that the imprinting temperatures must be aboveTg or Tm of the polymers which is sufficient for chain mobility and adequate for flowability at thenanoscale in AAO cavities. In Figure 3, PS1–VANs and PS2–VANs with lengths of 70 and 40 µmrespectively can be produced. The obtained VANs, having bundle–like structure, exhibit orderlyarray structures but the tendency to agglomerate. The appearance of agglomerated nanostructures isexplainable by surface adhesion induced between individual nanostructures. Generally, the qualityof agglomerated nanostructures was enhanced by using lower aspect ratios and higher interporedistances in AAO templates.

Polymers 2019, 11, 1039 6 of 18


decreased surface tensions between polymer melts and AAO templates [11–15]. However, the length 197 of VANs did not increase at temperatures above 180 °C because the polymer was easily ejected out 198 from the AAO nanocavities under the high imprinting pressure and relatively low viscosity. PS1–199 VANs and PS2–VANs of about 15–72 µm long were obtained depending upon the imprinting 200 conditions. 201

202 Figure 3. Effect of imprinting temperature on length of PS–VANs (SEM images refer to the 203 solid symbols). 204

The flowability at the nanoscale of the molten polymers in the AAO templates driven by the 205 imprinting process is in agreement with the Hagen–Poiseuille expression [15,50]. The length (H ) or 206 distance the polymer flowed through AAO can be estimated as shown below. 207

2 cos

232

P trH r

γ θ

η

+ =

(1)

This equation shows that the factors related to flowability in confined spaces are imprinting 208 pressure (P ), radius of the nanochannel ( ), surface tension of the molten polymer ( ), contact 209 angle of the polymer on AAO pore wall ( ), viscosity of the molten polymer ( ), and infiltration 210 time ( ). Our previous study revealed that the reduction of apparent viscosity estimated by Equation 211 (1) may be the result of shear rate, wall slip, and the pressure used in the imprinting process [15]. 212

To avoid the effects of confinement on the phase separation and transformation behaviors at the 213 nanoscale, therefore, we investigated this change in molecular free–volume by analyzing the thermal 214 properties of obtained PS–VANs as compared with PS bulk films. Specimens were obtained by 215 imprinting conditions of 180 °C, 5 MPa, for 30 min with different AAO pore size and Tg of polymers. 216 Figure 4 shows the DSC heating thermograms of the PS bulk films and PS–VANs. The results indicate 217 a change in molecular free–volume was found in the case of PS2–VANs when using AAO1 templates. 218 In the case of PS2, an increase in Tg of the PS2–VANs to 104 °C is found deviating from the PS2 bulk 219 film precursors having Tg of 100 °C. The increase in Tg of PS2 after the imprinting process using the 220 AAO1 template due to the nanoconfinement, influences the reduction of the free−volume or the 221 increase in polymer chain packing density caused by the compression, and shortens each polymer 222 chain distance. It also might be caused by the applied force from the imprinting pressure induced the 223 reduction of free–volume polymer in the AAO nanopore. Moreover, the polymer chains at the 224 interface have a strong interfacial interaction with the AAO pore walls that can induce the orientation 225 change of polymer chains [51,52] and also immobilize the polymer chain at the interface, especially 226

r γθ η

t

Figure 3. Effect of imprinting temperature on length of PS–VANs (SEM images refer to the solidsymbols).

As Figure 3, the difference in the length of VANs due to the different flowability of the PSprecursors in the range of 160–220 ◦C is shown, with fixed imprinting pressure and imprinting time of5.0 MPa and 30 min respectively. The longer length of PS1–VANs is compared with PS2–VANs, due tothe higher MFR and lower Tg of PS1. The lower value of Tg indicated higher chain mobility of polymerdue to larger free–volume of the polymer chain packing. The PS1 showed the higher flowability withinAAO1 templates as compared with PS2. The flowability of the polymers was enhanced by increasedimprinting temperature due to the reduction of polymer viscosities, and decreased surface tensionsbetween polymer melts and AAO templates [11–15]. However, the length of VANs did not increase attemperatures above 180 ◦C because the polymer was easily ejected out from the AAO nanocavitiesunder the high imprinting pressure and relatively low viscosity. PS1–VANs and PS2–VANs of about15–72 µm long were obtained depending upon the imprinting conditions.

The flowability at the nanoscale of the molten polymers in the AAO templates driven by theimprinting process is in agreement with the Hagen–Poiseuille expression [15,50]. The length (H) ordistance the polymer flowed through AAO can be estimated as shown below.

H = 2r

√√(P + 2γ cosθr

)t

32η(1)

This equation shows that the factors related to flowability in confined spaces are imprintingpressure (P), radius of the nanochannel (r), surface tension of the molten polymer (γ), contact angle ofthe polymer on AAO pore wall (θ), viscosity of the molten polymer (η), and infiltration time (t). Ourprevious study revealed that the reduction of apparent viscosity estimated by Equation (1) may be theresult of shear rate, wall slip, and the pressure used in the imprinting process [15].

To avoid the effects of confinement on the phase separation and transformation behaviors atthe nanoscale, therefore, we investigated this change in molecular free–volume by analyzing thethermal properties of obtained PS–VANs as compared with PS bulk films. Specimens were obtained byimprinting conditions of 180 ◦C, 5 MPa, for 30 min with different AAO pore size and Tg of polymers.Figure 4 shows the DSC heating thermograms of the PS bulk films and PS–VANs. The results indicatea change in molecular free–volume was found in the case of PS2–VANs when using AAO1 templates.In the case of PS2, an increase in Tg of the PS2–VANs to 104 ◦C is found deviating from the PS2 bulkfilm precursors having Tg of 100 ◦C. The increase in Tg of PS2 after the imprinting process usingthe AAO1 template due to the nanoconfinement, influences the reduction of the free−volume or theincrease in polymer chain packing density caused by the compression, and shortens each polymer

