Introduction

High performance polymers cover a wide spectrum of materials including polyimides, polyamides, poly(ether-ether ketone)s, polysulfones, polycarbonates etc., which play an increasingly important role in our daily life, compared with commodity polymers. They are valued, amongst other properties, for their temperature resistance, dimensional stability, strength and chemical resistance in several demanding applications [1, 2]. Thus, the high performance polymers have gained growing interests in aerospace, national defence and many industrial applications, such as electronics, medical or transportation [3–8]. Poly(ether-imide)s are derivatives of polyimide materials; they contain flexible ether groups which were introduced between rigid and highly polar imide units in order to improve the melt processability while maintaining the outstanding properties of polyimides. The better processability and superior properties make poly(ether-imide)s suitable replacement of polyimides in various applications, including aerospace, ground transportation as well as electronics. Several key performance features of poly(ether-imide)s include excellent dimensional stability at high temperatures under load, transparency, optical properties, and excellent aesthetics, high strength and modulus, good solvent resistance over a wide temperature range, high continuous-use temperature, high oxygen index and low smoke, medical- and food-contact compliance, good electrical properties with low ion content, excellent machinability etc. [9–17].

Aromatic poly(ether-imide)s have been materials of choice in high-performance microelectronic applications because of their high degree of ductility, inherently low thermal expansion coefficient and low dielectric constants [18, 19]. Their low loss behaviour and high glass transition temperature make them valuable alternative to some of conventional electrical insulation materials and dielectrics, i.e., polypropylene based thin film electrostatic capacitors [20, 21]. Previous studies on poly(ether-imide)s have indicated that the dielectric behaviour as a function of temperature and frequency were different in strong relation with the chemical structure [22] and the moisture content [23]. The behaviour of charge mobility and polarization determines the electrostatic charging of the film and how it can be dissipated during the production of capacitor devices [24].

Recently, the specifications for insulating films are that their dielectric constants should be ca. 3.0 and, within the next generation of integrated circuit production, devices may require for materials to have dielectric constants approaching or below 2.0. A low dielectric constant is one of the most attractive properties of poly(ether-imide) materials for microelectronics applications. In order to achieve a polymer structure with a low dielectric constant, repeating units with low polarity and low polarizability are needed [25].

On the other hand, aromatic poly(1,3,4-oxadiazole)s have been widely studied owing to their high thermal stability in oxidative atmosphere and specific properties determined by the electronic structure of 1,3,4-oxadiazole ring [26] Besides their outstanding resistance to high temperature, polyoxadiazoles are known to have a combination of many valuable properties, such as good hydrolytic stability, high glass transition temperature, tough mechanical properties, and low dielectric constants [27–32]. Polyoxadiazoles exhibit excellent fiber- and film-forming capability, which provides great potential for applications in fields of aerospace, military, protective garments, filtering fabrics and general industry. More recently, specific properties determined by the electronic structure of oxadiazole ring, particularly its electron-withdrawing character, stimulated an intensive research on these polymers with the aim to use them as electron-transporting materials or light-emitting layers in electroluminescent devices.

Poly(ether-imide)s are a particularly attractive class of polymers for structure-property investigation as the synthetic chemistry is flexible and offer a great potential for introducing numerous different structural moieties. Our aim was to investigate structural features which influence the photo-physical behaviour of new materials obtained by the incorporation of oxadiazole together with imide rings, flexible ether linkages and voluminous units as isopropylidene, hexafluoroisopropylidene, fluorene or cyclohexane into the polymer chains, expecting to provide a good combination of high performance properties and good processability, particularly in thin films and coatings. This paper is dealing with the study of the mechanical, optical and dielectric properties, as well as other physical processes taking place during heating in thin films made from poly(ether-imide)s containing 1,3,4-oxadiazole heterocycle.

Materials and methods

Synthesis and basic characterization of poly(ether-imide) powders

Poly(ether-imide)s P1–P4 containing 1,3,4-oxadiazole rings were synthesized as described elsewhere [33]. Briefly, the polycondensation reaction was carried out in 1-methyl-2-pyrrolidone (NMP) solution at high temperatures (200 °C), at a solid content of 12–14 %, without the isolation of the intermediate polyamidic acid (Scheme 1). The clear resulting poly(ether-imide) solutions were poured into water to precipitate the solid polymers, followed by washing with plenty of water and finally with ethanol.

