Influence of Trimethylsilyl Side Groups on the Molecular Mobility and

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Influence of Trimethylsilyl Side Groups on the Molecular Mobility and Charge Transport in Highly Permeable Glassy Polynorbornenes Huajie Yin, Pavel Chapala, Maxim Bermeshev, Brian Pauw, Andreas Schönhals, and Martin Böhning ACS Appl. Polym. Mater., Just Accepted Manuscript • DOI: 10.1021/acsapm.9b00092 • Publication Date (Web): 08 Mar 2019 Downloaded from http://pubs.acs.org on March 9, 2019

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ACS Applied Polymer Materials

Influence of Trimethylsilyl Side Groups on the Molecular Mobility and Charge Transport in Highly Permeable Glassy Polynorbornenes Huajie Yin†, Pavel Chapala‡, Maxim Bermeshev‡, Brian R. Pauw†, 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 KEYWORDS Polynorbornene, Dielectric Spectroscopy, Molecular Mobility, Conductivity, Gas Separation Membranes

ACS Paragon Plus Environment

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ACS Applied Polymer Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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ABSTRACT: Superglassy polymers with large fractional free volume emerge as novel materials with a broad range of applications, especially in the field of membrane separations. Highly permeable addition-type substituted polynorbornenes with high thermal resistance and chemical stability are among the most promising materials. The major obstacle for extending the practical membrane application is their strong tendency to physical aging, leading to a partial decline in their superior transport performance over time. In the present study, broadband dielectric spectroscopy with complementary X-ray scattering techniques were employed to reveal changes in microporous structure, molecular mobility and conductivity by systematic comparison of two polynorbornenes with different numbers of trimethylsilyl side groups. Their response upon heating (aging) was compared in terms of structure, dynamics and charge transport behavior. Furthermore, a detailed analysis of the observed Maxwell-Wagner-Sillars polarization at internal interfaces provides unique information about the microporous structure in the solid films. The knowledge obtained from the experiments will guide and unlock potential in synthesizing addition-type polynorbornenes with versatile properties.


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INTRODUCTION With respect to improving energy efficiency and addressing current environmental issues as well as the use of renewable energy sources, membrane technology must be considered an essential key technology.1 Especially for gas separation applications, polymers are still the preferred membrane materials due to their easy solution-based processing into thin permselective layers of different membrane configurations. Historically, the focus of research and industrial development in this field shifted from rubbers to glassy polymers which usually exhibit much lower permeabilities but significantly higher permselectivities. Subsequently, research efforts were directed towards glassy polymers with substantially higher fractional free volume, e.g. from conventional polysulfone to highperformance polyimides.2 A strong research stimulus in this field was the emerging of polymers with extremely high free volume, such as poly(trimethylsilylpropyne) (PTMSP)3 and poly(4-methyl-2-pentyne) (PMP)4,5 in the early 1980’s. These polyacetylenes provide gas permeabilities orders of magnitude higher than conventional rubbery or glassy polymers and PTMSP is still one of the most permeable polymers known. Meanwhile, other classes of such “superglassy” polymers were established,6 especially addition-type poly(norbornenes) and poly(tricyclononenes)

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as well as polymers of intrinsic

microporosity (PIMs)6,10,11,12 . All these classes of polymers exhibit a very rigid backbone in combination with either bulky substituents or a contorted structure induced e.g. by spiro-centers resulting in insufficient packing of the polymer chains in the solid state. This gives rise to very large free volume and in some cases to a continuous void phase13,14 becoming manifest by BET surface areas of several hundred m2/g. For conventional glassy polymers the influence of structural


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modifications on the gas transport properties has been investigated15. For these materials, also a direct relation of solution-diffusion based gas transport and mobility of the polymeric matrix is observed16 which no longer fully holds for most high-free volume superglassy polymers. On the other hand, the permselectivites are significantly higher than expected for a pure pore-flow mechanism. Depending on size and temporal persistence of the porous structures it is plausible to relate the actual transport mechanism to the transition region between solution-diffusion and poreflow.17 Another characteristic of superglassy polymers is that usually no glass transition can be detected by conventional thermal analysis such as DSC before the onset of thermal decomposition. However, there is evidence that despite their limited motional degrees of freedom of the polymer backbone such polymers can undergo a glass transition which can be detected using ultra-high heating rates decoupling the time-scales of glass transition and decomposition processes.18 Nevertheless, due to their high gas permeability, superglassy polymers generally provide a very attractive separation performance for light gases as depicted for relevant gas pairs in a so-called Robeson diagram,19,20,21 where ideal permselectivity  the ratio of pure gas permeabilities P1 and P2, is plotted against the permeability P1 of the more permeable gas. Several of these polymers are located very near to the limiting upper-bound representing the trade-off relation between 𝑃

