Structure – promoted high performance properties of triphenylmethane - containing polyimides and copolyimides

European Polymer Journal 108 (2018) 554–569

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European Polymer Journal journal homepage: www.elsevier.com/locate/europolj

Structure – promoted high performance properties of triphenylmethane containing polyimides and copolyimides

T

Ion Savaa, Mariana-Dana Damaceanua, , Catalin-Paul Constantina, Mihai Asandulesaa, Aleksandra Wolińska-Grabczykb, Andrzej Jankowskib ⁎

a b

“Petru Poni” Institute of Macromolecular Chemistry, Aleea Gr. Ghica Voda 41A, Iasi 700553, Romania Centre of Polymer and Carbon Materials, Polish Academy of Sciences, 41-819 Zabrze, M. Curie-Skłodowska Str. 34, Poland

ARTICLE INFO

ABSTRACT

Keywords: Polyimides Triphenylmethane Halogen Though films Dielectric behaviour Gas permeation

Novel triphenylmethane-based polyimides and copolyimides were synthesized by polycondensation reaction of 4,4′-(4,4′-isopropylidenediphenoxy)bis(phthalic anhydride) with a triphenylmethane-based diamine, or an equimolar mixture of 4,4′-oxydianiline and triphenylmethane-based diamine, via conventional two-step procedure. Three novel diamines were developed as well, being obtained from ortho-toluidine and one of the following aldehydes: benzaldehyde, 4-bromobenzaldehyde or 4-fluorobenzaldehyde. The relationship between the macromolecular structural motif and physical properties of the synthesized polymers was investigated in detail, with emphasis on thermal transitions and stability, dielectric behavior, molecular disorder, mechanical toughness and gas separation performance. The amorphous, free-standing, tough and defectless films prepared from these polymers showed excellent thermal stability, their decomposition starting above 425 °C. For all polymer films which were subjected to dielectric properties measurements, the variation of the real and imaginary parts of the dielectric permittivity with frequency and temperature was measured and discussed. The values of the dielectric constant and dielectric loss were measured at room temperature and in the frequency domain from 1 Hz to 1 MHz, and the obtained results proved the beneficial effect of fluoro and bromo graphting on the dielectric constant reduction. The dielectric spectroscopy data showed distinct γ and β subglass transitions at lower activation energies for copolyimides compared to polyimides, suggesting that the incorporation of the comonomer diamino-diphenyl ether allows faster motions of the small molecules, but hinders the mobility of charge carriers. The permeability of several gases through these membrane-forming materials were measured and discussed with respect to the structural variations in the polymer repeating unit.

1. Introduction

The extensive studies on structure-property correlations have already indicated few fundamental rules for design of polyimides with improved properties. It has been established that incorporation of ether or isopropyl groups into the main polyimide chains generally leads to lower glass transition temperatures (Tg), as well as to significant improvement of solubility and thermoplasticity of the polymers [4]. Additionally, introduction of bulky groups into the polymer main chains or attachment of bulky pendant groups can impart significant increase in Tg by restricting the segmental mobility, while providing a good solubility due to the decreased degree of packing and lower crystallization. Taking into account those facts, possibility to obtain materials with desired properties is closely related to the use of an appropriate molecular design. In case of polyimides, implementation of design results can be carried out by chemical structure modulation of diamine and dianhydride monomers.

Aromatic polyimides have aquired growing importance as high performance materials, due to their high thermal stability, ability to maintain dielectric, physical and mechanical properties over a wide temperature range, low thermal expansion, accessibility and easy proccesing to final products [1]. Commercial and synthesized dianhydrides and diamines of diverse chemical structures have been used as condensation monomers to develop novel aromatic polyimides. As a result, versatile soluble and processable polyimides have been reported and evaluated as materials useful in applications related to advanced technologies [2–4]. Although the unique properties of polyimides make them attractive in many areas, there is still a need for better understanding of structure-properties relationships to improve their characteristics and to advance their application in new commercial fields.

⁎

Corresponding author. E-mail address: [email protected] (M.-D. Damaceanu).

https://doi.org/10.1016/j.eurpolymj.2018.09.029 Received 19 July 2018; Received in revised form 7 September 2018; Accepted 16 September 2018 Available online 18 September 2018 0014-3057/ © 2018 Elsevier Ltd. All rights reserved.

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Scheme 1. Synthetic procedure to diamine monomers.

It has been generally known that penetration of small molecules is restricted through polymer matrices with strong interchain attraction forces and high degree of molecular packing, as it is the case of wholly aromatic classical polyimides. Thus, the main approach used to improve gas permeability is to disrupt chain packing, whereas that used to enhance selectivity is based on reducing chain mobility. Unfortunately, polymers with good selectivity exhibit low permeability and vice versa [5,6]. To minimize the trade-off between permeabilty and selectivity as well as that between processability and improved physical propeties, a specific tailoring of the polyimide molecular structure has to be carried out. Selection of adequate chemical structure to achieve the most favourable balance of transport properties can be considered a challenge for researchers. The majority of the chemical modification attempts have been directed to enhance diffusivity through an increment of the fractional free volume (FFV). From this point of view, fluorine-containing polyimides, mainly those containing the hexafluoroisopropylidene core, have received special importance as they provide a favourable balance of permeability and selectivity [7,8]. Another way to increase FFV and to improve gas separation properties of a polymer material is to introduce bulky groups and to increase at the same time the rigidity of the main chain, because the use of structures with high rigidity results in a strong size sieving ability [9]. The monomers containing bulky pendant groups, aromatic diamines or dianhydrides, could lead to aromatic polyimides with good solubility, preserving their thermal stability and mechanical properties [4]. The introduction of pendant phenyl ring in the structure of an aromatic diamine is an accesible way for the synthesis of processable polymers. Among various bulky diamines used in the synthesis of aromatic polyimides those containing 4,4′- triphenylmethane (TPM) core were rarely used. The pendant phenyl ring induces a free internal rotation of the triphenylmethane bridging group and makes these diamine excellent candidate for synthesis of processable polymers [10]. The literature mentions the existence of a wide variety of diamines derived from TPM using different substituted anilines and benzaldehydes [11], but only several aromatic polyimides have been reported to date [10,12–15]. Thus, the approach followed in this work consisted of using diamines with bulky side groups, such as phenyl or Br- / F-substituted phenyl units, conveniently placed to produce an increase in both FFV and rigidity, thus improving the gas permeation properties, beside others. Graphting a bromo substituent on the phenylene ring would make the polymer more bulky in nature, leading to lower interchain interactions and preventing the polymer chains from packing into tight structures. Furthermore, by virtue of restricted rotation about the bond joining the aromatic ring system, the Br-substituted polymers should have higher glass transition temperatures with respect to analogues polymers without this substituent [16]. On the other hand, the substitution of C-H bonds in polymers by C-F bonds generally endows polyimides with unique properties, such as reduction of dielectric constants, refractive indices or water absorption and increases of solubility, free volume, thermal stability, transparency, or gas permeability, among others [17].

