Molecular Mobility and Physical Aging of a Highly ... - ACS Publications

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Molecular Mobility and Physical Aging of a Highly Permeable Glassy Polynorbornene as Revealed by Dielectric Spectroscopy Huajie Yin,† Pavel Chapala,‡ Maxim Bermeshev,‡ Andreas Schönhals,† and Martin Böhning*,† †

Bundesanstalt für Materialforschung und−prüfung (BAM), Unter den Eichen 87, 12205 Berlin, Germany A.V. Topchiev Institute of Petrochemical Synthesis of Russian Academy of Science, Leninskii prospect, 29, 119991 Moscow, Russia

‡

S Supporting Information *

ABSTRACT: Polymeric membranes represent a cost- and energyefficient solution for gas separation. Recently superglassy polymers with high free volume outperform many conventional dense polymers in terms of gas permeability and selectivity. However, such polymers are prone to pronounced physical aging, resulting in a dramatic reduction in the gas permeability. Molecular mobility of polymer segments plays an important role in the physical aging and the gas transport performance of polymeric membranes. Molecular mobility and physical aging of a representative superglassy polynorbornene with very high gas permeability, PTCNSi2g, was monitored by using dielectric spectroscopy with state-of-the-art high-resolution analyzers. This work helps to shed some light on the structure−property relationship of superglassy polymers on a molecular level and to provide practical “design rules” for the development of high performance polymers for gas separation.

M

categories: chemical modification/cross-linking6,7 and mixed matrix materials.8,9 The influence of physical aging on the physical properties of polymeric membranes, like gas permeability,10 optical properties,11 and so on, were intensively studied. However, a deep understanding of local motions in glassy polymer membranes, considered being responsible for physical aging, and its relationship with gas transport properties is still missing. The gas transport in nonporous polymeric membranes, well described by the solution-diffusion model,12 is usually related to local motions of the polymer chains, which controls the presence of transient free volume elements enabling the diffusion of gas molecules as supported by fundamental transport models,13 molecular dynamics simulations,14 and experimental measurements.15 Gas separation performance is restricted by the well-known “Robeson upper-bound”.16,17 In recent years, especially superglassy polymers with microporous structure based on rigid chain structures, potentially capable of gas separation by molecular sieving, were reported to outperform this “upper bound”.18−21 Empirical conclusions were drawn that both the microporosity (i.e., a very high fractional free volume) and the local chain motion contribute to this outstanding separation performance. However, the correlation between the free volume distribution, local segment, and chain mobility and gas molecule diffusion based on a

embrane technology has emerged as a promising alternative to conventional energy intensive gas separation techniques such as amine capture of CO2 and cryogenic distillation for the removal of N2 from natural gas.1 It has been demonstrated to be a cost- and energy-efficient solution for gas separation. Polymeric membranes are being widely used for gas separation mainly because of their relatively low cost and easy processing into a hollow fiber configuration for large-scale industrial applications. Among them Sicontaining polymers have played an important role in the development of membrane materials for gas separation, e.g., polydimethylsiloxane (PDMS) and poly(1-trimethylsilyl-1propyne) (PTMSP).2 A new addition type poly(3,3-bis(trimethylsilyl)tricyclononene-7) (PTCNSi2g) with extremely high gas permeability was reported recently.3 This polymer belongs to the most permeable members of a novel class of membrane materials: Si-substituted addition polynorbornenes. These polymers have especially a great potential for separation of components of natural gases due to their solubility controlled selectivity for hydrocarbons. Like many other superglassy polymers with high/excess free volume (PTMSP,4 PIM-1,5 etc.), they are prone to pronounced physical aging, a continuous slow relaxation process toward thermodynamic equilibrium. The initial microporous structures approach a denser state via local segment and chain rearrangements that results in a dramatic reduction in the gas permeability. The practical utility of typical superglassy polymers for gas separation is limited by these aging effects. Many efforts have been made to suppress or control physical aging of such polymers. Approaches generally fall into two © XXXX American Chemical Society

