Eur. Phys. J. Appl. Phys. 80, 10901 (2017)© EDP Sciences, 2017DOI: 10.1051/epjap/2017160451

THE EUROPEANPHYSICAL JOURNAL

Regular Article

APPLIED PHYSICS

Mechanical, dielectric, and physicochemical propertiesof impregnating resin based on unsaturated polyesterimides★

Louiza Fetouhi1,2,3,*, Benoit Petitgas3, Eric Dantras4, and

Juan Martinez-Vega1,2

1 Université de Toulouse, UPS, INP, CNRS LAPLACE (Laboratoire Plasma et Conversion d’Energie), Toulouse, France2 CNRS, LAPLACE, 118 route de Narbonne, 31062 Toulouse Cedex 9, France3 NIDEC Corporation, Bd Marcellin Leroy, 16915 Angoulême Cedex 9, France4 Physique des Polymères, Institut Carnot CIRIMAT, Université Paul Sabatier, Toulouse, France

★ ContribuSymposium* e-mail: f

Received: 5 December 2016 / Received in final form: 23 May 2017 / Accepted: 7 August 2017

Abstract. This work aims to characterize the dielectric and the mechanical properties of a resin based on anunsaturated polyesterimide diluted in methacrylate reactive diluents used in the impregnation of rotatingmachines. The broadband dielectric spectrometry and the dynamic mechanical analysis were used to quantifythe changes in dielectric and mechanical properties of the network PEI resin, as a function of temperature andfrequency. The network characterizations highlight the presence of two main relaxations, a and a0, confirmed bythe differential scanning calorimetry analysis, showing the complexity of the chemical composition of this resin.The dielectric spectroscopy shows a significant increase in the dielectric values due to an increase of the materialconductivity, while the mechanical spectroscopy shows an important decrease of the polymer rigidity andviscosity expressed by an important decrease in the storagemodulus. The PEI resin shows a high reactivity whenit is submitted in successive heating ramps, which involves in a post-cross-linking reaction.

1 Introduction

The resins based on polyesters are widely used in theimpregnation industry to provide mechanical and electricalinsulating support to the windings of rotating machines.Unsaturated polyesters (UP) are used because of their goodmechanical and electrical properties, as well as hightemperature resistance [1]. With the technical develop-ments of recent decades, the constraints applied toelectrical devices (especially rotating machines such asmotors) have increased substantially. The componentssubjected to increasing temperatures involve the use ofvarnishes, which are increasingly resistant to heat. Thisrequirement has led to the establishment of thermal classescorresponding to the maximum operating temperature ofinsulating systems. This performance is meant to developproducts of higher thermal classes and has led to thedevelopment of new resins’ families, such as polyester-imide based ones. The challenge is to improve the heatresistance of the impregnating resins while considering theimplementation aspect. In particular, the final productmust be a single component with a viscosity at roomtemperature less than 2000mPa·s. Theymust also be cross-linked according to an industrially applicable thermalcycle. Because of energetic costs, the thermal cycle ispreferably of a few hours, with a temperature around

tion to the topical issue “Electrical Engineering(SGE 2016)”, edited by Adel Razek

etouhi@laplace.univ-tlse.fr

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150 °C. In addition, environmental constraints require thereduction of the volatile organic compounds (CoVs) toagree with the new REACH’s standards.

Among the methods used to provide more resistance topolyester resins (UP) is the addition of imide functions,either at the end of the molecules or in the backbone. In thefirst case, unsaturated polyesterimide oligomer disfunc-tionalized imide is diluted in styrene, resulting inunsaturated polyester resins with terminal functions ofimide. In this case, the UP structures are partially modifiedto improve their properties. The unsaturations located onthe end chains do not participate in the copolymerizationreaction with the reactive diluent. Consequently, these endchain unsaturations do not take part in building the 3Dnetwork, providing few improvements to the resin proper-ties. The other method is the modification of unsaturatedpolyester chains by incorporating imide structures in thebackbone, thus using directly unsaturated polyesterimideprepolymers. This method provides many more improve-ments in thermal and mechanical properties, but then anissue of imide homopolarization occurs. The formation ofpolyester-imide blocks with high molar masses tends toincrease the viscosity of the resin that constitutes aproblematic implementation.

The syntheses of polyesterimide resins are generallyperformed in three steps [2,3]. The first step consists of thesynthesis of the unsaturated polyester-imide prepolymer,which involves a polycondensation reaction betweenhydroxylamine and a heated anhydride. This is followedby a radical copolymerization reaction between theunsaturations of the prepolymer and those of the reactive

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diluent, which is generally styrene or vinyl toluene [4]. Thiscopolymerization step is needed to form the 3D cross-linkednetwork, which consists of transition of the polymer from aliquid state to a viscoelastic one. This phenomenon ofgelation formation occurs in two successive steps: forma-tion of microgels followed by their coalescence, resulting inthe formation of macrogels.