Scheme 1
scheme 1

Synthesis of poly(ether-imide)s containing oxadiazole rings

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The structure of polymers was identified by FTIR spectra that provided evidence for the formation of imide groups by characteristic imide ring absorptions in the range of 1770–1780 cm−1 (asymmetrical C = O imide stretching), 1710–1720 cm−1 (symmetrical C = O imide stretching) and 720–740 cm−1 (imide ring deformation). The strong absorption bands in the range of 1236–1247 cm−1 were linked to the presence of aromatic ether stretching, while characteristic absorption bands for oxadiazole ring were found at 970–960 cm−1 and 1020–1010 cm−1. Aliphatic linkages C-H in CH3 and C6H10 groups of polyimides P1 and P4, respectively, displayed absorption peaks at 2950 cm−1. In the FTIR spectrum of polymer P2 the hexafluoroisopropyllidene (6 F) group was identified by the absorption peak at 1210 cm−1. The solubility behaviour was tested qualitatively in various organic solvents. All these poly(ether-imide)s were easily soluble at room temperature in polar aprotic solvents such as N-methylpyrrolidone (NMP), dimethylformamide (DMF), dimethylsulfoxide (DMSO) and dimethylacetamide (DMAc) and even in less polar solvents such as tetrahydrofurane (THF) and chloroform, which are convenient and easy accessible solvents.

The molecular weight values Mw measured by gel permeation chromatography are in the range of 162,000–139,500 g/mol, Mn in the range of 107,800–93,700 g/mol and polydispersity Mw/Mn in the range of 1.30–1.50. The glass transition temperature values of the polymer powders taken from the second heating curves obtained in DSC plots were found to be in the range of 188–232 °C. The polymers showed excellent thermal stability, as expected in case of aromatic polyimides. They began to decompose in the range of 470–485 °C, as indicated by the temperature of 5 % weight loss in thermogravimetric analysis, indicating a high thermal stability. An in-depth study regarding the structural identification and the high performance properties of these polymers was published elsewhere [33].

Preparation of polymer films

Films of polymers P1–P4 were prepared by casting a 10 % polymer solution in chloroform onto glass plates, followed by heating from room temperature up to 60 °C for 10 h to remove the residual solvent. The resulting light-yellow, transparent coatings displayed strong adhesion to the glass support. The self-standing films were stripped off the plates by immersion in boiling water, followed by drying in oven at 110 °C. These films had the thickness in the range of 40–50 μm and were used afterwards for various measurements.

Measurements

Mechanical properties of the polymer films were determined by tensile testing using an Instron 5566 apparatus. The samples were used in the form of strips having the thickness of 0.05 mm, gauge length of 25 mm and width of 15 mm. Stiffness, tensile strength and elongation-at-break were evaluated at 10 mm/min cross-head speed. The dependencies of tensile stress (MPa) versus tensile strain (%) were recorded.

Dynamic mechanical analysis (DMA) was performed in tension mode on a Perkin Elmer Diamond instrument by using polyamide films with dimensions 10 × 10 × 0.05 mm. The temperature scan was run with 2 °C/min, at 1 Hz, from 0 °C up to the temperature where the sample became too soft to allow the experiment to be continued. The variation of the storage modulus (E’), loss modulus (E”) and loss tangent (tanδ) as a function of temperature was obtained.

Broadband dielectric spectroscopy (BDS) measurements were run in broad temperature and frequency range by Novocontrol Dielectric Spectrometer Concept 40 (Germany). Data were expressed as complex permittivity ε * (f) = ε ′ (f) − ε ″ (f), where ε’ and ε” are the relative permittivity and dielectric loss, respectively. The relative permittivity (ε’) of a material, also commonly known as dielectric constant, is its dielectric permittivity expressed as a ratio of absolute permittivity (SI units: F/m) to the permittivity of free space or vacuum (ε o = 8.85 F/m) and is dimensionless number. The dielectric loss (ε”) is quantized by a number known as the dielectric loss factor or dielectric loss index and is calculated as the product of the dielectric constant and the tangent of the dielectric loss angle. The measurements were recorded at a constant temperature by frequency scans between 10 Hz to 1 MHz, at every 5 °C between −120 °C and +250 °C, either at a constant frequency by temperature scan from −120 °C up to +250 °C. Polyimide films with thickness of about 50 μm placed between two circular steel electrodes with the upper electrode having 20 mm diameter were used for those measurements. The samples were placed inside a cryostat with controlled temperature in nitrogen atmosphere and the amplitude of AC applied voltage was 1 V.

