Molecular Mobility of the High Performance Membrane Polymer PIM-1

Letter pubs.acs.org/macroletters

Molecular Mobility of the High Performance Membrane Polymer PIM‑1 as Investigated by Dielectric Spectroscopy Nora Konnertz,† Yi Ding,† Wayne J. Harrison,‡ Peter M. Budd,‡ Andreas Schönhals,† and Martin Böhning*,† †

Bundesanstalt für Materialforschung und−prüfung (BAM), Unter den Eichen 87, 12205 Berlin, Germany School of Chemistry, University of Manchester, Manchester M13 9PL, U.K.

‡

S Supporting Information *

ABSTRACT: The increasing demand for energy efficient separation processes fosters the development of new high performance polymers as selective separation layers for membranes. PIM-1 is the archetypal representative of the class of polymers of intrinsic microporosity (PIM) which are considered most promising in this sector, especially for gas separations. Since their introduction, PIMs stimulated a vast amount of research in this field and meanwhile evolved to the state-of-the-art in membrane technology for gas separation. The major obstacle for extending the practical membrane application is their strong tendency to physical aging. For the first time, investigations by broadband dielectric spectroscopy (BDS) addressing molecular dynamics and conductivity in PIM-1 are presented. As chain packing during film formation from the casting solution and physical aging are key factors determining the separation performance of PIMs as membrane materials, characterization of the molecular mobility in such materials as revealed by BDS will provide valuable information for further development and optimization.

M

For conventional glassy polymers, a correlation between certain molecular motions of the solid polymeric matrix and the diffusion of small molecules within can be observed. This follows from fundamental transport models10,11 and molecular dynamics simulations12,13 and is also discussed based on experimental data.14,15 For the superglassy high free volume polymers such a correlation no longer holds as the transport mechanism is obviously different. Distinct differences are found in terms of the corresponding activation energies.16,17 This is also in agreement with a continuous free volume phase, found, e.g., in detailed atomistic molecular dynamics simulations of such polymers.18,19 An investigation addressing the molecular mobility of PIMs is of great importance, as two major phenomena which determine the practical performance in membrane application, i.e., the physical aging and the plasticization induced by highly sorbing penetrants, are directly related. Furthermore, the film formation during casting, i.e., the solidification of the polymer by solvent evaporation, is predominantly governed by the molecular mobility of the polymer matrix. Thus, systematic studies of various PIMs as well as corresponding nanocomposites will substantially contribute to further developments toward practical applications of PIMs in membrane

embrane technology is widely considered one of the key technologies for improving the energy efficiency of separation processes in the chemical industry as well as in energy supply. Especially for gas separation membranes, polymer materials are still the material of choice for the active separation layer due to their easy processability. In recent years, research in this area was stimulated particularly by a new class of glassy materials, the so-called polymers of intrinsic microporosity (PIM), introduced by Budd and McKeown.1,2 While previously studied polymers with very high fractional free volume and extremely high gas permeabilities, e.g., polyacetylenes such as PTMSP, show only poor selectivities, most PIMs exhibit an attractive combination of almost as high permeabilities with reasonable permselectivities. Therefore, PIMs today represent the current state-of-the-art in air separation and hydrogen recovery.3,4 All these superglassy polymers share a distinct tendency to physical aging, which is the major drawback with regard to practical membrane applications. After formation of the dense film (or thin selective layer)sometimes followed by a nonsolvent treatment in order to exchange residual casting solventthe initial permeability decreases significantly with time.5,6 Therefore, one of the current research topics in the field aims at the suppression of physical aging either by optimizing the chemical architecture of the polymers or by adding suitable (nano)fillers which stabilize their structural arrangement in the glassy state.5,7−9 © XXXX American Chemical Society

Received: March 15, 2016 Accepted: April 5, 2016

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DOI: 10.1021/acsmacrolett.6b00209 ACS Macro Lett. 2016, 5, 528−532