Since the development of unsaturated polyester resins,styrene is the most commonly used reactive diluent, mainlybecause of its high diluting capacity and its ability toradically copolymerize with the unsaturations of polyes-ters. Because of its high volatility and its hazardous,carcinogenic [5], and pollutant [6] nature, replacing thestyrene became a major issue in the impregnating resinsindustry. Therefore, other reactive diluents have beenidentified as potentially interesting [7,8]. They can begrouped into three distinct families: styrenic, (meth)acrylic, and allylic diluents. Although researches haveshown that acrylic or methacrylic diluents are both used inthe resin formulation as a reactive diluent (alone or as amixture with styrene) [9,10], they are not well studied. Dueto its compatibility with polyesters, as well as low viscosityand low volatility during polymerization (1–3%), theBDDMA (1,4-butanediol dimethacrylate) [11,12] canreplace styrene as reactive diluent. The choice of BDDMAappears to be potentially interesting because it producesresins with equivalent or better properties to those made instyrene. This is either in terms of implementation orthermal and mechanical resistance.

The resistance of varnishes to mechanical and electricalstresses depends not only on the operating temperatureparameter, but also on a solicitation frequency, such asvibration or AC voltage frequencies. Therefore, the studyof the characteristic relaxations based on the stressfrequency is one of the keys to understand the behaviorof these electrical insulators under severe thermal andstress constraints. In this study, we analyze the dielectricand mechanical properties of a thermal class H resin, whichhas a high operating temperature of 180 °C. This resin isbased on an unsaturated polyester-imide precursor andBDDMA and another methacrylate monomer, as reactivediluents. The different characterizations are performedunder a wide range of temperatures and frequencies inorder to assess the network behavior of the polymer.

2 Experimental

The sample used in dielectric analyses is prepared bycoating liquid varnish on Polyimide film (Kapton®). Thecoated films are then formed using a hydraulic press undera curing cycle usually applied in industry. Two heatingramps are used. The first one is soft and reaches an initialtemperature of 130 °C in order to eliminate volatiles andavoid trapping bubbles in the sample. In this step, acopolymerization mechanism between unsaturated polyes-ter-imide prepolymers and the reactive diluent starts andthe resin transitions from a liquid state to a viscoelasticone. This is called the gelation process [13]. The secondtemperature ramp is steeper and involves a cure of 3 h toobtain a solid sample. In the second step, the cross-linking

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progresses and tends to be completed. The disk-shapedsamples obtained have a radius of 15mm and a thickness of0.3mm. They are then coated with gold by sputtering oneach face before measurements. The sample used inmechanical spectroscopy is obtained by curing a block ofvarnish using the previous curing cycle. The final samplehas a bar shape with 0.5mm thickness, 50mm length, and1mm width and is obtained by mechanical polishing of thevarnish block.

The dielectric properties are measured by broadbanddielectric spectroscopy (BDS) using a Novocontrol®

spectrometer with a ZGS testing cell. Dielectric measure-ments are performed under a nitrogen gas jet by applying alow voltage of 1V in a temperature range from �50 °C to210 °C with a ramp of 3 °Cmin�1. The frequency range isfrom 0.1Hz to 106Hz. The data have been obtained in theform of the dielectric complex permittivity (e� = e0 � ie00),where e0 and e00 are the permittivity and the dielectriclosses, respectively. The loss factor (tan (d)) is derived fromthe ratio between the imaginary (e00) and the real (e0) partsof the complex permittivity.

The mechanical properties are measured by dynamicmechanical analysis (DMA) using an ARES TA® device.Mechanical spectroscopy measurements are also performedunder a nitrogen gas by applying a sinusoidal shear stress of0.1Hz to the parallelepiped sample in the same temperatureramp and range conditions as the dielectric spectroscopy.The data are obtained in the form of the shear complexmodulus (G� =G0 � iG00), where G0 and G00 are the storageand the dissipative modulus, respectively. The loss factor(tan (d)) is derived from the ratio between the imaginary(G00) and the real (G0) parts of the complex shear modulus.

The differential scanning calorimetry (DSC) analysis isperformed using a TA Instruments Q2000 DSC to find theresin Tg. A specimen of 9mg is taken from PEI resinobtained using the same curing cycle showed above. Twosuccessive scans have been performed by DSC in atemperature range from �50 °C to 210 °C with a ramp of10 °Cmin�1.

The samples were characterized by dielectric andmechanical spectroscopies in their solid form. They werepreviously subjected to 4 scans under temperature rampsand frequency. We initially present the results obtained inthe first scan by DMA and in the second scan by BDS toavoid dealing with surface capacitance that generallyhappens in the first scan of dielectric spectroscopy.Afterwards, we discuss the obtained results by analyzingthe four successive performed scans, which allowedpreceding the post-cross-linking reaction of the resin.

