-
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
-
Polymers 2019, 11, 1039 3 of 18
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
-
Polymers 2019, 11, 1039 4 of 18
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.
Polymers 2019, 11, x FOR PEER REVIEW 4 of 18
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).
-
Polymers 2019, 11, 1039 5 of 18
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.
Polymers 2019, 11, x FOR PEER REVIEW 5 of 18
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
Polymers 2019, 11, x FOR PEER REVIEW 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.
Polymers 2019, 11, x FOR PEER REVIEW 7 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 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
Polymers 2019, 11, x FOR PEER REVIEW 7 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.
Polymers 2019, 11, x FOR PEER REVIEW 8 of 18
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
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 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.
Polymers 2019, 11, x FOR PEER REVIEW 10 of 18
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.
Polymers 2019, 11, x FOR PEER REVIEW 11 of 18
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.
Polymers 2019, 11, x FOR PEER REVIEW 12 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 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
Polymers 2019, 11, x FOR PEER REVIEW 12 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.
Polymers 2019, 11, x FOR PEER REVIEW 13 of 18
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.
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 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.
References
1. Hu, S.; Xia, S. Rational Design and Nanofabrication of
Gecko-Inspired Fibrillar Adhesives. Small 2012, 8,2464–2468.
[CrossRef] [PubMed]
2. Sheng, X.; Zhang, J. Superhydrophobic Behaviors of Polymeric
Surfaces with Aligned Nanofibers. Langmuir2009, 25, 6916–6922.
[CrossRef] [PubMed]
3. Singh, V.; Bougher, T.L.; Weathers, A.; Cai, Y.; Bi, K.;
Pettes, M.T.; McMenamin, S.A.; Lv, W.; Resler, D.P.;Gattuso, T.R.;
et al. High Thermal Conductivity of Chain-Oriented Amorphous
Polythiophene. Nat.Nanotechnol. 2014, 9, 384–390. [CrossRef]
[PubMed]
4. Kim, Y.; Limanto, F.; Lee, D.H.; Fearing, R.S.; Maboudian, R.
Role of Counter-substrate Surface Energy inMacroscale Friction of
Nanofiber Arrays. Langmuir 2012, 28, 2922–2927. [CrossRef]
[PubMed]
5. Mijangos, C.; Hernández, R.; Martín, J. A Review on the
Progress of Polymer Nanostructures with ModulatedMorphologies and
Properties, Using Nanoporous AAO Templates. Prog. Polym. Sci. 2016,
54–55, 148–182.[CrossRef]
6. Hernández, J.J.; Monclús, M.A.; Navarro-Baena, I.; Viela, F.;
Molina-Aldareguia, J.M.; Rodríguez, I.Multifunctional
Nano-engineered Polymer Surfaces with Enhanced Mechanical
Resistance andSuperhydrophobicity. Sci. Rep. 2017, 7, 43450.
[CrossRef]
http://dx.doi.org/10.1002/smll.201200413http://www.ncbi.nlm.nih.gov/pubmed/22641471http://dx.doi.org/10.1021/la9002077http://www.ncbi.nlm.nih.gov/pubmed/19326870http://dx.doi.org/10.1038/nnano.2014.44http://www.ncbi.nlm.nih.gov/pubmed/24681778http://dx.doi.org/10.1021/la204078zhttp://www.ncbi.nlm.nih.gov/pubmed/22263534http://dx.doi.org/10.1016/j.progpolymsci.2015.10.003http://dx.doi.org/10.1038/srep43450
-
Polymers 2019, 11, 1039 16 of 18
7. Dickson, M.N.; Liang, E.I.; Rodriguez, L.A.; Vollereaux, N.;
Yee, A.F. Nanopatterned Polymer Surfaces withBactericidal
Properties. Biointerphases 2015, 10, 021010. [CrossRef]
8. Li, X.; Cheung, G.S.; Watson, G.S.; Watson, J.A.; Lin, S.;
Schwarzkopf, L.; Green, D.W. The NanotippedHairs of Gecko Skin and
Biotemplated Replicas Impair and/or Kill Pathogenic Bacteria with
High Efficiency.Nanoscale 2016, 8, 18860–18869. [CrossRef]
9. Xue, L.; Kovalev, A.; Dening, K.; Eichler-Volf, A.;
Eickmeier, H.; Haase, M.; Enke, D.; Steinhart, M.; Gorb,
S.N.Reversible Adhesion Switching of Porous Fibrillar Adhesive Pads
by Humidity. Nano Lett. 2013, 13,5541–5548. [CrossRef]
10. Xue, L.; Kovalev, A.; Eichler-Volf, A.; Steinhart, M.; Gorb,
S.N. Humidity-Enhanced Wet Adhesion onInsect-Inspired Fibrillar
Adhesive Pads. Nat. Commun. 2015, 6, 6621. [CrossRef]
11. Huang, L.-B.; Zhou, Y.; Han, S.-T.; Yan, Y.; Zhou, L.; Roy,
V.A.L. The Role of Nanoparticle Monolayer on theFlowing Behaviour
of Polymer Melts in Nanochannels. Nanoscale 2014, 6, 11013–11018.