Polymers 2019, 11, 1039 7 of 18

chain distance. It also might be caused by the applied force from the imprinting pressure inducedthe reduction of free–volume polymer in the AAO nanopore. Moreover, the polymer chains at theinterface have a strong interfacial interaction with the AAO pore walls that can induce the orientationchange of polymer chains [51,52] and also immobilize the polymer chain at the interface, especiallyunder rapid cooling conditions. The result shows that using AAO1 template induced the confinementeffect on the reduction of the free–volume of PS2. Hence, the use of PS1–AAO2 imprinting was selectedto investigate the phase separation and structural transformation behaviors at the nanoscale.


under rapid cooling conditions. The result shows that using AAO1 template induced the confinement 227 effect on the reduction of the free–volume of PS2. Hence, the use of PS1–AAO2 imprinting was selected 228 to investigate the phase separation and structural transformation behaviors at the nanoscale. 229

230 Figure 4. Differential scanning calorimetry (DSC) thermograms of polymer films and 231 nanostructures of (a) PS1 and (b) PS2. 232

3.3. Phase Separation and Structural Transformation of Polymer Blends at the Nanoscale 233 Phase separation and the structural transformation behaviors of PS1/PVA blends confined in 234

AAO2 nanochannels are experimentally investigated herein. The effects of domain size and the mass 235 ratios of PS1/PVA blends were investigated to clarify characteristic behaviors confined in the AAO2 236 template, which influence the obtained morphologies of PS1 nanostructures. Figure 5 (a, A and b, B) 237 show the SEM images of the cross–sectional PS1/PVA blends prepared by the advanced high–shear 238 process. PVA dispersed phases with the domain size in the range of nanodomain and microdomain 239 were prepared with the difference in PVA mass ratios of 10 and 30 wt% respectively. In the case of 240 PS1/PVA (90/10 wt%) blend, PVA nanodomain was formed having a diameter in the range of 10–30 241 nm. For the PS1/PVA (70/30 wt%) blend, PVA microdomain was produced showing the smallest and 242 average diameters of 300 nm and of 30 µm, respectively. Figure 5 (c, C) presents the mapping image 243 of the ATR of the PS1/PVA film precursors. These results indicate that the uniform dispersion of the 244 PVA in the PS1 matrix was able to be prepared by using the high-shear process. However, the non–245 uniform area indicated signal distortion due to scattering of Infrared radiation (IR). 246

247

Figure 5. Surface morphologies and surface chemical analysis of PS1/PVA blends; (a, b) 248 SEM images of PS1/PVA (70/30 wt%), (A, B) SEM images of PS1/PVA (90/10 wt%), (c) ATR 249

mapping of PS1/PVA (70/30 wt%) and (C) ATR mapping of PS1/PVA (90/10 wt%). 250

Figure 4. Differential scanning calorimetry (DSC) thermograms of polymer films and nanostructures of(a) PS1 and (b) PS2.

3.3. Phase Separation and Structural Transformation of Polymer Blends at the Nanoscale

Phase separation and the structural transformation behaviors of PS1/PVA blends confined inAAO2 nanochannels are experimentally investigated herein. The effects of domain size and the massratios of PS1/PVA blends were investigated to clarify characteristic behaviors confined in the AAO2

template, which influence the obtained morphologies of PS1 nanostructures. Figure 5a,A and b,B showthe SEM images of the cross–sectional PS1/PVA blends prepared by the advanced high–shear process.PVA dispersed phases with the domain size in the range of nanodomain and microdomain wereprepared with the difference in PVA mass ratios of 10 and 30 wt% respectively. In the case of PS1/PVA(90/10 wt%) blend, PVA nanodomain was formed having a diameter in the range of 10–30 nm. Forthe PS1/PVA (70/30 wt%) blend, PVA microdomain was produced showing the smallest and averagediameters of 300 nm and of 30 µm, respectively. Figure 5c,C presents the mapping image of the ATRof the PS1/PVA film precursors. These results indicate that the uniform dispersion of the PVA in thePS1 matrix was able to be prepared by using the high-shear process. However, the non–uniform areaindicated signal distortion due to scattering of Infrared radiation (IR).

Polymers 2019, 11, 1039 8 of 18


under rapid cooling conditions. The result shows that using AAO1 template induced the confinement 227 effect on the reduction of the free–volume of PS2. Hence, the use of PS1–AAO2 imprinting was selected 228 to investigate the phase separation and structural transformation behaviors at the nanoscale. 229

230 Figure 4. Differential scanning calorimetry (DSC) thermograms of polymer films and 231 nanostructures of (a) PS1 and (b) PS2. 232

3.3. Phase Separation and Structural Transformation of Polymer Blends at the Nanoscale 233 Phase separation and the structural transformation behaviors of PS1/PVA blends confined in 234

AAO2 nanochannels are experimentally investigated herein. The effects of domain size and the mass 235 ratios of PS1/PVA blends were investigated to clarify characteristic behaviors confined in the AAO2 236 template, which influence the obtained morphologies of PS1 nanostructures. Figure 5 (a, A and b, B) 237 show the SEM images of the cross–sectional PS1/PVA blends prepared by the advanced high–shear 238 process. PVA dispersed phases with the domain size in the range of nanodomain and microdomain 239 were prepared with the difference in PVA mass ratios of 10 and 30 wt% respectively. In the case of 240 PS1/PVA (90/10 wt%) blend, PVA nanodomain was formed having a diameter in the range of 10–30 241 nm. For the PS1/PVA (70/30 wt%) blend, PVA microdomain was produced showing the smallest and 242 average diameters of 300 nm and of 30 µm, respectively. Figure 5 (c, C) presents the mapping image 243 of the ATR of the PS1/PVA film precursors. These results indicate that the uniform dispersion of the 244 PVA in the PS1 matrix was able to be prepared by using the high-shear process. However, the non–245 uniform area indicated signal distortion due to scattering of Infrared radiation (IR). 246

247

Figure 5. Surface morphologies and surface chemical analysis of PS1/PVA blends; (a, b) 248 SEM images of PS1/PVA (70/30 wt%), (A, B) SEM images of PS1/PVA (90/10 wt%), (c) ATR 249

mapping of PS1/PVA (70/30 wt%) and (C) ATR mapping of PS1/PVA (90/10 wt%). 250

Figure 5. Surface morphologies and surface chemical analysis of PS1/PVA blends; (a,b) SEM images ofPS1/PVA (70/30 wt%), (A,B) SEM images of PS1/PVA (90/10 wt%), (c) ATR mapping of PS1/PVA (70/30wt%) and (C) ATR mapping of PS1/PVA (90/10 wt%).