3 Dielectric and mechanical results

3.1 Mechanical spectroscopy results

The a relaxation associated to the manifestation of theglass transition temperature could be pointed out using theDMA technique. Three different methods could be used:the extrapolated onset of the sigmoidal change in storagemodulus, the peak of the loss modulus signal, or the peak ofthe tangent delta signal. Figure 1 shows the evolution of thestorageG0 (in black) and the loss modulusG00 (in red) of the

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Fig. 2. Loss factor (tan (d)) of the PEI resin as function oftemperature obtained at the 1st run of DMA at 0.1Hz.

Fig. 3. Real part of the complex permittivity (e0) as function oftemperature and frequency obtained in 2nd run by BDS.

Fig. 1. Storage (G0) and dissipative (G00) modulus of the PEIresin as function of temperature obtained at the 1st run of DMAat 0.1Hz.

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PEI resin as function of temperature in the first run ofdynamic mechanical analyses. Figure 2 shows the corre-sponding loss factor (tan (d)).

In Figure 1, we can see a significant drop in the storagemodulus (G0) over a wide temperature range, from 20 °C to140 °C. The initial vitreous modulus with a value of almost1GPa decreases by about two decades to reach a value of20MPa. This strong drop in the polymer rigidity involvesan important change in thematerial viscosity. This shows asignificant decrease in the polymer viscosity when it crossesthe glass transition temperature domain. According to thesimplified Eyring model of viscous fluids [14], the polymerviscosity is proportional to the storage modulus and isinversely proportional to temperature.

The complexmodulus (G00) has amaximumof 60MPa at20 °C. This peak is related to the network chains mobilitylocated and comforting the presence of a Tg sample in a lowtemperature domain.

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The tan (d) spectrum shows two peaks in Figure 2. Thefirst one is located at 20 °C and the second one is at about80 °C. This finding could be related to the presence of twodistinct mechanical relaxations, therefore showing thepresence of two distinguished phases in the polymer blend.A study conducted by Cousinet et al. [15] highlighted thepresence of two peaks on tan (d) spectrum at 70 °C and165 °C, on UP resins dissolved in methacrylate compoundssuch as BDDMA and IBOMA (isobornyl methacrylate).The first peak at about 70 °C is from the UP-BDDMA partand is more specifically the phase rich in polyester resin.The second peak at about 165 °C is relative to the phase ofIBOMA, whose intensity increased with the concentrationof this compound in the overall mix.

3.2 Dielectric spectroscopy results

In this section, we present the dielectric results obtained inthe second run of the BDS performed on the PEI resin.Figures 3 and 4 present respectively the real (e0) and theimaginary (e00) parts of the complex dielectric permittivityas function of temperature obtained for the isofrequenciesof 0.1, 50, and 104Hz. Figure 5 displays the correspondingloss factor (tan (d)).

Firstly, one can see that the permittivity (Fig. 3)increases very slightly from �50 °C to 0 °C. Then, thisincrease begins to be more pronounced from 0 °C and seemsto take place two times. Indeed, for the isofrequency of0.1Hz, the permittivity increases and reaches the value of 6at 40 °C and exceeds a value of 10 at 100 °C. These couldcorrespond to the presence of two main relaxations (a anda0), which are related to the increase in polarized dipolesmobility. According to the Adams-Gibbs theory [16], themobility of these dipoles is allowed by the cooperativesegments motions, which is activated at the glass transitiontemperature (Tg). The possible presence of two dielectricmain relaxations (a and a0) stays in good agreement withthe DMA results (Fig. 2) that showed also the presence oftwo distinguished main relaxations also called a and a0.Polymers with immiscible blends usually show several

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Fig. 4. Imaginarypart of the complexpermittivity (e00) as functionof temperature and frequency obtained in 2nd run by BDS.

Fig. 5. Loss factor (tan (d)) as function of temperature andfrequency obtained in 2nd run by BDS.

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main a-relaxations, where each relaxation is related to aglass transition of each component [17,18]. On the samefrequencyspectrumof0.1Hzandfrom100 °C, the realpartofthe complex permittivity continues to increase rapidly andreaches the maximum value of 1050 at 160 °C. This could berelated to the electrode polarization (EP) phenomenon,generally observedat low frequencies,which is superimposedto the conduction phenomenon [19,20]. At this lowestfrequency (0.1Hz), a phenomenon of electrode polarization(EP)contributes to the important increase invalues, therebymasking the real dielectric response of the material. Themolecular origin of the EP phenomenon is the partialblocking of charge carriers at the sample/electrode interface.This leads to a separation of positive and negative charges,which give rise to an additional polarization. The increase intemperature makes it possible for the ionic charges to movefrom the bulk to the interface between the electrodes and theinsulator [21,22]. One can observe a slight decrease inpermittivity from160 °C for the same isofrequency spectrumof 0.1Hz. This could come from chemical changes in thepolymer network rather than from polarization dipolesphenomena. Themeasurement of the sample thickness afterthe second run of the dielectric spectroscopy confirms thenonvariation of the sample thickness. It can, therefore, beconcluded that the decrease in the observed permittivitycould be related to a post-cross-linking process or chemicalspecies removal such as residual humidity.

