Thermophysical Properties and Molecular

Relaxations in Cured Epoxy Resinþ PEO Blends:

Observations on Factors Controlling Miscibility

Ioannis M. Kalogeras,*1 Aglaia Vassilikou-Dova,1 Iraklis Christakis,1 Dorota Pietkiewicz,2 Witold Brostow2

1 Solid State Physics Section, Department of Physics, University of Athens, Panepistimiopolis, 157 84 Zografos, GreeceE-mail: ikaloger@phys.uoa.gr

2 Laboratory of Advanced Polymers & Optimized Materials (LAPOM), Department of Materials Science and Engineering,University of North Texas, P.O. Box 305310, Denton, Texas, USA

Received: February 2, 2006; Revised: March 10, 2006; Accepted: March 14, 2006; DOI: 10.1002/macp.200600052

Keywords: blends; crystallization; dielectric properties; glass transition; poly(ethylene oxide); thermosets

Introduction

Relation to Previous Work

Since the early 1960s, a variety of homopolymers, random

or block copolymers have been incorporated into epoxy or

novolac resins. The resulting blends exhibit microstructures

that depend on type of polymers, their concentrations and

curing programs (time and temperature).[1–3] The majority

of thermosetting polymer blends studied hitherto are

immiscible;[4–9] as a result of competition among kinetic

and thermodynamic factors. In cases where miscibility has

been found, attempts were made to relate it to structural

similarities and in particular to the presence of various

secondary interaction forces (e.g., hydrogen bonding

dH,[10,11] ion–dipole forces and electron—donor–acceptor

interactions). Poly(ethylene oxide) (PEO), a material with

promising application in high-energy density electrochem-

ical systems and biomedical devices (e.g., in electroche-

mical sensors, solid-state batteries, haemodialysis

membranes and drug delivery systems), is widely used as

an elastomeric modifier of thermosets for enhancement of

their mechanical and electrical properties, and as a mean to

Summary: Thermophysical properties and molecular relaxa-tions in aromatic amine-cured diglycidyl ether of bisphenol-A(DGEBA) epoxy oligomer and poly(ethylene oxide) (PEO)mixtures were determined by DSC and dielectric techniques(TSC, DRS). The binary blends were judged to be fully misciblein the amorphous state (wPEO< 40 wt.-%), as evidenced by thesingle composition-dependent glass transition temperaturesTgs. In the amorphous blends, negative deviations of dielectric/thermalTg-estimates from the linear mixing rule or the behaviorpredicted by the Fox equation reveal weaker intermolecularinteractions, compared to strong self-association of hydroxylsin the cured thermoset. Morphological changes in PEO-richblends (wPEO� 40 wt.-%) are in accordance with theircomplicated interface structure, previously reported to consistof amorphous PEO regions, branched epoxy resin chains and animperfect epoxy resin network located between PEO lamellae.In these blends, PEO crystallites exert steric hindrances in theamorphous regions, causing strong Tg upshifts. Changes in therelaxation dynamics of glyceryl segments (e.g., in the activationenergy barrier and relaxation strength) are in accordance withthe idea that the close matching between the molecularpolarities of PEO, epoxy resin and the cure agent significantlycontributes to the observed miscibility.

Plots of transition temperatures as function of blendcomposition.

Macromol. Chem. Phys. 2006, 207, 879–892 � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Full Paper DOI: 10.1002/macp.200600052 879

study crystallization processes, miscibility, compatibility,

and parameters controlling phase interactions.[1–4]

PEO is a proton-accepting polyether known to exhibit

strong interactions and miscibility with several polymers

possessing functional acid groups, such as poly(acrylic

acid) and poly(methacrylic acid) in which complexion

occurs from strong acid hydroxyls.[12] An important

example of miscible epoxy resin (ER)þ PEO mixtures is

the blend of PEO with poly(hydroxyl ether of bisphenol-A)

(phenoxy).[13,14] Epoxy resins in general, and specifically

diglycidyl ether of bisphenol-A (DGEBA)-type epoxies,

have a structure similar to phenoxy;[15,16] thus, we expect

them to show high compatibility with PEO. Nevertheless,

there is ample experimental evidence suggesting that the

presence of functional groups in thermosetting component

polymers cannot assure the formation of favorable

intermolecular specific interactions, nor can it ensure the

homogeneity of the resulting blends. For example, several

polyether-type thermoplastics, such as poly(propylene

oxide), poly(vinyl ethyl ether), polyacetal (polyoxymethyl-

ene), poly(2,6-dimethyl-1,4-phenylene oxide),[5,6] and

poly(ether imide)[8], when incorporated in cross-linking

ERs yield phase-separated morphology. Substantial bear-

ing on the morphology of the final (cured) product has also

the choice of the hardening agent as well as the cure

conditions. For example, PEO has been reported to exhibit

miscibility with DGEBA cured with 1,3,5-trihydroxybenz-

ene (THB),[11] aromatic amines [e.g. 4,40-methylenebis(2-

chloroaniline) (MOCA),[17] 4,40-diaminodiphenylmethane

(DDM),[18] 4,40-diaminodiphenyl sulfone (DDS),[19,20] and

4,40-methylenebis(3-chloro-2,6-diethylaniline)

(MCDEA)[21]], and aromatic anhydrides [e.g., pyromellitic

dianhydride (PMDA)[20] and phthalic anhydride (PA)[22]].