Figure 6A illustrates the top view images of PS1–VANs surface morphologies produced by thermalnanoimprint process. The obtained smooth surface of the PS1–VANs can be prepared using the neat PS1

film precursors. PS1–VAPNs can be produced using the PS1/PVA (90/10 wt%) precursor films followedby the selective removal of PVA segment as shown in Figure 6B,b. In the case of the PVA microdomain,the PS1–VANs with the nanopeapods–like form was replicated consisting of nanorods and nanospheres(See Figure 6C,c). The microholes on the imprinted film surface were observed because of the removalof the PVA segment. Transformation of PS1 nanostructures from nanocolumns into the nanorods andthe nanospheres was induced by the Rayleigh instability due to the interfacial interaction of PS1 andPVA. The reduction of dimensionality from 1D into 0D is the intention of decreasing the surface energyof the system. In this experiment, the influence of imprinting conditions (temperature and time) wereinvestigated to clarify the phase separation and structural transformation behaviors of the polymernanostructures confined in AAO2 templates.


Figure 6A illustrates the top view images of PS1–VANs surface morphologies produced by 251 thermal nanoimprint process. The obtained smooth surface of the PS1–VANs can be prepared using 252 the neat PS1 film precursors. PS1–VAPNs can be produced using the PS1/PVA (90/10 wt%) precursor 253 films followed by the selective removal of PVA segment as shown in Figure 6 (B, b). In the case of the 254 PVA microdomain, the PS1–VANs with the nanopeapods–like form was replicated consisting of 255 nanorods and nanospheres (See Figure 6 (C, c)). The microholes on the imprinted film surface were 256 observed because of the removal of the PVA segment. Transformation of PS1 nanostructures from 257 nanocolumns into the nanorods and the nanospheres was induced by the Rayleigh instability due to 258 the interfacial interaction of PS1 and PVA. The reduction of dimensionality from 1D into 0D is the 259 intention of decreasing the surface energy of the system. In this experiment, the influence of 260 imprinting conditions (temperature and time) were investigated to clarify the phase separation and 261 structural transformation behaviors of the polymer nanostructures confined in AAO2 templates. 262

263

Figure 6. Surface morphologies of (A, a) PS1–VANs (B, b) PS1–VAPNs and (C, c) PS1 264 nanopeapods. 265

Formation of the porous PS1 nanostructured surfaces was able to be generated by the surface–266 induced phase separation of the PS1/PVA (90/10 wt%) blend in AAO2 templates. Figure 7a–d 267 illustrates the influence of the imprinting temperatures (120–180 °C) on the formation of pores on the 268 nanostructured surfaces of PS1. The visible pores were able to be imaged at the imprinting 269 temperature at 180 °C. At this temperature, the formation of nanopores was caused by; (1) the 270 flowability of PVA at melting state (Tm = 161 °C) and (2) the surface–induced phase separation due to 271 the affinity of the hydrogen bond between AAO and PVA. Interfacial interaction between PVA-AAO2 272 also promotes the coarsening behavior owing to the hydrogen bond. Therefore, the PVA droplets in 273 the PS1 matrix were formed on the pore wall of the AAO2 template. Moreover, the reduction of PS1 274 viscosity with increasing imprinting temperature can promote the surface-induced phase separation. 275 Figure 7e illustrates the reduction of the polymer melt droplet angles of the PS1/PVA (90/10 wt%) 276 blend with increasing temperature. At the temperature of 180 °C, it was found that the polymer melt 277 droplets had a contact angle lower than 90° which indicates the oxophilicity with AAO2 templates. 278 Hence, the surface interaction between polymers and AAO becomes the important role in the phase 279 separation at the imprinting temperature of 180 °C. Furthermore, thermal-induced flowability of PVA 280 might enhance the formation of PVA droplets on the AAO2 pore wall due to the low viscosity at 281 above melting temperature. 282

Figure 6. Surface morphologies of (A,a) PS1—VANs (B,b) PS1—VAPNs and (C,c) PS1 nanopeapods.

Polymers 2019, 11, 1039 9 of 18

Formation of the porous PS1 nanostructured surfaces was able to be generated by thesurface–induced phase separation of the PS1/PVA (90/10 wt%) blend in AAO2 templates. Figure 7a–dillustrates the influence of the imprinting temperatures (120–180 ◦C) on the formation of pores on thenanostructured surfaces of PS1. The visible pores were able to be imaged at the imprinting temperatureat 180 ◦C. At this temperature, the formation of nanopores was caused by; (1) the flowability of PVA atmelting state (Tm = 161 ◦C) and (2) the surface–induced phase separation due to the affinity of thehydrogen bond between AAO and PVA. Interfacial interaction between PVA-AAO2 also promotesthe coarsening behavior owing to the hydrogen bond. Therefore, the PVA droplets in the PS1 matrixwere formed on the pore wall of the AAO2 template. Moreover, the reduction of PS1 viscosity withincreasing imprinting temperature can promote the surface-induced phase separation. Figure 7eillustrates the reduction of the polymer melt droplet angles of the PS1/PVA (90/10 wt%) blend withincreasing temperature. At the temperature of 180 ◦C, it was found that the polymer melt dropletshad a contact angle lower than 90