For the spectra obtained at 50 and 104Hz, thepermittivity shows, until 80 °C, a similar behavior thanthe spectrum of 0.1Hz. Above 80 °C, it seems that theincrease in the permittivity values are less important athigh temperatures for these scanning frequencies. This iscertainly due to the less pronounced or absence of electrodepolarization at higher frequencies. In addition, one canobserve the classical shift of the dielectric phenomena tohigher temperatures with the increasing frequency.

The imaginary part of the complex permittivity (e00)(Fig. 4) and the loss factor (tan d) (Fig. 5) display fourdifferent dielectric processes. The first one is located below0 °C for the isofrequency of 0.1Hz. One can observe a slight

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dropof thedielectric losses andthe loss factor valueswith theincreasing temperature.This couldcorrespondto theendofarelaxation peak, and more precisely to a secondaryb-relaxation, which generally occurs at low temperatures.Theb-relaxation is due to themotion of local noncooperativedipoles in the condensed phase. In the second temperaturedomain, located between 0 °C and 100 °C, one can observe amore pronounced increase in both the dielectric losses (e00)and the loss factor (tan d) values. We can observe the firstrelaxation (a) peak located at 40 °C in both e00 and tan dmagnitudes. Around 80 °C, the loss factor spectrum showsthe presence of a second peak that could be related to thesecondmaina0 relaxation.On the dielectric losses spectrum,it is more difficult to observe the second relaxation peak a0.Indeed, the high temperature conduction overlaps to thepolarization electrode (EP) that contributes to mask thedielectric response of the resin, and thus the secondrelaxation a0. The last temperature domain is located abovethemain relaxation area (beyond 100 °C). There, one can seean important increase in the dielectric losses and loss factorvalues. Several phenomena are superimposed, highlightingthe complexity of the mechanisms associated with eachdifferent magnitude. Indeed, e00 and tan (d) increase rapidlyand reach the maximum values of 2� 105 and 20respectively, forthe lowest frequencyof0.1Hz.Thisbehavioris typically due to anEPphenomenon overlapped to the hightemperature s-conduction. The electrode polarization isinduced by the relaxation of extremely mobile space-chargelayers in the polymer.Above theTg region, the free volume isincreased andmakes it possible for the ionic charges tomove[16]. This finding is widespread to other frequencies (50 and104Hz); however, the value of bothmagnitudes increase lessthan for the scanning frequency of 0.1Hz.

4 Discussion

4.1 Relaxation spectra

The results obtained bymechanical and dielectric spectros-copies show the presence of two distinguished relaxationsa and a0 in the studied range [20–140 °C]. The first

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Fig. 6. Heat flow spectrum as function of temperature obtainedby DSC.

Fig. 7. Comparisonof the factorof losses (tan (d))obtained ineachrun by DMA as function of temperature and for a frequency of0.1Hz.

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relaxation (a) is around the room temperature (20 °C) inthe mechanical spectroscopy, whereas it is located around40 °C in the dielectric spectroscopy. The second relaxation(a0) is around 80 °C, taking into account the relaxationsobserved on each loss factor spectrum obtained by DMAand BDS at the scanning frequency of 0.1Hz. The presenceof two distinguished phases is comforted by the continuousdecrease in the storage modulus G0 (cf. Fig. 1). Thecontinuous decrease in viscosity over the range [20–140 °C]could cover both relaxations observed at 20 °C and 80 °C.The presence of two relaxations in the polymer network ofthe resin suggests the existence of two possible glasstransitions. PEI resin is a mixture of at least two polymerswith varying levels of miscibility. The studied resin iscomposed of a polyesterimide precursor present in areactive diluent of BDDMA and another methacrylatemonomer. The partial miscibility of the polyesterimide inthe reactive diluent, the homopolymerization of meth-acrylates, or secondary reactions that take place during thesynthesis can induce the phase segregation in material [17].

The most commonly used method to evaluate themiscibility of components in a mixture is the analyses of theglass transition temperatures. The presence of two Tgindicates the presence of two phases in the polymer system.The presence of a single Tg is often attributed to perfectlymiscible polymer blends [18]. The results obtained during asecond scan of DSC on the PEI resin are presented inFigure 6. The DSC spectrum displays two endothermicvariations. The first one is at about 25 °C and the secondone is at 90 °C.

The measured Tgs on the resin based on PEI andmethacrylate monomers diluent are relatively low (25 °Cand 90 °C). The polyester-imide monomers are known fortheir high glass transition temperature thanks to the imidefunctions, which increase as a function of the number ofimide functions in the polyester matrix. As mentioned inSection 1, unsaturated polyester resins with terminal imidefunctions are assumed to provide few improvements to theresin properties. This is because the end chain unsatura-tions do not take part to build the 3D network. In addition,

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in this case, the Tg is related to the concentration of thereactive diluent based on methacrylate, especially theBDDMA that is in contrary known to have a low Tg. Theanalysis of theTgs gives an idea of the methacrylate diluentcontent. Studies show that the increase in styrene contentshifts the Tg to a higher temperature and stiffens the UPresin [23]. The use of methacrylate diluent in order toreplace styrene seems to be not only for health andenvironmental interest, but also to reduce the glasstransition temperature of the PEI resin. This decrease intheTg resin is advantageous in the impregnating field. Thissuggests that the resin operates in a viscous state, whichallows for absorbing vibrations better and mechanicalstresses during rotating machine operation.