permselectivity 𝛼 = 𝑃1 and permeability P1 for a certain gas pair, essentially based on the data 2

collection of a vast number of polymers. Some of these polymers even outperform the current state of the art represented by this upper bound.8,9,22-24 Furthermore, for superglassy polymers often a reverse selectivity is observed with respect to separation of hydrocarbons, where the permeability is no longer determined by a sieving mechanism, i.e. a permeability that is dominated by diffusivity


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decreasing with increasing penetrant size, but it is determined by penetrant solubility which increases with molecule size of the hydrocarbons8, 25 . This makes such polymers even more attractive for a number of technically relevant gas separation applications. Different from PIMs, whose insufficient packing in the solid state is often attributed to rigid ladderlike components with a site of contortion in their backbone structure, the formation of large free volume

and

microporosity

in

addition-type

Si-substituted

poly(norbornenes)

and

poly(tricyclononenes) is strongly dependent on the introduction of bulky side groups.9,26 It was reported that the introduction of the trimethylsilyl (SiMe3) group to the addition-type polynorbornene chain leads to an increase in gas permeability by a factor of 80–110.26,27 There is further evidence showing that gas permeability of Si-containing polymers is also sensitive to the number of SiMe3 groups.8,9,28,29 One of the polymers investigated in this study, PTCNSi2g (poly(3,3-bis(trimethylsilyl) tricyclononene-7), chemical structure given in Figure 1) with two SiMe3 side groups, is currently the most permeable member of this class of membrane materials 8,9,26,30 which also are reverseselective for higher hydrocarbon gases. Although superglassy polymers mostly exhibit a superior performance in gas separation, a practical drawback is their pronounced tendency to physical aging,12, 31 - 36 causing the initial microporous structure to approach a more dense state resulting in a significant reduction of permeability. For conventional non-porous polymers physical aging is thought to occur via molecular reorganizations through segmental fluctuations.

37

However, this is not so

straightforward for superglassy polymers in the context of their rigid backbone structures and lack of conformational degrees of freedom. Also, the fact that superglassy polymers with higher chain


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rigidity exhibit more pronounced and faster physical aging compared to conventional glassy polymers seems counterintuitive at first glance. But this behavior becomes clear, when the process of the previous film formation is taken into account: the rigid backbone leads to a premature solidification of the polymer which is then farther from equilibrium and this results in a greater driving force for physical aging as a process approaching this equilibrium. Measurements of gas permeability over different aging periods of time at room temperature or slightly above were carried out to track their aging behavior.8,9, 38 ,39 Despite their significant tendency to physical aging, addition-type polynorbornenes showed relatively slow rates of aging compared to other superglassy polymers with similar level of permeability. This is favorable for their long-term stability required in membrane applications.8,9 Due to the limited number of conformational degrees of freedom, it is difficult to develop a simple picture of physical aging in such systems related to local segmental motions. Alternatively, physical aging may be rationalized as a continuous slow relaxation towards thermodynamic equilibrium in terms of reorganization of free volume40 also governing rate and extent of this process. 41 Nevertheless, broadband dielectric spectroscopy (BDS) can be used to monitor this process, thus providing direct information improving the fundamental understanding of the longterm behavior of this novel class of polymers. On the other hand, knowledge about the dynamics of superglassy polymer helps to reconsider the connection between the diffusive transport of gas molecules in a polymeric matrix and the molecular dynamics on a certain length scale. This relationship is well established for conventional solution-diffusion membranes in which gas molecules diffuse through transient free volume elements or channels forming via thermally activated chain motions,17 which is supported by


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fundamental transport models,