In contrast, immiscibility has been reported for ERþPEO blends cured with aliphatic amines, such as diethyl-

ene triamine (DETA)[20] and tetraethylene pentamine

(TEPA).[4] Thus, compared with linear polymer blends,

miscibility and intermolecular specific interactions in

thermosetting blends are much more complicated; kinetic

factors, curing processes, and topological structures of

thermosetting systems are closely coupled with the

thermodynamics of mixing.

The common criterion for miscibility is the detection of a

single glass transition temperature Tg located between the

respective temperatures of the pure components. This is the

reason why differential scanning calorimetry (DSC) has

been one of the most widely used techniques in determining

miscibility. DSC, however, can only sense movements at

length scales larger than �10 nm, so that any nano-

heterogeneities are averaged out. Dielectric relaxation

probes have been claimed to be more sensitive for

measuring structural heterogeneity in polymer blends,

copolymers, and composites, due to their considerably

smaller probe length scales.[23] The study of molecular

dynamics in thermosetþ polymer blends by means of

thermally stimulated currents (TSC) and by its ac counter-

part dielectric relaxation spectroscopy (DRS) can provide

valuable information on phase interactions and their

miscibility. Hitherto, dielectric methods have been imple-

mented to monitor phase separation, gelation, and vitri-

fication processes during cure,[24] or to study local-chain

and segmental relaxation modes in several epoxy resins.[25–

28] Bulk PEO and its complexes with several salts (e.g.,

potassium thiocyanate and lithium perchlorate) have also

been extensively studied dielectrically.[29,30] However, a

survey of the literature pertaining to cured ERþ PEO

blends revealed a shortage of reports on their relaxation

dynamics studied by dielectric techniques.

Using PEO as a model polymer to be incorporated in an

epoxy network, DSC, TSC, and DRS are used here to probe

the influence of molecular structure, components polarity,

and interactions, on thermal properties and their dielectric

relaxation response, with the aim to acquire information on

phase compatibility in cured epoxyþ polymer networks.

The merit of performing comparative thermal, dielectric,

and electrical studies on the molecular dynamics of

ERþ PEO systems is supported by reports of enhanced

conductivity of PEO-based fast ion conductors induced by

the addition of novolac phenolic resins.[31,32] A lower

pseudoactivation energy in lithium transfer and compli-

cated interactions among PEO, resin, and lithium salt, have

been detected by Chu et al.[31] The high ionic conductivity

in PEOþLiClO4 presumably originates from effective salt

dissociation and the establishment of a new conduction

network formed by amorphous PEO and the phenolic resin.

Furthermore, structural modifications were reported to

extend both thermal stability and mechanical strength of the

solid polymerþ electrolyte composite.[31]

Experimental Part

Materials

Figure 1 shows the structures of the chemicals used in thisstudy. PEO (Sigma Chemical Co., USA) with molecularweight Mw ¼ 20 000 g �mol�1 was used without furtherpurification. The epoxy resin used in the study was DGEBA(commercially available from Shell Chemicals Corporation,USA, as EPONTM resin 828) with the degree of oligomeriza-tion n� 0.1. The number-average molecular mass isMnðERÞ ¼ 360 g �mol�1 with the epoxide equivalent mass�185 g. For curing DGEBAþPEO mixtures, we employed acommon aromatic amine curing agent (hardener), DDM(chemical grade, MnðDDMÞ ¼ 198.26 g �mol�1) from AldrichChemical Inc. (USA), in granular form. Acetone solutions ofthe curing agent were mixed at 90 8C (just above the curingagent’s melting point, at 87 8C) with acetone solutionscontaining DGEBA and predetermined quantities of PEO.The amount of hardener relative to the DGEBA resin was28 phr (i.e., 28 parts of curing agent per 100 parts of resin),which results in the stoichiometric amino hydrogen/epoxy

880 I. M. Kalogeras, A. Vassilikou-Dova, I. Christakis, D. Pietkiewicz, W. Brostow

Macromol. Chem. Phys. 2006, 207, 879–892 www.mcp-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

ratio [r]. Curing of the well-mixed ternary DGE-BAþPEOþDDM blends was performed at 90 8C (for 4 h),followed by post-curing at 120 8C (for 2 h). The concentrationof PEO in each blend is expressed in wt.-% (e.g., 10/90ERþPEO for samples with 10 wt.-% of DGEBAþDDM and90 wt.-% PEO).

Techniques

Dielectric Relaxation Spectroscopy (DRS)

The complex dielectric permittivity data, e�ðoÞ ¼ e0ðoÞ�je00ðoÞ, determined by DRS, with e0(o) the real part ofpermittivity, e00(o) the dielectric loss and o¼ 2pf the angularfrequency, are usually fitted by assuming that the dielectricresponse is a linear superposition of the contribution from dc-conductivity sdc and various polarization processes expressedby their complex dielectric constant ei*(o). To account for thenon-Debye character of the ac response, ei*(o) can bedescribed quite well by the empirical asymmetric Havriliak-Negami (H–N) function.[33] Thus

e�ðoÞ ¼ e1 þXi

Dei

1 þ jf

f0;i

� �1�ai" #bi

� jsdc

e0os

ð1Þwhere 0< s� 1 (s¼ 1 for ohmic conductivity), e0 is thepermittivity of free space, and e1 is e0(f) for f >> fmax¼ 1/2pt(t is the relaxation time). In the framework of the H–N model,the temperature-dependent parameters of each relaxationprocess are the mean relaxation time, the relaxation strengthDe, and the width a and asymmetry b of the relaxation inrelation to a single Debye process (a¼ 0, b¼ 1). Sub-glassyrelaxations exhibit thermally activated relaxation timesexpressed by the Arrhenius-type equation

tðTÞ ¼ t0 expE

kBT

� �ð2Þ

where t0 is related to the inverse of the characteristic vibrationfrequency at infinite temperature of a charge carrier trapped inits site, E is the apparent activation energy barrier, and kB is theBoltzmann constant. The relaxation parameters (E, t0) can bededuced respectively from the slope and the intercept of thelinear plot of ln(fmax) versus 1/T (Arrhenius plots).