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ACS Macro Letters

part, i = √−1) in a frequency range from 10−1 to 106 Hz. The samples were measured in parallel plate geometry. In order to ensure a good electrical contact, gold electrodes with a diameter of 20 mm were evaporated on both sites of the sample film. The measurements were carried out by isothermal frequency scans at the selected temperatures. A detailed temperature program with several heating and cooling cycles in the range of 173−523 K (−100−250 °C) was used (ΔT = 3 K) to determine the influence of the temperature treatment on structure and dynamics of the sample. The temperature was controlled by a Novocontrol Quatro cryo-system with stability better than 0.1 K. For more details see reference 27. During the whole temperature program the sample was kept in a dry nitrogen atmosphere. The relationship between the complex dielectric function ε* and the complex conductivity σ* is given by eq 1.

separationsespecially as BDS is complementary to investigations using X-ray scattering20,21 or addressing electronic properties of PIMs.22 In this study the first synthesized and most representative PIM-1 (see Figure 1) is investigated by broadband dielectric spectroscopy (BDS).

Figure 1. Chemical structure of the repeat unit of PIM-1.

PIM-1 was synthesized according to ref 23. SEC in chloroform against polystyrene standards gave Mw = 82800 g/mol and a polydispersity index of PDI = 2.8 (see also Supporting Information). The casting solution was prepared by dissolving PIM-1 in chloroform (1 g of PIM-1 per 12 mL of CHCl3) by shaking for 4−5 h, filtered using a 5 μm PTFE-filter, and finally treated in an ultrasonic bath for 5 min. The solution was then cast into a round Teflon mold of 8 cm diameter which was placed in a closed chamber saturated with chloroform vapor in order to reduce the solvent evaporation from the film. Overnight, a solid film had formed at room temperature. It was removed from the Teflon mold and stored for another 2 days in the vapor chamber at room temperature. Afterward the film was dried for 3 days at 75 °C in a vacuum oven (see below). It has to be noted that for PIM-1 very different casting/ drying protocols can be found in the literature, which lead to different states of the solid film (as revealed, e.g., by different gas permeabilities). Drying temperatures in the range from 40 °C up to 100 °C have been reported5,24,25the strong tendency to physical aging of PIMs suggests as low as possible temperatures as long as complete solvent removal is ensured. Furthermore, it is known that a methanol treatment for solvent exchange results in lower packing density (and high gas permeabilities), while contact with water (e.g., for removing the film from the casting plate) has an opposite effect. In order to obtain a representative PIM-1-film, a protocol without methanol treatment and without direct water contact was chosen. An optimal annealing temperature was determined by several casting/drying steps (temperature, time) with subsequent thermogravimetric analysis (TGA), respectively. After drying for 1 day in vacuum at 40 °C in the TGA, a remaining mass loss of about 2.6% at 200 °C was observed, which after 5 days at 40 °C was still 1.3%. For the second series of films, temperatures above the boiling point of the solvent used (chloroform, bp = 61 °C) were chosen: 75 and 100 °C. After 3 days, at 75 °C there was a mass loss at 200 °C of 0.44% in the TGA, which was the same as for 1 day at 100 °C. After 3 days at 100 °C only a further reduction to 0.33% was achieved. Therefore, 3 days at 75 °C was chosen for the film used in this study. Selected TGA curves are given in the Supporting Information (Figure S3). The molecular mobility was estimated by broadband dielectric spectroscopy (BDS). A high-resolution ALPHA analyzer interfaced to an active sample head (Novocontrol, Montabaur, Germany) was used to measure the complex dielectric permittivity ε*( f) = ε′( f) − iε″( f) ( f - frequency, ω angular frequency: ω = 2πf, ε′ - real part, ε″ - loss or imaginary

σ *(ω) = σ ′(ω) + iσ ″(ω) = iωε0ε*(ω)

(1)

where σ′ and σ″ are the real and imaginary part of σ*. They are defined as σ ′(ω) = ωε0ε″(ω)

(2)

and σ ″(ω) = ωε0ε′(ω)

(3)

ε0 is the permittivity of free space. In order to characterize the dielectric behavior and also consider reversible and irreversible effects on molecular mobility due to changes of packing density induced by heating and/or sorbed water or remaining traces of solvent, a protocol of two subsequent heating stages was used (depicted in Figure 2). The first heating starts at room

Figure 2. Heating/cooling cycles of the dielectric measurements on PIM-1. The step width between individual measurements was increased below room temperature due to available liquid nitrogen capacity for cooling.

