Simulation of mechanical properties of multilayered propylene–ethylene copolymer/ethylene 1-octene copolymer composites by equivalent box model and its experimental verification Jiabin Shen, Ming Wang, Jiang Li * , Shaoyun Guo * , Shuangxi Xu, Yuqing Zhang, Ting Li, Ming Wen The State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute of Sichuan University, Chengdu, Sichuan 610065, China article info Article history: Received 22 April 2009 Received in revised form 20 July 2009 Accepted 21 July 2009 Available online 28 July 2009 Keywords: Multilayered coextrusion Morphology Mechanical properties Equivalent box model abstract Multilayered propylene–ethylene copolymer (PPE)/ethylene 1-octene copolymer (POE) composites were prepared by a microlayered coextrusion system. Static and dynamic ten- sile results showed that yield strengths and storage moduli of the multilayered samples were distinctly larger than those of conventional blends. The equivalent box model (EBM) proposed by Kolarik was used to explain the effect of morphology on mechanical properties. Experimental and theoretical results indicated that the excellent mechanical properties of multilayered composites were ascribed to the phase continuity. The effects of interfacial layers in multilayered composites on mechanical properties were also dis- cussed. The existence of interfaces between POE and PPE layers not only led to the inval- idation of EBM for prediction of mechanical properties of multilayered composites with larger number of layers, but also induced a new absorbing peak in loss modulus-tempera- ture spectrum because of the shearing friction between POE and PPE layers through interfaces. Ó 2009 Elsevier Ltd. All rights reserved. 1. Introduction Most of the polymer blends were found to be immisci- ble and exhibited many kinds of heterogeneous morpholo- gies, which could be divided into three classes, i.e., dispersed, co-continuous and stratified morphologies [1]. The stratified structure was regarded as a special co-con- tinuous morphology, where each layer was parallel and continuous along a direction. Comparing with the dis- persed and co-continuous morphologies, the stratified morphology was not easily obtained by the conventional method under the limits of the composition, viscosity and processing conditions. Because of the extensive applications of polymer blends, the study on the prediction of their mechanical properties was considered to be necessary [2–7]. Obvi- ously, the rule of mixture, which meant that properties of homogeneous systems were the average of their ingredi- ents, was not available for the prediction of mechanical properties of heterogeneous blends. Their mechanical properties were dependent not only on the corresponding properties and the composition of individual phases, but also mostly on the spatial organization of each phase and the nature of interfaces. A common example was that the addition of compatibilizer into a sea-island system led to the increases of yield and impact strengths because of the decreases of size of dispersed phase and interfacial ten- sion [8]. Usually, an important aspect of spatial organiza- tion encompassed the size, shape, and distribution of dispersed phases and the coarseness and tortuosity of a continuous phase [9]. Many investigators [10–12] focused on the effective factors on mechanical properties and found that, in a droplet–matrix structure, the tensile mod- ulus was largely determined by that of the matrix phase. However, the modulus of the fibrous blend was deter- mined by that of the dispersed phase, and in the case of 0014-3057/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.eurpolymj.2009.07.013 * Corresponding authors. Tel./fax: +86 28 85405135. E-mail addresses: [email protected](J. Li), [email protected](S. Guo). European Polymer Journal 45 (2009) 3269–3281 Contents lists available at ScienceDirect European Polymer Journal journal homepage: www.elsevier.com/locate/europolj
Simulation of mechanical properties of multilayered propylene–ethylenecopolymer/ethylene 1-octene copolymer composites by equivalent boxmodel and its experimental verification
Jiabin Shen, Ming Wang, Jiang Li *, Shaoyun Guo *, Shuangxi Xu, Yuqing Zhang, Ting Li, Ming WenThe State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute of Sichuan University, Chengdu, Sichuan 610065, China
a r t i c l e i n f o
Article history:Received 22 April 2009Received in revised form 20 July 2009Accepted 21 July 2009Available online 28 July 2009
Keywords:Multilayered coextrusionMorphologyMechanical propertiesEquivalent box model
0014-3057/$ - see front matter � 2009 Elsevier Ltddoi:10.1016/j.eurpolymj.2009.07.013
Multilayered propylene–ethylene copolymer (PPE)/ethylene 1-octene copolymer (POE)composites were prepared by a microlayered coextrusion system. Static and dynamic ten-sile results showed that yield strengths and storage moduli of the multilayered sampleswere distinctly larger than those of conventional blends. The equivalent box model(EBM) proposed by Kolarik was used to explain the effect of morphology on mechanicalproperties. Experimental and theoretical results indicated that the excellent mechanicalproperties of multilayered composites were ascribed to the phase continuity. The effectsof interfacial layers in multilayered composites on mechanical properties were also dis-cussed. The existence of interfaces between POE and PPE layers not only led to the inval-idation of EBM for prediction of mechanical properties of multilayered composites withlarger number of layers, but also induced a new absorbing peak in loss modulus-tempera-ture spectrum because of the shearing friction between POE and PPE layers throughinterfaces.