For DRS measurements disk-like specimens were sand-wiched between brass or gold-coated brass electrodes and e*was determined as a function of the frequency (10�2–106 Hz)at selected temperatures between �140 and 150 8C (controlledto better than �0.1 K). A Novocontrol Alpha Analyzer wasused in combination with the Novocontrol Quatro Cryosystem.

Thermally Stimulated Currents (TSC)

As an alternative to DRS, TSC measures the thermallyactivated release of stored dielectric polarization. The currentthermogram [i(T)] corresponds to an e00(T) plot at the frequencyof �10�3 Hz. Typically, the TSC experiment consists ofpolarizing a sample at some sufficiently high temperatureTp byapplying a static electric field Ep for a sufficient period of timetp. When sample temperature is lowered to T0 with the electricfield applied and then shorted at T0, the non-equilibrium stateof the system is frozen-in. During the successive heating, at aconstant heating rate b¼ dT/dt, the recovery of the system ismonitored by measuring the depolarization current density

JðTÞ ¼ P0ðTPÞtðTÞ exp � 1

b

ðTT0

dT 0

tðT 0Þ

264

375 ð3Þ

where P0 is the saturation polarization. Deviations in the shapeof a TSC peak from that described by Equation (3) areinterpreted on the basis of discrete or continuous distributionsin the relaxation times t(T). Some features of the TSC methodmake it particularly attractive as a complementary technique toDRS: The method is fast allowing for a quick characterizationof the overall dielectric response of the material underinvestigation, it is characterized by high sensitivity and highpeak resolving power (due to its low equivalent frequency)and it offers special techniques to analyze experimentallycomplex relaxation mechanisms into approximately singleresponses.[34]

The TSC scans were performed in vacuum (10�2–10�4 Pa),usually in the range �243–77 8C (or to 180 8C for blends withwPEO� 20 wt.-%), Ep� 2� 106 V/m, tp¼ 5 min and b¼ 5 8C/min. In each case, Tp was kept below the DSC-determinedmelting point of the crystalline component. Samples of typicaldimensions (5� 5� 1) mm3 were used. In order to facilitatecomparisons between different samples, the TSC spectra aregiven in transient conductivity units s(T)¼ J(T)/Ep. In mostcases, spectra were collected with the MISIM electrode

Figure 1. Chemical structures of: (a) poly(ethylene oxide) (PEO), (b) 4,40-diaminodiphenyl-methane(DDM), and (c) diglycidyl ether of bisphenol-A (DGEBA) resin.

Thermophysical Properties and Molecular Relaxations in Cured Epoxy Resinþ PEO Blends: . . . 881

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configuration (M¼ copper electrode electrolytically coveredby chromium, I¼ thin insulating FEP or Teflon foil, S¼sample).

Differential Scanning Calorimetry (DSC)

A Perkin-Elmer Pyris 6 differential scanning calorimeter undera dry nitrogen atmosphere was used, calibrated with an indiumstandard. Tg values were determined as follows. The samples(with weights 3.3–7 mg) were rapidly cooled from roomtemperature to�120 8C and subsequently heated to 110 8C andheld there for 5 min to eliminate the thermal history. This wasfollowed by a second quenching to �120 8C. Finally, thesamples were reheated to 200 or 250 8C (second scan at aheating rate of 20 8C min�1). Glass transition temperatureswere estimated as the midpoint of the heat capacity change.The apparent melting enthalpy Hm,blend and melting tempera-tures Tm(2) were determined from the DSC endothermic peakson the second heating run, based on the area and the maximumof the endothermic peaks, respectively. Crystallization enthal-piesHc and temperatures Tc were derived respectively from thearea and the maximum of the exothermic peaks recorded in thesecond cooling step.

Thermal Properties of Cured ERþPEO Blends

As a result of curing, the initial mixture (DGE-

BAþDDMþ PEO) undergoes a series of structural

changes, e.g., growth of chains, branching of chains, and

gelation, which transform the initial ternary system to a

binary system consisting of epoxy resin and PEO.

ERþ PEO blends with less than 50 wt.-% PEO were

transparent. On the other hand, due to the formation of PEO

spherulites, the blends with higher PEO contents were

opaque. Figure 2 shows the second heating scans for

selected cured blends. Relatively weak glass transitions,

indiscernible in the scale of the plot, were observed for most

samples and their characteristics (Tg, DCp) are listed in

Table 1. Significant is the absence of cold-crystallization

transitions during the second heating run, which implies

that crystallization was rapid and occurred to completion

during the quenching. Crystalline PEO, evidenced by

exothermic crystallization peaks in the region between 35

and 41 8C, has been observed in samples with PEO weight

fractions wPEO> 40 wt.-%.

The values shown in Table 1 for the degree of crystallinity

were calculated as:

Xc;blend ¼ Hm;blend

H0m

ð4Þ

and

Xc ¼Hm;PEO

H0m

ð5Þ

where Hm,blend and Hm,PEO are the apparent melting

enthalpies per gram of blend and of PEO present on the

blend, respectively, and Hm0 is the enthalpy of melting per

gram of perfectly crystallized PEO (Hm0 ¼ 205 J/g). From

the values shown in Table 1 it is clear that the crystallinity of

the blends decreases with increasing concentration of the

non-crystallizable component (ER). This result could be

ascribed to the formation of the cross-linked epoxy network

and its strong topological influence on PEO crystallization.

