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Materials Science and Engineering A 480 (2008) 259–265

Creep behaviour of layered silicate/starch–polycaprolactoneblends nanocomposites

C.J. Perez, V.A. Alvarez, A. Vazquez ∗Research Institute of Material Science and Technology (INTEMA), National University of Mar del Plata (UNMdP),

Av. Juan B. Justo 4302, 7600 Mar del Plata, Argentina

Received 14 May 2007; received in revised form 2 July 2007; accepted 12 July 2007

bstract

The creep behaviour of biodegradable composites based on starch/polycaprolactone commercial blends reinforced with an organo-modifiedanoclay, processed by melt-intercalation was studied. Clay content and temperature effects were also analyzed. The experimental behaviour wasorrelated with several models. The master curves were built by means of the time–temperature superposition principle. Findley’s power law

orrectly predicted the strain–time curves and also the complete compliance curves. The Burgers’s model (four parameters) was also used allowinghe relationship between the creep behaviour and the composite morphology. All the results showed that the addition of clay to the neat matrixeads to a significant improvement of creep resistance. 2007 Elsevier B.V. All rights reserved.

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eywords: Creep; Clay; Nanocomposites; Biodegradable polymers; Modelling

. Introduction

In the last years, biodegradable polymers have been widelysed for packaging with the emphasis on the reduction ofnvironmental pollution [1,2]. Even tough biodegradable poly-ers mean ecological benefits, they are required to have

hysical–mechanical properties and cost similar to those of theraditional plastics. The interest on starch-based polymers is con-tantly increasing due to its low cost but it is well know that itroperties are poor [3,4]. In order to improve the mechanicalroperties, small quantities of layered silicates can be addedroducing materials with desired properties [5,6].

Creep behaviour is a very important property of a thermo-lastic composite that controls its dimensional stability andspecially in applications where the material has to support loadsor long periods of time [7–9]. This mechanism can drive to annacceptable deformation level and in due course to a structuralailure.

On this field, Vlasveld et al. [10] have investigated thereep and physical aging behaviour of polyamide six (PA6)anocomposites. They demonstrated that the creep compliance

∗ Corresponding author. Tel.: +54 223 4816600.E-mail address: anvazque@fi.mdp.edu.ar (A. Vazquez).

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921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2007.07.031

as diminished by layered silicate incorporation and related thisith the modulus increment, but the shape of the curves was sim-

lar to that of unfilled PA6. In addition, Pegoretti et al. [11] haveonsidered the effect of nanoparticles on the creep behaviour ofolyethylene terephthalate (PET). They showed that nanoparti-les improved the creep resistance of the neat matrix. Galgalit al. [12] have studied the creep properties of polypropy-ene (PP)/maleated polypropylene (PP-g-Ma) reinforced withlays. They established that clay produces an enlarged creepesistance and drive to higher shear viscosity. Ranade et al.13] have analyzed the creep performance of polyethylenemaleated and non-maleated)—montmorillenite layered silicatelown films. They found a decrease on the creep compliance andn increase on the modulus by clay addition. Yang et al. [14,15]ave characterized the tensile creep resistance of polyamide 66PA66) nanocomposites. They proved that the deformation ofhe nanocomposites were clearly lower than that of the pure

atrix showing their improved creep performance. They haveemonstrated that the nanocomposites exhibited higher elastic-ty and that the nanoreinforcements were able to improve theolymer elasticity. Chiou et al. [16] have evaluated the rheol-

gy of starch–clay nanocomposites by using several clays. Theyonfirmed that the decrease on the creep compliance was evi-ently related with the polymer–clay interaction being highern the case of higher compatibility. Martucci et al. [17] have

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60 C.J. Perez et al. / Materials Science

tudied creep behaviour of glutaraldehyde-crosslinked gelatine.hey demonstrated that chemical crosslinking drives to a loweriscous creep and an increment on the elastic part.