Received: June 23, 2017 Accepted: July 17, 2017

813

DOI: 10.1021/acsmacrolett.7b00456 ACS Macro Lett. 2017, 6, 813−818

Letter

ACS Macro Letters

trimethylsilyl- (SiMe3-) side groups. No glass transition was detected by DSC before the onset of the thermal decomposition above 600 K under nitrogen atmosphere.3 Therefore, the α-relaxation related to the cooperative segmental motion cannot be observed by BDS in the temperature range under investigation. Moreover, there is a steep increase of the dielectric loss above 550 K with decreasing frequency and increasing temperature. The contribution to this higher dielectric loss is due to the dielectric processes as well as the conductivity, which is better revealed by plotting the real part of the complex conductivity versus frequency at different temperatures (see Figure S3 in the Supporting Information). The process in this region corresponds to the β*2-process at relatively low frequencies. The DC-conductivity (σdc) is assumed to originate from the mobile ionic impurities, which are often present in nonpolar, dielectric polymers like polyethylene24,25 in ppm concentrations and can drift under the influence of the electric field within the polymer matrix.23 In rigid polymers, with insufficient chain packing,26,27 its loose packing structure enables the charge carriers to move through the polymer matrix even when the segmental dynamics is slow or frozen.28 Thermal decomposition of PTCNSi2g in the temperature range of the dielectric investigation can be excluded, as confirmed by TGA (see Figure S1 in the Supporting Information). In Figure 2 the dielectric loss is plotted versus temperature at a frequency of 6098 Hz for different heating and cooling cycles.

molecular-level understanding has not been established in detail so far. Therefore, exploring the molecular mobility of such polymeric systems is of high importance. Such work helps to build up the connection between structure and properties in the scope of dynamics, which will contribute to providing practical “design rules” for the development of polymers with optimal gas transport properties. In this study, the molecular mobility of a novel highly gas-permeable polynorbornene-based membrane material with intrinsic microporosity, PTCNSi2g (see inset in Figure 1), was investigated by broadband dielectric

Figure 1. 3D representation of the dielectric results for the second heating run for PTCNSi2g. Inset: Chemical structure of PTCNSi2g.

spectroscopy (BDS). BDS is proved to be a suitable method to obtain detailed information regarding the dynamics for the new class of polymers.5,22 According to our previous work,5 the dielectric behavior of such polymers of high free volume tends to evolve toward a more stable state during the first several heating/cooling cycles due to changes of packing density−evidencing the occurrence of physical aging. Therefore, a protocol of three subsequent heating stages was used for in situ monitoring of physical aging. The detailed temperature protocol is given in Figure S2 in the Supporting Information. Figure 1 shows a three-dimensional representation of the dielectric measurement for the second heating run (2 H), i.e., the dielectric loss versus frequency and temperature. The dielectric response of PTCNSi2g is weak because its repeating unit carries no intrinsic dipole moment. The detected dielectric loss, nevertheless directly connected to molecular motions of the polymer matrix, is probably due to the presence of impurities in the polymer and, moreover, a small number of polar carbonyl groups formed by a slight degree of oxidation.23 The FTIR spectra indicated the presence of carbonyl groups in the polymer (Figure S4 in the Supporting Information). It should be noted that the magnitude of the dielectric loss (ε″) ranging from 10−4−10−3 is extremely low and underlying processes are only detectable with state-of-the-art highresolution BDS analyzers. Four dielectrically active processes, denoted as β-relaxation (β1-, β2-relaxation at lower temperatures) and β*-process (β*1-, β*2-processes at higher temperatures), are observed, each indicated by a peak in the dielectric loss. Multiple dielectric processes might point to a heterogeneous structure of PTCNSi2g caused by the introduction of two bulky

Figure 2. Dielectric loss vs temperature of PTCNSi2g at a frequency of 6098 Hz for different heating/cooling cycles, with arrows indicating β1-relaxation and β2-relaxation.