Observation of single calorimetric glass transition peaks

(or equivalently of dielectric and mechanical segmental

relaxations) between the constituent values for the

amorphous components is regarded as a clear manifestation

of phase miscibility in binary cured thermosetþ polymer

systems.[18–22] The Tg–composition relationship observed

in our blends (Table 1 and Figure 3) has been described in

the left half of the diagram by an empirical equation of

Kwei[35]

Tblendg ¼

wPEOTPEOg þ kwERT

ERg

wPEO þ kwER

þ qwERwPEO ð6Þ

with k¼ 1 and q¼�203� 5. Equation (6) reduces to the

empirical equation developed by Jenckel and Heusch when

k¼ 1.[36] Apart from Equation (6) several other relations

exist,[37] such as the Gordon–Taylor equation[38]

Tblendg ¼

wPEOTPEOg þ kwERT

ERg

wPEO þ kwER

ð7Þ

and an empirical equation developed by Fox[39]

1

Tblendg

¼ wPEO

TPEOg

þ wER

TERg

ð8Þ

These equations were tested for their applicability in

describing the Tg-composition dependence. The DSC-

Figure 2. DSC curves of DDM-cured ERþ PEO blends (secondscans).

882 I. M. Kalogeras, A. Vassilikou-Dova, I. Christakis, D. Pietkiewicz, W. Brostow

Macromol. Chem. Phys. 2006, 207, 879–892 www.mcp-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

based estimate for the Gordon–Taylor parameter k equals

0.42� 0.02 (Figure 3). The first term in Equation (6) and (7)

represents the mixing contribution derived from the

additive rule of entropy and/or of the volume of the

blends, while the quadratic term (qwERwPEO) in

Equation (6) was introduced to compensate for the binary

interaction between the segments of the component

polymers. In the above equations, k is an adjustable

parameter related to the degree of curvature; q is supposed

to be related to the change in energy associated with the

formation of a contact pair between different segments in

the mixture.

The physical meaning of q has been discussed in some

detail by Kwei and co-workers [35] who claim that

q ¼ 2fd

3kB

E12 �E11 þ E22

2

� �ð9Þ

where fd is a dispersion factor related to the completeness of

the mixing process of the two polymers and E11, E22, E12

represent the contact energy of the segments of polymers 1

and 2 (cured ER and PEO). Therefore, q should be

indicative of the extent of intermolecular interactions

between blend components. The highly negative value for

the interaction parameter q suggests lower propensity of

chain segments to form intermolecular interactions. We

note that with the same k values, the q value determined for

the present system (q¼�203) is close to that determined by

Zheng et al.[40] for THB-cured ERþ poly(e-caprolactone)

(PCL) blends (q¼�187). Nevertheless, it is substantially

lower from that determined by Ni and Zheng[41a] for DDS-

cured ERþ PCL blends (q¼�33) and by de Juena and

Cortazar[41b] and Kuo et al.[41c] for phenoxyþ PCL blends

(q¼�128 and �100, respectively). Assuming that Kwei

explanation is correct, the strength of the intermolecular

hydrogen-bonding interactions in our blends is much

weaker than that in the latter systems.

Irrespective of the functional form used to fit the Tg

versus wt.-% PEO relationship we observe strong positive

deviations of the experimental Tgs from the fitting curves

above 50 wt.-% PEO. This agrees with several calorimetric

results,[4,11,21] and appears independent of hardener’s type

and content. The deviations are related to the compositional

enrichment of the amorphous region, created by crystal-

lization of PEO (i.e., increase of the ER concentration in the

residual amorphous PEO phase). Moreover, an increase of

the glass transition temperature is expected due to the

stiffening of the amorphous phase by reinforcement from

PEO spherulites.[10] We see the Gordon–Taylor curve in

Figure 3 provides a better fit to the midpoint concentration

(50% of each component) of the Kwei equation. Up-shift of

Figure 3. Plots of transition temperatures as function of blendcomposition. Data fitting lines, based on the Kwei, the Gordon–Taylor and the classical Fox equation are shown. The value of�57 8C was used for the glass transition temperature ofamorphous PEO.[13]

Table 1. Parameters of thermal transitions:Tg, melting temperature determined from the second heating scanTm(2), specific heat capacityjump DCp, Tc, Hc, and values of the crystallinity of the PEO component in the blends Xc and of the cured ERþ PEO blends Xc,blend. Theweight percentage of PEO in the amorphous phase (w0

PEO) is also included.

wPEO Tg DCp Tm(2) Hm,blend Tc Hc Xc Xc,blend w’PEO

wt.-% 8C J � g�18C�1 8C J � g�1 8C J � g�1 % % wt.-%

0 162.0 0.27 – – – – – – –10 123.0 0.33 – – – – – – –20 83.1 0.31 – – – – – – –30 54.9 0.27 – – – – – – –40 – – 60.6 11.38 35.4 �9.80 6 2 3950 9.0 0.07 53.2 20.30 – – 10 5 4760 7.3 0.18 60.8 54.64 36.9 �50.19 27 16 5270 – – 66.0 99.38 38.4 �93.27 48 34 5580 9.3 0.16 62.2 126.80 39.6 �124.10 62 49 6090 3.0 0.05 62.0 147.05 40.7 �145.46 72 65 72

100 �46.9 0.09 64.1 157.94 40.8 �152.94 77 77 100

Thermophysical Properties and Molecular Relaxations in Cured Epoxy Resinþ PEO Blends: . . . 883

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this experimental Tg value seems to be due to the formation

of a crystalline PEO phase.