Related with starch-based blends, Cyras et al. [9] and Alvarezt al. [18] have investigated the effect of natural fibres on thereep behaviour of such kind of blends. Both demonstrated thathe addition of fibres improves the creep resistance of poly-

eric matrices and this improvement strongly depends on thebre content, their dimensions, orientation, distribution, and thedhesion between the fibres and matrix.

The aim of this work was to investigate the influence oflay incorporation and clay content on the creep behaviourf starch/polycaprolactone blend. The creep behaviour of thisaterial is of crucial interest because it is used above its glass

ransition temperature. Several models and equations were usedo predict the long-time response of studied materials.

. Experimental

.1. Materials

The matrix used in this work was a commercialtarch/polycaprolactone blend called MaterBi Z [19], kindlyupplied by Novamont, Novara, Italy. An organoclay, under theommercial name of Cloisite 10A® (C10A) was used as nanor-inforcement. It was purchased from Southern Clay Productsnc., USA. The organoclay was a natural montmorillonite mod-fied by quaternary ammonium salt:

at a concentration: 125 mequiv/100 g of clay. This organo-odified clay was chosen due to their compatibility with the

iodegradable matrix, obtained in a previous work [20].

.2. Composite preparation

Composites were prepared by melt-intercalation followed byompression-moulding. An intensive Brabender type mixer withwo counter-rotating roller was used. Mixing temperature was00 ◦C; speed of rotation was 150 rpm and mixing time was0 min. The concentration of clay ranged from 0 to 7.5 wt.%.ompression moulding was carried out in a hydraulic press for0 min at 100 ◦C. The thickness of the samples was between 0.3nd 0.5 mm.

.3. Methods

Creep tests were conducted in a DMA 7-e Perkin-lmer under nitrogen atmosphere. Specimen dimensions were0 mm × 2 mm × 0.4 mm, according to ASTM 2990 standards.ests were carried out at five different temperatures: 3, 10, 20,5 and 30 ◦C, all of them above the glass transition tempera-

ure of the matrix (−66 ◦C). These temperatures were selectedecause these are typical temperatures of use. A 3 MPa stressas applied for 30 min and then it was removed and recoveryeasurements were recorded.

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ngineering A 480 (2008) 259–265

.3.1. Theoretical backgroundCreep performance is commonly represented by creep com-

liance, J(t) which is defined as the ratio between creep strain ε

t) and the applied stress (σ) as follows:

(t) = ε(t)

σ(1)

It is known that creep modulus; defined as the relationshipetween the applied stress and the deformation at a given time,ecreases as a function of time and temperature. William, Landelnd Ferry [21] have developed the WLF theory showing that, inhe linear viscoelastic range, a viscoelastic curve determined atn arbitrary temperature T1 can be carried to another one T2aking an appropriate translation on the time axis by applyingshift factor aT at a reference temperature Tr. When reference

emperature is far from the glass transition temperature (morehan 50 K) the Arrhenius law is preferred. Based on it, the shiftactor can be described as a function of temperature as follows:

og aT = − Q

2.303R

(1

T− 1

Tr

)(2)

here Q is the activation energy and R is the universal gasonstant.

.3.2. Constitutive modelThe viscoeslastic behaviour of thermoplastic composites has

een modelled by using a constitutive model based on a poweraw equation known as Findley power law [22]. It can bexpressed as follows:

(t) = ε0 + Atn (3)

here ε(t) is creep strain at time t, ε0 the instantaneous initialtrain, A the amplitude of transient creep strain and n is the timexponent.

The instantaneous strain, the amplitude of transient creeptrain, and the time exponent are defined as creep parameters.he relative creep strain, εr, is defined as the ratio between thereep strain value at each time, ε(t), and the instantaneous initialtrain value, ε0. Rearranging Eq. (3):

r = A′tn (4)

here εr is the relative strain creep in percentage and A′ is thelope of the power law. A′ and n are parameters that can bebtained by plotting log εr versus log t; log A′ is obtained fromhe intersection with log εr axe and n from the slope of the curve.