◦which indicates the oxophilicity with AAO2 templates. Hence, the

surface interaction between polymers and AAO becomes the important role in the phase separationat the imprinting temperature of 180 ◦C. Furthermore, thermal-induced flowability of PVA mightenhance the formation of PVA droplets on the AAO2 pore wall due to the low viscosity at abovemelting temperature.Polymers 2019, 11, x FOR PEER REVIEW 9 of 18

283

Figure 7. Effect of imprinting temperature on formation of nanoporous (left); (a) 120℃, (b) 284 140 ℃, (c) 160 ℃, (d) 140 ℃, and (e) PS1/PVA melt droplets on AAO surfaces (right). 285

Surface-induced phase separation at the macroscale has been investigated with the annealing 286 process of PS1/PVA (90/10 wt%) blend films covered by the AAO template (upper) and aluminum 287 plate (lower) at the temperature of 120–180 °C for 30 min. The surface chemical properties of the 288 annealed films were simulated using the ATR mapping mode. By the result, the surface film of the 289 polymer blend precursor has the composition of PVA of about ~10% (% hydroxyl group). After the 290 annealing process at 180 °C, the ratio of PVA on the blend film surface adhered on the AAO side was 291 reached about ~60–80% but showed the lower ratio on the aluminum side. The increase in the ratio 292 of PVA on the annealed PS1/PVA blend film surface indicated the strong surface interaction between 293 PVA and AAO above the melting temperature of PVA, as previously explained for the formation of 294 nanoporous results. This can be confirmed that the formation of the porous PS nanostructured surface 295 is caused by surface-induced phase separation at the imprinting temperature of 180 °C (See Figure 296 8). 297

298

Figure 8. Annealing behaviors of PS1/PVA (90/10 wt%) blends. 299

According to the coarsening behavior, influences of coarsening temperature (with the function 300 of viscosities and surface tensions) and coarsening times are the crucial parameters to the formation 301 of the porous nanostructures. In this part, the influence of the imprinting time on the formation of 302 the porous nanostructured surfaces will be investigated. The imprinting temperature and pressure 303 of 180 °C and 5 MPa were performed in the experiment. Figure 9 presents the surface appearances of 304

Figure 7. Effect of imprinting temperature on formation of nanoporous (left); (a) 120°C, (b) 140 °C,(c) 160 °C, (d) 140 °C, and (e) PS1/PVA melt droplets on AAO surfaces (right).

Surface-induced phase separation at the macroscale has been investigated with the annealingprocess of PS1/PVA (90/10 wt%) blend films covered by the AAO template (upper) and aluminum plate(lower) at the temperature of 120–180 ◦C for 30 min. The surface chemical properties of the annealedfilms were simulated using the ATR mapping mode. By the result, the surface film of the polymerblend precursor has the composition of PVA of about ~10% (% hydroxyl group). After the annealingprocess at 180 ◦C, the ratio of PVA on the blend film surface adhered on the AAO side was reachedabout ~60–80% but showed the lower ratio on the aluminum side. The increase in the ratio of PVA onthe annealed PS1/PVA blend film surface indicated the strong surface interaction between PVA andAAO above the melting temperature of PVA, as previously explained for the formation of nanoporousresults. This can be confirmed that the formation of the porous PS nanostructured surface is caused bysurface-induced phase separation at the imprinting temperature of 180 ◦C (See Figure 8).

Polymers 2019, 11, 1039 10 of 18


283

Figure 7. Effect of imprinting temperature on formation of nanoporous (left); (a) 120℃, (b) 284 140 ℃, (c) 160 ℃, (d) 140 ℃, and (e) PS1/PVA melt droplets on AAO surfaces (right). 285

Surface-induced phase separation at the macroscale has been investigated with the annealing 286 process of PS1/PVA (90/10 wt%) blend films covered by the AAO template (upper) and aluminum 287 plate (lower) at the temperature of 120–180 °C for 30 min. The surface chemical properties of the 288 annealed films were simulated using the ATR mapping mode. By the result, the surface film of the 289 polymer blend precursor has the composition of PVA of about ~10% (% hydroxyl group). After the 290 annealing process at 180 °C, the ratio of PVA on the blend film surface adhered on the AAO side was 291 reached about ~60–80% but showed the lower ratio on the aluminum side. The increase in the ratio 292 of PVA on the annealed PS1/PVA blend film surface indicated the strong surface interaction between 293 PVA and AAO above the melting temperature of PVA, as previously explained for the formation of 294 nanoporous results. This can be confirmed that the formation of the porous PS nanostructured surface 295 is caused by surface-induced phase separation at the imprinting temperature of 180 °C (See Figure 296 8). 297

298

Figure 8. Annealing behaviors of PS1/PVA (90/10 wt%) blends. 299

According to the coarsening behavior, influences of coarsening temperature (with the function 300 of viscosities and surface tensions) and coarsening times are the crucial parameters to the formation 301 of the porous nanostructures. In this part, the influence of the imprinting time on the formation of 302 the porous nanostructured surfaces will be investigated. The imprinting temperature and pressure 303 of 180 °C and 5 MPa were performed in the experiment. Figure 9 presents the surface appearances of 304

Figure 8. Annealing behaviors of PS1/PVA (90/10 wt%) blends.

According to the coarsening behavior, influences of coarsening temperature (with the functionof viscosities and surface tensions) and coarsening times are the crucial parameters to the formationof the porous nanostructures. In this part, the influence of the imprinting time on the formation ofthe porous nanostructured surfaces will be investigated. The imprinting temperature and pressure of180 ◦C and 5 MPa were performed in the experiment. Figure 9 presents the surface appearances ofPS1 nanostructures obtained from the imprinting time of 15–120 min. With the long imprinting timeof about 90–120 min, we found the breakup of the PS1 nanocolumns into nanofragments due to theagglomeration of the PVA droplets during and after the coarsening behaviors. The agglomeration ofPVA droplets depends on the time of diffusion and mass transfer of the PVA.