Two processes due to the localized fluctuations are observed and indicated as β1- and β2-relaxation. Both peaks are visible in first heating (1 H), first cooling (1 C), and second heating (2 H) runs, while only one peak at higher temperature is present in the second cooling (2 C) and third heating (3 H) runs. The decrease of dielectric loss and the elimination of a relaxation process are related to the thermal treatment of the film accompanying the measurement. Similar experimental results were reported for conventional polymers, like polycarbonate (PC),29 poly(ethylene terephthalate) (PET) and related polyesters,30 poly(methyl methacrylate) (PMMA),31 poly(vinyl ethylene) (PVE),32 and so on, in a less pronounced manner. One possible molecular origin is that the amplitude of angular fluctuations of dipoles attributed to the local relaxation processes in the glassy state is reduced.29 It should be further noted that a slight difference in the dielectric loss between 1 H and 1 C was revealed. As evaporation of residual solvent can be 814


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ACS Macro Letters

In addition to the β-relaxations, in the higher temperature region, β*-processes can be observed. For the ease of tracking the two peaks (β*1- and β*2-process), the dielectric loss is presented in the frequency domain in Figure 4 at T = 558 K.

practically excluded (see TGA, Figure S1 in the Supporting Information), this is most likely due to a structural change within the sample during the 1 H run up to 473 K. The spectra for the dielectric loss corresponding to 1 C and 2 H coincide in the temperature range of 303−473 K, indicating no remaining solvent and further structural change. In contrast, significant differences between the next two runs (1 C and 2 H) and the subsequent two runs (2 C and 3 H) were observed, indicating a further structural change of the film up to 593 K. For further analysis of the relaxation processes, the model function of Havriliak−Negami (eq 1) was fitted to the data. * (ω) = ε∞ + εNH

Δε (1 + (iωτHN)β )γ

(1)

ε∞ represents the real part ε′ with ε∞ = limω→∞ε′(ω), Δε is the dielectric strength, and τNH is the relaxation time corresponding to the frequency of maximal dielectric loss f max. β and γ are shape parameters, which describe the symmetric and asymmetric broadening of relaxation peaks. An example for that procedure is given in Figure S5a in the Supporting Information, where two HN functions are fitted to the data of the β1- and β2-relaxation of PTCNSi2g. Figure 3 gives the temperature dependence of the relaxation rate for both the β1- and β2-relaxation in the lower temperature

Figure 4. Dielectric loss vs frequency for PTCNSi2g at the temperature T = 558 K for the different heating/cooling cycles, with the arrows indicating the relaxation processes.

The intensity of the processes generally decreases with the measurement cycle as the whole spectra evolved to lower dielectric loss for further test cycles. An example of a HN-fit to the dielectric loss data at 558 K to determine the rate f max is given in Figure S5b in the Supporting Information. The relaxation map, including different measuring runs, is presented in Figure 5. The data for the 1 H and 1 C run are not

Figure 3. Relaxation rate f max,β vs inverse temperature in the low temperature range for the different measuring runs. The solid symbols correspond to β1-relaxation and the open symbols for β2-relaxation. The lines are fits of the Arrhenius equation to the corresponding data. The relevant activation energy was estimated. Figure 5. Rate f max,β* vs inverse temperature for PTCNSi2g in the high temperature region for the different measuring runs. The solid symbols correspond to the β*1-process, while the open symbols represent the β*2-process. The solid lines are fits of the Arrhenius equation to the corresponding data. The relevant activation energies are indicated.

region for different measuring cycles. The data sets corresponding to the two peaks obtained in 1 H, 1 C, and 2 H and the data sets for a single peak (β2-relaxation) in 2 C and 3 H processes located at higher temperatures obey the Arrhenius law33 given by eq 2. log fp = log f∞

ln 10EA − kBT

shown, as no β*-process could be observed below 473 K. For the 2 H run, one peak was present initially in the high frequency region (β*1-process), followed by the second peak at lower frequency (β*2-process). As shown in Figure 5, the temperature dependence of the rate f max is curved. This curvature is probably not due to a VFT behavior34−36 related to glassy dynamics. More likely its curvature is due to a continuous change in the sample structure during 2 H leading to a continuous slowing-down of the molecular fluctuations above 538 K, where the deviation from linear Arrhenius behavior becomes obvious. During the subsequent 2 C cycle,

(2)