� 2009 Elsevier Ltd. All rights reserved.
1. Introduction
Most of the polymer blends were found to be immisci-ble and exhibited many kinds of heterogeneous morpholo-gies, which could be divided into three classes, i.e.,dispersed, co-continuous and stratified morphologies [1].The stratified structure was regarded as a special co-con-tinuous morphology, where each layer was parallel andcontinuous along a direction. Comparing with the dis-persed and co-continuous morphologies, the stratifiedmorphology was not easily obtained by the conventionalmethod under the limits of the composition, viscosityand processing conditions.
Because of the extensive applications of polymerblends, the study on the prediction of their mechanicalproperties was considered to be necessary [2–7]. Obvi-
ously, the rule of mixture, which meant that properties ofhomogeneous systems were the average of their ingredi-ents, was not available for the prediction of mechanicalproperties of heterogeneous blends. Their mechanicalproperties were dependent not only on the correspondingproperties and the composition of individual phases, butalso mostly on the spatial organization of each phase andthe nature of interfaces. A common example was that theaddition of compatibilizer into a sea-island system led tothe increases of yield and impact strengths because ofthe decreases of size of dispersed phase and interfacial ten-sion [8]. Usually, an important aspect of spatial organiza-tion encompassed the size, shape, and distribution ofdispersed phases and the coarseness and tortuosity of acontinuous phase [9]. Many investigators [10–12] focusedon the effective factors on mechanical properties andfound that, in a droplet–matrix structure, the tensile mod-ulus was largely determined by that of the matrix phase.However, the modulus of the fibrous blend was deter-mined by that of the dispersed phase, and in the case of
Fig. 1. Schematic of microlayer coextrusion system: A, B-single screwextruder; C-connector; D-coextrusion block; E-layer multiplying ele-ments; F-die.
3270 J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281
co-continuous blends, both components could fully con-tribute to the modulus of the blend. On the other hand,the interfacial adhesion between phases also played animportant role in governing the mechanical properties ofpolymer blends, including the transfer of external force,the interfacial friction between phases and so on [13,14].
As described above, mechanical properties of polymerblends were strongly dependent on the composition andmorphology, which could be described by mass (or vol-ume) fraction and geometrical parameters, respectively.Based on these concepts, many theoretical models, suchas the equivalent box, Halpin–Tsai, Takayanagi, Daviesand Kerner models, etc. have been proposed to predictthe mechanical properties of polymer blends [2,9,15–17].In these theoretical models, the equivalent box model, pro-posed by Kolarik, got widely used due to its simple expres-sion. In this model, two immiscible phases were dividedinto parallel and series branches. Moreover, this modelgave the upper bound of the mechanical properties deter-mined by parallel model, which corresponded to the strat-ified structure, and the lower limit determined by seriesmodel. It indicated that the mechanical properties of dis-persed, co-continuous structure would be lower than thoseof stratified structure. Kolarik [13] ascribed this enhance-ment of mechanical properties of stratified morphologyto its best phase continuity in the direction of the actingforce, where the interfaces could not affect the transfer ofthe stress. Previously, Boyd has developed a similar modelto the equivalent box model to study the mechanical re-sponse of laminar semi-crystalline polymers in terms ofseparated crystalline- and amorphous-phase properties[18]. He considered that the mechanical response can beobtained by incorporating uniform in-plain strain andadditive forces from layer to layer of crystalline and amor-phous-phases and uniform stress and additive displace-ments normal to the lamellar surfaces.
To the best of our knowledge, the equivalent box modelwas mostly applied in sea-island and co-continuous sys-tems [15]. In despite of the potential advantage of themechanical properties of stratified materials predicted bythe equivalent box model, their experimental verificationwas hardly involved in actual investigations because itwas difficult to obtain through the conventional processingmethod. Moreover, some basic scientific studies on thevalidity of EBM, for example the effect of interfacial layerson mechanical properties, could not be carried out due tothe absence of laminar model materials with large numberof layers.