In Figure 3 we also show how the melting temperature of

the crystalline component Tm and Tc change as a function of

blend composition. Tms of PEO in the blends slightly

decrease with an increase of DGEBA concentration. The

above phenomenon is characteristic of miscible blends and

has been explained thermodynamically by Nishi and

Wang.[42] In the present study, however, the melting point

depression is neither strong nor consistent so as to provide a

clear-cut criterion for phase miscibility.

Dielectric Characterization ofCured ERþPEO Blends

Transient Conductivity Spectra

Figure 4 shows transient conductivitys(T) signals of typical

relaxation mechanisms in vacuum-dried thick films of

DDM-cured DGEBA (ER) and pure PEO. The spectra of

selected cured ERþ PEO blends are divided into two

groups. In Figure 4(a) we show spectra of blends containing

appreciable crystalline PEO fractions (wPEO > 40 wt.-%).

The location of the relaxation modes ascribed to the

constituting materials is indicated. Figure 4(b) shows

selected spectra of amorphous blends along with the

nomenclature used for the main relaxation modes. In all

cases, thin insulating FEP films were interposed between

the samples and the electrodes to separate relaxation

response from the intense (in some blends) high-T dc

conductivity signals.

The main relaxation mode in the TSC spectra of cured

epoxy resins is the glass-transition signal aER. As expected,

the signal is highly dependent on hardener’s type and content

as well as curing time and temperature. For aromatic amine-

cured DGEBA, the aER band peaks at Ta¼ 156 8C, close to

the calorimetric Tg estimate of 162 8C. This signal is outside

the temperature range of the spectra shown in Figure 4; a low

temperature component, peaking near the polarizing temper-

ature (Tp¼ 57 8C) is only present.

The secondary transition [bER, Figure 4(a)] is located

around �108 8C, in agreement with the TSC estimate

reported by Shin et al.[25] Based on dynamic mechanical

analysis (DMA) of DDM-cured ER and ERþ PEO

blends,[5] the bER mode has a loss factor (tan d) peak

around �55 8C (measured at f¼ 0.1 Hz). The lower

temperature of the peak maximum, Tmax, in the TSC

spectrum is in agreement with the lower equivalent

frequency of the method (feq� 10�3 Hz). Several other

cured ER systems show TSC bER processes in the same

temperature range. For example, DDS-cured tetraglycidyl-

4,40-diaminodiphenyl (TGDDM) methane shows bER

maxima between �100 and �80 8C and triethylenetetr-

amine (TETA)-cured DGEBA shows this sub-Tg relaxation

between �113 and �73 8C.[27] In terms of the amino cross-

linked epoxy systems, the bER transition is attributed to the

crankshaft motion of the hydroxyl-ether groups (glyceryl

segment I, Figure 5), which most of them form after the

amine-epoxy curing reaction. It has been also reported that

the diphenyl propane units contribute to the relaxation with

the ‘‘flip-flop’’ motion of the phenyl rings that would relax

near �110 8C.[43]

Earlier TSC and thermally stimulated creep (TSCr)

studies of amine-cured epoxy networks revealed that both

Figure 4. Transient conductivity s(T) spectra of cured ER, purePEO and cured ERþ PEO blends. To reduce the strong high-T dcconductivity signal of PEO and reveal submerged relaxations thespectra were recorded with the MISIM electrode configuration(I¼ FEP).

884 I. M. Kalogeras, A. Vassilikou-Dova, I. Christakis, D. Pietkiewicz, W. Brostow

Macromol. Chem. Phys. 2006, 207, 879–892 www.mcp-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

the magnitude and Tmax of the bER mode increases with

increasing cross-link density.[43] According to Mangion and

Johari,[26] during curing the height of the secondary

relaxation peak first increases to a maximum value and

then decreases (at post-cure periods) due to physical aging

effects. Based on the above results, the TSC observation by

Lee et al.[28] of a bER process near �12 8C for DDM-cured

DGEBA should be considered with caution. A strong peak

upshift would necessitate very high cross-link density

accompanied by strong dipole–dipole, van der Waals or

hydrogen bonding interactions of the relaxation entities

with their surroundings; yet, none of the above factors

seems sufficient to account for such a dramatic upshift. The

very weak signal around �25 8C, generally ascribed to

unreacted hardener molecules,[18] complicates further an

explanation.

The diffuse low-T gER mode is attributed to main-chain

motion, most likely of the glycidyl ether segment (segment

II). The corresponding TSC signal appears here at temper-

atures well below �140 8C (Tmax��113,[28] �123,[27]

�145,[25] and �155 8C[43]). Among others, Mangion and

Johari[26] observed that the concentration of unreacted

epoxy groups decreases with increasing cross-linked

density, reached at longer curing and post-curing (or aging)

programs. In addition, Boye et al.[43] reported a drop of the

gER TSC band with [r] increasing from 0.5 to 1 (i.e.,

increased cross-linking density). In the present case, the low

dielectric strength of the gER mode is consistent with the

high cross-linking density obtained for [r]¼ 1.