As part of our continuing efforts in developing easy processable, highly thermostable polymers for use as advanced materials, it appeared chalenging to us to synthesize novel aromatic polyimides and copolyimides containing halogen-substituted/unsubstituted triphenylmethane core and flexible isopropylidene units in the main chain. Our work surveys the correlations between the structural features and polymer characteristics, with a special concern on the effect of chemical structure on thermal transitions and stability, dielectric behavior, molecular disorder, mechanical toughness and gas separation performance. By studying those effects, additional insight into structureproperties relationship is expected to be obtained, which may develop strategies to better control polyimide properties as well as to improve predictions of their performance as advanced materials. 2. Experimental 2.1. Starting materials Benzaldehyde (≥99.5%), 4-fluorobenzaldehyde (98%), 4-bromobenzaldehyde (ReagentPlus®, 99%), o-toluidine (≥99%), 4,4′-(4,4′isopropylidenediphenoxy)bis(phthalic anhydride), (6HDA, 97%), 4,4′oxydianiline (97%), sodium hydroxide (granulated EMPLURA®), hydrochloric acid (37%, AR grade), 1-methyl-2-pyrrolidinone (HPLC grade; NMP), chloroform (anhydrous, 99%) and ethanol (analytical standard) were purchased from Sigma–Aldrich. 2.2. Monomers The synthetic pathway to triphenylmethane-based diamines, namely αα-bis(4-amino-3-methylphenyl)phenylmethane (AMPM), α,αbis(4-amino-3-methylphenyl)-4′-fluorophenylmethane (AMPM-F) and α,α-bis(4-amino-3-methylphenyl)-4′-bromophenylmethane (AMPMBr) is outlined in Scheme 1 and the details are described as follows. 2.2.1. α,α-Bis(4-amino-3-methylphenyl)phenylmethane (AMPM) The synthesis of AMPM was carried out by a slightly modified procedure reported for related diamines [13]. Shortly, o-toluidine (21.4 g, 0.2 mol) was heated at 120 °C in nitrogen, then benzaldehyde (10.6 g, 0.1 mol) dissolved in 15 mL (12 N) of hydrochloric acid was added dropwise over a period of 1.5 hrs. The reaction mixture was continously refluxed at ca. 120–124 °C for another 12–14 hrs. After cooled to room temperature, 7.92g of 20% aqueous solution of sodium hydroxide was added to yield a pale blue suspension, which was separated by filtration. The purification of the crude diamine involved the dissolution in chloroform and washing with acid water (pH = 1), when diamine chlorohydrate is formed and passed into water phase. The organic phase was removed and the water phase was washed three times with chloroform to eliminate the impurities. Then the pH was ajusted to 9, and the diamine was extracted with chloroform. The organic residue was concentrated under reduced presure, redissolved in small quantity of methanol, followed by pouring into alcaline solution 555

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of pH 9–9.5. After stirring for 12 h, the resulting precipitate was filtered, washed with water and dried at 60 °C under vacuum. Yield: 65.5%, Melting point: 127–129 °C. 1 H NMR (DMSO, 400 MHz), δ (ppm): 7.24–7.20 (t, 2H, H-12), 716–7.12 (t, 1H, H-13), 7.07–7.05 (d, 2H, H-11), 6.65 (s, 2H, H-5), 6.59–6.57 (d, 2H, H-7), 6.51–6.49 (d, 2H, H-8), 5.14 (s, 1H, H-9), 4.65 (s, 4H, H-2), 1.97 (s, 6H, H-1). 13 C NMR (DMSO, 400 MHz), δ (ppm): 146.25 (C-10), 144.93 (C-3), 132.71 (C-4), 131.01 (C-5), 129.36 (C-11), 128.41 (C-12), 127.38 (C-7), 126.01 (C-13), 121.18 (C-6), 114.21 (C-8), 55.20 (C-9), 18.07 (C-1).

equimolar quantities of 6HDA and a 1:1 mixture of 4,4′-diaminodiphenyl ether and one of the synthesized triphenylmetane-based diamine (Scheme 3). The following example presents the details of procedure used. In a 100 mL three-necked, round-bottomed flask, equipped with a mechanical stirrer and nitrogen inlet and outlet, 0.95 g (0.0025 mol) diamine AMPM-Br, 0.5 g (0.0025 mol) of 4,4′-diaminodiphenyl ether and 15.5 mL NMP as solvent were introduced under nitrogen stream. After complete solubilization of the diamines, a pale yellow solution was obtain. To this solution, 2.6 g (0.005 mol) of 6HDA was added. The reaction mixture was stirred at room temperature for 4 h under nitrogen flow, being acommpanied by the color change to brilliant violet. A part of the formed polyamidic acid solution was used to obtain copolyimide films. The other part was heated under stirring at 180 °C for 4 h under strong nitrogen stream to achive the copolyimide solution. By precipitation into water, a blue violet copolymer was obtained which was finally treated with ethanol in a Soxhlet apparatus for 1 day to remove the oligomers and the high boiling point solvent. Finally, copolyimide CPI-3 was obtained as a pale blue powder after drying in an oven, under vacuum, at 100 °C for 6 h. Any attempt to evaluate the copolymeric composition has failed. Although the reaction mixture was charged with 50% of each diamine, this does not implies that the copolymeric composition AMPM-Br:DDE:6HDA is 1:1:2, due to the different reactivity of each diamine in the polycondensation reaction with 6HDA. The inherent viscosity values of the synthesized polyimides and copolyimides in NMP solution were found in the range of 0.2–0.38 dL/ g. GPC measurements of polymers in CHCl3 solutions provided number average molecular weight (Mn) values between 12,000 and 18,500 Dalton and polydispersity index (Mw/Mn) between 1.3 and 1.8. However, these values have to be taken into consideration with precaution since calibration with polystyrene standards may result in questionable data since the polarity and backbone stiffness of the studied polymers deviate strongly from those of polystyrene.

2.2.2. α,α-Bis(4-amino-3-methylphenyl)-4′-fluorophenylmethane (AMPMF) The reaction was performed as described above, by using 12.4 g (0.1 mol) 4-fluorbenzaldehyde dissolved in 20 mL (12 N) of hydrochloric acid and 21.4 g (0.2 mol) of o-toluidine. Yield: 68.75%, Melting point: 137–139 °C. 1 H NMR (DMSO, 400 MHz), δ (ppm): 7.07–7.05 (d, 4H, H-11,12), 6.64 (s, 2H, H-5), 6.58–6.56 (d, 2H, H-7), 6.52–6.50 (d, 2H, H-8), 5.16 (s, 1H, H-9), 4.66 (s, 4H, H-2), 1.98 (s, 6H, H-1). 13 C NMR (DMSO, 400 MHz), δ (ppm): 162.01 and 159.60 (C-13), 145.02 (C-3), 142.43 (C-10), 132.57 (C-4), 131.02 (C-11), 130.93 (C-5), 127.32 (C-7), 121.26 (C-6), 115.16 and 114.95 (C-12), 114.26 (C-8), 54.31 (C-9), 18.06 (C-1). 2.2.3. α,α-Bis(4-amino-3-methylphenyl)-4-bromophenylmethane (AMPMBr) For the synthesis of AMPM-Br, 21.4 g (0.1 mol) o-toluidine and 18.5 g (0.1 mol) 4-bromobenzaldehyde disolved in 20 mL concentrated HCl were used. The reaction and also purification were carried out as reported for AMPM. Yield: 69%, Melting point: 137–140 °C. 1 H NMR (DMSO, 400 MHz), δ (ppm): 7.44–7.42 (d, 2H, H-12), 7.01–6.99 (d, 2H, H-11), 6.63 (s, 2H, H-5), 6.58–6.56 (d, 2H, H-7), 6.51–6.49 (d, 2H, H-8), 5.14 (s, 1H, H-9), 4.68 (s, 4H, H-2), 1.97 (s, 6H, H-1). 13 C NMR (DMSO, 400 MHz), δ (ppm): 145.82 (C-10), 145.10 (C-3), 132.10 (C-4), 131.57 (C-11), 131.28 (C-12), 130.92 (C-5), 127.34 (C-7), 121.29 (C-6), 119.10 (C-13), 114.26 (C-8), 54.48 (C-9), 18.06 (C-1).

