Stabilization of Poly(methyl methacrylate) Nanofibers with Core–Shell

Feb 15, 2017 - The glass transition behaviors of poly(methyl methacrylate) (PMMA) nanofibers confined in pristine and surface-modified AAO templates a...
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Stabilization of Poly(methyl methacrylate) Nanofibers with Core− Shell Structures Confined in AAO Templates by the Balance between Geometric Curvature, Interfacial Interactions, and Cooling Rate Chen Zhang, Linling Li, Xiaoliang Wang,* and Gi Xue Key Laboratory of High Performance Polymer Materials and Technology of Ministry of Education, Department of Polymer Science and Engineering, School of Chemistry and Chemical Engineering, State Key Laboratory of Coordination Chemistry, Nanjing National Laboratory of Microstructures, Nanjing University, Nanjing 210093, P. R. China S Supporting Information *

ABSTRACT: The glass transition behaviors of poly(methyl methacrylate) (PMMA) nanofibers confined in pristine and surfacemodified AAO templates are investigated by differential scanning calorimetry (DSC) and broadband dielectric spectroscopy (BDS). During an ultraslow cooling process (0.1 K/min) across the Tg, two glass transition temperatures (Tg,low and Tg,high) are clearly identified by DSC and BDS, which correspond to the core and shell, respectively. The Tg,high originates from the transition of the adsorbed layer and is mainly dominated by the geometric curvature radius of the nanopores rather than the chemical nature of the wall surface. A dramatic change in the glass transition behaviors is detected when the cooling rate is changed from 40 to 0.1 K/min, which reflects the inherent evolution between the shell and the core through a nonequilibrium interlayer. Furthermore, by studying the system before and after surface modification of the nanopores by silanization, we suggest that such evolution could be sped up through the benefit of the stronger interfacial interactions. Our findings provide insight into achieving stable glassy polymer structures confined in nanopores by balancing the geometric curvature, interfacial interactions, and cooling rate.

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INTRODUCTION

few reports on achieving stable nanostructures, which is very important to stabilized functionalization. Nanoporous aluminum oxide (AAO) templates have been widely used to provide a constrained environment, as it contains arrays of parallel, cylindrical nanopores that are uniform in length and diameter.40 When polymers infiltrate into such uniform nanopores, different glassy behaviors are observed, and interesting hypotheses have been presented from exploring nanoconfined structures.5,7,15−18,41−45 The Tg and dynamics of confined polymers have been reported to be

Upon reducing the dimension or altering the nanostructure, polymers manifest abnormal physical behaviors compared to those of the bulk,1−18 which play a key role in the application of functional materials, nanoprinting, sensors, and electronic devices.19−25 Hence, the study of polymers under confinement is essential for the design of nanostructured materials. Pioneering works into polymers confined at the nanoscale have mainly focused on thin films.1−3,9−11,26−32 With an increase in the confinement dimension, more complicated and interesting applications arise. Therefore, the glass transition behavior of polymers under higher geometrical confinements has attracted increasing attention.33−39 However, there are still © XXXX American Chemical Society

Received: November 14, 2016 Revised: January 22, 2017

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rate of the adsorbed layer should be affected by the strength of the interfacial interactions.