The optical transmission spectra of the polymer films were registered with a Specord M42 apparatus, by using as-prepared free standing films.

Results and discussions

All the studied poly(oxadiazole-ether-imide)s gave transparent and light-yellow free-standing and flexible films that resisted to repeated bending. The toughness of these films allowed their easy mechanical handling. Their elastic modulus, tensile strength and elongation to break have been determined as averages of three to five drawing experiments. The values of tensile strength were in the range of 86–102 MPa, elastic modulus in the range of 2.32–3.60 GPa and elongation to break in the range of 5.12–7.68 % [33]. The tensile tests of polymers P1–P4 are displayed in Fig. 1. Thus, the investigated polymers are suitable candidates for use as advanced materials and, therefore, they were subjected to dynamo-mechanical, dielectric behaviour and optical clarity investigations.

Fig. 1
figure 1

Tensile tests of poly(ether-imide)s P1–P4

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Dielectric properties

Recently, low dielectric constant (low-k) is one of the most desirable properties of next-generation electronic devices [34, 35]. Therefore, polyimides with low dielectric constants are required to increase the efficiency at which a microelectronic device transmits signals. Usually, polyimides containing few polarizable atoms or functional groups display low dielectric constants, in all frequency ranges, as expected in the case of poly(ether-imide) films P1–P4, as well. In the dielectric analysis, a sinusoidal voltage is applied to create an alternating electric field, and the measured current is expressed in a complex form (ε*), which can be divided into capacitive and conductive components, known as the permittivity (ε’) and dielectric loss (ε”), respectively. The relative permittivity (or dielectric constant) refers to the polarizability of a material and is strongly dependent on its chemical structure [36, 37]. The dielectric loss is a measure of the energy required for molecular motion in the presence of an electric field and consists of two contributions: energy losses due to the orientation of molecular dipoles, and energy losses due to the conduction of ionic species.

The dielectric properties of oxadiazole-containing poly(ether-imide)s films P1–P4 were evaluated on the basis of the dielectric constant and dielectric loss and their variation with frequency and temperature. Figure 2 shows the dependence of the real and imaginary part of the complex permittivity, on frequency, at low (−150 °C), room (25 °C) and high (150 °C) temperatures.

Fig. 2
figure 2

Relative permittivity (ε’) and dielectric loss (ε”) versus frequency, at different temperatures, for P1–P4

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At low, room and high temperatures, the relative permittivity (or dielectric constant) of those polymers increases very slowly, linearly, with decreasing the frequency, taking low values even at low frequencies (10 Hz). The dielectric constant of polymers decreases gradually with increasing frequency because the response of the electronic, atomic and dipolar polarizable moieties vary with frequency in the specific domain. The ability of polarizable units to orient fast enough as to follow the oscillation of the alternative electric field is reflected in the relative permittivity (ε’) values [38, 39]. With increasing frequency, there is not enough time for the dipoles and ions to orient themselves in the alternating electric fields and ε’ decreases. Figure 2 shows that the increasing rate of the relative permittivity with decreasing frequency for P1–P4 polymer films at all investigated temperatures are quite similar, demonstrating their likewise polarization capability. A higher increase rate of the relative permittivity can be observed at high temperature (150 °C) due to the mobility of the charge carriers. However, P1–P4 exhibit low relative permittivity values over the whole domain of frequencies (10–106 Hz), even at 150 °C. Dielectric materials with stable dielectric constant and dissipation factor values across large frequency and temperature ranges are highly preferred for majority of electronic applications.

The relative permittivity values of the poly(ether-imide) films P1–P4 at 10, 103 and 105 Hz and room temperature (25 °C) are listed in Table 1.