temperature (298 K/35 °C) and ends at 473 K (200 °C), while the second heating starts at 173 K (−100 °C) and goes up to 523 K (250 °C). During cooling between the two heating runs, also BDS measurements were performed at the same temperatures. It has to be noted that for PIM-1, as for PIMs in general, no thermal glass transition is observed below the decomposition at about 640 K (370 °C),1 so we do not expect to find a related dynamic glass transition (segmental dynamics, α-relaxation) in the applied temperature range. In Figure 3 the temperature-dependent dielectric loss log ε″ at a fixed frequency of 1000 Hz is plotted for the two heating runs and the cooling run between them. A significant difference between first heating and the two subsequent runs is revealed. The dielectric loss curve of the first heating shows one distinct peak and a shoulder. The peak at higher temperatures, around 460 K (187 °C), can be assigned to a molecular relaxation 529


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film. For further analysis of the relaxation process, the Havriliak−Negami function (eq 4) was fitted to the data. * (ω) = ε∞ + εHN

Δε (1 + (iωτHN)β )γ

(4)

ε∞ represents the real part ε′ with ε∞ = limω→∞ε′(ω); Δε is the dielectric strength; and τHN 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. Conductivity effects are treated in the usual way by adding a power law to the dielectric loss: ε″cond = σ0/ωs·ε0; here ε0 is the permittivity of free space, and σ0 is related to the specific DC conductivity of the sample. The parameter s (0 < s ≤ 1) describes for s = 1 Ohmic and for s < 1 non-Ohmic effects in the conductivityfor details see refs 27 and 28. An example is given in Figure S4 in the Supporting Information. In Figure 5 the determined frequencies of maximal dielectric loss are plotted in Arrhenius coordinates. At first glance, the

Figure 3. Comparison of dielectric spectra (log ε″ vs temperature) at a fixed frequency of f = 1000 Hz for the different heating and cooling runs.

process which is found similarly during the first cooling and second heating. It has to be noted that the dielectric loss in the first heating run, up to about 400 K (127 °C), is in general significantly higher than in the subsequent measurements. This is probably caused by the remaining solvent molecules. Together with the adumbrated shoulder around 313 K (40 °C) this indicates rather the evaporation process than a higher molecular mobility due to plasticization by traces of solvent. Also the influence of water, sorbed from the atmosphere, might be taken into account, for which a strong effect on the gas transport properties due to partial blocking of the accessible free volume is well-known.16,26 The curves of the dielectric loss measured during first cooling and second heating are almost identicalindicating that no further changes occur in this temperature rangeand also show the distinct relaxation peak around 460 K (187 °C). The observed onset at lower temperature points at a second relaxation peak below 200 K (−73 °C), but this can unfortunately not be fully characterized within the temperature range accessible for the BDS measurements. Figure 4 shows a three-dimensional representation of the BDS measurement, i.e., the dielectric loss log ε″ versus frequency log f and temperature T. The peak of the molecular relaxation of PIM-1 is clearly seen, and as expected it shifts to higher frequencies with increase of temperature. The increase of dielectric loss with decreasing frequency (and increasing temperature) is due to conductivity, which is related to the drift motion of charge carriers in the

Figure 5. Plot of the frequency of maximal dielectric loss f max,β* in Arrhenius coordinates. Inset: DC conductivity σDC for PIM-1 in dependence of the inverse temperature. The lines are fits of the Arrhenius equation to the data.

linear behavior seems to indicate a localized molecular relaxation process, usually denoted as β-relaxation, which is typically found for most polymers. In contrast to that, cooperative segmental relaxation processes usually follow a Vogel−Fulcher−Tammann (VFT) law, distinctly deviating from the linear Arrhenius behavior. An activation energy of EA = 86.1 kJ/mol is obtained for the PIM-1 under investigation. This value is relatively high for commonly observed local βprocesses in polymers. For that, usually activation energies between 40 and 60 kJ/mol are expected. A similar phenomenon was also found for example in poly(ethylene naphtalate) (PEN). The activation energy of the β*-relaxation (EAβ*) for PEN is in the similar range as the value obtained for PIM-1. For PEN this is ascribed to the formation of intermolecular sandwich-like agglomerates of aromatic moieties of the polymer chain with strong interaction of the respective π-systems29−31 This leads to higher activation energies and a more cooperative character of the molecular relaxation process. Moreover, recently a similar behavior was found for Matrimid, a commercially available polyimide also considered as potential material for gas separation membranes.32 Although for PIMs generally a limited molecular mobility is assumed due to their rigid structure, which leads to the low packing density and the exceptional gas transport properties, a