EA is the activation energy, kB is the Boltzmann constant, and log f∞ denotes the limiting dielectric relaxation time at infinite temperature. The molecular origin of the β-relaxation is probably related to carbonyl groups (CO) due to minor oxidation of the SiMe3-group. The activation energy for all processes is estimated and given in the Supporting Information (Table S1). 815


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ACS Macro Letters f max versus 1/T follows an Arrhenius-like behavior also in the high-temperature region before it is no longer detectable below 550 K. Activation energies are shown in Table S2 in the Supporting Information. With further heating, 3 H, the peak for β*1-process was also too weak to be analyzed, in contrast to the β*2-process. Its rate follows an Arrhenius-like behavior with a slightly lower activation energy. It is worth mentioning that the activation energies for the two β*-processes of PTCNSi2g are extremely high (170−240 kJ/mol) compared to those of commonly observed local fluctuations in polymers (40−60 kJ/ mol). The observation of such high activation energies for β* needs further investigations with respect to its molecular mechanism: possibly a ring fluctuation or a small-angle jiggling of the cyclobutane ring, sterically restricted by the presence of the two bulky side SiMe3-groups in PTCNSi2g. Pure methyl group rotations could not be made responsible for fluctuations with such a high activation energy.37 Physical aging in polymer glasses is generally accompanied by changes in different physical properties (enthalpy, volume, etc.).38 It was demonstrated that the temperature dependence of the dielectric strength Δε follows structural changes in a sample.39 The correlation between the dielectric strength of the local relaxation and thermodynamic properties has been, for instance, described by Power et al. in a dielectric study of glassforming D-sorbitol.40 For further analysis, the dielectric strength Δε is obtained in addition to the relaxation rate from the fit of HN functions to the data. The Debye theory of dielectric relaxation generalized by Kirkwood and Fröhlich gives for the dielectric strength as a function of the temperature T (eq 3).33 Δε =

1 μ2 N g 3ε0 kBT V

Figure 6. Dielectric strength vs inverse temperature for β* process obtained from different measuring cycles: with solid symbols for β*1process and with open symbols for β*2-relaxation. The dashed lines are the guides to the eyes.

dependence compared to that of 2 H, different from the predicted decrease upon cooling per eq 3. This might indicate a weaker or perhaps ceased physical aging, but it has to be noted, that it is already difficult to analyze this process due to its small dielectric strength. On the other hand, this would be supported by the linear Arrhenius-like behavior shown in Figure 5. The open symbols in Figure 6 show the distinctly higher dielectric strength for the β*2-process. During 2 H, a similar temperature dependence of the dielectric strength to that of the β*1-process was observed. First, the dielectric strength is increasing with increasing temperature up to about 563 K, followed by an opposite temperature dependence until 588 K. During the subsequent second cooling run, the dielectric strength still decreases with decreasing temperature, reaching a constant value. In the third heating, the dielectric strength decreased to even lower values with slight dependence on the temperature indicating that there is still an ongoing collapse (i.e., physical aging) of the microporous structure. Evidently, the temperature dependence of dielectric strength is the result of an interplay between the thermal energy (kBT), causing a change in the number or the fluctuation angle of contributing dipoles and the densification-induced change in the volume (V), which are correlated with each other. However, the change in the temperature dependence of dielectric strength for PTCNSi2g is clearly related to an overall structural change of the sample during the measurement related to physical aging. In conclusion, a highly permeable glassy polymer (PTCNSi2g) was investigated by dielectric spectroscopy. βRelaxation (β1- and β1-relaxation) and β*-process (β*1- and β*2-process) have been observed, which showed time evolution of f max and dielectric strength. Such observations can be explained by the impact of structural relaxation on the dielectric response of PTCNSi2g. Activation energy for different processes further pointed to an overall densification (physical aging) of the polymer matrix. No α-relaxation was found, while conductivity due to ionic impurities was observed in the glassy state where the loose microporous packing structure enables the charge carriers to move through the polymer.