The conventional coextrusion was a valid method toconstruct the laminar polymer materials, but only sev-eral-layer composites were obtained. In recent years, anadvanced processing technology, multilayered (or micro-layered) coextrusion, has been developed to prepare thestratified composite sheets with alternating polymer lay-ers, where each layer was continuous along the extrusiondirection, [19] and the number of layers could amount tothousands with the increase of number of layer multiply-ing elements. Some studies showed that the multilayeredstructure endowed the blends (composites) with the out-standing barrier [20,21], optical [22,23], electrical [24],especially mechanical [25–28] properties. In addition,
these multilayered composites could also be treated aslaminar model materials for the experiment verificationof EBM and studies of layered interfaces.
In this paper, the multilayered polypropylene–ethylenecopolymer/ethylene 1-octene copolymer (PPE/POE) com-posites with alternating PPE and POE layers were preparedby multilayered coextrusion technology developed in ourlab. The equivalent box model was introduced to predictand explain their yield strengths and dynamical mechani-cal properties.
2. Experimental
2.1. Materials
The polymers used in this work were PPE and POE. ThePPE was EPS30R from Qi Lu Petroleum Chemical Company,Ltd. (China) consisting of 6.5 wt.% ethylene; the POE elasto-mer was Engage 8150 from Dupont-Dow Chemical consist-ing of 25 wt.% octene. The densities of them were0.897 � 103 and 0.868 � 103 kg/m3, respectively.
2.2. Specimen preparation
PPE and POE were, respectively, dried in an oven at 85 �Cand 50 �C for 12 h prior to the processing. Each stratifiedsample was coextruded as a sheet about 1.6 mm thick and40 mm wide using the multilayered coextrusion system de-signed by our lab, the schematic of which was illustrated inFig. 1. Two polymers were simultaneously extruded fromdifferent extruders, and combined as 2-layer melt in thecoextrusion block, then flowed through a series of layermultiplying elements (LMEs). In a LME, the melt was slicedinto two left and right sections by a divider, and then recom-bined vertically as shown in Fig. 2. An assembly of n LMEscould produce a laminar composite with 2(n+1) layers. In thiswork, 2-, 16-, 32-, 64- and 128-layer samples were extrudedwith 0, 3, 4, 5 and 6 LMEs, respectively. The ratio of two-components was varied by adjusting the screw speed ofextruders. The temperatures of extruders for PPE and POEwere 180 �C and 200 �C, respectively, and the temperature
Fig. 2. Schematic of multiplying elements.
Table 1The content of POE in blended and laminated samples.
J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281 3271
of LMEs was 190 �C so that viscosities of PPE and POE werematched as they entered into LMEs.
For comparison, conventional blended PPE/POE sampleswere prepared as sheets with a similar dimension to strat-ified ones. Firstly, the dried PPE and POE were mixed in ahigh-speed mixer, and then extruded using an extruderof multilayered coextrusion system in order to keep thesame heat history as corresponding stratified composites.The temperature at the die was 190 �C. The neat PPE andPOE sheets were also prepared, respectively, by an extru-der, with the die temperature of 190 �C. The LMEs werenot used for producing these samples.
2.3. Density measurement
The densities of samples were tested in a densitometer,the measuring range of which was 0.800–0.900 g/cm3, tocalculate the volume fraction of PPE or POE according tothe additive principle. Firstly, the specimen was immersedinto water, and then the density of the solvent was ad-justed by adding gradually the isopropanol into water untilthe sample suspended in it. At this time, the density ofsample was equal to that of mixture solvent.
The volume fraction of POE (VPOE) could be calculatedfrom the following formula:
VPOE ¼qC � qPPE
qPOE � qPPE
where, qC is the density of the composite; qPPE and qPOE
represent the densities of PPE and POE, values of whichare 0.897 � 103 and 0.868 � 103kg/m3, respectively.
The sample codes and their volume fraction of POE arelisted in Table 1.
2.4. Tensile tests
Tensile tests were performed using an Instron 4302 ten-sion machine (Canton, MA, USA) at a strain rate of 0.02, 0.1,0.2, 1, 2, 20 min�1 with a experimental temperature of23 �C in accordance with ASTM D638. The yield strengthswere determined from stress–strain curves. At least of fivespecimens for each sample were tested and the average va-lue was calculated.