PS1 nanostructures obtained from the imprinting time of 15–120 min. With the long imprinting time 305 of about 90–120 min, we found the breakup of the PS1 nanocolumns into nanofragments due to the 306 agglomeration of the PVA droplets during and after the coarsening behaviors. The agglomeration of 307 PVA droplets depends on the time of diffusion and mass transfer of the PVA. 308

309

Figure 9. Effect of imprinting time on the obtained PS1 nanostructures. 310

Jiun Tai Chen et al [19] have explained the surface–induced phase separation behaviors by the 311 coarsening process between PMMA and THF confined in AAO templates. The explanation reveals 312 that the coarsening process progressed by the three possible mechanisms consisting of the Ostwald 313 ripening, coalescence, and hydrodynamic flow. The Ostwald ripening was widely used to clarify the 314 phase separation of the polymer solutions. The driving force of this mechanism is to minimize the 315 interfacial energy between the polymer–rich phase and solvent-rich phase or the PS1–rich phase and 316 PVA–rich phase for this case. The molecules of the tiny droplets of the dispersed phase dissolve due 317 to the higher curvatures and precipitate on the surface of the large droplets. The reduction of the total 318 interfacial area results in the decrease in the interfacial energy. The asymptotic power–low relation 319 was proposed to explain the domain size of the droplets as in the following equation, 320

( ) 3131 tDd ξ≈ , (2)

where d is domain size of droplet, D is the diffusion coefficient,ξ is the correlation length, and t321 is the coarsening time. 322

The second mechanism of the coarsening process is caused by the coalescence of the two droplets 323 to form the larger droplet. The impinging of two droplets caused by the translational diffusion 324 promotes the mass transfer, and then the larger single droplet is formed. The droplet domain size 325 relates to time as in Equation (3) 326

13 1

3Bk Td tη

≈

(3)

where Bk is the Boltzmann constant, T is the temperature, and η is the viscosity. In this study, 327 the viscosity of molten polymers used was reduced with increasing temperature referred to the 328 viscosity Arrhenius model [11]. 329

The third coarsening mechanism caused by the hydrodynamic flow of the fluid mixtures. The 330 gradients of pressure along the axis of the cylindrical pore play the role of the flowability from a 331 narrow to a wide region that induces the bi–continuous of the two polymers formed during the phase 332

Figure 9. Effect of imprinting time on the obtained PS1 nanostructures.

Jiun Tai Chen et al. [19] have explained the surface–induced phase separation behaviors by thecoarsening process between PMMA and THF confined in AAO templates. The explanation revealsthat the coarsening process progressed by the three possible mechanisms consisting of the Ostwaldripening, coalescence, and hydrodynamic flow. The Ostwald ripening was widely used to clarify thephase separation of the polymer solutions. The driving force of this mechanism is to minimize theinterfacial energy between the polymer–rich phase and solvent-rich phase or the PS1–rich phase andPVA–rich phase for this case. The molecules of the tiny droplets of the dispersed phase dissolve due tothe higher curvatures and precipitate on the surface of the large droplets. The reduction of the total

Polymers 2019, 11, 1039 11 of 18

interfacial area results in the decrease in the interfacial energy. The asymptotic power–low relationwas proposed to explain the domain size of the droplets as in the following equation,

d ≈ (Dξ)1/3t1/3, (2)

where d is domain size of droplet, D is the diffusion coefficient, ξ is the correlation length, and t is thecoarsening time.

The second mechanism of the coarsening process is caused by the coalescence of the two dropletsto form the larger droplet. The impinging of two droplets caused by the translational diffusionpromotes the mass transfer, and then the larger single droplet is formed. The droplet domain sizerelates to time as in Equation (3)

d ≈(

kBTη

) 13

t13 (3)

where kB is the Boltzmann constant, T is the temperature, and η is the viscosity. In this study, theviscosity of molten polymers used was reduced with increasing temperature referred to the viscosityArrhenius model [11].

The third coarsening mechanism caused by the hydrodynamic flow of the fluid mixtures. Thegradients of pressure along the axis of the cylindrical pore play the role of the flowability from anarrow to a wide region that induces the bi–continuous of the two polymers formed during the phaseseparation. The growth of the droplet domain is related linearly with time as in the following equation.Figure 10 simulates the phase separation caused by the coarsening behavior.

d ≈ ση

t (4)

where σ is the surface tension.


separation. The growth of the droplet domain is related linearly with time as in the following 333 equation. Figure 10 simulates the phase separation caused by the coarsening behavior. 334

d tση

≈ (4)

where σ is the surface tension. 335

336

337 Figure 10. Formation of polymer droplets caused by surface−induced phase separation and 338 coarsening behaviors. 339

In the case of the PVA microdomain, Figure 7 (C, c) illustrates the morphologies of PS1 340 nanostructured surfaces using PS1/PVA (70/30 wt%) blend precursor. After the removal of the PVA 341 segment, we found PS1 nanostructure arrays with microholes. The appearance of microholes is 342 caused by the PVA microdomain size of the film precursor. At the microholes area, we found PS1 343 nanostructures with nanorods and nanospheres (See Figure 11). The transformation of the PS1 344 nanostructures is the result of the surface interaction between PS1 and PVA in the AAO2 templates. 345 The PS1 encapsulated by the PVA–rich phase (30 wt%) adjoined with the AAO2 pore wall, and the 346 interfacial interaction between PS1 and PVA lead to the structural transformation of the polymer 347 nanostructures, which was driven by the Rayleigh instability within AAO2 templates. Rayleigh 348 instability is the common transformation phenomena when the columns of water are falling and 349 breakup into water droplets. The instability is the result of the undulation of the liquid cylinder 350 surface at low surface tension. Herein, the influences of imprinting temperatures (120–180 °C) and 351 imprinting time (30–120 min) have been investigated. It was found that the number of the spherical 352 structures increased when increasing the imprinting temperature and imprinting time. With 353 increasing the imprinting temperatures, Rayleigh instability caused by the decrease in viscosities that 354 enhances the degree of structural transformation of the PS1 nanocolumns and converted to the 355 nanorods and nanospheres. The increased imprinting temperature also resulted in the increasing rate 356 of mass transfer and the kinetic phenomenon of the transformation. The increase in the immiscibility 357 of the polymer blends due to the rise in the temperature that influences the increasing heat of mixing. 358 Therefore, it was found that the number of nanorods and nanospheres increased when elevating the 359 imprinting temperatures. 360

Figure 10. Formation of polymer droplets caused by surface−induced phase separation andcoarsening behaviors.