Table 1 Dielectric data of oxadiazole-containing poly(ether-imide) films at selected frequencies and 25 °C
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Poly(ether-imide)s P1–P4 display relative permittivity (or dielectric constant) values at room temperature and in the frequency range of 10 Hz–1 MHz between 2.87 and 3.23, being lower or comparable with that of Kapton HN polyimide film, whose dielectric constant values, measured under the same conditions, range from 3.13 to 3.24 [40]. The relative permittivity values of the investigated polymer films depend on their chemical structure, chain conformation and molecular packing. The lowest value of the relative permittivity was achieved for P3 and the highest one for P1 film. This behaviour was expected for P3 since it has the highest content of aromatic units in the dianhydride segment and the most voluminous moiety - fluorene, which decreases the chain packing and increases the free volume, thus decreasing the relative permittivity of the polymer film. On the other hand, with decreasing the aromatic component and increasing the aliphatic segment, the relative permittivity increases for P1 most probably due to a denser packing which leads to a decrease of the free volume of polymer. The introduction of CF3 groups in the structure of poly(ether-imide) P2 results in the decrease of the relative permittivity compared with P1 since flexible hexafluoroisopropylidene groups decreases the chain packing and increases the free volume of the polymers, lowering the polarization by decreasing the number of polarizable groups per volume unit and thus decreasing the relative permittivity [41]. A similar effect was observed in the case of P4 by introduction of a cycloaliphatic moiety with low polarity in the dianhydride segment, its relative permittivity values being very close to that of P2 on the entire temperature and frequency range.

Beside lower dielectric constant, low dielectric loss is also desirable since it mirrors the efficiency of a dielectric material. At room and high temperatures, P1–P4 show a maximum on the dielectric loss vs. frequency plot, being centred at approx. 105 Hz at 25 °C and 103–104 Hz at 150 °C. These maxima correspond to the relaxation processes associated with small or large chain segments mobility [42]. Since the value of glass transition temperatures for these polymers are with approximately 40 °C higher than the set temperature (see Table 2), we may attribute these peaks to the secondary relaxations that appear in the glassy state, γ and β relaxations, respectively, being known that they shift to higher temperatures with increasing frequency. At very low temperatures, all polymer films, except P2, showed a very low constant dielectric loss on the whole range of the measured frequencies. P2 showed a sharp increase of the dielectric loss with increasing frequency only in the high frequency range (105–106Hz) due to the mobility of charge carriers. The dielectric loss values of the studied films at 10, 103 and 105 Hz and room temperature (25 °C) are collected in Table 1. All the films presented low values of dielectric loss, comparable with that of Kapton HN [40]. The obtained low values of the dielectric loss (below 0.012) reflect a minimal dissipation of energy in the form of heat in the dielectric material.

Table 2 Dynamic-mechanical data of poly(ether-imide)s P1–P4
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Relaxation behaviour

Viscoelastic relaxation behavior

Since plastic materials are used in many engineering designs, the dynamic stress-strain response of these materials is of considerable interest. The dynamic response of amorphous polymers requires a careful examination of both the temperature rise due to rapid plastic deformation and the molecular rearrangement of polymer chains. The dynamic mechanical analysis (DMA) offers a valuable help in understanding some physical phenomena tacking place during heating, being sensitive to molecular motion for determination of transitions and transformations in polymers [43]. Table 2 presents the main results taken from the DMA curves of polymers P1–P4.

The drops in E’ curves and the peaks on E” and tan δ curves indicate the physical transitions in polymers. Usually, the transition temperatures are recorded at the maximum rate of turndown of the storage modulus E’ or at the maximum loss modulus E” or of tan δ peaks. At low temperature, the values of E’ are over 109 Pa, while the E” lies at about 108 Pa, which is typical of glassy polymers. This is the glassy region where the macromolecules are in the “frozen” state. With rising temperature, a small increase of E” and tanδ occurred as a result of β relaxation. The β relaxation being associated with rotational vibrations of different groups coming from the diamine segment either from the dianhydride unit extends on a broad temperature range, until the beginning of the α relaxation, except for P3 (Fig. 3). A qualitative comparison of the β relaxations of the four poly(ether-imide)s showed one shoulder-like peak centred at 42, 15 and 75 °C for P2, P3 and P4, respectively. P1 presented two more intense β relaxation peaks, at 29 and 130 °C, suggesting that the local motions of the small groups and fragments of the monomeric units are much favoured compared with the other polymers. The β relaxation is less resolved in the case of P3 as a result of more restricted local motions associated with this transition.