suppressed, increased, or remain constant compared with the bulk value, depending on the diameter of the nanopores, the interfacial interactions, and the sample preparation. Using scattering techniques and DSC, Shin and Russell et al. first observed an intermolecular enhancement in the chain mobility with reduced entanglement but unchanged Tg for polystyrene (PS) nanofibers confined in AAO nanopores, when the pore size is smaller than the chain dimensions (2Rg) of PS.5 Duran et al. reported that the Tg of synthesized poly(γ-benzyl-Lglutamate) nanofibers in AAO templates was reduced by as much as 50 K.45 Almost at the same time, Martin and coworkers detected restricted segmental dynamics associated with the polymer−alumina interfacial region for confined semicrystalline poly(vinylidene fluoride) (PVDF).7 In general, a remarkable broadening of the distribution in relaxation times for polymers confined in AAO templates has been reported in previous relevant works.14,18,46 Hu et al. found a relationship between the reduced Tg and the molar mass of confined polymers that depends on the diameter of the nanopore.42 According to current studies, “size effects” and “interfacial effects” have mainly been suggested to account for the variations in the glass transition behavior of polymers confined in AAO nanopores. The geometric curvature radius of the nanopores has been found to be a main factor altering the degree of confinement, and interfacial interactions can be tuned by chemical modification of the hydroxylated pore walls. In our previous works,15−17,44 both nanotubes and nanofibers of polymers confined in AAO nanopores were investigated. Using fluorescence nonradiative energy transfer (NRET) methods, the interchain proximity and Tg of polymer chains in confined nanotubes were proven to depend on the curvature radius of the AAO nanopores.44 Double Tgs were observed for poly(methyl methacrylate) (PMMA)15 and poly(n-butyl methacrylate) (PBMA)16 nanofibers located inside AAO templates with pore diameters much larger than the their 2Rgs. A two-layer model was proposed in which the polymers in the adsorbed layer attached to the pore walls show an increased Tg due to interfacial interactions, while for the chains in core volume, a reduced or unchanged Tg was observed compared to the bulk value. Attempts to achieving a stable nanostructure by balancing the geometric curvature, interfacial interactions, and thermal annealing conditions are still being made. As Martin noted in a recent review, studies considering the combination of these factors for nanostructured polymers confined in AAO templates are hitherto scarce in the literature.47 In this work, we mainly aim to study the impacts of the geometric curvature and interfacial interactions on the segmental dynamics of PMMA nanofibers confined in AAO nanopores. During an ultraslow cooling process (0.1 K/min) across the Tg, the glass transition temperatures and dynamics of PMMA confined in pristine AAO templates with different diameters (25−250 nm) are investigated by differential scanning calorimetry (DSC) and broadband dielectric spectroscopy (BDS). In addition, the evolution of the core−shell structure of the confined polymer nanofibers is studied by cooling rate experiments (40−0.1 K/min), which reveals the dynamic exchange between chains in the adsorbed layer and core volume. Then, the surface of AAO is silanized by hexamethyldisiloxane (HMDSO). By comparing the results of DSC and BDS for PMMA samples confined in pristine and surface modified AAO templates, we believe that the geometric curvature defined by the pore size is a dominant factor affecting the variation in the Tgs in two-layer model, and the formation

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EXPERIMENTAL SECTION

Materials. A monodisperse PMMA sample with a number-average molecular weight of Mn = 6 kg/mol (stereoregular composition of 4% isotactic, 37% atactic, and 59% syndiotactic, polydispersity index of 1.1, and Rg of 2.5 nm) was purchased from Polymer Source Inc. (Dorval, Canada). As a classic material for nanotechnology applications, PMMA was chosen because it has been confirmed to form core−shell structures when confined in AAO templates. In addition, the interaction between PMMA and AAO nanopores can be easily modulated by altering the chemical nature of the wall surface. Anodic aluminum oxide (AAO) templates with a narrow distribution of pore sizes, which are open on both ends, were purchased from HeFei PUYUAN Nano Ltd. The thickness of the AAO templates is approximately 70 μm, and the average pore sizes are 25, 55, 130, and 250 nm, which are much larger than the Rg of PMMA. Before use, the templates were rinsed with ethanol and water and then dried at 150 °C for 12 h under vacuum to remove volatile impurities on the surface of the nanopores. Sample Preparation/Infiltration Procedure. PMMA was dissolved in toluene, and films with a thickness of approximately 100 μm were prepared by solution-casting onto clean cover glass substrates. After drying under ambient conditions for several days, the films were then dried under vacuum for 24 h at 150 °C. Then, the AAO template was sandwiched between two cover glasses with PMMA melt on the surface, and the whole system was maintained at 180 °C under vacuum (100 mbar) for 24 h to allow the PMMA melt to fill the nanopores via capillary forces. After slowly cooling to room temperature, excess bulk polymer on the surface of the AAO template was carefully removed with a surgical blade. The pristine surface of AAO nanopores is hydrophilic, but here we change it to hydrophobic by surface modification. Before modification, pristine AAO templates were cleaned with 30% H2O2 solution at 100 °C to ensure enough −OH groups are present for the subsequent modification reaction. Then, a liquid phase reaction was carried out in 10% HMDSO/ chloroform solution at room temperature for 24 h. Finally, the residual modifier was removed by washing with chloroform several times.48 Characterization. SEM Measurements. The nanostructure morphologies of the PMMA-filled AAO samples were recorded using a scanning electron microscope (HITACHIS-4800) with an acceleration voltage of 10 kV. To confirm the morphology of PMMA nanofibers, the alumina templates were dissolved in 5 wt % H3PO4 at room temperature.25 All samples were coated with several nanometers of Au before SEM measurement. Contact Angle Measurements. The surface contact angles of water on pristine and modified AAO templates were determined by an optical contact angle measuring device (OCA30, Dataphysics Instruments GmbH) at room temperature. Thermal Analysis. Calorimetric measurements were carried out by a Mettler-Toledo DSC1 STARe differential scanning calorimeter under dry nitrogen atmosphere (50 mL/min). Temperature and enthalpy calibrations were performed before the experiments using indium and zinc standards. An approximately 20 mg portion of the PMMA-filled AAO sample (2 mg bulk PMMA was used) was sealed in an aluminum pan (40 μL) for DSC measurement. The mass fractions of PMMA filled in the AAO templates with different pore sizes were determined by thermogravimetric analysis (TGA). The samples were heated from 25 to 700 °C at 10 K/min under a nitrogen atmosphere using a PerkinElmer TGA-Pyris system, and the results are shown in Figure S1. Then, the PMMA-filled AAO sample was heated from 30 to 180 °C at 10 K/min, and the sample reached a final thermal steady state after cooling to 30 °C at 0.1 K/min. The glass transition temperatures and ΔCp values were determined from the second heating thermographs. In addition, the cooling rate dependence was investigated for the PMMA-filled AAO samples. The samples with different pore diameters were first heated to 180 °C to eliminate the previous thermal history. After the cooling across the Tg at different cooling B