2.4. Preparation of polymer free-standing films The polyamidic acid solutions in NMP obtained during the polycondensation reactions were used to prepare free-standing polymer films by casting the solutions onto dust-free glass plates. These films were gradually heated from room temperature up to 70–80 °C for 4 h and then at 100, 150, 200 and 250 °C (20 min at each temperature) to remove the residual solvent, and after that kept at 250 °C for 2 h to complete the imidization in solid state. The films were stripped off the plates by immersing into water and dried under vacuum at 100 °C. All polymer films were flexible and mechanically resistant, and they were used for different measurements.

2.3. Polymers Three novel polyimides were synthesized by the solution polycondensation reaction of equimolar amounts of an aromatic diamine (AMPM, AMPM-F, AMPM-Br) with 4,4′-(4,4′-isopropylidenediphenoxy)bis(phthalic anhydride) (6HDA) by a two-step pathway (Scheme 2). The following example illustrates the general procedure. In a 100 mL three-necked, round-bottomed flask, equipped with a mechanical stirrer and nitrogen inlet and outlet, 1.51 g (0.005 mol) diamine AMPM and 15 mL NMP as solvent were introduced under nitrogen flow. After complete solubilization of the diamine, a pale-yellow solution was obtained. To this solution 2.6 g (0.005 mol) dianhydride 6HDA was added. The color of the reaction mixture changed to blue-violet. The reaction mixture was stirred at room temperature for 4 h with the formation of intermediate polyamidic acid. A part of the polyamidic acid solution was used to prepare polyimide films. The other part of solution was heated under strong nitrogen stream up to 180 °C and stirred at this temperature for 3–4 h, then it was gradually cooled to room temperature. The formed polyimide solution was poured into water to precipitate the solid polymer. The polymer was washed with plenty of water and finally treated with ethanol in a Soxhlet apparatus for 1 day in order to remove the oligomers and the high boiling point solvent. Finally, polyimide PI-1 was obtained as a cream-coloured powder after drying in an oven, under vacuum, at 100 °C for 6 h. Following the same procedure, a series of three copolyimides (CP-1 – CP-3) was synthesized as well. They were obtained by using

2.5. Measurements The NMR spectra were registered with a Bruker Avance III 400 spectrometer, operating at 400.1 and 100.6 MHz for 1H and 13C nuclei, respectively. 1H and 13C chemical shifts are reported in δ units (ppm) relative to the residual peak of the solvent, dimethyl sulfoxide (DMSO‑d6). FTIR spectra were recorded with a FT-IR VERTEX 70 (Bruker Optics Company), with a resolution of 0.5 cm−1 in ATR mode, by using freestanding films. XRD spectra were recorded on a Bruker D8 Avance diffractometer, using the Ni-filtered Cu-Kα radiation (λ = 0.1541 nm). The working conditions were 36 kV and 30 mA. All diffractograms were registered in the range of 2–40 (2θ degrees), at room temperature. The initial samples for X-ray measurements were polymer films. All diffractograms are reported as registered. The d-spacing values which allow the estimation of interchain spacing distances were calculated from the diffraction peak maximum, by using the Bragg equation: 556

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Scheme 2. Synthetic route to triphenylmethane-based polyimides.

Scheme 3. Synthetic route to copolyimides based on triphenylmethane.

d = /2 sin

(1)

Ubbelohde viscometer with k = 0.05002. The average molecular weights were measured by gel permeation chromatography (GPC) using a ParSEC Chromatography Ver. 5.67 Brookhaven Instruments Corp. apparatus provided with refraction and UV detectors and PL Mixed C Column. Measurements were carried out with polymer solutions of 0.2% concentration in CHCl3 as solvent.

where λ is the wavelength of the radiation and 2θ is the angle of maximum intensity in the amorphous halo exhibited by the polymer. The inherent viscosities of the polymers were determined at 20 °C, by using NMP-polymer solutions of 0.5 g/dL concentration, with an 557

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Differential scanning calorimetry (DSC) analysis was performed using a Mettler Toledo DSC 1 (Mettler Toledo, Switzerland) operating with version 9.1 of STARe software. The samples (6–9 mg) were encapsulated in aluminium pans having pierced lids to allow escape of volatiles. The heating rates of 10 °C min–1 and nitrogen purge at 100 mL min–1 were employed. To determine the glass transition temperature, heating–cooling-heating waves were registered in the temperature range from 25 to 300 °C. Heat flow vs. temperature scans from the second heating run were plotted and used for calculation of the glass transition temperature. The inflection-point of the curve resulting from the typical second heating was assigned as the glass transition temperature of the respective materials. Permeation measurements were made at 7 bar of applied gas pressure, and at 35 °C using a constant-pressure/variable-volume apparatus [18]. The membrane thickness was in the range of 115–175 µm for homopolymers and 90–110 µm for copolymers and it was measured with a digital micrometer readable to ± 1 µm. The membrane diameter was 5 cm. Pure helium (He), oxygen (O2), nitrogen (N2), and carbon dioxide (CO2) were used for permeation experiments, and those gases were measured in the given sequence. The gas permeability, P, expressed in Barrer units, was determined according to the Eq. (2).

AMPM

AMPM-F

P = 10

10

q· l (p1 p2 ) A·t

(2) 3

where q is the quantity of a permeant [cm (STP)] passing through the membrane in time t [s], A is the effective membrane area [cm2], l is the membrane thickness, and p1 and p2 are the upstream and downstream pressure, respectively. The error of the permeability coefficient associated with this system was determined to be 10–15%, the error being larger the lower the gas permeability. From single gas permeation experiments, and by using the Eq. (3), the ideal selectivity (α) for gases A and B was also calculated.

= AMPM-Br

PA PB

(3)

The tensile properties were recorded on a Shimadzu AGS-J deformation apparatus at ambient temperature and at a rate of deformation of 1 mm/min with a load cell capable of measuring forces up to 1 kN. For each data point, three film samples of 20 mm × 5 mm size were tested and the average value was taken. 3. Results and discussion 3.1. Structural characterization of monomers and polymers Three aromatic diamines containing triphenylmethane group, namely α,α-bis(4-amino-3-methylphenyl)phenylmethane (AMPM), α,α-bis(4amino-3-methylphenyl)-p-fluoro-phenylmethane (AMPM-F), and α,α-bis (4-amino-3-methylphenyl)-p-bromo-phenylmethane (AMPM-Br) were successfully prepared in a facile way, by a single step procedure, as shown in Scheme 1. These diamines were characterized by IR, 1H- and 13C NMR spectroscopy, which proved their correct chemical structure. In the FTIR spectra, AMPM, AMPM-F and AMPM-Br (Fig. S1,ESI) showed characteristic absorption bands for N-H stretching at 3355 cm−1 and 3460 cm−1, NH deformation at 1622 cm−1 and CeN stretching at 1278 cm−1, confirming the presence of primary amine groups in the products. The absorption bands from 2855 cm−1 and 2925 cm−1 characteristic for aliphatic CeH stretching were associated with the presence of methyl units in all three diamines. The absorption bands characteristic for aromatic CeH linkage were found at 3018 and 700 cm−1, while those characteristc for C]C linkage were registered at 1500 cm−1. The presence of CeBr linkage in the diamine AMPM-Br was evidenced by the absorbtion band at 748 cm−1, while the C-F stretching in AMPM-F was assigned to the absorption band at 1154 cm−1 [19]. Fig. 1 display the 1H NMR spectra corresponding to AMPM, AMPMF and AMPM-Br. Their 13C NMR spectra are shown in Fig. S2, ESI. The

Fig. 1. 1H NMR spectra of AMPM, AMPM-F and AMPM-Br.