(3)

The Onsager factor is omitted for sake of simplicity, ε0 is the permittivity of the vacuum, N/V is the number density of dipoles involved in the relaxation process, and μ is the mean dipole moment of the process under consideration. g denotes the so-called Kirkwood−Fröhlich factor, which describes static correlation between the dipoles. Following eq 3, one would expect a decreasing Δε with increasing temperature, which is usually observed for αrelaxations in polymers. In contrast to that, generally Δε increases with increasing temperature for localized β-relaxations of bulk polymers. This behavior can be explained by a dominating increase of the number of contributing dipoles (N) in combination with a possible increase of the angular magnitude of the fluctuations with temperature. The Kirkwood−Fröhlich factor g will not change with temperature for polymers in the glassy state and the impacts of the thermal energy (kBT) and volume change are minor.33 At first, the dielectric strength Δε for the β*1-process increases with temperature as expected up to about 538 K (2 H; see Figure 6). In the same temperature region, the rate f max for the β*1-process follows a typical Arrhenius-like behavior, as indicated by the dashed line in Figure 5. In contrast, further heating above 538 K results in a decrease of Δε, which might be attributed to a decrease in the fluctuation angle and/or number density of dipoles involved in the β*1-process because of densification due to physical aging which overcompensates the impacts of both the structurally and thermally induced increase in the number density of dipoles (N/V). In the following cooling run (2 C), the dielectric strength for β*1-process remains at the lower value with negligible temperature 816


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Polymers of Intrinsic Microporosity. Macromolecules 2015, 48 (18), 6553−6561. (11) Huang, Y.; Paul, D. R. Physical Aging of Thin Glassy Polymer Films Monitored by Optical Properties. Macromolecules 2006, 39 (4), 1554−1559. (12) Wijmans, J. G.; Baker, R. W. The Solution-diffusion Model: A Review. J. Membr. Sci. 1995, 107 (1−2), 1−21. (13) Alentiev, A. Y.; Yampolskii, Y. P. Meares Equation and the Role of Cohesion Energy Density in Diffusion in Polymers. J. Membr. Sci. 2002, 206 (1−2), 291−306. (14) Hofmann, D.; Fritz, L.; Ulbrich, J.; Schepers, C.; Böhning, M. Detailed-atomistic Molecular Modeling of Small Molecule Diffusion and Solution Processes in Polymeric Membrane Materials. Macromol. Theory Simul. 