Equation (3) can be also expressed in terms of creep compli-nce given by the following equation:

(t) = J0 + J1tn (5)

here J0 is the time-independent compliance, J1 is the coeffi-ient of time-dependent term and n is a constant, generally lowerhan one, independent on the applied stress.

Differentiating Eq. (3), it is possible to obtain the power lawreep rate, i.e. εPL, as follows:

˙PL = nA

t1−n(6)

ce and Engineering A 480 (2008) 259–265 261

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ε

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Ag

ε

3

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Fig. 1. Strain (percentage)–time curves for the neat matrix (MaterBi Z) at severaltemperatures.

Fn

copenhanced creep performance. On the other hand, n was around0.113 ± 0.015 and it was not affected by nanoclay incorpora-tion. Similar behaviour was observed by Yang et al. [15] in the

Table 1Average parameters obtained from Findley power law for the matrix (MaterBiZ) and their nanocomposites with clay

Clay content (wt.%) A′ n

0.0 0.661 ± 0.084 0.1010 ± 0.006

C.J. Perez et al. / Materials Scien

.3.3. Four parameters modelFour-parameters (or Burgers) model [23], gives the relation-

hip between the morphology of the composites and their creepehaviour. This model is a series combination of the Maxwellnd Kelvin–Voigt model [24], the total strain ε is given by theeneral equation:

= ε1 + ε2 + ε3 (7)

= σ

E1+ σ

E2

[1 − exp

(− tE2

η2

)]+ σ

η1t (8)

here ε1 and ε3 are the elastic and viscous strains representedy Maxwell model and ε2 is the viscoelastic strain representedy the Kelvin–Voigt model. E1 and E2 are elastic moduli, η2 and1 are viscosities, σ the applied stress and t is the creep time.his model has only one relaxation time, τ2 = η2/E2. When aonstant load is applied, the initial deformation takes place inhe spring with the modulus E1. Later deformation comes fromhe spring E2 and dashpot η2, in parallel, and from the dashpotith the viscosity η1. In the recovery test, after all the load is

emoved at 30 min, the creep is all recoverable except for theow that occurred in the dashpot with viscosity η1.

From this model, it is also possible to estimate the Burger’sreep rate, i.e. εB, by differencing Eq. (8). In this case, the creepate can be calculated as:

˙B = σ

η1+ σ

η2exp

(−tE2

η2

)(9)

t very long time, when creep rate get to a constant value, it isiven by:

˙B = σ

η1(10)

. Results and discussion

Fig. 1 shows the strain (percentage)–time curves for the neatatrix at several temperatures. In these curves, the creep stages

instantaneous deformation, primary and secondary creeps) arevident. On the other hand, there is no evidence of tertiary creep,.e. creep rupture, which would require longer times. As it wasxpected, strain increases with temperature. Similar behaviouras displayed by the nanocomposites. Fig. 2 shows the same

urves but at room temperature for the matrix and nanocom-osites with different clay contents. It is clear from this figurehat the strain of nanocomposites is lower than that of the neat

atrix and this implies that the creep behaviour was improvedy the presence of organoclay nanoparticles. In order to showhe effect of nanoparticles on the deformation of the pure blendt different temperatures, Fig. 3a and b summarizes the initialinstantaneous) and total (after 30 min of applied stress) defor-ations. Both deformations are noticeably less in the case of

anocomposites and they are reduced when clay concentrations increased from 1.0 to 7.5 wt.%. In addition the instantaneous

eformation as well as the total deformation exhibited almostn exponential variation with the temperature.