Polystyrene standards of known molecular weight were used for calibration. Dielectric spectroscopy measurements were performed with a Broadband Dielectric Spectrometer (Concept 40, GmbH Germany) in wide frequency range (from 1 Hz to 1 MHz) and temperature (from −150 °C to 280 °C). The alternative field was applied with an Alpha-A High Performance Frequency Analyzer and the temperature was controlled with a high accuracy Novocontrol Quatro Cryosystem. The samples were placed between two gold coated plate electrodes and the analyses were carried out in dry nitrogen atmosphere, avoiding the water absorption. Thermogravimetric analysis (TGA) was performed on a Mettler Toledo model TGA/SDTA 851 under nitrogen flow (20 cm3 min−1), at a heating rate of 10 °C/min from 25 to 900 °C, by using polymer films. The initial mass of the samples was 6–9 mg. 558

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assignments of 1H and 13C chemical shifts were made according to the numbering scheme presented in Fig. 1. These assignments are based on 2D NMR heteronuclear correlations HMQC (Heteronuclear MultipleQuantum Correlation) that allows to obtain a 2D heteronuclear chemical shift correlation map between directly-bonded 1H and 13C –heteronuclei (Fig. S3, ESI). In the 1H NMR spectra, the aromatic amine protons appeared at 4.65, 4.66, and 4.68 ppm, while the aliphatic protons of the triphenylmethane core were identified at 5.14, 5.16, and 5.14 ppm for AMPM, AMPM-F and AMPM-Br, respectively. In the aliphatic region, the synthesized diamines exhibited a singlet at 1.97, 1.98, and 1.97 ppm for the same diamine order, which is associated with the protons of CH3 groups. The 13C NMR spectra showed the signals at 18.07, 18.06, and 18.06 ppm corresponding to the aliphatic carbons attached to the benzene rings, while the aliphatic carbons of the triphenylmethane unit appeared at 55.20, 54.31, and 54.48 ppm. The carbon atoms linked directly to the amine groups were identified at 144.93, 145.02, and 145.10 ppm for AMPM, AMPM-F and AMPM-Br, respectively. Solution polycondensation reaction of 4,4′-(4,4′-isopropylidenediphenoxy)bis(phthalic anhydride) with one of the aromatic triphenylmethane-based diamines, AMPM, AMPM-F or AMPMBr enabled novel aromatic polyimides PI-1 – PI-3 to be obtained which contain voluminous pendent phenyl, p-fluorophenyl or p-bromophenyl units, as seen in Scheme 2. The two-step polycondensation reaction of the same dianhydride with a 1:1 mixture of 4,4′-oxydianiline and one of the triphenylmethane-based diamines, AMPM, AMPM-F or AMPM-Br, in NMP solutions, yielded copolyimides CP-1 - CP-3, as shown in Scheme 3. The FTIR spectra of the synthesized polyimides and copolyimides proved their correct structure, judging by the main characteristic absorptions bands listed in Table 1. Characteristic imide ring absorptions appeared in the range of 1776–1777 cm−1 (asymmetrical C]O imide stretching), 1715–1718 cm−1 (symmetrical C]O imide stretching) and 745–747 cm−1 (imide ring deformation), while no amide group absorption from the intermediate polyamidic acids, corresponding to NH stretching (wide bands) and carbonyl stretching (amide I and amide II) was observed. Such features proved the total conversion of the intermediate polyamidic acid into final polyimide or copolyimide structures by polycondensation in thin films at high temperature. The incorporation of the methyl groups in the diamine core of polymer chains was evidenced by the strong absorption peaks in the range of 1443–1445 cm−1 associated with CH3 bending. The absorption band at 2966–2968 cm−1 atributed to the CeH stretching from the triphenylmethane structure is overlapped with that of aliphatic CeH stretching from CH3 units. CeH and C]C linkages in aromatic rings showed absorption peaks at 3058–3066 and 1498–1503 cm−1,

respectively, while the strong absorption bands in the range of 1235–1236 were associated with aromatic ether stretchings. The polyimides and copolyimides containing aditional halogen atoms showed characteristic absorption peaks in the range of 1158–1159 cm−1 (fluorine-containing polymers) and 820–829 cm−1 (bromine-containing polymers). Reprezentative FTIR spectra are displayed in Fig. S4, ESI. The new polyimides and copolyimides could be easily dissolved in aprotic polar solvents such as NMP, DMSO, DMAc, and some of them in more common organic solvents, such as THF and chloroform (Table S1, ESI). The dissolution in these solvents yielded homogeneous polymer solutions which were stable at room temperature without phase separation, gelation or precipitation. The solubility behaviour of polyimides is generally influenced by the chain packing density and intermolecular interactions, which are ussually promoted by the polymer chains rigidity. The enhanced solubility of these polymers can be explained by the presence of bulky pendent phenyl units in the triphenylmethane core which led to an increase of polymer chain distances and a decrease of physical bonding. The sp3 hybridization of the tertiary carbon from both diamine and dianhydride segments induces an increase in the freedom of the chains movements and, due to the generation of non-coplanar structures, facilitates the penetration of small solvent molecules between polymer chains, thus solubilizing them. As a consequence, due to their good solubility, the presented polymers are appropriate candidates for spin-coating and casting processes and the resulting thin and thick films could be used in advanced applications. 3.2. Thermal properties The thermal features of polyimides PI-1 - PI-3 and copolyimides CPI-1 - CPI-3 were evaluated by means of differential scanning calorimetry (DSC), and thermogravimetric analysis (TGA) using small pieces of polymer films as samples for all measurements. The results of these investigations are presented in Table 2. Fig. S5, ESI section, shows the DSC curves of the synthesized polyimides and copolyimides. The Tg values of the polymers taken from the DSC second heating scan were found in the range between 206 and 234 °C (Table 2). There are two factors that appear to dictate the Tg values of both polyimides and copolyimides: chain rigidity and molecular packing state. Those polyimides contain the same dianhydride component and, therefore, we assume that the difference between Tg values are given by the diamine component. It was observed that the introduction of both fluorine and bromine substituents led to the increase of glass transition temperature of PI-2 and PI-3 with respect to PI-1 which does not contain any substituent on the pendent phenyl ring. It can be expected that Br atoms restrict the rotational motion within the PI-3 chains due to their bulky nature, and consequently rise the chain rigidity and Tg [20,21]. On the other hand, the fluorine atom combines the properties of its large electronegativity and small size so that the polyimide PI-2 exhibits higher Tg with respect to polyimide PI-1. The stronger interchain

Table 1 FTIR (films, cm−1) analysis of the synthesized polyimides and copolyimides. Band assign

PI-1

PI-2

PI-3

CPI-1

CPI-2

CPI-3

aromatic imide, C]O asymmetric stretching aromatic imide, C]O symmetric stretching CeN bending aromatic CeH stretching aromatic C]C stretching aromatic CeF stretching aromatic CeBr stretching aromatic CeOeC symmetric stretching CeN stretching aliphatic CeH stretching CH3 bend

1777

1777

1776

1776

1776

1776

1716

1717

1718

1715

1715

1716

747 3058 1502 – – 1236

746 3066 1503 1159 – 1235

747 3059 1501 – 820 1236

745 3058 1499 – – 1236

745 3061 1499 1158 – 1236

745 3065 1498 – 829 1236

1369 2967 1444

1369 2968 1443

1370 2967 1443

1370 2967 1444

1369 2967 1445

1369 2966 1445

Table 2 Thermal data of triphenylmethane-based polyimides and copolyimides. Polymer