42

molecular dynamics simulations,

43

and experimental

measurements.16, 44 In contrast to that, superglassy polymers exhibit extremely low or even negative effective activation energies for permeation which cannot be understood on a simple physical basis, but rather reflect the complex interplay of diffusivity and solubitity.9,38 As mentioned above, the actual gas transport mechanism in such polymers may be related to the transition of solution-diffusion and pore flow showing features of both, i.e. typical permselectivities of the former and extremely high permeabilities due to a continuous void phase equivalent to a pore network. Wijmans and Baker 17 postulate, that the temporal stability (i.e. the lifetime) of the pore network might be a key factor determining the transport mechanism in this context. Therefore, the investigation of the dynamics of superglassy polymer might also contribute to a better understanding of their outstanding gas separation properties. Moreover, in our recent work a relatively high conductivity was observed for different superglassy polymers even deep in the glassy state.36, 45 This observation was ascribed to the intrinsic microporous structure facilitating the transport of charge carriers.36 The latter can move easily through the loose structure even when segmental motions are frozen. Such phenomena were also found for more fragile polymers with stronger frustration in chain packing and superionic glasses with excess free volume in their structure. 46 , 47 In contrast, in conventional non-microporous polymers, charge carriers can normally diffuse through the dense structure only when segmental mobility is active, i.e. above the glass transition temperature. Following earlier work characterizing the dielectric behavior of PIM-1 33 and an addition-type poly(tricyclononene)

36

the present study focusses on the comparison of two poly(tricyclononenes)

exhibiting one and two trimethylsilyl (SiMe3) substituents, respectively. This allows for a more


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detailed analysis of the influence of the bulky SiMe3-groups on the dynamics and relaxation behavior of these superglassy polymers in the solid state. Although polynorbornenes possess no dipole moment, relaxation behavior, dynamic properties and conductivity can be addressed by broadband dielectric spectroscopy (BDS) as minor dipolar constituents of the polymer matrix (carbonyl groups formed by partial oxidation or impurities like residual catalyst) can be monitored due to the high sensitivity of the experimental equipment. With this method it is possible to cover a wide range of frequencies and temperatures. 48 The measurements are complemented by X-ray scattering experiments to correlate effects of physical aging not only to dynamic properties, but also to identify respective structural changes occurring during heating/cooling cycles applied in the course of the BDS measurements, especially related to the microporosity of the polymers.


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EXPERIMENTAL SECTION Materials. The two polymers under investigation belong to the class of addition-type polynorbornenes. In order to obtain sufficient reactivity of monomers with several bulky side groups with respect to the addition-polymerization, the side groups were located farther from the double bond of the norbornene structure 28. The resulting polytricyclononenes may still be denoted as polynorbornenes, so both terms can be used for these polymer structures. PTCNSi1 and PTCNSi2g were synthesized following the procedures reported in references.8,9 The chemical structures of the materials are given in Figure 1, PTCNSi1 bearing one trimethylsilyl side group and PTCNSi2g two.

Figure 1. Chemical structure of PTCNSi1 (a) and PTCNSi2g (b).

Both polymers exhibit quite good film forming properties due to molar masses of 550000 g/mol (PTCNSi1) and 350000 g/mol (PTCNSi2g) – see also the Supporting Information. Some physico-chemical and gas transport properties of these polymers are reported by Chapala et al. in ref.8,9 for films cast from toluene solution. Both polymers exhibit distinct microporous characteristics with BET surface areas of 790 m2/g (PTCNSi2g) and 610 m2/g (PTCNSi1),


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respectively. Gas permeation coefficients at room temperature are significantly higher for PTCNSi2g (PO2=4750 Barrer, PCH4=6900 Barrer, PCO2=19900 Barrer) compared to PTCNSi1 (PO2=990 Barrer, PCH4=1010 Barrer, PCO2=5300 Barrer), but both on the high level expected for superglassy polymers. Sample Preparation. The solutions with a concentration of ca. 5.4 wt% were prepared for film casting by dissolving 0.25 g PTCNSi1 and PTCNSi2g in toluene (5 ml), respectively, followed by shaking for 5 h. Toluene was selected as solvent because the reported physico-chemical and gas transport properties reported above were obtained for samples prepared from toluene solutions. The solutions were filtered with a 0.2 µm PVDF-filter and cast into a round Teflon mold with a diameter of 4 cm. The mold was placed in a closed chamber with saturated toluene vapor at room temperature to control the evaporation of the solvent from the film. A film with a thickness of ca. 60 µm was formed after 3 days and was removed from the Teflon mold. The film was subsequently dried for 3 days at 393 K (120 °C) in a vacuum oven (oil-free). From this procedure films were obtained which are completely clear and transparent. Photographs of the films are shown in the Supporting Information (see Figure S1). Thermogravimetric analysis (TGA) of the prepared films showed nearly no mass loss for PTCNSi1 and PTCNSi2g up to temperatures of 603 K (330 °C), respectively, indicating a negligible amount of residual solvent (see Figure S2 in the Supporting Information). A glass transition temperature (Tg) could not be detected in the temperature range below the thermal decomposition in nitrogen atmosphere above 600 K.8,9 The dielectric measurements were carried out in parallel plate geometry. Aluminum electrodes with a diameter of 10 mm were evaporated onto both sides of the polymer film in an ultrahigh vacuum (10−5 mbar), to ensure a good electrical contact. Aluminum was chosen instead of gold to