[CrossRef] [PubMed]
12. Yung, K.-L.; Kong, J.; Xu, Y. Studies on Flow Behaviors of
Polymer Melts in Nanochannels by Wetting Actions.Polymer 2007, 48,
7645–7652. [CrossRef]
13. Ali, S.; Tian, W.; Ali, N.; Shi, L.; Kong, J.; Ali, N.
Polymer Melt Flow through Nanochannels: From Theoryand Fabrication
to Application. RSC Adv. 2015, 5, 7160–7172. [CrossRef]
14. Mei, S.; Feng, X.; Jin, Z. Polymer Nanofibers by
Controllable Infiltration of Vapour Swollen Polymers
intoCylindrical Nanopores. Soft Matter 2013, 9, 945–951.
[CrossRef]
15. Muanchan, P.; Suzuki, S.; Kyotani, T.; Ito, H.
One-dimensional Polymer Nanofiber Arrays with High AspectRatio
Obtained by Thermal Nanoimprint Method. Polym. Eng. Sci. 2017, 57,
214–223. [CrossRef]
16. Wu, H.; Su, Z.; Takahara, A. Characterization of An
Isotactic Polystyrene/Poly(2,6 dimethylphenylene oxide)Nanorod
Blend with Gradient Composition and Crystallinity. RSC Adv. 2012,
2, 8707–8712. [CrossRef]
17. Wu, H.; Su, Z.; Takahara, A. Gradient Composition
Distribution in Poly(2,6-dimethylphenyleneoxide)/Polystyrene Blend
Nanorods. Soft Matter 2011, 7, 1868–1873. [CrossRef]
18. Wu, H.; Su, Z.; Takahara, A. Molecular Composition
Distribution of Polycarbonate/Polystyrene Blends inCylindrical
Nanopores. Polym. J. 2011, 43, 600–605. [CrossRef]
19. Wei, T.-H.; Chi, M.-H.; Tsai, C.-C.; Ko, H.-W.; Chen, J.-T.
Porous Polymer Nanostructures Fabricated by theSurface-Induced
Phase Separation of Polymer Solutions in Anodic Aluminum Oxide
Templates. Langmuir2013, 29, 9972–9978. [CrossRef]
20. Chi, M.-H.; Chang, C.-W.; Ko, H.-W.; Su, C.-H.; Lee, C.-W.;
Peng, C.-H.; Chen, J.-T. Solvent-Induced Dewettingon Curved
Substrates: Fabrication of Porous Polymer Nanotubes by Anodic
Aluminum Oxide Templates.Macromolecule 2015, 48, 6241–6250.
[CrossRef]
21. Chen, D.; Chen, J.-T.; Glogowski, E.; Emrick, T.; Russell,
T.P. Thin Film Instabilities in Blends under
CylindricalConfinement. Macromol. Rapid Commun. 2009, 30, 377–383.
[CrossRef] [PubMed]
22. Lee, C.-W.; Wei, T.-H.; Chang, C.-W.; Chen, J.-T. Effect of
Nonsolvent on the Formation of PolymerNanomaterials in the
Nanopores of Anodic Aluminum Oxide Templates. Macromol. Rapid
Commun. 2012, 33,1381–1387. [CrossRef] [PubMed]
23. Chen, J.-T.; Lee, C.-W.; Chi, M.-H.; Yao, I.-C.
Solvent-Annealing-Induced Nanowetting in Templates: TowardsTailored
Polymer Nanostructures. Macromol. Rapid Commun. 2013, 34, 348–354.