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Macromolecules rates (40, 20, 10, 5, 2, 1, and 0.1 K/min), subsequent heating curves were obtained at heating rates of 10 K/min. Broadband Dielectric Spectroscopy (BDS). Dielectric measurements were performed in a temperature range of 30−180 °C for bulk PMMA and the PMMA-filled AAO samples using a Novocontrol Alpha dielectric spectrometer (Concept80, Novocontrol Technologies GmbH & Co, KG) over a frequency range of 10−2−106 Hz at atmospheric pressure. For the PMMA-filled AAO samples, the template was capped between two gold-plated copper electrodes (20 mm diameter); for bulk PMMA, a PTFE ring (area of 59.69 mm2 and thickness of 0.5 mm) was used as a spacer. To erase the pervious thermal history, both the bulk PMMA and PMMA-filled AAO samples were first heated to 180 °C and then cooled to room temperature at an ultraslow cooling rate. Measurements were corrected for spacer capacitance, stray capacitance, and edge compensation.

other is much higher (denoted Tg,high). The overshoots accompanying the glass transition are enthalpy relaxation peaks, which appeared after slow cooling across the glass transition temperature and could be further used to prove the corresponding glass transition. The deviation between the two Tgs increases over the range of 30−60 °C as the pore diameter decreases from 250 to 25 nm (as shown in Figure 2b). Similar glass transition behaviors have been reported for small molecule systems by both experiments and simulations.49−54 However, few studies have shown a double glass transition behavior for polymers under nanoconfinement.15,55,56 A core−shell model can be proposed to interpret this phenomenon, which describes a strongly constrained, adsorbed layer with an increased Tg and a core volume with a bulk-like Tg. Figure 2b reveals the effect of pore diameter on the glass transitions of the confined PMMA nanofibers. An approximately linear relationship between the double Tgs and the geometric curvature radius d (d is the diameter of the AAO nanopores) is achieved. Clearly, the Tg,high increases with a decrease pore size, but there is no significant change to Tg,low. For the adsorbed layer, Napolitano et al. proposed that the local free volume at the buried polymer− substrate interfaces should influence the glass transition behavior of polymers under confinement.29,57 Because of the presence of pore walls with a certain geometric curvature, the polymers near the interface may exhibit strong spatial ordering and sharp density variations. In our previous work, using fluorescence nonradiative energy transfer (NRET) and calorimetric methods, the Tg of polymer chains confined in nanotubes was proven to depend on the curvature radius of the AAO nanopores. The interchain proximity of polystyrene nanotubes in AAO templates decreased with increasing geometric curvature radius d of the nanopores.44 For confined nanofibers, the average interchain proximity of the polymer was also detected to be closer compared with that of the bulk.16 Thus, we believe the trend in Tg,high should be connected to the geometric curvature of the substrate. For the PMMA chains in the adsorbed layer, increased deviation from the bulk behavior may arise from the closer interchain proximity or reduced interfacial free volume with a decrease in nanopore diameter. However, considering the PMMA in the core volume, the impact of the geometric curvature imposed by the hard wall can be ignored during ultraslow cooling across the Tg. Since the influence of the substrate on the glass transition temperature significantly decreases with increasing distance from the substrate,3,58−60 the chain behavior returns to that of the bulk. To further confirm the results obtained from the calorimetry measurements, complementary BDS studies were performed. Dielectric relaxations of PMMA confined in the AAO templates with different pore sizes were analyzed, and systematic differences were observed between the structural relaxations of the confined samples and that of the bulk. For all the spectra, the dielectric loss peaks shift toward higher frequencies with increasing temperature due to a decrease in viscosity and enhancement in mobility. This phenomenon is characteristic of thermally activated processes, as shown in Figure S2 (bulk PMMA and 55 nm PMMA-filled AAO samples, for example). The complex dielectric permittivity ε*(ω) = ε′ − iε″ is a function of frequency ω and temperature T, where ε′ is the real part and ε″ is the imaginary part. To determine the structural relaxation time, the obtained spectra were analyzed using the Havriliak−Negami (HN) function.61