4.2 Post-cross-linking effects on mechanical properties

To study the effect of the PEI resin cross-linking rate onthe dielectric and mechanical properties of the resin, wecarried out additional scans on samples, those hadpreviously been submitted to two dynamic mechanicaland dielectric scans.

The peaks of tan (d) initially observed in the first runtend to extenuate and homogenize in the 2nd, 3rd, and 4thscans (Fig. 7). This suggests that the two distinctrelaxations, a and a0, previously highlighted tend to mergeas the system advances in its cross-linking. The heteroge-neity of the polymer network is characterized by theincrease in the half-depth of tan (d) [24]. However, weobserve in the case of PEI resin shrinkage of the area underthe curve of tan (d), as the successive scans under a ramptemperature are performed. This leads to two hypotheses.At first we can assume that the blends contained in the PEIresin become increasingly miscible as the cross-linkingreaction proceeds. Then, as we mentioned in theintroduction, the radical copolymerization between thereactive diluent and the functional units of polyesterimideprecursor leads to the formation of 3D chemical gel thatrigidify under the influence of heat. During the polymeri-

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Fig. 8. Comparison of storage modulus (G0) obtained in each runby DMA as function of temperature and for a frequency of 0.1Hz.

Fig. 9. Comparison of dissipative modulus (G00) obtained in eachrun by DMA as function of temperature and for a frequency of0.1Hz.

Fig. 10. Real part of the complex permittivity (e0) as function oftemperature obtained on the 4 successive runs for the frequency of0.1Hz by BDS.

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zation, the cross-linked infusible network is formed. Thenin continuing the cross-linking of the network, thecopolymerization would continue [25].

The storage modulus (G0), shown in Figure 8, evolvesprogressively with the performed scans under temperatureramps. We observe an increase of this modulus on the 4thrun of about 5MPa in a rubbery state. This is directlyrelated to the post-cross-linking phenomenon effect [23].Indeed, this irreversible phenomenon allows the polymerdensity to increase and consequently increases its rigidity.This results in the recovery of the cross-linking reactioninitiated by the heating ramps.

In Figure 9, we observe a shift of the dissipativemodulus (G00) peaks to higher temperatures as thesuccessive runs are performed. However, the maximumof the 3rd scan that appears to be shifted slightly to lowertemperatures than the 2nd one is probably due the intervalof time between the two scans. The amplitude of G00decreases on the 2nd and 3rd transitions due to a post-

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cross-linking effect. As mentioned above, the loss modulusG00 characterizes the chain network mobility of thepolymer. The post-cross-linking of the network duringsuccessive scans under temperature has resulted inincreased network density and thus reduced the chainsmobility. However, the dissipative modulus G00 shows anincrease in amplitude of 20MPa in the last scan. Thisincrease of mobility of the chains might be explained by theinitiation of the chemical degradation mechanisms of thepolymer network (chain scission), because the networkbecomes more rigid and denser. The application of fourramping temperature of 3 °Cmin�1 in a temperature rangeto 210 °C could induce the scission of certain chemicalfunctional groups, such as the covalent band C–O–C,which are known to be the weak point of ester functions.Montaudo and Puglisi [26] located the initial degradationtemperatures of polyesters between 230 °C and 300 °Cunder inert atmosphere.

4.3 Post-cross-linking effects on dielectric properties

The dielectric properties of the PEI resin obtained in the3rd and 4th runs are significantly reduced. The permittivi-ty (e0) and the dielectric losses (e00), shown in Figures 10 and11 respectively, decrease by about one decade between thefirst and the last runs for the scanning frequency of 0.1Hz.Both the number of dipoles and the conduction phenomenadue to the electronic and ionic charges motions aresignificantly reduced on the annealed sample. Thisreduction in charge carrier mobility and density could berelated to the increase in the density network. Indeed, onecan observe in both Figures 11 and 12 a shift of the maina-relaxations initially located at 20 °C and 80 °C to highertemperatures of 40 °C and 120 °C, respectively. The slightfluctuation of the signal intensity, related to the main arelaxations, is attenuated as the heating scans areperformed. This is also related to the post-reticulationthat occurs in the PEI resin during the different measure-ments, which provides less mobility to the polymer

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Fig. 11. Imaginary part of the complex permittivity (e00) asfunction of temperature obtained on the 4 successive runs for thefrequency of 0.1Hz by BDS.