2000, 9 (6), 293−327. (15) Inoue, R.; Kanaya, T.; Masuda, T.; Nishida, K.; Yamamuro, O. Relationship between the Local Dynamics and Gas Permeability of Para-Substituted Poly(1-chloro-2-phenylacetylenes). Macromolecules 2012, 45 (15), 6008−6014. (16) Robeson, L. M. The Upper Bound Revisited. J. Membr. Sci. 2008, 320 (1−2), 390−400. (17) Freeman, B. D. Basis of permeability/selectivity tradeoff relations in polymeric gas separation membranes. Macromolecules 1999, 32 (2), 375−380. (18) Du, N.; Park, H. B.; Robertson, G. P.; Dal-Cin, M. M.; Visser, T.; Scoles, L.; Guiver, M. D. Polymer Nanosieve Membranes for CO2capture Applications. Nat. Mater. 2011, 10 (5), 372−375. (19) Carta, M.; Malpass-Evans, R.; Croad, M.; Rogan, Y.; Jansen, J. C.; Bernardo, P.; Bazzarelli, F.; McKeown, N. B. An Efficient Polymer Molecular Sieve for Membrane Gas Separations. Science 2013, 339 (6117), 303−307. (20) Wang, Z.; Wang, D.; Zhang, F.; Jin, J. Tröger’s Base-Based Microporous Polyimide Membranes for High-Performance Gas Separation. ACS Macro Lett. 2014, 3 (7), 597−601. (21) Swaidan, R.; Ghanem, B.; Pinnau, I. Fine-Tuned Intrinsically Ultramicroporous Polymers Redefine the Permeability/Selectivity Upper Bounds of Membrane-Based Air and Hydrogen Separations. ACS Macro Lett. 2015, 4 (9), 947−951. (22) Konnertz, N.; Ding, Y.; Harrison, W. J.; Budd, P. M.; Schönhals, A.; Böhning, M. Molecular Mobility and Gas Transport Properties of Nanocomposites Based on PIM-1 and Polyhedral Oligomeric Phenethyl-silsesquioxanes (POSS). J. Membr. Sci. 2017, 529, 274−285. (23) van den Berg, O.; Sengers, W. G. F.; Jager, W. F.; Picken, S. J.; Wübbenhorst, M. Dielectric and Fluorescent Probes to Investigate Glass Transition, Melt, and Crystallization in Polyolefins. Macromolecules 2004, 37 (7), 2460−2470. (24) Huang, Y.; Jämbeck, J.; Unge, M. Understanding the Ionic Conduction in Dielectric Polymers at High Electric Fields Using Molecular Dynamics Simulations. ACS Macro Lett. 2017, 6, 571−574. (25) Sasabe, H.; Saito, S. Relationship between Ionic Mobility and Segmental Mobility in Polymers in the Liquid State. Polym. J. 1972, 3 (5), 624−630. (26) Kunal, K.; Robertson, C. G.; Pawlus, S.; Hahn, S. F.; Sokolov, A. P. Role of Chemical Structure in Fragility of Polymers: A Qualitative Picture. Macromolecules 2008, 41 (19), 7232−7238. (27) Stukalin, E. B.; Douglas, J. F.; Freed, K. F. Application of the Entropy Theory of Glass Formation to Poly(α-olefins). J. Chem. Phys. 2009, 131 (11), 114905. (28) Heres, M.; Cosby, T.; Mapesa, E. U.; Sangoro, J. Probing Nanoscale Ion Dynamics in Ultrathin Films of Polymerized Ionic Liquids by Broadband Dielectric Spectroscopy. ACS Macro Lett. 2016, 5 (9), 1065−1069. (29) Cangialosi, D.; Wübbenhorst, M.; Groenewold, J.; Mendes, E.; Schut, H.; van Veen, A.; Picken, S. J. Physical Aging of Polycarbonate Far below the Glass Transition Temperature: Evidence for the Diffusion Mechanism. Phys. Rev. B: Condens. Matter Mater. Phys. 2004, 70 (22), 224213. (30) McGonigle, E.-A.; Daly, J. H.; Jenkins, S. D.; Liggat, J. J.; Pethrick, R. A. Influence of Physical Aging on the Molecular Motion