Polymers 2019, 11, 1039 12 of 18

In the case of the PVA microdomain, Figure 7C,c illustrates the morphologies of PS1 nanostructuredsurfaces using PS1/PVA (70/30 wt%) blend precursor. After the removal of the PVA segment, wefound PS1 nanostructure arrays with microholes. The appearance of microholes is caused by the PVAmicrodomain size of the film precursor. At the microholes area, we found PS1 nanostructures withnanorods and nanospheres (See Figure 11). The transformation of the PS1 nanostructures is the resultof the surface interaction between PS1 and PVA in the AAO2 templates. The PS1 encapsulated by thePVA–rich phase (30 wt%) adjoined with the AAO2 pore wall, and the interfacial interaction betweenPS1 and PVA lead to the structural transformation of the polymer nanostructures, which was driven bythe Rayleigh instability within AAO2 templates. Rayleigh instability is the common transformationphenomena when the columns of water are falling and breakup into water droplets. The instability isthe result of the undulation of the liquid cylinder surface at low surface tension. Herein, the influencesof imprinting temperatures (120–180 ◦C) and imprinting time (30–120 min) have been investigated.It was found that the number of the spherical structures increased when increasing the imprintingtemperature and imprinting time. With increasing the imprinting temperatures, Rayleigh instabilitycaused by the decrease in viscosities that enhances the degree of structural transformation of the PS1

nanocolumns and converted to the nanorods and nanospheres. The increased imprinting temperaturealso resulted in the increasing rate of mass transfer and the kinetic phenomenon of the transformation.The increase in the immiscibility of the polymer blends due to the rise in the temperature that influencesthe increasing heat of mixing. Therefore, it was found that the number of nanorods and nanospheresincreased when elevating the imprinting temperatures.


361 Figure 11. SEM images of PS1 nanostructured arrays obtained by PS1/PVA (70/30 wt%); 362 (upper; imprinting time 30 min) with the imprinting temperatures of (A) 120 (B) 140 (C) 160 363 and (D) 180 °C, lower; imprinting temperature 180 °C with imprinting times of (a) 15 (b) 30 364 (c) 60 and (d) 120 min). 365

In terms of the influence of the imprinting time, we found the similarity in the structural 366 transformation behaviors when increasing the imprinting temperatures. The number of nanorods 367 and nanospheres increased with increasing the length of imprinting time as shown in Figure 11. 368 However, the result reveals that the agglomeration of the PS1 segment has not much effect in this case. 369 Jiun Tai Chen et al. explained clearly the transformation of the nanostructures driven by the Rayleigh 370 instability with the characteristic time scale that agrees with their experimental results. The Rayleigh 371 instability was able to explain this by Equations (5) and (6) as below [24,25]. Figure 12 simulates the 372 phase separation caused by the coarsening behavior. 373

374 Figure 12. Transformation of polymer nanostructures driven by the Rayleigh instability. 375

In the case of fluid viscosity is neglected, the break up time as the characteristic time scale driven 376 by the Rayleigh instability is 377

( ) 210 γρτ R= (5)

where ρ is density of the fluid and 0R is the initial radius of fluid cylinder. 378

In the case of viscoelastic materials, the characteristic time is given by 379

γητ 0R= (6)

Figure 11. SEM images of PS1 nanostructured arrays obtained by PS1/PVA (70/30 wt%); (upper;imprinting time 30 min) with the imprinting temperatures of (A) 120 (B) 140 (C) 160 and (D) 180 ◦C,lower; imprinting temperature 180 ◦C with imprinting times of (a) 15 (b) 30 (c) 60 and (d) 120 min).

In terms of the influence of the imprinting time, we found the similarity in the structuraltransformation behaviors when increasing the imprinting temperatures. The number of nanorods andnanospheres increased with increasing the length of imprinting time as shown in Figure 11. However,the result reveals that the agglomeration of the PS1 segment has not much effect in this case. JiunTai Chen et al. explained clearly the transformation of the nanostructures driven by the Rayleighinstability with the characteristic time scale that agrees with their experimental results. The Rayleighinstability was able to explain this by Equations (5) and (6) as below [24,25]. Figure 12 simulates thephase separation caused by the coarsening behavior.

Polymers 2019, 11, 1039 13 of 18


361 Figure 11. SEM images of PS1 nanostructured arrays obtained by PS1/PVA (70/30 wt%); 362 (upper; imprinting time 30 min) with the imprinting temperatures of (A) 120 (B) 140 (C) 160 363 and (D) 180 °C, lower; imprinting temperature 180 °C with imprinting times of (a) 15 (b) 30 364 (c) 60 and (d) 120 min). 365

In terms of the influence of the imprinting time, we found the similarity in the structural 366 transformation behaviors when increasing the imprinting temperatures. The number of nanorods 367 and nanospheres increased with increasing the length of imprinting time as shown in Figure 11. 368 However, the result reveals that the agglomeration of the PS1 segment has not much effect in this case. 369 Jiun Tai Chen et al. explained clearly the transformation of the nanostructures driven by the Rayleigh 370 instability with the characteristic time scale that agrees with their experimental results. The Rayleigh 371 instability was able to explain this by Equations (5) and (6) as below [24,25]. Figure 12 simulates the 372 phase separation caused by the coarsening behavior. 373

374 Figure 12. Transformation of polymer nanostructures driven by the Rayleigh instability. 375

In the case of fluid viscosity is neglected, the break up time as the characteristic time scale driven 376 by the Rayleigh instability is 377

( ) 210 γρτ R= (5)

where ρ is density of the fluid and 0R is the initial radius of fluid cylinder. 378

In the case of viscoelastic materials, the characteristic time is given by 379

γητ 0R= (6)

Figure 12. Transformation of polymer nanostructures driven by the Rayleigh instability.