Tg (°C)

Tonset (°C)

T10% (°C)

Tmax (°C)

W900 (%)

PI-1 PI-2 PI-3 CPI-1 CPI-2 CPI-3

222 234 226 206 227 230

427 466 447 457 470 476

499 504 526 519 494 516

514 516 503 520 524 523

64.2 62.3 60.2 60.5 58.5 62.3

Tg = glass transition temperature. Tonset = onset temperature on the TG curve. T10% = temperature of 10% weight loss on the TG curve. Tmax = temperature of maximum rate of decomposition. W900 = char yield at 900 °C, under nitrogen. 559

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Fig. 3. XRD spectra of the investigated polymer films.

between 13 and 39 °C being registered with respect to the unsubstituted polymers. Very interesting, the incorporation of diamino-diphenylether core into the main chains of CPI-1 and CPI-3 copolymers, for which polar interchain interactions are less probable than for CPI-2, resulted in an increase of the initial decomposition temperature by approx. 30 °C compared with analogue polymers, PI-1 and PI-3. On the contrary, no significant difference in the thermal stability of PI-2 and CPI-2 was observed since the fluorine atoms promoted interchain interactions to such extent that balanced the effect of diamino-diphenylether segments on polymer thermal stability. The anaerobic char yield value at 900 °C (W900) for each polymer is also summarized in Table 2. The remaining char yields were all above 58.5% at 900 °C under nitrogen atmosphere, demonstrating the high thermal stability of the investigated polymers. 3.3. X-ray diffraction studies X-ray diffraction patterns registered for films of polyimides PI-1 – PI-3 and copolyimides CPI-1 - CPI-3 are illustrated in Fig. 3. XRD investigations revealed that all the synthesized polymers show only one broad halo centered between 16.63° and 17.92° (2θ), being a proof of their amorphous character. Most probably, the molecular disorder was induced by both the tetrahedral carbon atom being present in the triphenylmethane and isopropylidene units of the polymer chains, as well as the pendent side groups, disrupting the polymer planarity and hence the possibility of order. The most prominent peak in the WAXD spectra of amorphous glassy polymers is often used to estimate the average interchain spacing distance (d-spacing) by using the Bragg equation (Table 3). There are obvious differences between the d-spacing values within the group of polyimides, on one hand, and between polyimides and copolyimides, on the other hand. This can be readily explained by comparing the polyimide and copolyimide backbone segments derived from the diamine core. PI-3 contains voluminous Br atoms substituted on phenyl ring in the diamine segment that hinder the rotational motion within the chain, consequently enhance the rigidity and allow a denser packing with respect to the PI-1 polymer chains. A similar effect

Fig. 2. TGA-DTG curves of triphenylmethane-based polyimides (a) and copolyimides (b).

interactions generated by the polar nature of F atoms can partially balance its small sizing effect [22]. The same judgment could be made for copolymer series, in which copolyimide CPI-1 has lower Tg value compared to CPI-2 and CPI-3. The structural difference between polyimides on one hand, and copolyimides on the other hand, is given by the diamine segment, containing additional ether groups in the case of copolyimides. Generally, the presence of ether linkages increases the flexibility of the polymer backbones by deacreasing the rotational barrier of the main chains. This is clearly reflected in the decreased Tg values of the CPI-1 and CPI-2 copolymers compared to the analogue polymers. However, the bromine substituted polymers exhibit different behaviour. In this case, CPI-3 displays a small increase in Tg value instead of its reduction, when compared to PI-3, indicating that the ether linkages did not improve the rotational motion within the copolymer chains. Moreover, other factors such as the molecular weight should be considered responsabile for a higher Tg of CPI-3 compared to PI-3. The results of TGA measurements for the investigated polyimides and copolyimides are listed Table 2. Representative TGA curves are shown in Fig. 2. The polymers showed excellent thermal stability, as expected in case of aromatic polyimides. The initial decomposition temperature (IDT) defined by the onset value on the TG curve was found in the range of 427–476 °C, while the temperature of 10% gravimetric loss (T10%), an important parameter for evaluation of thermal stability, ranged between 494 and 526 °C. The temperature corresponding to the maximum rate of decomposition (Tmax) was in the range of 503–523 °C. The incorporation of F or Br atoms to the pendent polymer side groups proved to be beneficial for the thermal stability, an increase

Table 3 XRD and tensile test data of the investigated polymers.

560

Polymer

2θ(°)

dspacing (Å)

Young’s Modulus (GPa)

Tensile strength (MPa)

Elongation to break (%)

PI-1 PI-2 PI-3 CPI-1 CPI-2 CPI-3

16.63 17.73 17.92 17.28 17.74 17.45

5.33 5.00 4.95 5.13 4.99 5.08

0.76 0.83 0.72 1.67 1.12 1.36

53.53 ± 7.11 54.38 ± 6.66 25.02 ± 1.4 97.54 ± 23.1 105.6 ± 0.84 91.9 ± 12.83

7.65 ± 1.94 6.41 ± 0.74 1.6 ± 0.54 5.42 ± 1.5 9.45 ± 1.39 6.53 ± 0.1

± ± ± ± ± ±

0.18 0.02 0.17 0.03 0.15 0.23

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is induced by the F atoms in PI-2, only that the denser packing can be explained by the polar interchain interactions promoted by its strong electronegativity. The d-spacing values of the copolyimide series suggest that the incorporation of diamino-diphenylether segments in the main polymers chains have an intriguing effect on chain packing ability of copolymers having different substituents on the pendent phenyl ring: the packing is more pronounced when no substituent is present (CPI-1), remain similar in the case of F-substituted copolymer (CPI-2), and is slightly reduced for Br-substituted copolymer (CPI-3). Therefore, we can assert that the competition between several factors, such as the rigidity of the main polymer chains, intermolecular forces, and barriers to rotation along the main polymer chains leads to those different values of d-spacing, without following a general rule. However, the XRD features are consistent with the previously discussed DSC results.

The values of tensile strength were found in the range of 25.02–105.6 MPa, elastic modulus in the range of 0.72–1.67 GPa and elongation at break in the range of 1.6–9.45 %. Fig. 4b displays some representative tensile tests of the films of polyimide and copolyimide series. These values are slightly lower than those reported for related polyimides based on the same dianhydride but with different aromatic diamines [23]. By comparing the tensile properties of polyimide films without and with halogen substituents we can see that F units improved the tensile properties due to a better packing of polymer chains promoted by this atom. In the case of copolymer series, the introduction of diphenylether units in the macromolecules was beneficial for the mechanical properties as well, leading to higher both tensile stress and tensile strain. Thus, all the tensile data presented in Table 3 demonstrate that films with satisfactory mechanical properties (toughness and modulus) can be achieved from the studied polymers. Accordingly, the investigated polymers are suitable candidates for use as advanced materials and, therefore, they were subjected to dielectric behaviour investigations and tested as gas separation membranes.

3.4. Tensile tests of free-standing films The free-standing polyimide films having the thickness in the range of 115–175 μm were obtained from polyamidic acid solutions casted onto glass plates and gradually heated up to 250 °C to eliminate the solvent and complete the imidization in solid state. The resulting films were defect-free, flexible, tough, and maintained their integrity after repeated bending (Fig. 4a). The toughness of these films allowed their easy mechanical handling. The tensile properties of these materials are collected in Table 3. Elastic modulus, tensile strength and elongation at break have been evaluated on the basis of three drawing experiments.