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RESULTS AND DISCUSSION Because of the lower surface energy, the polymer melt tends to wet the surface of AAO spontaneously. By annealing at high temperature for a long enough time, the PMMA melt gradually infiltrates into the cylindrical nanopores, and polymeric nanofibers can be fabricated. Figure 1 shows the SEM

Figure 1. SEM micrographs of the PMMA-filled AAO templates from the top view: (a) 250, (b) 55, and (c) 25 nm. SEM micrographs of the same PMMA nanofibers after removing the AAO templates: (d) 250, (e) 55, and (f) 25 nm.

micrographs of the PMMA-filled AAO samples and the PMMA nanofibers after removing the alumina templates. As can be seen, in all cases, we obtained PMMA nanofibers with uniform diameters and good quality. Geometric Curvature Effect. The DSC technique was used to investigate the glass transition behavior of the constrained PMMA nanorods. All samples were first cooled form 180 to 30 °C at 0.1 K/min to ensure the complete formation of the core−shell structure and were then heated at 10 K/min. As shown in Figure 2, two distinct Tgs are detected for PMMA confined in pristine AAO nanopores: one is close to the bulk value at approximately 90 °C (denoted Tg,low), and the C

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Figure 2. (a) Normalized DSC heating traces of PMMA filled in the pristine AAO templates with different pore sizes. All samples were first cooled from 180 to 30 °C at 0.1 K/min to ensure the complete formation of the core−shell structure and then heated at 10 K/min. The DSC curves were normalized on the basis of the TGA results. (b) The double Tgs of confined PMMA as a function of nanopore diameter d.

Figure 3. Frequency dependence of the normalized dielectric loss ε″ for bulk PMMA and PMMA confined in pristine AAO templates with different diameters (25, 55, 130, and 250 nm) at (a) 105 °C and (b) 150 °C. The black circles represent the measurement data, the red solid lines represent the HN fitting curves, and the dashed lines represent the individual processes (the olive dashed lines are the conductivity contributions).

ε*(ω) = ε∞ +

Δε(T ) [1 + (iωτHN(T ))m ]n

Figure 3 shows the contributions of the different structural relaxations (α′, α, and β relaxation) for bulk and PMMA confined in AAO templates with different pore sizes in isothermal plots, with the respective HN fittings. In the frequency window studied, the dielectric response of the bulk sample is characterized by two main relaxation processes labeled α and β at temperatures slightly higher than Tg, such as 105 °C. At higher temperatures, such as 150 °C, the α- and β-relaxations are less separated in frequency and finally completely merge into one apparent relaxation peak.62 For confined PMMA, closer inspection of the dielectric spectra reveals the presence of another relaxation process, besides the dc conductivity, and merged structural relaxations (α and β relaxation) at high temperatures called α′ relaxation. Therefore,

(1)

where ε∞ is the high-frequency limit of the permittivity, Δε(T) is the relaxation strength related to the amount of effective dipoles involved in the relaxation process at a certain temperature, τHN(T) is the central relaxation time of the equation, m and n (m > 0, mn ≤ 1) are the shape parameters characterizing the symmetrical and asymmetrical broadening of the dielectric loss curve, and ω is the angular frequency. The relaxation time at maximum loss (τmax) is calculated by ⎛ πm ⎞ −1/ m⎛ πmn ⎞ τmax = τHN sin−1/ m⎜ ⎟ sin ⎜ ⎟ ⎝ 2(1 + n) ⎠ ⎝ 2(1 + n) ⎠

(2) D

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Figure 4. (a) Characteristic relaxation time τmax as a function of reciprocal temperature for α′, α, and β relaxations of bulk PMMA and PMMA confined in pristine AAO templates with different pore sizes. The dashed lines correspond to α′ VFT fittings, dash-dotted lines to α VFT fittings, and solid lines to β Arrhenius fittings. (b) The low-frequency slope m corresponds to the segmental relaxation of the bulk (black squares) and the confined PMMA in pristine AAO templates.