The structure of the TSC spectrum is highly dependent on

PEO’s crystallinity, molecular weight, degree of hydration,

and preparation conditions. The bPEO transition

(Tmax ��40 to �21 8C) is attributed to long-range

segmental chain reorientations directly correlated with the

amount of amorphous phase in the amorphous–crystalline

two-phase structure of PEO.[44] Association of the bPEO

mode with a cooperative glass transition phenomenon in

disordered regions is supported by the disappearance of this

signal in single-crystal laminate solution-grown PEO (with

Mw > 105). The presence of the bPEO band well above the

frequently reported calorimetricTg of PEO (e.g.,�57 8C for

high–molecular-weight material) is attributed to the high

degree of crystallinity. The secondary gPEO process

(Tmax ��170 8C) is assigned to local twisting movements

(crank-shaft or kink motion) in the main chains in both

defective regions within crystallites and non-crystalline

regions,[44] including amorphous segments in the fold

structure on crystal surfaces. Note that in earlier TSC

studies of uncomplexed PEO the gPEO process has been

reported to peak at about �155 8C, showing (non-system-

atic) high-T shifts in the presence of salts and also great

sensitivity to the nature of electrodes.[45,46] The latter effect

was also observed in the present study (e.g., see Figure 6).

Both the location and strength of the peak recorded above

bPEO (e.g., at 20 8C in the homopolymer) is also highly

dependent on the selection of the electrode material.[47]

Blending induces modifications in the segmental and

local-chain dynamics of the constituting phases. The

relaxation bands attributed to local-chain motions (gPEO

and gER) activate in the same temperature region forming a

complex band (denoted gblend), which locates between

�190 and �140 8C. The main component of this signal

peaks near the temperature of the gPEO signal and is

relatively strong (higher smax(g)) and high-temperature

Figure 5. Section of cross-linked DDM-cured DGEBA. Thechain segments responsible for the bER (segment I) and gER

(segment II) relaxation modes of the cured epoxy resin areindicated.

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shifted (higher Tmax(g)) for blends with PEO contents

between 40 and 60 wt.-%.

The addition of PEO affects also the position and the

strength of the bER peak, in relation to the signal recorded in

the thermoset. The spectral information shown in Figure 7

and its insert cast doubt on the interpretation of earlier DMA

data. Zheng et al. report that DDM-cured blends show bER-

related signals in the range �100 to 0 8C, with loss factor

(tan d¼ e00/e0) peaks around �59 and �80 8C, for pure ER

and 50/50 ERþ PEO, respectively (measured at f¼ 0.1

Hz).[18] Based on our results involving the same composi-

tions, the TSC signal shows a clear upshift, irrespective of

the blocking-electrode system (e.g., from �107 to �99 8Cusing FEP films). According to Zheng et al., a high-T shift of

the DMA bER signal was observed only in the blend with

10 wt.-% PEO.[18] They explained this by association of

hydroxyl-ether units in the epoxy with oxygen atoms in

PEO, hindering the rotational motion of the hydroxyl ether

groups. Further, Zheng and co-workers suggested that

increasing concentration of the linear polymer component

should create, and progressively make dominant, plasticiz-

ing of the bER mode in the epoxy. We have studied more

blends than Zheng et al. and we find that plasticization in not

supported by simple DTmax considerations: a low-T shift

was absent for bblend modes in blends with 40–50 wt.-%

Figure 6. Dependence of the s(T) spectrum of pure PEO(Mw ¼ 2� 104 g �mol�1) on the electrode configuration: (a)copper electrodes electrolytically covered by chromium (ohmicelectrodes), and blocking electrode MISIM configurations, withI¼ FEP (b) or Teflon (c) films.

Figure 7. Plots of the transient conductivity signal at peak position (smax) of thebblend signal as a function of the PEO content in the blends. The inset shows thevariation of the band temperature maximum. Spectral parameters obtained withdifferent blocking electrodes are shown. The dashed lines are guide for the eye.

886 I. M. Kalogeras, A. Vassilikou-Dova, I. Christakis, D. Pietkiewicz, W. Brostow

Macromol. Chem. Phys. 2006, 207, 879–892 www.mcp-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

PEO. Conceivably, plasticization may be invoked to explain

an increase in the relaxation strength of bER (Figure 7) in

blends with 30 and 40 wt.-% PEO.

The position of the single prominent glass transition

mode of each blend (ablend) changes irregularly with the

change of the PEO/ER weight fraction. Its shape is also

dependent on this ratio (Figure 8). Broadening is highly

dependent upon the actual electrode configuration, since

this alters contributions from satellite peaks in a rather

irregular way. However, it is clear that broadening is

significant in blends with 20 wt.-% of PEO – when the

polyether is fully amorphous.

Frequency Dependence of Dielectric Modes

Representative frequency dependence of dielectric permit-

tivity e0 and loss factor e00 for the blend with wPEO¼ 90 wt.-

% are shown in Figure 9. The spectra in Figure 9(a) refer to a

dry blend (recorded at T¼ 50 8C) and in Figure 9(b) to a

hydrated one (weight increase due to water absorption

h¼ 4.14%, T¼ 25 8C). The main phenomena seen are the

glass transition (in the 1–105 Hz range) and dc conductivity

signals at lower frequencies. We also provide the electric

modulus, M� ¼ 1=e� ¼ M0 þ jM00.[48] The real M0 and

imaginary M00 parts of the electric modulus were obtained

by transforming the permittivity data as follows:

M0ðf Þ ¼ e0ðf Þe02ðf Þ þ e002ðf Þ ð10Þ

and

M00ðf Þ ¼ e00ðf Þe02ðf Þ þ e002ðf Þ ð11Þ

Using the modulus formalism the large contribution of the

non-local relaxation at low frequencies is suppressed, while

theM00(f) spectrum preserves the general shape and position

of the ablend dispersion. The temperature dependence of the

maximum of the low-frequency M00(f) band and its position

in the frequency scale below the glass transition region

suggests that the signal is space-charge related. A ‘‘con-

ductivity relaxation’’ (CR) mechanism is the most probable

cause;[49] however, considering the heterogeneities found in

this blend, a Maxwell–Wagner–Sillars (MWS) polariza-

tion phenomenon most likely overlaps with CR. Strong

overlap of CR and MWS signals is common in the DRS

studies; a recent example is the structure of the M00(f)spectra of microphase-separated bioactive polyur-

ethanes.[50] We note that with the samples in the hydrated

state both mechanisms shift to higher frequencies.