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minimize the heat impact onto the film to avoid a possible damage of the polymer, and to reduce the diffusion of metal atoms into the film with microporous structure.49 Dielectric Relaxation Spectroscopy. To monitor the molecular dynamics of the investigated materials, broadband dielectric spectroscopy (BDS) was employed. The set-up consisted of an ALPHA analyzer with high resolution, equipped with an active sample head (Novocontrol, Montabaur, Germany). In the frequency range from 10-1 Hz to 106 Hz the complex dielectric function 𝜀 ∗ (𝑓) = 𝜀 ′ (𝑓) − 𝑖𝜀 ′′ (𝑓) (f – frequency, ω – angular frequency: ω=2πf, ε' – real part, ε"loss or imaginary part, 𝑖 = √−1) was measured. The frequency range was extended (10-2 to 106 Hz) for selected temperatures – see below. A QUATRO cryo-system (Novocontrol) was used for temperature control better than ± 0.1 K. A pure and dry nitrogen atmosphere was maintained during the complete measurement. For more details see reference.50 It is worth to note that the samples remain clear and completely transparent after the dielectric measurements. Also, the measured films remained completely dissolvable. Additionally, the complex conductivity σ* can be obtained related to the complex dielectric function by Equation 1. 𝜎 ∗ (𝜔) = 𝜎 ′ (𝜔) + 𝑖𝜎 ′′ (𝜔) = 𝑖𝜔𝜀0 𝜀 ∗ (𝜔)

(1)

where σ' and σ" are the real and imaginary part of σ*; ε0 is the permittivity of free space. It holds:

𝜎 ′ (𝜔) = 𝜔𝜀0 𝜀′′(𝜔)

(2)

and


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𝜎 ′′ (𝜔) = 𝜔𝜀0 𝜀 ′ (𝜔)

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(3) .

Thermogravimetric Analysis. Thermogravimetric analysis (TGA) was performed in the temperature range from 300 to 700 K to estimate the onset decomposition temperature of PTCNSi1 and PTCNSi2g under N2 atmosphere and to confirm the complete removal of the casting solvent (heating rate 10 K/min, N2 flow rate 200 ml/min) – see Fig. S2. Small-Angle and Wide-Angle X-ray Scattering. For small-angle and wide-angle X-ray scattering (SAXS/WAXS) measurements a Xenocs Nano-inXider-SW, with a microfocus X-ray tube (copper target) was used. X-rays (=0.154 nm) were focused and monochromatized using a multilayer optic and then collimated by means of two scatterless pinholes. A Pilatus 100k singlephoton counting detector module, located at 93 mm distance from the sample was employed as WAXS detector. This covered a scattering vector (q) range of 3.1 nm-1 < q < 4.2 nm-1. An identical module was mounted 933 mm from the sample behind a semitransparent beamstop, allowing for a practical SAXS q range of 0.06 nm-1 < q < 3.6 nm-1. Data processing for both detectors was performed using the DAWN software package51 according to the latest available procedures.52 Data treatment comprised masking and corrections for counting time, dark-current, transmission, primary beam flux, background, thickness and solid angle followed by a final scaling to absolute units. The treatment was followed by azimuthal integration step.