[CrossRef] [PubMed]
24. Huang, Y.-C.; Fan, P.-W.; Lee, C.-W.; Chu, C.-W.; Tsai,
C.-C.; Chen, J.-T. Transformation of Polymer Nanofibersto
Nanospheres Driven by the Rayleigh Instability. ACS Appl. Mater.
Interfaces 2013, 5, 3134–3142. [CrossRef][PubMed]
25. Chen, J.-T.; Wei, T.-H.; Chang, C.-W.; Ko, H.-W.; Chu,
C.-W.; Chi, M.-H.; Tsai, C.-C. Fabrication of PolymerNanopeapods in
the Nanopores of Anodic Aluminum Oxide Templates Using a
Double-Solution WettingMethod. Macromolecules 2014, 47, 5227–5235.
[CrossRef]
26. Tsai, C.-C.; Chen, J.-T. Rayleigh Instability in Polymer
Thin Films Coated in the Nanopores of AnodicAluminum Oxide
Templates. Langmuir 2013, 30, 387–393. [CrossRef] [PubMed]
27. Feng, X.; Jin, Z. Spontaneous Formation of Nanoscale Polymer
Spheres, Capsules, or Rods by Evaporation ofPolymer Solutions in
Cylindrical Alumina Nanopores. Macromolecules 2009, 42, 569–572.
[CrossRef]
28. Ko, H.-W.; Chi, M.-H.; Chang, C.-W.; Su, C.-H.; Wei, T.-H.;
Tsai, C.-C.; Peng, C.-H.; Chen, J.-T. Fabrication ofMulticomponent
Polymer Nanostructures Containing PMMA Shells and Encapsulated PS
Nanospheresin the Nanopores of Anodic Aluminum Oxide Templates.
Macromol. Rapid Commun. 2015, 36, 439–446.[CrossRef]
http://dx.doi.org/10.1116/1.4922157http://dx.doi.org/10.1039/C6NR05046Hhttp://dx.doi.org/10.1021/nl403144whttp://dx.doi.org/10.1038/ncomms7621http://dx.doi.org/10.1039/C4NR03002Hhttp://www.ncbi.nlm.nih.gov/pubmed/25132623http://dx.doi.org/10.1016/j.polymer.2007.11.013http://dx.doi.org/10.1039/C4RA14787Ahttp://dx.doi.org/10.1039/C2SM27026Ahttp://dx.doi.org/10.1002/pen.24403http://dx.doi.org/10.1039/c2ra21570ehttp://dx.doi.org/10.1039/C0SM00968Ghttp://dx.doi.org/10.1038/pj.2011.25http://dx.doi.org/10.1021/la401797fhttp://dx.doi.org/10.1021/acs.macromol.5b01207http://dx.doi.org/10.1002/marc.200800730http://www.ncbi.nlm.nih.gov/pubmed/21706613http://dx.doi.org/10.1002/marc.201200200http://www.ncbi.nlm.nih.gov/pubmed/22605615http://dx.doi.org/10.1002/marc.201200640http://www.ncbi.nlm.nih.gov/pubmed/23238887http://dx.doi.org/10.1021/am400029jhttp://www.ncbi.nlm.nih.gov/pubmed/23514621http://dx.doi.org/10.1021/ma500568jhttp://dx.doi.org/10.1021/la4041002http://www.ncbi.nlm.nih.gov/pubmed/24380368http://dx.doi.org/10.1021/ma8024283http://dx.doi.org/10.1002/marc.201400662
-
Polymers 2019, 11, 1039 17 of 18
29. Feng, X.; Mei, S.; Jin, Z. Wettability Transition Induced
Transformation and Entrapment of PolymerNanostructures in
Cylindrical Nanopores. Langmuir 2011, 27, 14240–14247.
[CrossRef]
30. Tsai, C.-C.; Chen, J.-T. Effect of the Polymer Concentration
on the Rayleigh-Instability-Type Transformationin Polymer Thin
Films Coated in the Nanopores of Anodic Aluminum Oxide Templates.
Langmuir 2015, 31,2569–2575. [CrossRef]
31. Chen, J.-T.; Zhang, M.; Russell, T.P. Instabilities in
Nanoporous Media. Nano Lett. 2007, 7, 183–187.