Fig. 12. Loss factor (tan (d)) as a function of temperatureobtained on the 4 successive runs for the frequency of 0.1Hz byBDS.

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network. In addition, processed polymers are usually in anonequilibrium state at the end of the shaping process;thus, heating occurs during measurements operates as anannealing that allows the network tending toward anequilibrium state with a structural recovery [27]. Themacroscopic properties of the thermosets depend on thecross-link density [15]. The heating increases the averagenumber of cross-links and reduces the average lengthbetween them, thus increasing the polymer density.

The increase in dielectric values is partly explained by adecrease in viscosity shown by the mechanical spectrosco-py. In general, the material conductivity is considered to beinversely proportional to its viscosity. Consequently, thepost-cross-linking reactions should affect the polymerconductivity directly. In order to analyze more preciselythe s-conduction obtained in each dielectric-scanning run,the AC conductivity is presented as function of frequencyusing equation (1) for temperatures above the glasstransition domain (>120 °C) for each run. We considerthe measured real part of the AC conductivity (s0

AC)following the universal power law given by:

s0AC ¼ ve0e00ðvÞ ¼ s þ Avs; ð1Þ

where v is the angular frequency, e0 is the vacuumpermittivity, e00 is the measured dielectric losses, A is atemperature-dependent parameter, and s is the exponent ofthe “universal” power law. s is considered as the DCconductivity obtained at the lowest frequency of 0.1Hzthrough the AC conductivity (s0

AC). The angular frequencyexponent s is 0< s<1 and is related to a non-ohmicconductivity [28].

This representation (cf. Fig. 13) allows a plateau todisclose independently on frequency and that corresponds tothe DC conductivity magnitude (s). The DC conductivity(s) is calculated by extrapolating the conductivity curves tozero frequency for different temperatures. In the higherfrequency region, conductivity sAC becomes dependent andincreases linearlywith frequency exhibiting a slope close to1.

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In the first run, one can observe a progressive shift of theplateau toward high frequencies when the temperatureincreases. These changes correspond not only to thethermally activated character of the AC conductivity butalso to the increase of the DC bulk conductivity in the PEI.

The sAC value around 10�8V�1 cm�1 at 210 °C in thefirst run decreases to about 10�9V�1 cm�1 in the last 4thrun. As far as the scanning heat is performed, theconduction magnitude is reduced. The annealing allowsthe resin to improve its dielectric behavior, showing thereduction of the s-conduction of nearly one order ofmagnitude between the first and the last runs. The Tgincreases with the annealing and allows to better controlthe free volume expansion responsible of charges motions inthe bulk. The material is more stable and the chargecarriers are trapped in the thermosetting polymer network.Furthermore, one can see that the reduction in s-conduc-tion after the first run is more important. This comes fromnot only from the residual water and volatiles eliminationand the reduction in the surface capacity, but also from thepostlinking progression that happens during the firstscanning heat. In the 2nd, 3rd, and the 4th scans, thesephenomena tend to be completed.

Figure 14 presents the reciprocal temperature depen-dence of the DC conductivity (s) extracted from the ACconductivity values obtained at the lowest frequency of0.1Hz for each scanning run. The conductivity being athermoactivated process, the Arrhenius model written bythe following form, fits well DC conductivity (s) valuesversus the inverse of the temperature:

s ¼ s0 exp�Ea

kBT

� �; ð2Þ

where s is the extrapolated DC conductivity from the ACform, s0 is the pre-exponential factor, Ea is the activationenergy, T is the temperature, and kB is the Boltzmannconstant (kB=8.62� 10�5 eVK�1).

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Fig. 13. AC conductivity (sAC) as function of the frequency obtained in each dielectric run for each temperature of measurements.

Fig. 14. Extrapolated DCconductivity (s) of PEI resin extractedfromthe lowfrequencyACconductivity as functionof the inverseoftemperatureobtainedforeachrunabovetheTg region.Thered linesrepresent the linear fits using the Arrhenius law.

Table 1. Activation energies obtained for each dielectricscan under ramping temperature.

Scanning run Ea (eV)

1st run 0.5102nd run 0.6973rd run 0.7164th run 0.734

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Above the second Ta0 region, the mechanisms follow the

Arrhenius law. The Ta0 region is shifted to higher temper-

atures, which is explained by a cross-linking reactionprogression resulting in the different sample’s annealing.

The activation energies (Ea) extracted from theArrhenius fit are reported in Table 1. The activationenergy increases from 0.51 eV obtained in the first run toreach 0.734 eV obtained in the last run of the dielectricspectroscopy. This clearly shows that the physicochemicalmechanisms are different between the first and last runs.

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The activation energy increases with the applied heating.Moreover, this indicates an increase in the potential barrierthat the charge carriers have to overcome to jump betweensites [29]. The thermal heating induces the post-cross-linking of the polymer network that could stabilize thestructures of the PEI resin.