In the case of fluid viscosity is neglected, the break up time as the characteristic time scale drivenby the Rayleigh instability is

τ = (ρR0/γ)1/2 (5)

where ρ is density of the fluid and R0 is the initial radius of fluid cylinder.In the case of viscoelastic materials, the characteristic time is given by

τ = ηR0/γ (6)

3.4. Surface Properties

Surface properties of polymers were evaluated using the water droplet angle measurement,macroscale friction test, and nanoindentation. The five types of polymer films used to investigatewere; (1) neat PS1, (2) neat PVA, (3) PS1/PVA (90/10 wt%) blend, (4) PS1–VANs, and (5) PS1–VAPNs.The imprinted samples were obtained using the imprinting condition of 180 ◦C, 5 MPa for 30 min(nanostructures length 70 ± 10 µm with the remained film thickness 40 ± 10 µm).

The results of the water droplet angle measurement was presented in Figure 13. The flat surfacesof PS1 and PVA films with the droplet angle of water about ~89.8◦ and ~40◦–50◦ (initial droplet) wereobtained. The water–soluble properties of the PVA influence the instability of the water droplet shapeand its angle as shown on the shadow side of the water droplet. The water droplet angle of PS1/PVAfilm with a value of about ~65.6◦ can be measured. PS1–VANs surface shows the water droplet anglereach ~130.1◦. The great increase in water droplet angle as compared to the flat films is the result of thephysical roughness of the nanostructures and the reduction of the contact area between polymer andwater droplet. Moreover, the PS1–VAPNs surface has a slight increase in the water droplet angle to~140.1◦ as compared with the PS1–VANs due to roughness at the nanoscale.


3.4. Surface Properties 380 Surface properties of polymers were evaluated using the water droplet angle measurement, 381

macroscale friction test, and nanoindentation. The five types of polymer films used to investigate 382 were; (1) neat PS1, (2) neat PVA, (3) PS1/PVA (90/10 wt%) blend, (4) PS1––VANs, and (5) PS1–VAPNs. 383 The imprinted samples were obtained using the imprinting condition of 180 °C, 5 MPa for 30 min 384 (nanostructures length 70 ± 10 µm with the remained film thickness 40 ± 10 µm). 385

The results of the water droplet angle measurement was presented in Figure 13. The flat surfaces 386 of PS1 and PVA films with the droplet angle of water about ~89.8° and ~40°–50° (initial droplet) were 387 obtained. The water–soluble properties of the PVA influence the instability of the water droplet shape 388 and its angle as shown on the shadow side of the water droplet. The water droplet angle of PS1/PVA 389 film with a value of about ~65.6° can be measured. PS1–VANs surface shows the water droplet angle 390 reach ~130.1°. The great increase in water droplet angle as compared to the flat films is the result of 391 the physical roughness of the nanostructures and the reduction of the contact area between polymer 392 and water droplet. Moreover, the PS1–VAPNs surface has a slight increase in the water droplet angle 393 to ~140.1° as compared with the PS1–VANs due to roughness at the nanoscale. 394

395 Figure 13. Water droplets on the polymer surfaces. 396

Surface friction properties at macroscale demonstrated the same tendency with results of the 397 water droplet angle measurement. The measured friction coefficients was placed in Figure 14. The 398 PS1–VAPNs surfaces exhibit the higher friction coefficient (~0.53–0.55) as compared with PS1–VANs 399 surface (~0.37–0.42). The results indicate the roughness at the nanoscale influences the changes in 400 surface friction at the macroscale, and the surface energy in the water droplet angle measurement as 401 in the explanation above [53–56]. In addition, the friction coefficient of PS1–VAPNs can be enhanced 402 when it was switched to the wetting condition. This result is due to the water bridge–mediated 403 contact formation induced by the solid–solid contact between the contact elements and the porous 404 surface of PS1–VAPNs [9,10]. 405

Figure 13. Water droplets on the polymer surfaces.

Polymers 2019, 11, 1039 14 of 18

Surface friction properties at macroscale demonstrated the same tendency with results of thewater droplet angle measurement. The measured friction coefficients was placed in Figure 14. ThePS1–VAPNs surfaces exhibit the higher friction coefficient (~0.53–0.55) as compared with PS1–VANssurface (~0.37–0.42). The results indicate the roughness at the nanoscale influences the changes insurface friction at the macroscale, and the surface energy in the water droplet angle measurement asin the explanation above [53–56]. In addition, the friction coefficient of PS1–VAPNs can be enhancedwhen it was switched to the wetting condition. This result is due to the water bridge–mediated contactformation induced by the solid–solid contact between the contact elements and the porous surface ofPS1–VAPNs [9,10].Polymers 2019, 11, x FOR PEER REVIEW 14 of 18

406 Figure 14. Friction coefficients of polymer surfaces 407

By the results of the nanoindentation test, the PS1–VAPNs show a lower surface hardness and 408 surface modulus as compared with PS1–VANs due to the scaling down or the size reduction effect of 409 the material structures after the removal of the PVA segment [57]. Hence, the size reduction has the 410 result of a decrease in the energy absorption capacity and the destruction of the nanostructures 411 becomes easier. The surface hardness of 0.6 and 0.8 MPa and surface modulus of 0.03 and 0.08 GPa 412 of the PS1–VAPNs and PS1–VANs can be respectively obtained using the indentation test. Figure 15 413 presents the nanoindentation test results of the imprinted nanostructures. 414