Table 1. VFT Parameters of α′ and α Relaxation for PMMAFilled AAO Templates

both the DSC and BDS measurements show two kinds of segmental dynamics for PMMA confined in AAO templates after ultraslow cooling. In the extreme case of the 25 nm sample, a remarkable α′ process appears at low frequencies, indicating greatly frustrated dynamics compared to the α relaxation of the bulk at the same temperature, and the usual α and β relaxations still can be observed in the high-frequencies regions. We ascribe the α′ relaxation to the partially immobilized polymer chains adsorbed to the pore walls. For the larger nanopore samples, such as 250 nm, although it is a challenge to accurately separate the α′ relaxation process from the primary data due to the dc contribution and low dielectric intensity, we can also detect a similar interfacial relaxation processes. Figure 4a shows the temperature dependence of the relaxation times for bulk PMMA and PMMA confined in the pristine AAO templates. The relaxation times of the interfacial α′ and segmental α relaxation processes follow the Vogel− Fulcher−Tammann equation (VFT): ⎛ DT0 ⎞ τmax = τ0 exp⎜ ⎟ ⎝ T − T0 ⎠

α relaxation bulk 250 nm 130 nm 55 nm 25 nm

α′ relaxation

D

T0 (K)

Tg (°C)

D

T0 (K)

Tg (°C)

6.1 3.6 1.8 1.3 0.9

297 310 329 333 334

78.9 82.0 82.3 82.4 79.8

2.0 2.1 2.0 2.7

355 355 357 358

114.0 116.6 119.5 129.6

with the results of the calorimetry experiment. Obviously, there are some discrepancies between the Tg values determined by DSC and BDS, which can result from the different types of molecular motions in thermal and dielectric relaxation.64 The fragility strength D and low-frequency slope m of the segmental relaxations of confined PMMA chains were found to be reduced compared to the bulk values. D could be associated with the cooperativity of the segmental relaxation of the polymers.65,66 A glass transition requiring higher cooperativity would depart strongly from Arrhenius behavior and exhibit lower D values. From the bulk to the 25 nm sample, the D of the α relaxation process gradually decreases from 6.1 to 0.9, which reflects an enhancement in the confinement effect due to the decreased pore size. Furthermore, at higher temperature, the segmental dynamics of PMMA chains in the core volume would suffer long-range effects from the adsorbed layer when confined in the AAO nanopores, which was not reflected in the DSC results above. The variations in m corresponding to the two segmental relaxations of PMMA confined in pristine AAO templates with varying pore diameters are shown in Figure 4b. Similar results were obtained by Floudas et al.18,46 As the pore size decreases, a reduction in m for both α′ and α relaxation processes is observed, which also provides strong evidence for our assumption. Compared to the bulk sample, confinement effects could broaden the distribution of the relaxation times of confined polymer chains not only in the adsorbed layer but also in the core volume, which may be connected with the slight broadening of the Tg regions observed in the DSC curves for PMMA confined in AAO nanopores. These results indicate the presence of different confined environments and possibly interactions between the PMMA dipoles at the time scales of

(3)

where τ0 is the relaxation time in the limit of very high temperature, D is a fitting parameter that reflects the fragility strength, and T0 is the “ideal” glass transition or Vogel temperature which is 30−70 K below the Tg.63 In contrast, the β relaxation times obey an Arrhenius temperature behavior. The β relaxation times of the bulk PMMA and confined PMMA samples are similar, which reveals that the β relaxation process is almost unaffected by nanoconfinement. However, for the α′ and α relaxation processes of the confined samples, lower D and higher T0 values are observed compared with those of the bulk. The fitting results are compiled in Table 1, and the glass transition temperatures for both the adsorbed layer (Tg,high) and core volume (Tg,low) are determined using the same definition: Tg = T(τ = 100 s). Tg,high, which is related to the structural relaxation dynamics of the molecules in the adsorbed layer, continuously shifts to a higher value with decreasing pore diameter. On the other hand, Tg,low is almost constant and stays around the value of the bulk Tg, and slight differences are possibly derived from experimental errors. These results agree E