2.5. Differential Scanning Calorimetry (DSC)
The melting points and the enthalpy values of thecomposites were investigated using DSC machine (NetzschDSC 204 made in Germany). The samples were heated from20 to 200 �C at a rate of 10 �C/min under a nitrogenatmosphere.
2.6. Scanning Electron Microscope (SEM)
In order to study the interfacial interaction between PPEand POE phases during tensile process, the neat PPE, 2-, 64-and 128-layer samples were stretched to 200% at the strainrate of 2 min�1, then unloaded the force and relaxed at23 �C for 24 h. All these samples were cryofractured inliquid nitrogen and SEM was performed in the directionperpendicular to the application of the tensile force on aJSM-5900LV (Japan).
2.7. Dynamic Mechanical Analysis (DMA)
The dynamic mechanical properties of PPE/POE blendsand 2-, 32-, 64-layer samples were tested by DMA (Q800,TA Instrument) at a heating rate of 3 �C/min from �80 to60 �C using a tensile model. The tested frequency was 10 Hz.
3. Equivalent box model (EBM)
As proposed in many literatures, mechanical propertiesof polymer blends were strongly dependent on the compo-
3272 J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281
sition and morphology and could be predicted by mathe-matic methods. At present, there existed several models[2,9,15–17] describing the elastic behavior of two-compo-nent systems. However, because of the manifold morphol-ogies, most theoretical models were applicable only withinlimited ranges. J. Kolarik proposed a predictive schemewhich was based on the combination of a two-parametermechanically equivalent box model (EBM) and the conceptof phase continuity [13,15].
EBM was improved from Takayanagi model and was ad-vanced for analyzing the contribution of each componentto the mechanical properties of multiphase composites. Itwas generally considered that the upper and lower boundsof mechanical properties of isotropic heterogeneous mate-rials could be represented by the parallel and series mod-els, respectively, [9]. In the parallel model, as all phaseswere continuous in the direction of the acting force, i.e.,the force did not cross any interface, the mechanical prop-erties could be fit by additive principle. In the series model,all components were discontinuous in the direction of theacting force, so that the mechanical properties were dom-inated by the component with the smallest mechanicalproperties and the interfacial adhesion between phases,and could be described by anti-additive principle. There-fore, in the case of isotropic heterogeneous materials, itwas considered that the mechanical properties could bepredicted by combining the parallel and series models. Amodel of the EBM was shown in Fig. 3, where the sub-scripts 1 and 2 indicate the phase A and phase B; the sub-scripts p and s denote parallel and series branch; the Vp andVs represent the volume fraction of parallel and seriesbranch in the composite; the Vp,1 and Vp,2 are the volumefraction of phase A and phase B in the parallel branch;the Vs,1 and Vs,2 are the volume fraction of phase A andphase B in the series branch.
The relation of these volume fractions could be ex-pressed as follows:
Vp;1 þ Vp;2 ¼ Vp Vs;1 þ Vs;2 ¼ Vs
Vp;1 þ Vs;1 ¼ V1 Vp;2 þ Vs;2 ¼ V2
V1 þ V2 ¼ Vp þ Vs ¼ 1ð1Þ
Fig. 3. The schematic of equivalent box model: Vp,1, Vp,2-volume fractionof phase A and phase B in the parallel branch; Vs,1, Vs,2-volume fraction ofphase A and phase B in the series branch.
Therefore, the mechanical properties of isotropic heter-ogeneous materials, Mc, could be expressed by,
Mc ¼ MpVp þMsVs
¼ M1Vp;1 þM2Vp;2� �
þ V2s
Vp;1
M1þ Vp;2
M2
� ��ð2Þ
when Vp = 1, Vs = 0 (parallel model, additive principle),
Mc ¼ Mp ¼ M1V1 þM2V2 ð3Þ
when Vp = 0, Vs = 1 (series model, anti-additive principle),
Mc ¼ Ms ¼ 1V1
M1þ V2
M2
� ��ð4Þ
where, Mp and Ms denote the mechanical properties in par-allel and series branches, respectively; M1 and M2 repre-sent the mechanical properties of phase A and phase B.Obviously, Mc in Eq. (2) is intervenient between the valuescalculated by additive principle and anti-additive principle.