415 Figure 15. Nanoindentation test results. 416

4. Conclusions 417 In summary, phase separation and the structural transformation behaviors of the PS1/PVA blend 418

in AAO template (100 nm diameter) for fabricating the PS1–VAPNs has been investigated. The 419 replication process used is the thermal nanoimprint lithography. In this study, the different behaviors 420 of the phase separation and the structural transformation were the result of the domain size and the 421 mass ratios of PVA precursors. 422

In the case of the PVA nanodomain (PS1/PVA 90/10 wt%), formation of the PVA droplets on the 423 AAO pore wall due to the surface–induced phase separation behavior, resulting in the porous surface 424 which occurred after the removal of PVA. The formation of pores on the PS1 nanostructured surface 425 occurred at imprinting temperatures above the melting temperature of PVA. This result indicated the 426 high flowability of the molten PVA has an important role in the mass transfer of the phase separation. 427

Figure 14. Friction coefficients of polymer surfaces

By the results of the nanoindentation test, the PS1–VAPNs show a lower surface hardness andsurface modulus as compared with PS1–VANs due to the scaling down or the size reduction effectof the material structures after the removal of the PVA segment [57]. Hence, the size reduction hasthe result of a decrease in the energy absorption capacity and the destruction of the nanostructuresbecomes easier. The surface hardness of 0.6 and 0.8 MPa and surface modulus of 0.03 and 0.08 GPaof the PS1–VAPNs and PS1–VANs can be respectively obtained using the indentation test. Figure 15presents the nanoindentation test results of the imprinted nanostructures.


406 Figure 14. Friction coefficients of polymer surfaces 407

By the results of the nanoindentation test, the PS1–VAPNs show a lower surface hardness and 408 surface modulus as compared with PS1–VANs due to the scaling down or the size reduction effect of 409 the material structures after the removal of the PVA segment [57]. Hence, the size reduction has the 410 result of a decrease in the energy absorption capacity and the destruction of the nanostructures 411 becomes easier. The surface hardness of 0.6 and 0.8 MPa and surface modulus of 0.03 and 0.08 GPa 412 of the PS1–VAPNs and PS1–VANs can be respectively obtained using the indentation test. Figure 15 413 presents the nanoindentation test results of the imprinted nanostructures. 414

415 Figure 15. Nanoindentation test results. 416

4. Conclusions 417 In summary, phase separation and the structural transformation behaviors of the PS1/PVA blend 418

in AAO template (100 nm diameter) for fabricating the PS1–VAPNs has been investigated. The 419 replication process used is the thermal nanoimprint lithography. In this study, the different behaviors 420 of the phase separation and the structural transformation were the result of the domain size and the 421 mass ratios of PVA precursors. 422

In the case of the PVA nanodomain (PS1/PVA 90/10 wt%), formation of the PVA droplets on the 423 AAO pore wall due to the surface–induced phase separation behavior, resulting in the porous surface 424 which occurred after the removal of PVA. The formation of pores on the PS1 nanostructured surface 425 occurred at imprinting temperatures above the melting temperature of PVA. This result indicated the 426 high flowability of the molten PVA has an important role in the mass transfer of the phase separation. 427

Figure 15. Nanoindentation test results.

Polymers 2019, 11, 1039 15 of 18

4. Conclusions

In summary, phase separation and the structural transformation behaviors of the PS1/PVA blend inAAO template (100 nm diameter) for fabricating the PS1–VAPNs has been investigated. The replicationprocess used is the thermal nanoimprint lithography. In this study, the different behaviors of the phaseseparation and the structural transformation were the result of the domain size and the mass ratios ofPVA precursors.

In the case of the PVA nanodomain (PS1/PVA 90/10 wt%), formation of the PVA droplets on theAAO pore wall due to the surface–induced phase separation behavior, resulting in the porous surfacewhich occurred after the removal of PVA. The formation of pores on the PS1 nanostructured surfaceoccurred at imprinting temperatures above the melting temperature of PVA. This result indicated thehigh flowability of the molten PVA has an important role in the mass transfer of the phase separation.The increase in imprinting time promotes the agglomeration of PVA droplets conductive to the breakupof PS1 columns into nanofragments.

With the PVA microdomain precursor (PS1/PVA 70/30 wt%), the results demonstrated thetransformation behaviors of PS1 nanostructures; PS1 nanocolumns transform into the nanopeapodsconsisting of nanorods and nanospheres. The structural transformation from 1D to 0D is induced byRayleigh instability due to the surface interaction between PS1 and PVA. The influences of the increasesin the imprinting time and the imprinting temperature resulted in an increased number of nanorodsand nanospheres due to enhancing the rate and duration of the mass transfer, and the decrease ofpolymer viscosities.

For the surface properties, the formation of pores on nanostructured surfaces of PS1 has theresult of increasing both the water droplet angle and the friction coefficient due to the increase insurface roughness at the nanoscale. Moreover, the friction coefficient of VAPNs can be increased whenswitchable to the wetting condition. However, we found the reduction in the surface hardness andsurface modulus caused by the scaling down effect.

Author Contributions: P.M. conceived, designed, performed the experiments and wrote the paper; H.I. and T.K.are advisors of this research.

Funding: This research was financially supported by the Japan Society for the Promotion of Science (JSPS)Grant−in−Aid for Scientific Research (C), project ID 16K06740.Acknowledgments: The authors gratefully acknowledge the support of the Japan Government (MEXT) andJSPS in cooperation with the Innovative Flex Course for Frontier Organic Material Systems (iFront) of YamagataUniversity to promote full doctoral scholarship for Paritat Muanchan.

Conflicts of Interest: The authors declare no conflict of interest.

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