Fig. 3
figure 3

β Relaxation region of oxadiazole-containing poly(ether-imide)s P1–P4

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The chemical composition and conformational features dictates the α relaxation of the polymers that is associated with the glass transition region. Table 2 presents the temperature of the α-relaxation as the value of E’ drop onset, the peak of E” and tanδ. Figure 4 shows the elastic component - storage modulus (E’), the viscous component - loss modulus (E”), and the loss factor tangent (tanδ = E”/E’) dependences on temperature for the films of poly(ether-imide)s P1 and P4.

Fig. 4
figure 4

The variation of E’, E” and tanδ with temperature for poly(oxadiazole-ether-imide)s P1 and P4

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In the glassy state, below the α-relaxation, tan δ < 0.1 for all polymers which means that the elastic behaviour is dominant. As the α-relaxation occurs, E” becomes higher than E’ (tan δ > 1) in a certain range of temperature for linear amorphous polymers and the film elongates. Tanδ drops back to pretransition values and the polymer recovers its original length when the α-relaxation is complete [44]. P1 and P4 followed such standard behaviour (Fig. 4) and displayed two cross points between E’ and E” (201 °C and 212 °C for P1 and 222 °C and 243 °C for P4) during α-relaxation and one cross point on the rubbery plateau region (at 235 °C and 243 °C for P1 and P4, respectively) after that the α-relaxation was completed. After α-relaxation, poly(ether-imide)s P1 and P4 gained some extra-mobility since at the temperature higher than at 235 °C for P1 and 243 °C for P4, the viscous behaviour is again dominant (E” > E’). It can be assigned to the flowing process [45], when the polymer chains start to slip past one another to some extent along the main macromolecular axis. The aliphatic flexible units, namely isopropylidene and cyclohexylidene, introduced in the main chains of polyimides P1 and P4, respectively, appear to be responsible for this behaviour.

The fluorene-containing polyimide P3 displayed only one cross point between E’ and E” at 245 °C and a pronounced tanδ increase during α-relaxation, the maximum value of tanδ reaching to 4. It seems that the bulky fluorene moiety improves the flexibility of the main P3 macromolecules and the flowing becomes more intense during α-relaxation, the E’ remaining lower than E” even at the end of the α-relaxation. With increasing temperature a continuous flowing occurs being accomplished by a continuous steep increase of tanδ and an abrupt decrease of E’ and E” being stopped by the rupture of the film. In the case of P2 containing voluminous hexafluoroisipropylidene units, the sample became too soft immediately after reaching α-relaxation to allow the experiment to continue and consequently no cross point between E’ and E” was registered.

The Tg values determined by DMA measurements were slightly higher than those recorded by DSC measurements, except P2 for which an accurate evaluation of Tg was not possible since the film broken ripped during α-relaxation. Usually, the values obtained from these two methods can differ up to 25 degrees from each other in the same material. DMA is a measurement of the dynamic moduli in an oscillatory mechanical deformation experiment during a programmed temperature scan at controlled frequency. The peak of the tan δ is a particularly discriminatory measure of Tg, although this is the center of the α relaxation, whereas in the DSC experiment the onset temperature of the α relaxation (Tg) is usually reported. In such a case the Tg measured by DSC will be lower than the one obtained by DMA by an amount that varies with the specific polymer. There is, in addition, a frequency effect which puts the mechanical Tg (measured at ca.1 Hz) higher than that for a DSC Tg (measured at ca. 0.0001 Hz) for an assumed activation energy of 400 KJ/mole which is typical for polymer Tg [46, 47].

The magnitude of the tanδ at the Tg represents a measure of the energy-damping characteristics of the polymer film and is related with the impact strength of the material. The results listed in Table 2 show that the energy-damping characteristics of the investigated polymers are very similar, very close to 2, with only small differences by a factor of 0.07 up to 0.17. However, it appears that the P4 film has the highest impact strength, since the magnitude of tanδ is superior to that of the other polymers, whereas the P1 proves to be more fragile in the series. These findings are in total agreement with the results of the tensile tests of these polyimide films [33] that revealed the highest values of the tensile strength for P4 (103 MPa) and the lowest one for P1 (77 MPa).