The parameters of Findley power law obtained for the matrixnd nanocomposites are showed on Table 1. A parameter was

1257

ig. 2. Strain (percentage)–time curves for the neat matrix (MaterBi Z) and theiranocomposites at room temperature and different clay contents.

alculated from A′ and ε0 and the obtained values are resumedn Fig. 4. It is apparent that the parameter increased with tem-erature and decreased with clay content which indicates an

.0 0.596 ± 0.127 0.1303 ± 0.025

.5 0.610 ± 0.058 0.1046 ± 0.010

.0 0.599 ± 0.130 0.1225 ± 0.013

.5 0.600 ± 0.070 0.1260 ± 0.014

262 C.J. Perez et al. / Materials Science and Engineering A 480 (2008) 259–265

Fig. 3. (a) Initial or instantaneous deformation (ε0) for matrix (MaterBi Z) andtheir nanocomposites as function of temperature, for different clay content. (b)Final, after 30 min of applied stress (ε30 min) deformation for matrix (MaterBi Z)and their nanocomposites as a function of temperature, for different clay content.

Fig. 4. A parameter of matrix (MaterBi Z) and their nanocomposites as a functionof temperature for different clay content.

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ig. 5. E1 parameter of matrix (MaterBi Z) and their nanocomposites as aunction of temperature for different clay content.

ase of polyamide 66 nanocomposites. ε0 parameter, which iselated with the instantaneous deformation, decreased with clayncorporation, showing an improvement on the elastic part.

Experimental curves were fitted by means of Burger’s model.he non-linear curve fit function of the OriginPro 7.5 softwareas used and the four parameters (E1, E2, η2 and η1) were

stimated. The correlation between the experimental data andodel prediction was really good. Fig. 5 summarizes the values

f E1 parameter for the matrix and nanocomposites as a func-ion of temperature. This parameter (associated to the Maxwellpring) established the instantaneous creep strain that would beecovered after stress elimination. It is evident that nanocom-osites have higher E1 values (increasing with clay content). E1f each kind of material displayed a decreasing trend with theemperature. This is related with the softening of the materialt high temperatures that drives to a decrease on the stiffness.or the range of temperatures studied in the present work, aseudo-linear relationship between E1 and temperature can bestablished as it is shown on Table 2. The results show that,lthough E1 was higher for the nanocomposites in the com-lete range of temperatures showing that C10A particles wereble to improve the elasticity of the neat matrix, the effect ofanoparticles on the stiffness (or elastic part) became lower as

emperature is raised. At lower temperatures, big changes in E1ere observed when organoclay was incorporated to starch/PCLlend, which in turns indicates that instantaneous elasticity mod-lus was improved by the presence of the nanofillers. The last

able 2elationship between E1 and the temperature for the matrix (MaterBi Z) and

heir nanocomposites with clay

lay content (wt.%) Modulus as a function of temperature r

.0 E (MPa) = 300.4 − 4.24 T(◦C) −0.980

.0 E (MPa) = 733.7 − 16.9 T(◦C) −0.996

.5 E (MPa) = 873.6 − 20.9 T(◦C) −0.993

.0 E (MPa) = 879.2 − 17.5 T(◦C) −0.977

.5 E (MPa) = 1222.1 − 23.6 T(◦C) −0.996

= regression coefficient.

ce and Engineering A 480 (2008) 259–265 263

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C.J. Perez et al. / Materials Scien

esult is in accordance with the increase in Young’s modulus ashe clay concentration increased previously reported [20]. Sim-lar trend on increasing E1 indicating the enhancement of theensile modulus was found by other authors [13,15]. It is impor-ant to note that all selected temperatures were higher to the glassransition temperature and, so that, the stiffness of the materials very low.

The retardant elasticity (E2), which is related to the stiffnessf amorphous polymer chains in the short time, decreased as aunction of temperature, as is shown on Fig. 6. This implies thathe deformation in the Kelvin unit turn into higher by the effect ofhe temperature. Another important effect was the reinforcementf the nanofillers on the Kelvin unit.