[CrossRef][PubMed]
32. Schlitt, S.; Greiner, A.; Wendorff, J.H. Cylindrical Polymer
Nanostructures by Solution Template Wetting.Macromolecules 2008,
41, 3228–3234. [CrossRef]
33. Chu, C.-J.; Cheng, M.-H.; Chung, P.-Y.; Chi, M.-H.; Jeng,
K.-S.; Chen, J.-T. Reversible Morphology Control
ofThree-Dimensional Block Copolymer Nanostructures by the Solvent
Annealing-Induced Wetting in AnodicAluminum Oxide Templates. Int.
J. Polym. Mater. 2016, 65, 695–701. [CrossRef]
34. Zhang, Z.; Ahn, D.U.; Ding, Y. Instabilities of PS/PMMA
Bilayer Patterns with a Corrugated Surface andInterface.
Macromolecules 2012, 45, 1972–1981. [CrossRef]
35. Fan, P.-W.; Chen, W.-L.; Lee, T.-H.; Chen, J.-T. Annealing
Effect on Electrospun Polymer Fibers and TheirTransformation into
Polymer Microspheres. Macromol. Rapid Commun. 2012, 33, 343–349.
[CrossRef][PubMed]
36. Tseng, H.-F.; Cheng, M.-H.; Jeng, K.-S.; Li, J.-W.; Chen,
J.-T. Asymmetric Polymer Particles with AnisotropicCurvatures by
Annealing Polystyrene Microspheres on Poly(vinyl alcohol) Films.
Macromol. Rapid Commun.2016, 37, 1825–1831. [CrossRef] [PubMed]
37. Mead-Hunter, R.; King, A.J.C.; Mullins, B.J. Plateau
Rayleigh Instability Simulation. Langmuir 2012, 28,6731–6735.
[CrossRef]
38. Fan, P.-W.; Chen, W.-L.; Lee, T.-H.; Chiu, Y.-J.; Chen,
J.-T. Rayleigh-Instability-Driven MorphologyTransformation by
Thermally Annealing Electrospun Polymer Fibers on Substrates.
Macromolecules 2012, 45,5816–5822. [CrossRef]
39. Wang, H.; Chang, T.; Li, X.; Zhang, W.; Hu, Z.; Jonas, A.M.
Scaled Down Glass Transition Temperature inConfined Polymer
Nanofibers. Nanoscale 2016, 8, 14950–14955. [CrossRef]
40. Tan, A.W.; Torkelson, J.M. Poly(methyl methacrylate)
Nanotubes in AAO Templates: Designing NanotubeThickness and
Characterizing the Tg-Confinement Effect by DSC. Polymer 2016, 82,
327–336. [CrossRef]
41. Li, L.; Zhou, D.; Huang, D.; Xue, G. Double Glass Transition
Temperatures of Poly(methyl methacrylate)Confined in Alumina
Nanotube Templates. Macromolecules 2014, 47, 297–303.
[CrossRef]
42. Higuchi, T.; Tajima, A.; Yabu, H.; Shimomura, M. Spontaneous
Formation of Polymer Nanoparticles withInner Micro-Phase Separation
Structures. Soft Matter 2008, 4, 1302–1305. [CrossRef]
43. Hou, P.; Fan, H.; Jin, Z. Spiral and Mesoporous Block
Polymer Nanofibers Generated in ConfinedNanochannels.
Macromolecules 2015, 48, 272–278. [CrossRef]
44. Chu, C.-J.; Chung, P.-Y.; Chi, M.-H.; Kao, Y.-H.; Chen,
J.-T. Three-Dimensional Block CopolymerNanostructures by the
Solvent-Annealing-Induced Wetting in Anodic Aluminum Oxide
Templates. Macromol.Rapid Commun. 2014, 35, 1598–1605. [CrossRef]
[PubMed]
45. Xue, J.; Xu, Y.; Jin, Z. Interfacial Interaction in Anodic
Aluminum Oxide Templates Modifies Morphology,Surface Area, and
Crystallization of Polyamide-6 Nanofibers. Langmuir 2016, 32,
2259–2266. [CrossRef][PubMed]
46. Ko, H.-W.; Cheng, M.-H.; Chi, M.-H.; Chang, C.-W.; Chen,
J.-T. Selective Template Wetting Routes toHierarchical Polymer
Films: Polymer Nanotubes from Phase-Separated Films via Solvent
Annealing.Langmuir 2016, 32, 2110–2116. [CrossRef] [PubMed]
47. Yan, N.; Sheng, Y.; Liu, H.; Zhu, Y.; Jiang, W. Templated
Self-Assembly of Block Copolymers and MorphologyTransformation
Driven by the Rayleigh Instability. Langmuir 2015, 31, 1660–1669.