Figure 4. 3D representation of the dielectric loss log ε″ versus frequency and temperature for the second heating run. 530


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from yellow to brownish. A comparison of FT-IR spectra for the virgin and the BDS-measured sample (Figure S6) does not show any changes. Furthermore, the strongly annealed material regains its original yellow color after being redissolved in chloroform. So, no indications for changes in the chemical structure were found, and an only reversible change of the amorphous packing in the solid film may be concluded. In conclusion, a molecular relaxation process with Arrhenius behavior and an unusually high activation energy, denoted β*, has been observed in PIM-1, together with a significant conductivity in the glassy state. As expected, no α-relaxation (related to a glass transition) was found. Both the β*-relaxation as well as the conductivity are explained with the formation of local intermolecular agglomerates due to interaction of πelectrons in aromatic moieties of the polymer backbone (π−πstacking). Up to now, this has not been taken into account when film formation, chain dynamics, free volume, and especially gas transport properties of PIMs were discussed. Although the reported findings have to be further investigatedalso considering more PIMs in a systematic manner they represent a new important aspect of this innovative class of polymers.

somewhat similar reason for the higher activation energy should be considered. Preliminary wide-angle X-ray measurements of our film sample show several broad reflections which evidence such aggregates with a molecular spacing of about 3.4 and 10 Å33 (see Figure S7 in the Supporting Information), which is in agreement with earlier investigations of PIM-1 reported in refs 20 and 21. Due to the resulting partial cooperative character, the observed molecular relaxation process of PIM-1 is denoted β*. For a further analysis of the dielectric behavior of PIM-1, also the conductivity is consideredespecially because of the postulated π−π-stacking of aromatic moieties of the polymer chain, which may have a distinct influence on the conductivity mechanism. The frequency dependence of the real part of the complex conductivity (Figure 6) shows the typical behavior expected for

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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/acsmacrolett.6b00209. Details of synthesis and characterization of PIM-1 as well as photographs and FT-IR spectra characterizing the changes of the PIM-1 film during the BDS measurements (PDF)

Figure 6. Frequency dependence of the real part of the complex conductivity σ′ at the indicated temperatures.

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semiconducting polymeric materials. For high frequencies the real part σ′ decreases with decreasing frequency with a power law down to a characteristic frequency fc, where a plateau value is reached. The plateau value corresponds to the DC conductivity.27 As can be seen from Figure 6, this DC conductivity increases with increasing temperature. The inset of Figure 5 shows a linear behavior for the plateau values σDC in dependence of the reciprocal temperature; i.e., the DC conductivity follows an Arrhenius relation. An apparent activation energy of EA,σDC = 101 kJ/mol was estimated, which is significantly higher than the activation energy of the β*relaxation. It is necessary to note that the conductivity contribution was observed in the glassy state. For most conventional polymers the transport mechanism of charge carriers is due to mobility of charge carriers (impurities) connected with segmental dynamics of the polymer matrix above the glass transition, and therefore the DC conductivity follows VFT-behavior. For PIM-1 no glass transition is observed before decomposition. Therefore, it is concluded that for PIM-1 the conductivity is not directly related to segmental dynamics. As discussed above wide-angle X-ray scattering measurements show evidence of agglomerates probably of stacked phenyl groups by π−π interaction. Due to the overlapping π-systems charge transport in PIM-1 is enhanced. It has to be noted that the PIM-1 sample shows a distinct change in color after the temperature cycle of the BDS measurement (Figure 2). As can be seen from the photographs in Figure S5 (see Supporting Information), the color changes

AUTHOR INFORMATION

Corresponding Author

*Tel.: +49 30 8104-1611. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors would like to thank T. Rybak and D. Neubert for experimental help and S. Wellert (TU-Berlin) for the WAXD measurements. N.K. is grateful for support of the PhD-program of BAM. W.J.H. is supported by EPSRC grant EP/K016946/1.

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Molecular Mobility of the High Performance Membrane Polymer PIM-1

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