3.5. Dielectric behaviour 3.5.1. Primary assessment of the dielectric behaviour of triphenylmethanebased polyimides and copolyimides Dielectric relaxation spectroscopy is a widely used technique to attain information about the storage and dissipation components of the complex permittivity, namely dielectric constant (ε′) and dielectric loss

Fig. 4. (a) Combined images showing the flexibility of the studied polymer films;(b) Stress-Strain curves of the polymers. 561

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Fig. 5. The evolution of dielectric constant versus frequency at various temperatures for the investigated polymers.

(ε”) of polymer films in a broad frequency and temperature range. Also, from their variation, thermal transitions correlated with dipolar movements can be pointed out. Dipolar relaxation processes related to thermal transitions are evidenced by the step increase in dielectric constant waves and dielectric loss peaks, which shift to higher temperature with increased frequency [24–26]. According to literature data reported for polyimides [27–29], and generally for polymers, the dielectric constant, ε′, drops with increasing frequency. The dielectric constant being determined by the ability of polarizable units to align to the alternating electric field decreases with frequency enhancement since the orientation of dipole moments demands longer time than that corresponding to the applied frequency [30]. Electro-insulating features of the films obtained from triphenylmethane-based polyimide films were assessed on the basis of the dielectric constant and dielectric loss and their variation with frequency and temperature. The dielectric constant and dielectric loss as a function of frequency for polyimide PI-1 - PI-3 and copolyimide films CPI-1 - CPI-3 are presented in Figs. 5 and 6. The values of dielectric constants, ε′, and dielectric losses, ε” were analyzed in the 1 ÷ 106 Hz frequency range at temperatures between -150 °C and 280 °C. The dielectric constant increases with frequency decreasing for all polyimide films at very low, room and high temperature (Fig. 5a–c). The rate increase is higher at low frequencies and high temperatures. A strong increase of the dielectric constant can be observed for PI-3 film at high temperature (150 °C) and in the frequency range 1–102 Hz, due to the mobility of the charge carriers (Fig. 5c). The dielectric constant values of the investigated polyimide films at 1 Hz, 103 Hz and 106 Hz, at 24 °C, are listed in Table 4. At a temperature of −90 °C (Fig. 6a), the dielectric loss versus frequency dependences reveal a dielectric peak, around 10 Hz, that is attributed to the γ-relaxation process. At room temperature (Fig. 6b), the ε”(f) dependencies contain two dielectric maxima specific to secondary relaxation-type

processes. The β-relaxation appeared at low frequencies as a broad dielectric signal, while an end-part of γ-relaxation is revealed at high frequencies as a narrower dielectric peak. Fig. 6c shows the dielectric loss evolution with frequency for polyimide and copolyimide series recorded at 150 °C. At low frequencies, the dependencies are dominated by the conductivity signal, especially for PI-3 sample, while in the high frequency spectral region, the end-part of β-relaxation is slightly visible. The dielectric constant values of these triphenylmethane-based polymers at 24 °C and in the frequency range between 1 Hz and 106 Hz are in the range of 2.94–4.09, being comparable with those of Kapton HN polyimide film [26]. Lower dielectric constant values were obtained for the F- and Br-substituted polymers PI-2 and PI-3 as compared to the unsubstituted analogue PI-1, due to the high hydrophobicity of halogens that diminishes the moisture absorption and, as a consequence, reduces the polarizability of the products. On the other hand, lower dielectric constant values were registered for copolyimide series. The decrease rate depends on the polymer structural motif, demonstrating that there are several factors with a cumulative effect on the ability of the polarizable units to orient fast enough to keep up with the oscillation of the alternative electric field [30]. The incorporation of diamino-diphenylether core in the main chains of copolymers disturbs the close packing of polymer chains, leading to decreased dielectric constants compared with those of polyimide series. Moreover, it was observed that the introduction of both bromine and fluorine substituents in the copolyimide chemical structures had a similar effect as in the case of polyimide series, leading to decreased dielectric constants with respect to the copolymer which do not contains any substituent on the pendent phenyl ring CPI-1. On the other hand, the dielectric constant of polyimide films was found to be depended on the temperature, as can be observed in the large temperature domain, from -150 °C to 280 °C. From a dielectric 562

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Fig. 6. The evolution of dielectric loss versus frequency at various temperatures for the investigated polymers.

point of view, all polyimide films reveal a high thermal stability, since the dielectric constant is not influenced by the temperature until 150 °C, as can be seen in Fig. 7a. By introduction of the second diamine in the macromolecular chains, the dielectric constant remained constant until 170 °C, increasing the thermal stability of the copolymer series. As temperature increases, the dielectric constant raises strongly because the chemical dipoles absorb sufficient thermal energy and could follow the oscillations of the external alternative field. The presence of the dielectric peak (e.g., for PI-3 sample, the peak maximum appears around 165 °C) during the increase of dielectric constant might be attributed to a physical process (e.g. removal of solvent traces that have been included in the bulk sample), while with further increase of temperature, ε′ increases considerably, especially due to high mobility of charge carriers. The position of the dielectric peak observed on the ε′(T) dependencies seems to be dependent on the chemical structure of the polyimide. While the presence of diamino-diphenylether core in the chemical structure of copolymers shifts the peak maxima to higher temperatures, the presence of bromine from polyimides and

copolyimides has the opposite effect, promoting the peak appearance at lower temperatures. Fig. 7b displays the temperature dependencies of dielectric loss for the polyimide and copolyimide series, revealing the presence of dipolar relaxation-type processes. At low temperatures, an intense dielectric peak is visible from -150 °C to -50 °C that might be associated with γrelaxation. At higher temperatures, β-relaxation appears as a broad peak (e.g., for CPI-2 sample, at a frequency of 10 Hz, β-relaxation appears around 50 °C) that is overlapped by the dielectric signal of charge carriers mobility. Fig. 8 shows exemplarily the temperature evolution of dielectric loss for CPI-2 sample at selected frequencies. On increasing temperature, the isochronal plots reveal three thermal transitions: γ-, β- and α-relaxations. The primary α-relaxation is closely connected with the glass transition and appears as a shoulder with reduced intensity, especially at high frequencies. This process is strongly overlapped by the dielectric signal associated with the conductivity of charge carriers. On the other hand, it is important to note that the maxima of the dielectric peak

Table 4 Dielectric data of the triphenylmethane-based polymers at 25 °C. Polymer

PI-1 PI-2 PI-3 CPI-1 CPI-2 CPI-3

Dielectric constant at 3

Dielectric loss at 6

Conductivity, S/cm at 3

1 Hz

10 Hz

10 Hz

1 Hz

10 Hz

10 Hz

1 Hz

103 Hz

106 Hz

4.09 3.70 3.81 3.36 3.02 3.38

4 3.64 3.70 3.33 2.97 3.33

3.88 3.56 3.56 3.30 2.94 3.28

2.7 × 10−2 2 × 10−2 2.9 × 10−2 1.3 × 10−2 1.8 × 10−2 1.6 × 10−2

1.3 × 10−2 8.2 × 10−3 1.8 × 10−2 2 × 10−3 4.8 × 10−3 3.7 × 10−3

4 × 10−2 2.8 × 10−2 6.3 × 10−2 3.8 × 10−3 1.3 × 10−2 5.6 × 10−2

1.5 × 10−14 1.1 × 10−14 1.6 × 10−14 7.1 × 10−15 1.0 × 10−14 8.8 × 10−15

7.2 × 10−12 4.6 × 10−12 1.0 × 10−11 1.1 × 10−12 2.7 × 10−12 2.1 × 10−12

2.2 × 10−8 1.5 × 10−8 3.5 × 10−8 2.1 × 10−9 7.2 × 10−9 3.1 × 10−8

563

6

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Fig. 7. The evolutions of dielectric constant (a) and dielectric loss (b) versus temperature, at 1 Hz, for polyimide and copolyimide series.

localized around 171 °C are not changing with increasing frequency and, consequently, cannot be attributed to a relaxation-type process. The representative 3D version of the dielectric loss of CPI-2 as a function of frequency and temperature allows a better visualization of relaxation-type processes.