A representative example of the temperature variation of

the dielectric response is shown in Figure 10. At low

temperatures the M00(f) spectrum of the blend with

wPEO¼ 50 wt.-% shows a gblend relaxation at high

frequencies (e.g., fg,blend� 105 Hz at �70 8C), a small

contribution to the overall relaxation response.[51] The

dielectric strength Deg of the DRS gblend mode attains its

maximum value in the 50/50 ERþ PEO blend,[47] con-

sistent with the TSC estimate. We note that many TSC

spectra reveal a bimodal nature for the gblend signal, while

the high-frequency DRS response can be described by a

single relaxation mechanism. In both cases, however, the

gblend signal is broad (e.g., the half-width of the M00 signal is

Figure 8. Normalized plots (s/smax vs. T/Tmax) for the ablend

TSC signal ascribed to the glass transition mechanism of theERþ PEO blends. All spectra were recorded in the MISIMelectrode configuration, with ‘‘I’’ referring to FEP [Figure 8(a)] orTeflon [Figure 8(b)] films.

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much greater than that of Debye peak, 1.14 decades)

confirming that this mechanism is characterized by a broad

distribution of relaxation times. Analogous observations

apply for the neighboring signal (bblend mode), which enters

the frequency window when temperature increases above

�90 8C. The broad range of the b relaxation signal

illustrates the heterogeneous structure of the material.[27]

The Havriliak–Negami model was used to describe the

dielectric behavior and the Arrhenius rate law was found to

fit well sub-glassy relaxations (Figure 11). In all blends for

which data analysis was possible, the (re)orientation of the –

OH groups is controlled by an apparent activation energy

barrier of Eg,blend¼ 34� 2 kJ �mol�1. This value lies in the

range �20–50 kJ �mol�1 for typical local relaxations in

polymer matrices.[52] Arrhenius-type fits for the bblend

mode give estimates for Eb,blend that gradually decrease

with the addition of PEO. For example, Eb,blend¼ 90� 2

kJ �mol�1 for the blend with 20 wt.-% PEO, and it reduces

with increasing polyether content to 86� 4 kJ/mol (50 wt.-

% PEO), 71� 1 kJ �mol�1 (70 wt.-% PEO) and 65� 4

kJ �mol�1 (90 wt.-% PEO).

Preliminary results on highly hydrophilic PEO-rich

blends show that certain characteristics of the space-charge

related (CR and MWS) and segmental (ablend) relaxations

dependent strongly on the equilibrium hydration level (h).

Their high-frequency shift with increasing equilibrium

hydration level (ranging from 0 to 4.14% with relative

humidity environments from 0 to 75%) can be explained by

plasticization caused by occlusion of water molecules.

Further, sdc increases and the glass transition signal

broadens with water absorption (Figure 12).

Conclusion

Perturbations in the relaxation dynamics of the constituting

phases induced by blending provide information on

intermolecular interactions as well as structural character-

istics of the cure products. A characteristic example is the

downshift of the single glass-transition-related TSC signals

(quantified by the temperature of the TSC peak maximum,

Ta) with increasingwPEO (Figure 13). The shift is inherently

characteristic for the system and appears regardless of the

blocking electrode material. It is also in agreement with

DSC results (Figure 3). Satisfactory fits of the TSC data

were obtained using the Kwei equation with k¼ 1 and

parameter q in the range �277� q��228 [Figure 13(a)]

and to a lower degree using the Gordon–Taylor equation

[0.32� k� 0.39, Figure 13(b)]. The DSC estimate for the

Gordon–Taylor equation parameter is only slightly higher

(k¼ 0.42).

It has been proposed that the adjusting parameter k, used

in the Gordon–Taylor equation, can be taken as a semi-

quantitative measure of miscibility and strength of the

intermolecular interaction between components of polymer

blends. For instance, in blends of PCL with poly(vinyl

chloride) and chlorinated poly(vinyl chloride)s, the DSC-

based estimates for the Gordon–Taylor parameter k

increase from 0.56 to 0.76 with increasing degree of

chlorination,[53] which gives rise to increased amount of

intermolecular hydrogen-bonding between the a-hydrogen

of chlorinated polymers and the carbonyl of PCL. The

present estimate of a relatively low k value may be a

consequence of a decreased enthalpy contribution, which in

Figure 9. Dielectric permittivity (e0), dielectric loss (e00), andelectric modulus (M0, M00) spectra of a representative ERþ PEOblend (wPEO¼ 90 wt.-%): (a) dry blends (h¼ 0.00%, T¼ 50 8C)and (b) hydrated blend (h¼ 4.14% at 75% RH, T¼ 25 8C).

888 I. M. Kalogeras, A. Vassilikou-Dova, I. Christakis, D. Pietkiewicz, W. Brostow

Macromol. Chem. Phys. 2006, 207, 879–892 www.mcp-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Figure 10. Isothermal M00(f) spectra for the blend with 50 wt.-% PEO, recordedat selected temperatures. The spectrum recorded at �70 8C is shown with athicker line.