5 Conclusion

This work has highlighted a complementary relationshipbetween the mechanical and dielectric properties of PEIresin conducted by dielectric and mechanical spectros-copies and correlated to its physicochemical characteristicsobtained by DSC. The different spectroscopies allowanalyzing the polymer network. That gives us preciousinformation on the miscibility of blends contained in theresin. In addition, this study allows evaluating thedetermining role of curing on the functional properties ofthe PEI resin.

Both DMA and BDS identify two main relaxations onthe PEI resin at about 40 °C and 80 °C. DSC analyses showendothermic jumps associated with the presence of Tg inthe neighboring of these temperatures and in a wide rangeof temperature [20–140 °C]. The presence of two glasstransitions in the polymer network suggests the existence ofblends with varying levels of miscibility. This implies anonmiscibility or a nonperfect compatibility between theprecursor polyesterimide with the reactive diluent consist-ing of BDDMA and another methacrylate monomer.However, it is important to note that the peaks associatedwith relaxations observed on the 3rd and 4th runs on thetan (d) spectrum obtained by DMA seem to homogenize(cf. Fig. 9). This initially implies an increase in themiscibility of the polyesterimide in the BDDMA. On theother hand, the post-cross-linking phenomenon observedon successive runs appears to continue the radicalcopolymerization reaction between the reactive diluentand polyesterimide. This could improve the differentblends miscibility and avoid the previous phase segregationobserved in the first run.

The post-cross-linking phenomenon was observed inboth mechanical and dielectric properties of the PEI resin.This pointed out by both T 0

gs shifts to higher temperaturesand the increase in the viscous modulus. The post-cross-linking seems to also improve the insulating properties ofthe resin with reference to the decrease in dielectric andconductivity values. Otherwise, even the evolution of thedissipative modulus (G00) in the first three runs comfortsthe post-cross-linking progression in the polymer network.Its increase in the 4th run of mechanical spectroscopyassumes the initiation of chains polymer. This gives ageneral idea of the several problems that are related to thecure reaction, where copolymerization and degradationreactions are competing. Many publications [30–32] dealwith these issues and still do not completely elucidate allaspects of formulation and cure of the impregnating resins.

The continuous and significant reduction in theconservative mechanical modulus (G0) observed in a widetemperature range, in the Tgs domain, indicates asignificant dilatation of the material above the Tg. This

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could induce two distinct effects. The first effect is relatedto the decrease of the rigidity of thematerial, which gives tothe material more flexibility and could help to absorb thedifferent vibration and/or mechanical stresses to which thevarnish is subjected during operation. The second is relatedto the presence of the polymerTgs in the temperature rangeof the machine operation, that is 20–140 °C. This involves acontinuous transition of the material between the vitreousand the rubbery states [25]. Therefore, the material isconstantly subjected to internal stresses of expansion andcompression, due to the important dilatation of the resin,which could then contribute to ageing acceleration of theresin and potentially induce microcracks. Many studiesdeal with microcracking induced by mechanical fatigue(mechanical cycling) [33–36] at the origin of increase in theDP probability of occurrences.

This study shows that a thermal treatment of the resinimproves the electrical insulation properties of the resin.The resin appears to be stabilized with the application ofthermal annealing to complete its cross-linking and toevacuate the volatile compounds responsible for thereactivity of the resin. Industrially, the application ofannealing requires input energy, which would requireconsiderable financial support.

The authors acknowledge the Leroy-Somer Company for theirfinancial support and their co-supervision in the PhD project. Theauthors are also thankful to the ANRT France for their financialsupport.

References

1. K.-W. Lienert, Poly(ester-imide)s for Industrial Use (BeckElektroisolier-Systeme, Hamburg, 1999)

2. H.V. Boenig, Polyesters, unsaturated, in Encyclopedia ofPolymer Science and Technology (1964), Vol. 11, pp. 129–168

3. K. Schmidt, D. Wille, U.S. Patent 3,562,219, 19714. Anonymous (Dow Chemicals Co.), Vinyltoluene-unsaturat-

ed interpolymers and methods of making the same,GB731111, 1952

5. Environmental Protection Agency (2003), Vol. 686. US Department of Health and Human Services, Public

Health Services, Report on Carcinogens, National ToxicologyProgram, 12th edn. (2011)

7. P.B. Zetterlund, R.G. Gosden, W. Weaver, A.F. Johnson,New aspects of unsaturated polyester resin synthesis. Part 2.Reactant sequence distribution and its effects on curekinetics, Polym. Int. 52, 749 (2003)

8. Anonymous (Wyandotte Chemical Corp.), Polyester,GB1062756, 1962

9. Anonymous (General Electric), Improvements relating toelectrical insulating structures, GB818038, 1954

10. E.N. Peters, Z. Lin (General Electric), Method for electricalinsulation and insulated electrical conductor,US20070066710, 2005

11. J.J. Bozell, G.R. Petersen, Green Chem. 12, 539 (2010)12. Top value added chemicals from biomass: vol I – results of

screening for potential candidates from sugars and synthesisgas, US Department of Energy, Technical report office ofscientific and technical information, N.R.E.L., 2004