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Figure 5. Comparison of the characteristic relaxation time for α′ and α relaxations of PMMA confined in pristine AAO (red) and modified AAO-R (blue) templates.

radius d, rather than the interfacial interactions, when confined in the AAO nanopores. However, according to BDS, the pronounced effect of interfacial interactions on the segmental dynamics of polymer in the adsorbed layer (α′ relaxation process) is revealed, as shown in Figure 5. The stronger interfacial interactions originate directly from the extensive hydrogen bonding between the pristine walls of the AAO template and the PMMA molecules, which should further suppress the α′ segmental dynamics compared to those of the AAO-R samples. However, this suppression effect decreases as the geometric curvature radius d of the nanopores increases, and no significant influence on the dynamics of polymer in the core volume (α relaxation process) is observed. Although altering the chemical nature of the wall surface imparts little change in the associated double Tg values (defined by DSC and BDS, respectively) of the confined PMMA nanofibers with stable core−shell structures, the interfacial interaction plays a very important role in the evolution of the core−shell structure as the cooling rate changes, which will be discussed in the following section. Cooling Rate Dependence. In our previous works, variations in the Tg and the conformations of confined polymers in AAO were linked to the thermal annealing conditions.15−17 However, the combined effect of interfacial interactions and cooling rate on the evolution of confined polymers with core−shell structures in AAO nanopores was underestimated. For polymer confined in AAO, we achieved a cooling rate of approximately 120 K/s36 by using liquid nitrogen, resulting in a glass with one single glass transition temperature around Tg,bulk,15 and two distinct Tgs were observed by ultraslow cooling across the Tg at 0.1 K/min. Then, the detailed evolution of the core−shell structure was monitored by DSC by examining the cooling rate across the Tg from 40 to 0.1 K/min, as shown in Figure 6. Dramatic changes in the glass transition behavior could be observed as the cooling rate across the Tg gradually decreases from 40 to 0.1 K/min.

the segmental processes, and the polymer chain relaxations in both fractions are retarded under confinement. Interfacial Modification. To investigate the effect of interfacial interactions on the dynamics of polymer confined in nanopores, the hydrophilic surface of the pristine AAO template was modified. The AAO templates modified with an alkylated pore surface are denoted AAO-R. The contact angles of water on the AAO templates before and after modification are shown in Figure S3 and Table SI. For the 55 nm AAO template, the contact angle changes from 32.0° to 90.8°, which demonstrates the successful transformation of the AAO surface chemistry from hydroxylation to alkylation. The diameter change caused by the surface modification can be neglected since the diameters of the nanopore are much larger than the thickness of the grafted alkylsilane. PMMA nanofibers confined in AAO-R were also obtained and investigated by DSC and BDS using the same protocol, and the corresponding results are shown in Figures S4−S6 and Table SII. Surprisingly, according to DSC, there are almost no differences between the glass transition temperatures of PMMA confined in pristine AAO and the modified AAO templates, as shown in Figure S4. In our previous research, the Tg,high of PBMA nanofibers confined in 300 nm pristine AAO disappeared when the surface of AAO was modified by trichloro(1H,1H,2H,2H-perfluorooctyl)silane (FOTS) (AAO-F).16 Here, we consider the order of the interactional strength between the polymer and the AAO surface: AAO-OH > AAO-R > AAO-F. The AAO-R surface can be considered to be a more weakly attractive substrate, while the repulsive force is dominant in AAO-F. Thus, we propose that with an attractive substrate the immobilization of confined polymer chains adjacent to the pore walls should not be essentially affected by the strength of the interfacial interaction. Polymer nanofibers with core−shell structures in this part were achieved during ultraslow cooling across the Tg, and the values of the double Tgs mainly depended on the geometric curvature F