Figure 11. Representative Arrhenius plots for mechanismsisolated in pure PEO and DDM-cured ERþ PEO blends (90, 70,50, and 20 wt.-% PEO). The lines are fits of the experimental data toappropriate equations for the temperature dependence for therelaxation frequencies.

Figure 12. Modification of the room temperature M00(f) spec-trum of the 10/90 ERþ PEO blend as result of (equilibrium) waterabsorption (RH%: relative humidity; h: equilibrium hydrationlevel).

Thermophysical Properties and Molecular Relaxations in Cured Epoxy Resinþ PEO Blends: . . . 889

Macromol. Chem. Phys. 2006, 207, 879–892 www.mcp-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

turn indicates weakened intermolecular-specific phase

interactions in the cross-linked structures. Analogous

assessments have been reported for miscible DDS-cured

(k¼ 0.42,[20]) and THB-cured (k¼ 0.26,[14]; k¼ 0.31,[40])

DGEBAþ PEO blends, with additional verification pro-

vided by FTIR spectroscopy results.

The observation of subtle changes in the local-chain

relaxation dynamics provides additional verification of our

arguments. It is generally accepted that the relaxation

behavior of the secondary relaxation in pure resins is

tailored by self-associations of –OH groups. A moderate

strengthening of the dielectric bER mode in several blends

(insert in Figure 7) suggests that the addition of the linear

polyether facilitates local motion of glyceryl segments. The

DRS result of a decrease in the apparent activation energy

barriers with the addition of PEO supports the above

scenario. Such ‘‘plasticization’’ effects indicate a reduction

in the number and/or the strength of intermolecular

interactions in the cross-linked epoxy network, accompa-

nied by some changes in the structural environment of the

relaxing entities. The role of the above effect is further

exemplified by the observation of Boye et al. that imperfect

cross-linking of the epoxy drastically depresses the bER

peak, i.e., due to the internal antiplasticization of the

mechanism by unreacted epoxy groups.[43] In our case, the

band intensification is readily observable for PEO contents

between 20 and 40 wt.-% (Figure 7), where the polyether

component is completely amorphous and the blend is in a

completely miscible state.[18] Considering that the strength

of the bblend signals shown in Figure 7 for PEO-rich blends

should be corrected to account for the actual content of ER

in the blend, it appears that polyether addition eventually

facilitates local motion of hydroxyl units in the epoxy.

Observations concerning the primary (a) and the

secondary (b) modes can be related to mechanisms

inducing miscibility of the components. The basic inter-

action force between the constituent chain segments is

considered to be hydrogen bonding between group

pairs, such as the hydroxyl (–OH) of epoxy and the ether

group (–O–) of PEO, which act antagonistically to the self-

association of hydroxyls in pure ER. As the hydrogen bond

is a directional attractive interaction between electron-

deficient hydrogen and a region of high electron density, its

strength is directly related to all the elements affecting the

acidity of the proton donor, the basicity of the proton

acceptor, and the accessibility of the donor and acceptor.

Among these, chemical and stereo structures of the donors

and acceptors are essential in determining the strength of

the hydrogen bonds. Therefore, factors such as steric

shielding (the screening effect) and low accessibility of

functional groups are likely to reduce specific interactions

between PEO and ER chains due to the formation of a cross-

linked network.[11] Surprisingly, similar observations have

been reported even for uncured binary blends.

Based on the present results, and in agreement with

earlier FTIR studies by Horng and Woo[20] and by Hu

et al.[11] in cured DGEBAþ PEO mixtures, the intermo-

lecular hydrogen-bonding interactions are much weaker

than the self-association of hydroxyls of ER. This behavior

Figure 13. Plots of TSC-based glass-transition temperaturesTa(�Tg) as a function of blends’ composition. TSC scans forsamples with wPEO � 20 wt.-% were conducted in the MSMelectrode configuration, since their relatively low dc conductivitydoes not obscure the isolation of the glass-transition related currentsignal. Note that the Ta values of blends containing PEOcrystallites (wPEO � 50 wt.-%) were not used in the fits.

890 I. M. Kalogeras, A. Vassilikou-Dova, I. Christakis, D. Pietkiewicz, W. Brostow

Macromol. Chem. Phys. 2006, 207, 879–892 www.mcp-journal.de � 2006 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

is in marked contrast to the enhanced strength of the

intermolecular interactions in cured phenoxyþ PEO (mis-

cible blend).[11] In cases where dH interactions are possible,

cured blends may be either miscible (e.g., DDS-cured

ERþ PEO) or phase-separated (e.g. DETA-cured

ERþ PEO).[20] A homogeneous structure has been reported

even for blends cured with aromatic anhydride (PMDA), for

which structural similarity with the cure components is

present but hydrogen bonding is not expected.[20] In

contrast, aliphatic agents fail to produce miscible cured

thermosetþ polyether blends. Combining the above obser-

vations, we infer that the presence of hydrogen-bonding

sites between specific functional groups is neither a

sufficient nor a necessary condition for phase homogeneity.

It is the structural similarity of the components (resin,

hardener and polyether) and the extent of inter-associated

hydrogen bonds (rather than the strength of the bond itself)

that seems to govern the miscibility or compatibility of the

particular structures.

Acknowledgements: The present research has been funded byE.P.E.A.K. 2 (Operational Programme for Education and InitialVocational Training, Athens) in the framework of PYTHAGORAS(Project 70/3/7362).

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