The other central parameter is η1 which represents therrecoverable creep strain. This parameter is plotted as a func-ion of temperature on Fig. 7. Cyras et al. [9] have shown forhe same matrix that the viscosity η1 increased with sisal fibreontent and lower flow was occurred at the dashpot and theermanent deformation decreased. Above the glass transitionemperature, as in this case, the amorphous chains have notewor-hy mobility that encourages irreversible creep deformations.ther authors [15] have pointed out that η1 could be relatedith the damage of crystalline polymer or oriented noncrys-

alline regions. This process includes chain’s pulling out, crystallipping and irreversible deformation from amorphous regions.n improved deformation resistance of nanocomposites, i.e.,

he decrease on the permanent creep deformation by nanofillersddition, should be associated with an alteration of the crys-alline region morphology of the neat matrix, as it was observedy the Avrami’s exponent of nanocomposites as compared withhe pure starch/PCL blend in a previous work [25] and also withhe immobilization of polymer chains previously establish inhe same work. Increasing temperature, a diminished η1 wasbserved. At higher temperatures, polymer chains were ther-

ally activated and large deformation of crystallized regions

s well as irreversible transitions of amorphous regions tooklace. Creep data provides information about zero-shear viscos-ty (associated with η1) which was higher for nanocomposites

ig. 6. E2 parameter of matrix (MaterBi Z) and their nanocomposites as aunction of temperature, for different clay content.

(twCtCppwgst

TRp

C

01257

ig. 7. η1 parameter of matrix (MaterBi Z) and their nanocomposites as aunction of temperature, for different clay content.

han for the neat matrix. There exist two possible explanationsor this behaviour: the larger barrier for flow of confined chainsr the frictional interactions between the anisotropic clay crys-allites.

As it is known, elastic materials exhibit no residual defor-ation after the stress is removed. Opposite, the viscoelastic

ne, display a relaxation time. This time is defined as the timeeeded to produce 63.2% of the total deformation in the Kelvinnit (1 − e−1). Table 3 shows the average relaxation time and thealue at room temperature (20 ◦C) for the matrix and nanocom-osites; η2 is also included in this table. Both parameters, τ and2 seems to decrease with clay incorporation and then, with clayontent. Neither the relaxation time nor the retardant viscosityisplayed a clear tendency with the temperature.

The power law (Findley) and the four parameters modelBurgers) were suitable to reproduce the strain data. Besidesheir parameters, another characteristic factor is the creep ratehich also determines the dimensional stability of materials.reep rates (directly fitted from the linear part of experimen-

al curves in the secondary creep stage) are shown on Fig. 8.reep rate constantly increased with temperature because moreolymer segments can be thermally activated at elevated tem-eratures but, nevertheless the temperature, creep rate decreasedith clay content indicating that the used organoclay is a

ood option to enhance the dimensional stability of the puretarch/polycaprolactone blend. Zhang et al. [26] have observedhe same tendency for PA66-TiO2 nanocomposites. The compar-

able 3elaxation times and η2 parameter of matrix (MaterBi Z) and their nanocom-osites with clay

lay content (wt.%) τaverage (s) τroom T (s) η2 (GPa s)

.0 1061 ± 200 1680 966

.0 1020 ± 300 1368 820

.5 1000 ± 79 1091 787

.0 722 ± 90 938 633

.5 674 ± 163 617 616

264 C.J. Perez et al. / Materials Science and Engineering A 480 (2008) 259–265

Fig. 8. Creep rates of matrix (MaterBi Z) and their nanocomposites as a functionof temperature, for different clay content.

Table 4Creep rates (experimental and models) of matrix (MaterBi Z) and their nanocom-posites with different content of clay

Clay content (wt.%) εexperimental (s−1) εB (× 10−6 s−1) εPL(× 10−7 s−1)

0.0 1.1 × 10−6 3.17 8.691.0 5.4 × 10−7 3.13 7.122.5 8.8 × 10−7 2.98 7.0157

imgtε

ce(n

wtgsut

TApo

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Fig. 9. (a) Master curves of matrix (MaterBi Z) and their nanocompositesobtained by using the time–temperature superposition principle and (b) com-pn

ufmt

.0 7.2 × 10−7 2.24 6.91

.5 6.5 × 10−7 1.25 5.48

son of experimental values and thus predicted from the two usedodels (at room temperature) are summarized on Table 4. As a

eneral rule, εB was higher than εPL for all kind of materials andhe experimental values were closer to power law values, εPL.˙B was higher than that directly obtained from the strain–timeurve indicating that this model was not able to reproduce thexperimental data of all studied materials but the power lawFindley’s model) was applicable to materials that exhibited aon-pronounced second creep stage [15].