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cooling rates such as 5−1 K/min, the PMMA nanofibers confined in 55 nm nanopores showed three Tg values. The previous broadened glass transition decomposed into two parts: the transition of polymer in core volume with Tg,low resembling the Tg of the bulk and the transition of polymers in the thermodynamic nonequilibrium interlayer between the adsorbed layer and the core volume. In addition, the heat capacity in the Tg,low and Tg,high regions gradually increased with decreasing cooling rate; meanwhile, that in the Tg,inter region continually decreased. However, at the ultraslow cooling rate of 0.1 K/min, the system has enough time to complete the chain transfer process, thus leading to the disappearance of the Tg,inter region and finally the formation of nanofibers with a core−shell structure. As a consequence, the DSC heating traces show two distinct Tg values with strong enthalpy relaxation peaks. A similar experimental phenomenon can be observed for PMMA confined in AAO and AAO-R with other pore diameters, as shown in Figures S7 and S8. To gain more information, we analyzed the results of the cooling rate experiments in depth, and the effect of the interfacial interaction began to unfold. Herein, we focus on the adsorbed layer of confined PMMA since it could intuitively reflect the evolution of the core−shell structure. As depicted in Figure 7a,b, as the cooling rate decreases, the Tg,high value gradually decreases, which can be interpreted as the change in the adsorbed polymer layer on the wall upon controlled thermal treatment.29 Interestingly, for the cooling rate of 40 K/ min, the Tg,high of PMMA confined in pristine 25 nm AAO is approximately 6 °C higher than that of the AAO-R sample with the same pore diameter. Nevertheless, this difference in value gradually approaches zero as the cooling rate decreases to 0.1

Figure 6. DSC heating traces of PMMA confined in 55 nm pristine AAO templates suffering different cooling rates across the Tg. The dashed and dotted lines are guides for the eye.

Two glass transitions were detected when fast cooling rates were applied, such as 40−10 K/min. Combined with the temperature derivative of the heat capacity curves shown in Figure S7b, it is quite clear that a primitive adsorbed layer with a Tg,high of approximately 150 °C had already formed. In addition, a broadened glass transition with a T g of approximately 105 °C was simultaneously observed as the major process, which indicates that most of confined PMMA chains are in a thermodynamic nonequilibrium state. We can describe this phenomenon using another “two-layer model”, in which the core volume is pristine and unstable. For moderate

Figure 7. Tg,high evolution of the (a) PMMA-filled AAO and (b) PMMA-filled AAO-R samples suffering different cooling rates across the Tg. In addition, the percent ΔCp in the interfacial layer, Φads, as a function of cooling rate for PMMA confined in (c) pristine AAO and (d) modified AAOR templates with different pore diameters. G

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Although the quantitative estimation of tads is difficult for PMMA confined in AAO templates, we suggest a qualitative description of the relationship between the evolution of the adsorbed layer and the cooling rate. When the cooling rate is relatively fast, such as 40−10 K/min, the growth of the adsorbed layer seems to be less pronounced in almost all samples with decreased cooling rate. The rapid thickening of the adsorbed layer is observed only if the cooling rate is reduced to an appropriate degree. Thermal treatment with longer annealing time corresponds to the time scale of the absorption transition, and dynamic molecular exchange at a certain temperature occurs in the thermodynamic nonequilibrium system. It is worth noting that there are clear differences in the growth rate of the adsorbed layer of the AAO and AAO-R samples. Compared to the AAO-R samples, the percent of the adsorbed layer in the pristine AAO samples increase faster, especially for the samples with the smallest nanopores, which indicates a faster growth rate of the adsorbed layer in the pristine AAO samples. Considering that the deviation in the α′ relaxation times between PMMA confined in pristine AAO and modified AAO templates increases as the geometric curvature radius d of the nanopores decreases, as shown in Figure 5, a tentative explanation can be proposed that a greater deviation in the relaxation times between the adsorbed layer and the core volume may provide a more powerful driving force for the growth of the adsorbed layer. In addition, as the pore size decreases, the interfacial effect should dominate due to the increased surface-to-volume ratio. Thus, the difference in the growth rate of the adsorbed layer in the smaller nanopores becomes even larger. Meanwhile, compared to the AAO-R samples, the percent of the adsorbed layer in the pristine AAO samples is slightly higher. This phenomenon can be explained by the long-range effect of interfacial interactions, which are also considered to be responsible for the distribution of the glass transition as a function of distance from the polymer layer to the substrate.60 Thus, it should be understood that a stronger interfacial interaction between the PMMA and AAO pore surface could propagate farther in the radial direction. The distinct content of the adsorbed layer, due to the different strength interfacial interactions, constantly accumulates and eventually appears during ultraslow cooling. In view of the above results, for polymer confined in AAO nanopores, we suggest that a smaller geometric curvature radius d and stronger interfacial interactions should be helpful for the formation of nanofibers with a core−shell structure.