Dielectric relaxation behaviour

Dielectric relaxation techniques have been widely applied to establish transition temperatures, relative relaxation intensity, and time-temperature characteristics of the motional transitions encountered in these materials. Usually, the thermal transitions associated with dipolar relaxations are evidenced from the frequency and temperature dependence of the relative permittivity and dielectric loss of a material.

Dielectric data for γ, β and α transitions of oxadiazole-containing poly(ether-imide)s P1 and P4 are presented in Fig. 5 as representative examples. Figure 5 displays the variation of relative permittivity and dielectric loss with temperature at several frequencies between 1 and 106 Hz, in the first heating cycle.

Fig. 5
figure 5

Relative permittivity and dielectric loss versus temperature for P1 and P4 at various frequencies (the inset shows a magnified view of the subglass transition region)

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A step increase in a small range between 0.1 and 0.2 units in the dielectric relative permittivity at the temperature of γ and β relaxations was observed for all polymers. The segmental mobility allows for alignment of the dipoles, which results in an increase of the relative permittivity and causes the material to become less insulating. The relative permittivity also decreases with increasing temperature in the temperature range from 30 °C to 80 °C. This process is detectable only for higher frequencies (inset in Fig. 5) because, at lower frequencies, it is overlapped by the conductive contributions. It is possible that some polar molecules (like water) adsorbed during storage are removed from the polyimide films on heating. Similar processes were observed in other aromatic polyimides or polyamides [39, 48]. In the second heating cycle, the relative permittivity is no longer decreasing with temperature and the γ relaxation intensity is reduced to some extent. With rising temperature, the relative permittivity values increased gradually up to the glass transition temperature. Below Tg the thermal energy absorbed by internal dipoles as well as dipolar groups of P1–P4 is still small and only a reduced number of dipoles can be oriented in the direction of the applied electric field. After Tg, the magnitude of ε’ sharply advances with further increase in the temperature since the dipoles of polymers gained enough energy to orient themselves in the direction of the applied field, thus leading to a high polarization of the investigated sample. The relative permittivity enhance with the temperature can be discussed as well in terms of disentanglement of the molecular chains that becomes easier at higher temperatures due to the temperature induced molecular vibrations [49].

The dielectric loss data also show distinct sub-glass and glass relaxations as clear peaks on ε” versus temperature curves for all polymers. The temperature corresponding to the centre of each relaxation is gathered in Table 3. A representative 3D plot showing γ, β and α relaxations of these polymer series is illustrated in Fig. 6.

Table 3 Dielectric relaxation data of poly(ether-imide)s P1–P4
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Fig. 6
figure 6

3D plot showing the subglass and glass relaxations of poly(ether-imide) P2

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Usually the γ transition of polyimide films is associated with phenyl ring oscillations and is influenced by moisture absorption content, aging history and morphology [50, 51]. Moreover, it was demonstrated that lower interchain interactions, as is the case of P1–P4, result in looser chain packing, thus facilitating the molecular process that generates the γ-relaxation [52]. The temperature at which γ relaxation appears is very close for all investigated polymers, respectively between −88 and −101 °C at 1Hz. At room temperature, the process is followed by other relaxation (β) with a reduced intensity, observed as a peak or shoulder. It appears around 58, 85, 61 and 69 °C for P1, P2 P3 and P4, respectively, at 1Hz. β Transition of polyimides was assigned to the segmental motion in diamine or dianhydride moiety [53, 54], but may comprise larger portions of the structural unit that respond in a correlated manner [55, 56] or can ultimately involve the entire repeating segment, as it is the case of rigid systems. The intensity of γ and β relaxations of P1–P4 decreases with temperature reflecting an apparent loss of net dipolar correlation with increasing thermal energy. For lower frequencies, on the low temperature side, β relaxation of P4 is overlapped by the effects of removal of polar molecules, since the latter is no longer visible in the second heating cycle. With increasing temperature, ε” undergoes a single relaxation peak for all polymer samples. It is reasonable to attribute this peak to the α relaxation process. It was found at 205, 215, 250 and 213 °C for P1, P2 P3 and P4, respectively, at 1Hz, in a perfect correlation with those found by DMA at similar frequency. The peak characteristic for α relaxation shifts to higher temperature with increasing frequency because the segmental chain mobility allows for the molecular dipoles to respond to changes in frequency.