In order to characterize the secondary relaxation processes, each ε”(f) dependency was analysed in detail using a sum of the model function introduced by Havriliak-Negami [34]:

3.5.2. Havriliak-Negami (HN) parameters Fig. 9 shows the representative frequency – dependent dielectric loss spectra for PI-2 and CPI-2 samples, at various temperatures. The detailed ε”(f) dependencies recorded at low temperatures (between −100 °C and -50 °C) (Fig. 9a for PI-2 and Fig. 9b for CPI-2 samples) exhibit prominent dielectric peaks assigned to γ-relaxation, that shifts to higher frequencies as temperature increases. According to literature, γ-relaxation is generally activated by localized fluctuations of small polymer segments from side chain macromolecules and is influenced by moisture absorption from atmosphere [27,31]. At temperatures around +100 °C (Fig. 9c for PI-2 and Fig. 9d for CPI-2 samples) a broad peak was registered, being associated with β-relaxation. β Transition of polyimides was attributed to the segmental motion in diamine or dianhydride units, but may include larger portions of the polymer repeating unit that respond in a correlated manner, or can ultimately imply the entire repeating segment, as in case of rigid polymers [31–33]. Usually, the aromatic polyimides show a supplementary relaxation process, α-relaxation, which is associated with the glass transition. This thermal transition takes place at higher temperatures than secondary relaxations occur and it is related to the large-scale cooperative chain motions. In our case, α-relaxation is hardly visible due to overlapping with conductivity dielectric signal.

where Δε = εs − ε∞ is the relaxation intensity, ω = 2πf represents the angular frequency, f is the frequency, τHN is the HN relaxation time for each process associated with peak maxima, and a and b represent the broadening and skewing parameters, respectively. The fitting procedure was performed with WinFIT software package provided by Novocontrol. Fig. S6, ESI displays examples of HN fitting procedure for γ- (a and b)) and β-relaxation (c and d) dipolar processes for PI-2 and CPI-2 samples. Similar operations were applied for PI-1, PI-3, CPI-1, and CPI3 systems. According to Fig. S6a and b, ESI the HN fit (displayed with black line) encloses a single HN term associated with γ-relaxation (displayed with a red line). Furthermore, the ε”(f) dependencies of βrelaxation are processed with one HN term for β-dipolar relaxation and an additional term employed to separate the conductivity signal. Because of overlapping with conductivity of charge carriers, the fit procedure for β-relaxation cannot be performed for all polyimide and copolyimide samples. For temperatures higher than 50 °C, the ε”(f) dependencies of PI-3 and CPI-1 samples are dominated by conductivity and, as a consequence, the secondary β-relaxation is unclear. The numerical values of the HN parameters obtained from Eq. (4) for γ- (a) and β- (b) secondary relaxations are represented in Fig. S7, ESI as function of temperature. For both dipolar processes, the highest

=

i

=

+

s

[1 + (i

HN )

a]b

Fig. 8. Variation of ε” with temperature at selected frequencies (a) and 3D view of ε” as a function of frequency and temperature (b) for CPI-2 sample. 564

(4)

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Fig. 9. Frequency dependencies of the dielectric loss at various temperatures for PI-2 (a and c) and CPI-2 (b and d) samples.

relaxation intensity is displayed for polyimide samples, indicating a more intense dipolar activity. The temperature dependence of the obtained relaxation times was found to respect Arrhenius equation:

(T ) =

(5)

0 exp(Ea/ KT )

where τ is the relaxation time associated with the peak maxima, Ea is the apparent activation energy of the process, k is Boltzmann’s constant, τ0 is the pre-exponential factor and represents the relaxation time at infinite temperature. The activation energy of dielectric relaxation processes is related to the potential energy barrier of electric dipoles that follows the alternative field. Table 5 exhibits the values of the activation energy for γ and β transitions obtained from the Arrhenius fit (Fig. 10). The values of activation energy for γ-relaxation are relatively low, confirming that the dipolar motions associated with this process could be localized and non-cooperative [27,35]. The obtained results showed that after introduction of diamine-diphenylether together with AMPM in the macromolecular chain the values of the activation energy slightly decreased

Fig. 10. Representation of secondary relaxation processes in the Arrhenius plots.

from 50.5 kJ/mol (PI-1) to around 44.2 kJ/mol (CPI-1). We can conclude that the presence of diphenylether units within the main chains of copolymers modifies the local interchain packing, allowing faster motions of the small chain sequences or groups. According to Table 5, the activation energy of the β-relaxation is relatively high as compared with that of γ-relaxation. This confirms the cooperativity of molecular motions for β-process [27].

Table 5 Activation energy values of the investigated polymers and copolymers. Sample

γ-relaxation τ0 (s)

PI-1 PI-2 PI-3 CPI-1 CPI-2 CPI-3

β-relaxation Eγ (kJ/mol)

−17

3.8 × 10 5.4 × 10−16 2.5 × 10−16 9.9 × 10−15 4.7 × 10−16 3.6 × 10−15

50.5 46.6 48.3 44.2 47.4 44.4

τ0 (s) −14

1.5 × 10 3 × 10−15 – – 4.3 × 10−15 9.5 × 10−16

σ-relaxation Eβ (kJ/mol)

Eσ (eV)

67.5 77.3 – – 70.6 85.1

0.59 0.92 0.7 1.83 1.21 1.94

3.5.3. Conductivity behaviour The ε′(f, T) and ε”(f, T) dependencies allow the identification of secondary relaxation processes, primary α-relaxation and a heavily dielectric signal that is generally characteristic to electrode polarization 565

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Fig. 11. Temperature dependence of M” at selected frequencies (a) and 3D representation of M” as function of frequency and temperature (b) for PI-3 sample.

and mobility of charge carriers. Employing the dielectric modulus formalism, the electric polarization is avoided and the ‘true’ conductivity signal might be furnished. Fig. 11a displays the imaginary component M” of dielectric modulus as function of temperature, at different frequencies, for PI-3 sample. Apart from the dielectric loss representation, the isochronal regime of M” behaviour exhibits an additional peak, localized at high temperatures (e.g. at f = 10 Hz, the dielectric peak appears at 138 °C) and attributed to conductivity contribution. The 3D version of M” as function of frequency and temperature (Fig. 11b) allows visualisation of the conductivity signal as a well-defined dielectric peak that shifts to higher frequencies as temperature increases. The measured conductivity, σm, was calculated with the equation:

=

0

signal is, hence, characteristic to long-range charge carriers. It is important to remark that, as temperature increases, the non-dependent frequency plateau characteristic to σDC is considerably enlarged to higher frequencies, providing that the conductivity of charge carriers is thermally activated. Comparative dependencies of measured conductivity as a function of frequency (the dependencies are selected at T = 250 °C, where both σAC and σDC regimes are revealed) for polyimide and copolyimide series are shown in Fig. 12b. The higher values of conductivity are found for polyimide series, providing that the incorporation of the second diamine in the chemical structure of copolyimides decreases the planarity and conjugation degree, and as a consequence, the conductivity. The numerical values of σm are displayed in Table 4, at 1 Hz, 103 Hz and 106 Hz. The activation energy of conductivity, Eσ, follows the Arrhenius equation:

(6)

",

where ε0 is the permitivitty of the free space and ω is the angular velocity [36]. As generally known, the total measured conductivity is a sum of AC conductivity, σAC, and DC conductivity, σDC [37]: m (f ,

T) =

AC (f ,

T) +

DC (T )

=

0 exp

E , kT

(8)

where k is the Boltzmann constant, and T is the absolute temperature [37]. The evolution of σm as function of inverse temperature (1/T), in the 200–280 °C interval is represented in Fig. 13. The Arrhenius plot representation reveals that the activation energy of polyimides is lower than for copolyimides, confirming that the incorporation of diphenylether core in the main chains of copolymers hinders the mobility of charge carriers.