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RESULTS AND DISCUSSION Aging Induced Changes in the Microstructure of addition-type Polynorbornenes. Because of their strong tendency to physical aging it is expected that the polynorbornenes under investigation undergo structural changes during the heating/cooling cycles in the course of BDS measurements, especially with respect to microporosity. Therefore, the effect of thermal annealing was investigated by SAXS/WAXS measurements. McDermott et al. carried out detailed investigation on the physical aging of the archetypal superglassy polymer of intrinsic microporosity PIM-1 by employing SAXS/WAXS with support from molecular dynamics simulations.35 Changes in PIM-1 scattering patterns as a function of storage temperature, time, and film thickness were consistent with physical aging. It was demonstrated that scattering methods provide a convenient method to characterize the structure of PIM films and the broad SAXS feature around 0.1 Å-1 was identified as being sensitive to microporosity.35 In our recent publication, neutron scattering investigations comparing freshly prepared and aged PIM-1 samples suggested a shrinkage of the microporous structure leading to a denser structure due to physical aging.34 In the present study, SAXS/WAXS measurements were performed on the polynorbornene films before and after annealing. Based on the heating cycles during BDS measurements (see below), annealing was performed at 523 K (250 °C) for 30 min followed by a further step at 573 K (300 °C) for 20 min in N2-atmosphere. Figure 2 depicts the SAXS/WAXS patterns of PTCNSi2g and PTCNSi1 before annealing in comparison with results obtained earlier for PIM-1, a polymer of intrinsic microporosity (inset of Figure 2).34 Scattering patterns from PTCNSi2g film exhibit three main features: power-law (I ~ q-3) scattering at low q values, a broad SAXS/WAXS feature with weak maximum centered at q values of ca. 0.13 Å-1, and a group of peaks at 0.4 Å-1 and larger scattering vectors. The main features are analogous to


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those of PIM-1, where the intensity decreases following a power law with an exponent of -3.0 or -3.2 at low q-values,34,35 a broad scattering feature with weak maximum at ca. 0.29 Å-1 and several peaks at q>0.5 Å-1.34 The low-q power-law is found for most amorphous polymers with no microporosity but with an exponent of 4 (Porod law). The exponent found for PTCNSi2g with a value of 3 points to a structure with a fractal surface (exponent 3.2) in agreement with the microporosity or to a dense mass fractal (exponent 3). The broad feature centered at 0.13 Å-1 corresponds to a Bragg d-spacing of 4.8 nm (48 Å), much larger than any pore size measured for PTCNSi2g (dominant pore-diameter of 7 Å) 8 and might therefore rather be related to a characteristic distance between micropores. At higher q, 0.39 Å-1 and 0.96 Å-1, two sharp peaks are observed. The corresponding Bragg distances are estimated to be 1.6 nm (16 Å) and 0.7 nm (7 Å), respectively. The characteristic scattering within the scattering vector range from 0.1 to 1 Å-1, i.e. the peaks at 0.13 Å-1 and also those at 0.39 Å-1 and 0.96 Å-1, are most probably arising from microporosity. The two latter scattering peaks can be even considered as an evidence for loose chain packing also discussed in ref. 9. These peaks are also sensitive to physical aging as discussed below which supports this assignment. PTCNSi1 has SAXS and WAXS patterns similar to that of PTCNSi2g, indicating no big difference in their microstructure. It exhibits a power law with an exponent of -3.6 at low q-values. This value is closer to the Porod exponent of 4 found for conventional polymers suggesting a denser structure. The scattering close to q=0.1 Å-1 is weak and similar behavior was observed for dense Matrimid.34 Such behavior also indicates a less porous structure of PTCNSi1. The scattering pattern is shifted to somewhat higher q-values with maxima at q=0.45 Å-1 and 1.08 Å-1 corresponding to Bragg d-spacings of 1.4 nm (14 Å) and 0.6 nm (6 Å) - smaller as for PTCNSi2g, being in agreement with smaller BET surface area and pore sizes of PTCNSi1.8,9 Also the lower absolute scattering intensity of PTCNSi1 compared to PTCNSi2g,


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points to less microporosity 35. The positions of the characteristic scattering peaks in the WAXS range are in good agreement with those found in earlier studies for PTCNSi1 9 and PTCNSi2g 8

I~q-3.0

4

3

log (Intensity / a.u.)

(see Table S1 in the Supporting Information).

log (Intensity / cm-1)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


I~q-3.0

PTCNSi2g I~q-3.2

PIM-1

0.01

PTCNSi2g

2

I~q-3.6

0.1

1

q / Å-1

PTCNSi1

1 0.01

0.1

1

q / Å-1 Figure 2. X-ray scattering data for a PTCNSi1 and PTCNSi2g for freshly prepared samples. The data include combined SAXS and WAXS patterns. The inset compares data for PIM-1 taken from reference34and PTNSi2g.