The molecular origins of the sub-Tg relaxations are clarified by using data from the dielectric sub-glass transition relaxations in an Arrhenius activation energy analysis. The magnitude of the activation energy (Ea) of a relaxation depends on rotation potential energy barriers, e.g. internal friction, the volume and the environment of the moving repeat units. The activation energy of γ, β and α relaxations of each polyimide was calculated by means of Arrhenius equation:

$$ f=A\; exp\left(-Ea/RT\right) $$

using the plot of the frequency logarithm vs. reciprocal temperature of relaxation:

$$ lnf=lnA-Ea/RT $$

where f is the frequency (Hz), A is the pre-exponential factor, R is the gas constant (8.314 J mol−1 K−1), T is the absolute temperature (K) and Ea is the activation energy (kJ mol−1). Table 3 exhibits the values of the activation energy for γ, β and α transitions obtained from the Arrhenius fit (Fig. 7).

Fig. 7
figure 7

Arrhenius plot for γ, β and α relaxations of poly(ether-imide)s P1–P4

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The γ relaxation of poly(ether-imide) films P1–P4 is characterized by relatively small activation energy values that vary between 40.9 and 48.37 kJ/mol and are consistent with previous studies on aromatic polyimides [40, 57], suggesting a common molecular origin for γ relaxation of all these polymers. The motions associated with this relaxation may be considered localized and non-cooperative [58].

Dielectric activation energies for the β relaxation of aromatic polyimides have been reported in the range of 95–180 kJ/mol. The β relaxation Ea values for P1–P4 ranged between 75.52 and 112.12 kJ/mol. The differences in the activation energy values of β relaxation of P1–P4 are mainly caused by the motion in the dianhydride moiety, since all polymer chains contain the same diamine segment. The introduction of a flexible or voluminous unit in the dianhydride segment leads to a more flexible structure and, therefore, the length of mobile segments is considerably increased, in strong correlation with its nature. The lowest activation energy for β relaxation was found for P3 that contains the most voluminous moiety - fluorene - causing a decrease chain packing and a higher free volume. Thus, for this system, the origin of the β relaxation may be due to large amplitude movements of the dianhydride segment. With increasing frequency, β-relaxation shifts toward higher temperatures in accordance with Arrhenius law. Sensitivity of β-relaxation to frequency is due to its lower activation energy.

Within Fig. 6, the slopes of the lines for the α-relaxation are steeper than those for the γ and β relaxations. The high slope indicates a much higher activation energy and a different volume and environment of the moving segments. The Arrhenius activation energies were calculated using the slopes of the best-fit lines and were found for P1–P4 in the range of 460.06–736.29 kJ/mol, being in agreement with those values reported for related polyimides [44, 59]. The Ea values of α-relaxation are much higher than for the local, sub-glass relaxations. This is characteristic of the thermal energy needed to activate the highly cooperative chain motions involved in the α-relaxation [60]. The highest Ea value was found for P3 since its segmental polymer chain motion is more restricted due to the bulky fluorene units. These results can be understood in terms of increase of free volume and motion constrains imposed by this moiety.

Optical transparency

Colorless polymers with high thermal and mechanical stability as well as good optical transparency have widely been studied as prospective flexible substrates for electronic and micro optical devices in the field of displays, memory, lighting, solar cell, sensor, and waveguide etc., in order to replace the classical glass substrate. In particular, aromatic polyimide films display high glass-transition and decomposition temperature, low water uptake, good transmission and low optical loss, and therefore, they are good candidates for use as transparent flexible substrates. Aromatic polyimides generally exhibit pale yellow to brown colour with characteristic absorption tailings in the visible region caused by the intra-and inter-molecular charge transfer complexes (CTC) of the polyimide backbones [61]. One way to overcome this yellowish colour is the introduction of controlled steric hindrance in the main chain caused by voluminous atoms, like fluorine or chlorine or bulky moieties. Since poly(ether-imide)s P1–P4 were modified in this respect, we have made a study on the optical transmission of the films made from these polymers by UV-vis spectroscopy.