The creep strains measured at different temperaturesere superposed rearranging the time scale by using the

ime–temperature superposition principle. The activation ener-ies are given on Table 5 and the obtained master curves are

hown on Fig. 9a. Master curves were obtained shifting the mod-lus/compliance data at different temperature, taking 20 ◦C ashe reference temperature because the materials are commonly

able 5ctivation energies for creep process and Findley’s power law parameters (com-liance) of matrix (MaterBi Z) and their nanocomposites with different contentf clay

lay content (wt.%) Q (kJ/mol) J0 (GPa−1) J1 ( × 10−4 GPa−1)

.0 133.2 2.31 × 10−3 8.66

.0 189.7 1.05 × 10−3 7.08

.5 256.1 7.45 × 10−4 6.16

.0 236.8 5.94 × 10−4 6.03

.5 211.1 5.56 × 10−4 4.09

adtEviteAclwwTc

liance as a function of corrected time for matrix (MaterBi Z) and theiranocomposites, for different clay content.

sed at this temperature. The shape of the curves was similaror un-reinforced and reinforced materials. The quality of theaster curve superposition is acceptable whereas the shift fac-

or exhibited a good linear dependence with the temperaturend the activation energy values showed that nanocompositesisplayed higher creep resistance under the effect of tempera-ure; i.e., that the nanofillers immobilize the polymer chains.q. (5) was used to model the complete curve. An averagealue of 0.18 ± 0.03 was obtained for n. The values of the time-ndependent compliance and the coefficient of time-dependenterm are summarized on Table 5 and the correlation betweenxperimental data and model prediction are shown on Fig. 9b.s expected, both parameters, J0 and J1, decreased with clay

ontent showing the enhancement on the creep stiffness. Theower initial creep compliance is related to the higher modulus

ith increasing layered-silicate concentration. Similar resultsere observed by Vlasveld et al. [10] for PA6 nanocomposites.he rise of compliance with the time was clearly diminished bylay incorporation. In addition, from the last graph, it is quite

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C.J. Perez et al. / Materials Scien

vident that the power law correctly predicts the complianceehaviour of the matrix and nanocomposites for long times.

. Conclusions

The creep behaviour of polycaprolactone/starch blends andts nanocomposites with organo-modified layered silicates wasnalyzed.

The creep compliance was decreased by the incorporationf nanofillers. This effect was related with the enhancement ofhe modulus and it was related with the mechanical propertiestudied in a previous work. In addition, the reduction in thereep strain due to clay addition drives to a material with higherimensional stability. Other creep parameters, such as creep rate,lso demonstrated this tendency.

The parameters of Burgers model also confirmed this trend:he instantaneous and retardant modulus increased and the per-

anent viscosity and relaxation time decreased as a functionf clay content. The effect of temperature on such parame-ers was established showing the lower stiffness and the highereformability of matrix and nanocomposites as a function of it.

The efficiency of nanofillers to get better the creep per-ormance of neat matrix was confirmed by the parametersf Findley’s power law which properly represent the realtrain–time curves of matrix and nanocomposites in the selectedange of temperatures.

The master curves displayed a tolerable superposition andhe activation energy for the shift factor was also an evidenceor the creep resistance development.

All the results discussed in the present work indicate thatanofillers contributed to an improvement on the creep resis-ance which is an important result for the application point ofiew.

It is important to point out that, for the range of clay con-ent studied here, the creep resistance is continuously improvedogether with the increase of the clay concentration.

cknowledgements

Authors acknowledged to CONICET (PIP 6254) andNPCyT, Project FONCyT-PICT No. 12-15074 and PICT 22-5766 for the financial support of this project.

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d Engineering A 480 (2008) 259–265 265

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