(7)

The conductivity evolution with frequency at specific temperatures is presented in Fig. 12a for PI-3 sample. At low temperatures, σm increases lineary in the integral frequency domain. This behavior is specific to σAC regime and is assigned to the bounded charges. Following the increase of temperature, σm became relatively independent by frequency, especially at low frequencies, revealing the presence of DC conductivity signal, contributing nothing to dielecric polarization. This

Fig. 12. The evolution of σm with frequency for PI-3 sample (a) and a comparative evolution of σm with frequency for polyimide and copolyimide series. 566

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with other results concerning gas permeation in glassy polymers found in the literature as well as by us earlier [38–40]. As can be seen from the permeability data, the presence of F and Br substituents in the pendent phenyl ring in the macromolecular chain of polyimides and copolyimides slightly increases their permeability that is accompanied by a slight decrease in selectivity. On the other hand, the introduction of the co-monomer - diamino-diphenylether reduces somewhat permeability and increases selectivity in comparison with pristine polyimides (Table 6). The both PI-2 and CPI-2 with F substituent show the highest permeability within the given series, and the lowest selectivity in accordance with the trade-off behavior [6]. Although the differences in transport properties among the polymers and copolymers are very small, the observed trends are clearly visible. When correlating transport parameters of the polymers with their Tg values (Table 2), which are the measure of the chain stiffness, it can be seen that higher permeability goes together with higher chain stiffness of the polymer, which is the opposite than expected. It may be explained as due to the inability of the polymers with rigid chain to pack effectively that results in a more open structure and, in turn, higher permeation rate. Interestingly, the introduction of the co-monomer with additional flexible ether linkages, which enhance segmental mobility, reduces gas permeability. This result indicates that flexible linkages allow for closer packing and that packing density is the parameter which mainly controls gas permeability through the investigated polymers. This conclusion is consistent with the results of X-ray diffraction studies concerning unsubstituted PI-1 and CPI-1, where the dspacing value of CPI-1 is lower than that of PI-1. However, any such a correlation has been noticed for both the F and Br substituted PI and CPI. These polymers show higher permeation rates than the unsubstituted ones despite their lower d-spacing values suggesting more complex relationship between structure and gas permeation. In Fig. 14, the results of gas permeability measurements from the present study and two literature examples for common polyimides are presented together with the Robeson’s 2008 upper bounds [6] for O2/ N2, CO2/N2, and He/N2 gas pairs, the separation of which can be of

Fig. 13. Arrhenius presentation of σ-relaxation for polyimide and copolyimide series. Table 6 Gas permeation and separation properties of the polyimide films at 35 °C. Polymer

P (N2) [Barrer]

P (O2) [Barrer]

P (He) Barrer]

P (CO2) [Barrer]

α O2/N2

α CO2/N2

α He/N2

PI-1 PI-2 PI-3 CPI-1 CPI-2 CPI-3

0.06 0.07 0.07 0.05 0.07 0.06

0.38 0.42 0.41 0.35 0.39 0.37

6.12 7.84 6.91 5.93 6.46 6.07

1.81 1.92 1.89 1.50 1.74 1.60

6.02 5.73 5.74 6.57 5.89 6.17

28.7 25.9 26.3 28.3 26.4 26.7

102 112 98.7 118.6 92.3 101.2

3.6. Gas transport properties Gas permeation results for the polyimide films prepared in this work are given in Table 6. The observed order of increasing permeability for the tested gases is as follows: N2 < O2 < CO2 < He, which is the order of their decreasing kinetic diameter, and that is in accordance

Fig. 14. Robeson’s upper bound correlations for (a) O2/N2, (b) CO2/N2, and (c) He/N2 separations. 567

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commercial interest. As can be seen from this figure, all of the synthesized polyimides and copolyimides show a relatively low performance, which places them well below the respective upper bounds, particularly in case of CO2/N2 gas pair. However, compared with the both commercial polyimides, the studied films exhibit better selectivity than Matrimid, however at the expense of permeability, as well as better permeability than PMDA-4,4′-ODA accompanied by only slightly lower selectivity. Overall, the performance of those materials, though not attractive from the applications point of view, may help to indicate a direction for future material design.

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4. Conclusions In summary, triphenylmethane-based polyimides and copoyimides with improved solubility and nice balance of properties were successfully synthesized and characterized. The reported polymers are based on three novel triphenylmethane-based diamines which combine some structural elements, like two methyl groups and fluoro/bromo atoms substituted to the benzene rings. Two-step polycondensation reaction of 4,4′-(4,4′-isopropylidenediphenoxy)bis(phthalic anhydride) with one of the newly obtained diamines or with an equimolar mixture of 4,4′oxydianiline with a triphenylmethane-based diamine led to triphenylmethane-based polyimides and copolyimides. Due to the bulky pendent phenyl units in the triphenylmethane core and generation of non-coplanar structures, the polymers were soluble in organic solvents, thus being appropriate for processing into free-standing films with good mechanical properties. The polymer films had good thermal stability, their decomposition starting above 425 °C, and relatively low glass transition temperature. The dielectric constant measured between 1 Hz and 1 MHz, at room temperature, ranged between 2.94 and 4.09, lower values being obtained for F- and Br-substituted polymers compared to unsubstituted analogues. At low and moderate temperature, the studied polymer films showed γ- and β-relaxations connected with local movements of polymer chain segments. The activation energy for γ- and β-relaxation processes was found in the range of 44–50 kJ/mol and 67.5–85.1 kJ/mol, respectively, with lower values for copolyimide films. Interestingly, the polyimide series exhibits higher values of conductivity than those found for copolyimide series, which endow the polymers with semconducting properties at high temperature and frequencies. For both series, homo- and copolymers, Br- and F- substitution increased slightly the gas permeability that was accompanied by a small decrease in selectivity. Incorporation of the diphenylether units into the triphenylmethane-based polyimide chains had the reverse effect, of a slight permeability reduction combined with a small selectivity increase suggesting the existence of a more complex relationship between polymer structure and gas permeation. Owing to the appealing characteristics mediated by the triphenylmethane core, the studied polyimides and copolyimides may find applications as new types of high-temperature polymeric materials, e.g. as dielectric films of high thermal stability in electronic devices or membranes for specific gas separation applications. Acknowledgements This work was supported by a grant of Ministry of Research and Innovation, CNCS-UEFISCDI, code PN-III-P4-ID-PCE-2016-0708, no. 66/2017, within PNCDI III. Appendix A. Supplementary material Supplementary data to this article can be found online at https:// doi.org/10.1016/j.eurpolymj.2018.09.029. 568

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