K/min. In principle, the magnitude of the confinement effect on the Tg due to the attractive substrate depends on the strength of such interfacial interactions between the polymer and the substrate.34 Although the thin, primitive adsorbed layer achieved by fast cooling is effected, as the cooling rate decreases, the effect of the interfacial interaction on the chain conformation would decay as the thickness of the adsorbed layer increases. Similar results have been observed for PMMA nanotubes confined in oxide nanopores with different surface chemistry67 and PS films adsorbed on substrates with different strength interfacial interactions.32 The percent of adsorbed polymer confined in the nanopores can be estimated by assuming that the volume of the polymer in a given layer is proportional to the heat capacity of that layer. The percent of the adsorbed polymer layer can be calculated by eq 4: Φads =

ΔCp ,high ΔCp ,high + ΔCp ,inter + ΔCp ,low

(4)

In the two-layer model, ΔCp,inter in eq 4 can be taken as 0. The results of the calculations have been shown in Figure 7c,d. These results provide evidence for the evolution of the adsorbed layer and the inherent dynamic exchange among the polymer chains in the core volume, interlayer, and adsorbed layer. It is reasonable to assume that only a thin layer adjacent to the pore walls with surface-restricted molecules exists at high temperature. Because of the diversity of the polymer chain relaxations in different parts, the dynamic molecular exchange may occur.68,69 We can assume that this process should involve the polymer chains transferring from the center region toward the interfacial region. Therefore, the thickness of the adsorbed layer would increase as the temperature was reduced due to the larger cooperative motions. Therefore, the growth of the adsorbed layer is strongly dependent on the cooling rate. Based on the concept of “adsorption transition”, polymer chains would undergo an “adsorption transition” when encountering an impenetrable solid substrate only if the temperature is higher than the adsorption temperature, Ta.70−72 It follows that as a thermodynamic-like process, the adsorption transition is inevitably influenced by the treatment temperature. Considering that the dynamic molecular exchange processes also need to overcome certain energy barriers, as the temperature decreases, molecular migration from the center region to the interfacial region becomes irreversible, which induces the thickening of the adsorbed layer. In our previous work, we investigated the influence of annealing time on the Tg evolution of PMMA nanofibers confined in AAO templates, and equilibration of the polymer dynamics at the nanoscale was found to require a long time scale by thermal experiments.15 Similar conclusions have been reported by Napolitano et al. involving the time evolution of the thickness of an irreversibly adsorbed polystyrene (PS) layer on aluminum oxide.57 They defined a dimensionless parameter t* = tANN/tads (the ratio between the annealing time and the time scale of adsorption) to describe the exponential growth of the adsorbed layer with increasing annealing time.73 For t* ≪ 1 and t* ≫ 1, the polymer properties did not change with respect to annealing time. In contrast, for t* ∼ 1 the conformations of the polymer chains in the interfacial region changed continuously with annealing time as well as with other physical properties. In our work, a slower cooling rate corresponds to a longer annealing time (tANN) at temperatures higher than Ta.

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CONCLUSIONS In summary, we systematically investigated the effects of the geometric surface radius, interfacial interactions, and cooling rate on the glass transition behavior and evolution of the core− shell structure for PMMA nanofibers confined in AAO cylindrical nanopores. For confined PMMA, both thermal and dielectric experiments reveal a stable glassy structure, and two glass transition processes could be achieved during ultraslow cooling across the Tg. The Tg of the stable adsorbed layer substantially deviates from that of the bulk and essentially depends on the geometric curvature radius. An approximate linear relationship between the Tg,high and diameter d of the nanopores was observed. The formation of core−shell structures by the confined PMMA nanofibers was studied by cooling rate experiments. The physical absorption of polymer chains could account for the evolution of the adsorbed layer. H

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Macromolecules

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Furthermore, we also studied the glass transition behaviors of PMMA confined in surface modified AAO templates. Although the interfacial interaction between PMMA and the pore wall has no pronounced effect on the two Tgs after ultraslow cooling across the Tg, it may influence the formation of nanofibers with core−shell structures by altering the cooling rate. Thus, we suggest that the evolution of nanofibers with core−shell structure could be sped up through the benefit of the smaller geometric curvature radius d and stronger interfacial interactions. This work facilitates the understanding of the glass transition temperature and dynamics of confined polymer and provides an important theoretical guidance for achieving stable glassy polymer structures under uniform confinement by balancing the geometric curvature, interfacial interactions, and cooling rate.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b02469. Figures S1−S8 and Tables SI and SII (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (X.W.). ORCID

Xiaoliang Wang: 0000-0001-7820-4706 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors appreciate the financial support of the National Natural Science Foundation of China (51133002, 51673094, 21574063, 21274060, and 21404055). This work was also supported by the Program for Changjiang Scholars and Innovative Research Team in University (PCSIRT).

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