To further investigate the annealing effects on the microporous structure of PTCNSi2g and PTCNSi1, X-ray scattering patterns of both polymers before annealing and after annealing are compared (see Figure S3 in the Supporting Information). For both polymers, PTCNSi2g and PTCNSi1, changes in the absolute scattering intensity were observed in the corresponding q ranges (0.1 to 1.5 Å-1) at the unchanged peak maxima positions which strongly supports the association of this scattering feature with microporosity. We further quantitatively estimated the percentage of reduction in maximal scattering intensity 35 (see Supporting Information) revealing that largest


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pores diffused out of the film most quickly according to the free volume holes diffusion model.53,54 In the case of PTCNSi1, the reduction in scattering intensity is stronger. The changes of smallest pores were more pronounced and accompanied by a slight shift in the peak position, indicating an additional shrinkage in the pore size. It should be noted that comparison of the reduction level between the two polymers must be discussed carefully with respect to sample preparation and treatment history. The film thickness is considered resulting in absolute values of scattered intensity.35,55,56 Impact of Trimethylsilyl Side Groups on Molecular Mobility of addition-type Polynorbornenes. The temperature protocol for the dielectric measurements, depicted in Figure S4 in the Supporting Information, consisted of three consecutive heating cycles and therefore allowed for a basic monitoring of physical aging effects induced during the measurements. Results of the dielectric measurements in the second heating cycle are shown for PTCNSI1 and PTCNSI2g in Figure 3, i.e. a three-dimensional representation of the respective dielectric loss vs. temperature and frequency. Generally, the dielectric response is weak for both polymers as they do not possess an intrinsic dipole moment in the structure of their repeat unit. The dielectric response, i.e. the dielectric loss, is ascribed to a small amount of polar carbonyl groups due to a slight degree of oxidation. 57 The presence of carbonyl groups (C=O) was experimentally confirmed by FT-IR, where characteristic peaks in the wavenumber region of 1790―1720 cm-1 were observed for both PTCNSi1 and PTCNSi2g (Figure S5 in the Supporting Information). As it can be estimated from Figure S5, the degree of oxidation is low and will not affect the morphological properties of the films. Moreover, the contribution of impurities like traces of residual catalyst to the dielectric loss cannot be excluded. For PTCNSi1, two dielectrically active processes are observed, indicated by a peak in the dielectric loss and further


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denoted as β*- and β-process. The peak of the β-relaxation of PTCNSi1 is clearly seen, while the β*-relaxation is relatively weak and partially masked by conductivity phenomena. In contrast, as already reported in the context of an earlier study36, four dielectrically active processes could be identified for PTCNSi2g denoted as β-relaxation and β*-process. The β1- and β2-relaxation is located at higher frequencies and lower temperatures, whereas the β*1- and β*2-processes are found at lower frequencies and higher temperatures. So, in this case, each type of process (β and β*) becomes manifest in two respective peaks in the dielectric loss. Thus, introducing the two bulky trimethylsilyl- (SiMe3-) side groups to a polynorbornene gives rise to a more heterogeneous structure, most probably being the reason the occurrence of multiple dielectric peaks. It is worth to mention that the dielectric loss increases with decreasing frequency and increasing temperature for both polynorbornenes, which is attributed to conductivity contributions. It is an unusual behavior because conductivity is observed far below the glass transition temperature (which could not be measured by conventional thermal analysis as discussed above). Similar but mechanistically different behavior was also observed for another membrane polymer with intrinsic microporosity, PIM-1.45 For PIM-1, the conductivity contribution to the dielectric loss is related to the drift motion of charge carriers in the film. The interaction of π-electrons in aromatic moieties of the PIM-1 backbone (π- π stacking) results in the formation of local intermolecular agglomerates, which further enhances the conductivity in PIM-1 together with the microporosity. In contrast to PIM-1, polynorbornenes have no overlapping π-systems being able to contribute to charge transport. The conductivity is therefore not dominated by charge transport facilitated by π- π stacking. Instead, it is more related to other structural features. More detailed discussion about the observed anomalous behavior will be given in the following section.