[CrossRef]
48. Giussi, J.M.; Blaszczyk-Lezak, I.; Cortizo, M.S.; Mijangos,
C. In-situ Polymerization of Styrene in AAONanocavities. Polymer
2013, 54, 6886–6893. [CrossRef]
49. Sanz, B.; Blaszczyk-Lezak, I.; Mijangos, C.; Palacios, J.K.;
Müller, A.J. New Double-Infiltration Methodologyto Prepare PCL-PS
Core-Shell Nanocylinders inside Anodic Aluminum Oxide Templates.
Langmuir 2016, 32,7860–7865. [CrossRef]
50. Liu, Z. One-step Fabrication of Crystalline Metal
Nanostructures by Direct Nanoimprinting below MeltingTemperatures.
Nat. Commun. 2017, 8, 14910. [CrossRef]
http://dx.doi.org/10.1021/la2030632http://dx.doi.org/10.1021/la504901hhttp://dx.doi.org/10.1021/nl0621241http://www.ncbi.nlm.nih.gov/pubmed/17212461http://dx.doi.org/10.1021/ma071822khttp://dx.doi.org/10.1080/00914037.2016.1157801http://dx.doi.org/10.1021/ma2026836http://dx.doi.org/10.1002/marc.201100734http://www.ncbi.nlm.nih.gov/pubmed/22271584http://dx.doi.org/10.1002/marc.201600417http://www.ncbi.nlm.nih.gov/pubmed/27611838http://dx.doi.org/10.1021/la300622hhttp://dx.doi.org/10.1021/ma300964yhttp://dx.doi.org/10.1039/C6NR04459Jhttp://dx.doi.org/10.1016/j.polymer.2015.11.054http://dx.doi.org/10.1021/ma4020017http://dx.doi.org/10.1039/b800904jhttp://dx.doi.org/10.1021/ma501933shttp://dx.doi.org/10.1002/marc.201400222http://www.ncbi.nlm.nih.gov/pubmed/25098757http://dx.doi.org/10.1021/acs.langmuir.5b04569http://www.ncbi.nlm.nih.gov/pubmed/26886176http://dx.doi.org/10.1021/acs.langmuir.5b04746http://www.ncbi.nlm.nih.gov/pubmed/26831764http://dx.doi.org/10.1021/la504672xhttp://dx.doi.org/10.1016/j.polymer.2013.10.045http://dx.doi.org/10.1021/acs.langmuir.6b01258http://dx.doi.org/10.1038/ncomms14910
-
Polymers 2019, 11, 1039 18 of 18
51. Jin, K.; Torkelson, J.M. Tg-confinement Effects in Strongly
Miscible Blends of Poly(2,6-Dimethyl-1,4-Phenyleneoxide) and
Polystyrene: Roles of Bulk Fragility and Chain Segregation. Polymer
2017, 11, 85–96. [CrossRef]
52. Napolitano, S.; Glynos, E.; Tito, N. Glass Transition of
Polymers in Bulk, Confined Geometries, and NearInterfaces. Rep.
Prog. Phys. 2017, 80, 036602. [CrossRef] [PubMed]
53. Zhang, E.; Liu, Y.; Yu, J.; Lv, T.; Li, L. Fabrication of
Hierarchical Gecko-Inspired Microarrays Using aThree-dimensional
Porous Nickel Oxide Template. J. Mater. Chem. B 2015, 3, 6571–6575.
[CrossRef]
54. Zhang, E.; Wang, Y.; Lv, T.; Li, L.; Cheng, Z.; Liu, Y.
Bio-inspired design of hierarchical PDMS microstructureswith
tunable adhesive superhydrophobicity. Nanoscale 2015, 7, 6151–6158.
[CrossRef] [PubMed]
55. Jiang, W.; Lei, B.; Liu, H.; Niu, D.; Zhao, T.; Chen, B.;
Yin, L.; Shi, Y.; Liu, X. Fabrication of DirectionalNanopillars
with High-Aspect-Ratio by Stretching Imprint with Microcavity Mold.
Nanoscale 2017, 9,2172–2177. [CrossRef] [PubMed]
56. Spear, J.C.; Custer, J.P.; Batteas, J.D. The Influence of
Nanoscale Roughness and Substrate Chemistry o