Figure 8 shows the optical transmission spectra of the polyimide films with a thickness of ~50 μm prepared via thermal imidization. Table 4 summarizes the cut-off wavelength values (λ0) estimated by the point where the optical transmission curve intersects a bisected line drawn through the intersection of the extrapolations of the two slopes.

Fig. 8
figure 8

Optical transmission spectra of poly(ether-imide) P1—P4 films

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Table 4 Optical properties of poly(ether-imide) P1–P4 films
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Poly(ether-imide)s P1–P4 reveal low λ0 and high optical transparency. The values of cut-off wavelengths are very close, between 360 and 366 nm and decrease in the order P2 > P1 > P3 > P4. The percentage of light transmittance is higher than 66 % at 450 nm, whereas the average transmittance in the visible spectral range is higher than 76 %. As shown in Fig. 8, the optical transmission spectra of polymers P3 and P4 show the highest transparency in the visible region, over 80 %. The introduction of voluminous fluorene and cyclohexylidene moieties was highly effective in reducing of the intermolecular charge-transfer complex CTC formation between alternating electron-donor (diamine) and electron-acceptor (dianhydride) moieties. These bulky units have dropped the CTC formation between polymer chains mainly through the steric hindrance, thus lowering intermolecular forces. Unexpectedly, the optical clarity of P2 is lower in the series; being known that the bulky and electron-withdrawing CF3- group is effective in reducing the CTC formation between polymer chains, there might be other factors that affect optical properties. Previous studies proved that the colour of a polyimide film, consequently its optical transparency, is affected not only by CT interactions which are driven by the polymer chain structures, but is also sensitive to processing conditions like monomer purity, kind of solvent, film thickness and cure conditions [61]. Moreover, the use of NMP as solvent for polyimide synthesis results in more intensely coloured films due to the traces of oxidized residual solvent that is retained in the virtue of hydrogen bonds. Thus, any of the mentioned factors can be responsible for a reduced transparency of the film prepared from fluorinated poly(ether-imide) P2 compared to others. However, it displays still low cut-off wavelength and high optical transparency, with an average transmittance in visible region of 76 %, as to allow its use in optical applications.

Conclusions

By introduction of oxadiazole together with imide rings, flexible ether linkages and voluminous moieties as isopropylidene, hexafluoroisopropylidene, fluorene or cyclohexane into the polymer chains, products with substantially improved solubility and processability into thin flexible and tough films were obtained. These polymer films displayed relative permittivity values in the range of 2.84–3.26 and low dielectric loss, below 0.012, at room temperature and in the frequency domain of 10–106 Hz. The low values for both relative permittivity and dielectric loss, comparable with those of Kapton HN film, indicate a low electrical signals loss in the dielectric medium and, therefore, these polymers are potential candidates for interlayer dielectric applications. The dielectric spectroscopy data showed distinct subglass transitions connected with the local motion in the poly(ether-imide) chains: γ relaxation with the activation energy between 40.9 and 48.37 kJ/mol, and β relaxation with the activation energy of 75.52–112.12 kJ/mol. The highly cooperative chain motions involved in the α-relaxation have been also studied, the highest activation energy being found for fluorene-containing polyimide, for which the segmental chain motion is more restricted due to the bulky fluorene units. The local motions of the small groups and fragments of the monomeric units were much favoured for the isopropylidene-containing poly(ether-imide) compared with the others, as it was evidenced by dynamic mechanical analysis. The relationship between the chemical composition and elastic / viscous components ratio in the α-relaxation region with increasing temperature was discussed in detail. During the α-relaxation the trend of storage modulus, loss modulus and loss factor tangent with temperature revealed the sequence of flowing and elongation phenomena for two of the polymers. After α-relaxation, all polymers gained increased mobility and the flowing process was dominant. The values of the tan δ magnitude which is related to the impact strength of the material are very similar and close to 2 for all investigated polymers, with only small differences, by a factor of 0.07 up to 0.17, in a good relation with the tensile test results. Each polyimide film was transparent in the visual light range and revealed low cut-off wavelengths and high optical transparency, with a percentage of light transmittance higher than 66 % at 450 nm and average transmittance in the visible region between 76 and 82 %. The useful properties of these poly(ether-imide)s can now be exploited in applications as dielectric transparent flexible substrates.