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsmacrolett.7b00456. Additional supporting details (PDF).

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Andreas Schönhals: 0000-0003-4330-9107 Martin Böhning: 0000-0001-9753-345X Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors would like to thank T. Rybak, D. Neubert, and A. M. Elert for experimental help. REFERENCES

(1) Bernardo, P.; Drioli, E.; Golemme, G. Membrane Gas Separation: A Review/State of the Art. Ind. Eng. Chem. Res. 2009, 48 (10), 4638− 4663. (2) Yampolskii, Y. Gas and Vapor Transport Properties of SiContaining and Related Polymers. In Membrane Materials for Gas and Vapor Separation; Yampolskii, Y., Finkelshtein, E., Eds.; John Wiley & Sons, Ltd, 2017; pp 271−306. (3) Chapala, P. P.; Bermeshev, M. V.; Starannikova, L. E.; Belov, N. A.; Ryzhikh, V. E.; Shantarovich, V. P.; Lakhin, V. G.; Gavrilova, N. N.; Yampolskii, Y. P.; Finkelshtein, E. S. A Novel, Highly Gas-Permeable Polymer Representing a New Class of Silicon-Containing Polynorbornens As Efficient Membrane Materials. Macromolecules 2015, 48 (22), 8055−8061. (4) Dorkenoo, K. D.; Pfromm, P. H. Accelerated Physical Aging of Thin Poly[1-(trimethylsilyl)-1-propyne] Films. Macromolecules 2000, 33 (10), 3747−3751. (5) Konnertz, N.; Ding, Y.; Harrison, W. J.; Budd, P. M.; Schönhals, A.; Böhning, M. Molecular Mobility of the High Performance Membrane Polymer PIM-1 as Investigated by Dielectric Spectroscopy. ACS Macro Lett. 2016, 5 (4), 528−532. (6) Kelman, S. D.; Rowe, B. W.; Bielawski, C. W.; Pas, S. J.; Hill, A. J.; Paul, D. R.; Freeman, B. D. Crosslinking Poly[1-(trimethylsilyl)-1propyne] and its Effect on Physical Stability. J. Membr. Sci. 2008, 320 (1−2), 123−134. (7) Li, F. Y.; Xiao, Y.; Chung, T.-S.; Kawi, S. High-Performance Thermally Self-Cross-Linked Polymer of Intrinsic Microporosity (PIM-1) Membranes for Energy Development. Macromolecules 2012, 45 (3), 1427−1437. (8) Lau, C. H.; Nguyen, P. T.; Hill, M. R.; Thornton, A. W.; Konstas, K.; Doherty, C. M.; Mulder, R. J.; Bourgeois, L.; Liu, A. C. Y.; Sprouster, D. J.; Sullivan, J. P.; Bastow, T. J.; Hill, A. J.; Gin, D. L.; Noble, R. D. Ending Aging in Super Glassy Polymer Membranes. Angew. Chem., Int. Ed. 2014, 53 (21), 5322−5326. (9) Bushell, A. F.; Budd, P. M.; Attfield, M. P.; Jones, J. T. A.; Hasell, T.; Cooper, A. I.; Bernardo, P.; Bazzarelli, F.; Clarizia, G.; Jansen, J. C. Nanoporous Organic Polymer/Cage Composite Membranes. Angew. Chem., Int. Ed. 2013, 52 (4), 1253−1256. (10) Swaidan, R.; Ghanem, B.; Litwiller, E.; Pinnau, I. Physical Aging, Plasticization and Their Effects on Gas Permeation in “Rigid” 817


Letter

ACS Macro Letters and Structural Relaxation in Poly(ethylene terephthalate) and Related Polyesters. Macromolecules 2000, 33 (2), 480−489. (31) Boucher, V. M.; Cangialosi, D.; Alegrìa, A.; Colmenero, J.; González-Irun, J.; Liz-Marzan, L. M. Physical Aging in PMMA/silica Nanocomposites: Enthalpy and Dielectric Relaxation. J. Non-Cryst. Solids 2011, 357 (2), 605−609. (32) Casalini, R.; Roland, C. M. Aging of the Secondary Relaxation to Probe Structural Relaxation in the Glassy State. Phys. Rev. Lett. 2009, 102 (3), 035701. (33) Schönhals, A. Molecular Dynamics in Polymer Model Systems. In Broadband Dielectric Spectroscopy, Kremer, F., Schönhals, A., Eds.; Springer: Berlin, 2003; pp 225−293. (34) Vogel, H. The Law of the Relation Between the Viscosity of Liquids and the Temperature. Phys. Z. 1921, 22, 645. (35) Fulcher, G. S. Analysis of Recent Measurements of the Viscosity of Glasses. J. Am. Ceram. Soc. 1925, 8 (6), 339−355. (36) Tammann, G.; Hesse, W. Die Abhängigkeit der Viscosität von der Temperatur bei unterkühlten Flüssigkeiten. Z. Anorg. Allg. Chem. 1926, 156 (1), 245−257. (37) Schönhals, A.; Schick, C.; Huth, H.; Frick, B.; Mayorova, M.; Zorn, R. Molecular Dynamics in Glass-forming Poly(phenyl methyl siloxane) as Investigated by Broadband Thermal, Dielectric and Neutron Spectroscopy. J. Non-Cryst. Solids 2007, 353 (41), 3853− 3861. (38) Hodge, I. M. Physical Aging in Polymer Glasses. Science 1995, 267 (5206), 1945−1947. (39) Boucher, V. M.; Cangialosi, D.; Alegria, A.; Colmenero, J.; Gonzalez-Irun, J.; Liz-Marzan, L. M. Accelerated Physical Aging in PMMA/Silica Nanocomposites. Soft Matter 2010, 6 (14), 3306−3317. (40) Power, G.; Johari, G. P.; Vij, J. K. Relaxation Strength of Localized Motions in D-sorbitol and Mimicry of Glass-softening Thermodynamics. J. Chem. Phys. 2003, 119 (1), 435−442.

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