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Figure 3. Three-dimensional representation of the dielectric loss versus frequency and temperature for the second heating run (2 H) of PTCNSi1 (a) and PTCNSi2g (b). In Figure 4 the dielectric loss is further plotted as a function of temperature at a fixed frequency of 6 kHz for different heating and cooling runs for PTCNSi1. The two dielectrically active processes are also clearly observed in this representation of the spectra. The peak in the low temperature region present in all heating/cooling runs is the β-relaxation caused by the localized fluctuations. The second process located at ca. 525 K (252 °C) corresponds to the β*-process. This process is probably attributed to another polarization process taking place on different time and spatial scales. The peak intensity of the β*-process is decreasing in the course of consecutive measurement cycles 2 H to 3 H while the intensity of β-relaxation increased slightly, and its peak position is shifted to higher temperatures, i.e. to lower frequencies, as a consequence of exposure to the highest temperature of 573 K (300 °C). Thus, the BDS measurement, i.e. the inevitable thermal treatment connected to it, significantly influences position and intensity of the peaks of both processes. The dielectric loss curves obtained for the first heating (1 H), the first cooling (1 C) as well as the second heating (2 H) up to 473 K


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(200 °C), are overlapping, which indicates the absence of residual solvent in the film, consistent with the TGA measurements (Figure S2). In contrast, a significant difference between the first three runs (1 H, 2 C, 2 H) and the following runs (2 C, 3 H,3 C) is found, which is a clear indicator for structural changes of the sample during the second heating cycle up to 573 K (300 °C). The dielectric spectrum at f=6 kHz, is compared to that at f=210 Hz for the second heating in the inset of Figure 4. Both peaks, β and β*, shift to higher temperatures with increasing frequency as expected. It is worth to note that the dielectric loss at high temperatures (above ca. 500 K (225 °C)) increases with decreasing frequency due to conductivity, in agreement with the findings shown in Figure 3.

Figure 4. Dielectric loss vs. temperature for PTCNSi1 at the frequency f=6 kHz for the different heating/cooling cycles. Inset: Comparison of dielectric loss vs. temperature for PTCNSi1 for cycle 2 H at frequencies of 210 Hz and 6 kHz.

As revealed in the three-dimensional plot shown in Figure 3b PTCNSi2g exhibits dielectrically active processes in similar temperature/frequency ranges as PTCNSi1 pointing to a similar origin


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of these processes. Interestingly, both processes, β and β*, are split in two contributions for PTCNSi2g, indicating a distinct heterogeneity in this polymer in comparison to PTCNSi1. In Figure 5 the β1- and β2-relaxation observed for PTCNSi2g are examined in more detail. In this plot, the two β*-processes appearing in the higher temperature region are not clearly discernible as their position with respect to temperature are very close. Therefore, in the inset the dielectric loss is presented in the frequency domain, enabling a distinct identification of the β*1- and β*2-process. The evolution of the dielectric loss signal regarding its intensity and peak position during the consecutive heating/cooling runs accompanying the dielectric measurements points to structural changes of the material corresponding to physical aging. A detailed discussion was already given in reference 36.

Figure 5. PTCNSi2g: Dielectric loss vs. temperature at the frequency f=6 kHz for the different heating/cooling cycles. The inset shows the dielectric loss as function of frequency at a constant


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temperature of T = 558 K for different measurement cycles – arrows indicate the position of the two dielectrically active processes β*1- and β*2.

The Havriliak-Negami (HN) function, given in Equation 4 was fitted to the data in order to further analyze the relaxation processes.58

∗ (𝜔) 𝜀𝑁𝐻 = 𝜀∞ +

∆𝜀 (1 + (𝑖𝜔𝜏𝐻𝑁 )𝛽 )𝛾

(4)

ε∞ is the real part ε´ with 𝜀∞ = lim 𝜀´(𝜔); Δε represents the dielectric strength; the relaxation 𝜔→∞

time τHN corresponds to fmax (the frequency of maximal dielectric loss). The shape parameters β and γ describe the symmetric and asymmetric broadening of relaxation peaks. If present, conductivity effects are considered in the usual way, i.e. by adding a power law to the dielectric ′′ loss: 𝜀𝑐𝑜𝑛𝑑 = 𝜎0 /𝜔 𝑠 ∙ 𝜀0 ; ε0 represents the permittivity of free space, whereas σ0 is related to the

specific DC-conductivity of the investigated sample. Furthermore, s is a parameter (0 < s ≤ 1) which describes Ohmic (s=1) and non-Ohmic (s

Influence of Trimethylsilyl Side Groups on the Molecular Mobility and

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