-p9

10 L. Fetouhi et al.: Eur. Phys. J. Appl. Phys. 80, 10901 (2017)

13. Y.S. Yang, L. Suspene, Curing of unsaturated polyesterresins: viscosity studies and simulation in pre-gel state,Polym. Eng. Sci. 31, 321 (1991)

14. H. Eyring, The activated complex in chemical reactions, J.Chem. Phys. 3, 107 (1935)

15. S. Cousinet et al., Toward replacement of styrene by bio-based methacrytes in unsaturated polyester resins, Eur.Polym. J. 67, 539 (2015)

16. J.H. Gibbs, E.A. DiMarzio, Nature of the glass transition andthe glassy state, J. Chem. Phys. 28, 373 (1958)

17. E.M.S. Sanchez,C.A.C.Zavaglia,M.I.Felisberti,Unsaturatedpolyester resins: influence of the styrene concentration on themiscibility andmechanical properties, Polymer 41, 765 (2000)

18. L. Xian,R.A.Weiss, Relationship between the glass transitiontemperature and the interaction parameter of miscible binarypolymer blends, Macromolecules 25, 3242 (1992)

19. S. Diaham, M. Locatelli, Novel high glass transition polyam-ide-imide: Tg influence on electrical conductivity at hightemperature, IEEE Trans. Dielectr. Electr. Insul. 22, 3053(2015)

20. A.K. Jonscher, Dielectric relaxation in solids, J. Phys. D:Appl. Phys. 32, R57 (1999)

21. P.B. Ishai, M.S. Talary, A. Caduff, E. Levy, Y. Feldman,Electrode polarization in dielectric measurements: a review,Meas. Sci. Technol. 24, 102001 (2013)

22. S. Diaham, M.-L. Locatelli, Time and frequency domains dcconductivity analysis in thin dielectric films at hightemperature, J. Phys. D: Appl. Phys. 44, 105402 (2011)

23. N.M. Alves et al., Viscoelastic behavior of poly(methylmethacrylate) networks with different cross-linking degrees,Macromolecules 37, 3735 (2004)

24. R. Ley, Structural heterogeneities and mechanical propertiesof vinyl/dimethacrylate networks synthesized by thermalfree radical polymerisation, Polymer 43, 4375 (2002)

10901

25. L. Fetouhi, B. Petitgas, E. Dantras, J. Martinez-Vega,Dielectric and mechanical properties correlated to physico-chemical characteristics of a polyester-imide resin used inrotating machines insulations, in International Conferenceon Dielectric ICD 2016, July 3–7, Montpellier, France (2016)

26. G. Montaudo, C. Puglisi, Thermal degradation mechanismsin condensation polymers, Dev. Polym. Degrad. 7, 48 (1987)

27. J.R. White, Polymer ageing: physics, chemistry or engineer-ing? Time to reflect, C. R. Chimie 9, 1396 (2006)

28. G.W. Bak, Dielectric relaxation in p-terphenyl crystals,Mater. Sci. Eng. 80, L37 (1986)

29. A. Bouzidi, K. Omri, L. El Mir, H. Guermazi, Preparation,structural and optical investigations of ITO nanopowder andITO/epoxy nanocomposites, Mater. Sci. Semicond. Process.39, 536 (2015)

30. K. De La Caba, P.Guerrero, A. Eceiza, I.Mondragon,Kineticand rheological studies of two unsaturated polyester resinscured at different temperatures, Eur. Polym. J. 33, 9 (1997)

31. A. Yousefi, P.G. Lafleur, R.J. Gauvin, The effects of cobaltpromoterandglassfiberson the curingbehavior ofunsaturatedpolyester resin, J. Vinyl Addit. Technol. 3, 157 (1997)

32. Y.S. Yang, L.J. Lee, Rheokinetic studies of unsaturatedpolyester resins, Polym. Proc. Eng. 5, 327 (1987–1988)

33. F. Awaja, S. Zhang, M. Tripathi, A. Nikiforov, N. Pugno,Cracks, microcracks and fracture in polymer structures:formation, detection, autonomic repair, Prog. Mater. Sci. 83,536 (2016)

34. J.C. Radon, Fatigue crack growth in polymers, Int. J. Fract.16, 533 (1980)

35. A. Wöhler, Wöhler’s experiments on the strength of metals,Engineering 2, 160 (1867)

36. E.H. Andrews, Cracking and crazing in polymeric glasses, inThe Physics of Glassy Polymers, edited by R.N. Haward(Springer, Netherlands, 1973), pp. 394–453

Cite this article as: Louiza Fetouhi, Benoit Petitgas, Eric Dantras, Juan Martinez-Vega, Mechanical, dielectric, andphysicochemical properties of impregnating resin based on unsaturated polyesterimides, Eur. Phys. J. Appl. Phys. 80, 10901 (2017)

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