Kolarik considered that if the mechanical properties ofphase A were smaller than those of phase B (i.e.,M1 < M2), the Eq. (2) could be simplified by:
Mc ¼ ðM1Vp;1 þM2Vp;2Þ þ A �M1 � Vs ð5Þ
where, A is the interfacial adhesion between phase A andphase B. If the adhesion was too weak (A = 0), the seriesbranch could not contribute to the mechanical propertiesof composites. If the adhesion was strong enough (A = 1),the mechanical properties of composites were contributedby both of the parallel and the series branches. In addition,the percolation theory [29] was also introduced to calcu-lated Vp,1 and Vp,2 as described by:
VP;1 ¼V1 � V1;cr
1� V1;cr
� �T1
VP;2 ¼V2 � V2;cr
1� V2;cr
� �T2
ð6Þ
where, Vi,cr is the critical volume fraction (the percolationthreshold); Ti is the critical universal exponent, the valuesof which are 0.16 and 1.8 for a three-dimensional lattice[13,15].
On the other hand, in order to explain the equivalentmechanical model clearly, the concept of phase continuity,C, was also introduced to describe the level of phase con-nectivity in the direction of the acting force. A larger phasecontinuity meant that there were more parallel couplingcontributing to the mechanical properties of compositesand advantageous for the stress transfer. In two-compo-nent systems, C could be acquired by,
C1 ¼ VP;1=V1 ¼V1 � V1;cr
1� V1;cr
� �T1,
V1
C2 ¼ VP;2=V2 ¼V2 � V2;cr
1� V2;cr
� �T21,
V2 ð7Þ
It was indicated that the phase continuity was related tothe proportion of parallel branch in each phase. In the par-allel and series models, the phase continuity was consid-ered equal to unity and zero, respectively.
Thus, according to EBM, it could be considered that themechanical properties of an isotropic heterogeneous mate-rial were mainly related not only to the mechanical
Fig. 4. SEM micrographs of PPE/POE conventional blend and 64-layer samples.
J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281 3273
properties of each component, but also to the level ofphase continuity in the direction of the acting force.
4. Results and discussion
4.1. Phase morphology
Fig. 4 showed SEM micrographs of conventional blendsand 64-layer composites with a POE content of 13 wt.%. Itwas illustrated that the blends had typical sea-island mor-phology and the POE phases were evenly distributed in thePPE matrix. On the other hand, the multilayered samplehad a laminar morphology as described by the parallelmodel, where alternating continuous PPE and POE layersaligned parallel to the extrusion direction. According tothe mechanical models proposed above, it was assumed
00.0
0
10
20
30
Yie
ld s
tren
gth,
σC (
MP
a)
Volume fractio
σC=28.0*[(V
PPE- 0.16)/(1 - 0.16)]1.8
Fig. 5. The effect of volume fraction of POE on the yield strengths of compositesadditive principle and anti-additive principle; solid line (—) was calculated from
that the phase continuity of the latter should be larger thanthe former and might approach to unity. In order to furtherstudy the effect of phase morphologies on the mechanicalproperties, results of tensile tests were discussed in thenext part.
4.2. Static tensile properties
Fig. 5 showed the effect of volume fraction of POE (VPOE)on the yield strengths of the composites. Apparently, theyield strengths of laminar composites were larger thanthat of conventional blends at the same content of POE,which was consistent with the theoretical prediction de-scribed above. As demonstrated in Fig. 5, the data ofblended samples were intervenient between upper andlower bounds decided by additive and anti-additive princi-
1.0.5
n of POE, VPOE
Blended samples
2-layer samples
16-layer samples
64-layer samples
: dash line (– –) and short dash line (- - -) were theoretical prediction byEqs. (5) and (6).
100 120 140 160 180
Hea
t flo
w, E
ndo
Temperature
B1
B3B5B6
16-116-3
128-1
128-3
128-4
64-3
64-1
Fig. 6. Comparison of DSC curves of blended and multilayered PPE/POE composites at a heating rate of 10 �C/min.
Table 2DSC results of blended and multilayered PPE/POE composites.
3274 J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281
ple (Eqs. (3) and (4)), respectively, indicating that theiryield strengths could be simulated by combining the paral-lel and series models. According to the improved simulat-ing Eqs. (5) and (6) by Kolarik, the correlation of the yieldstrength and VPPE was as follow:
rC ¼ 28:0 � VPPE � 0:161� 0:16
� �1:8
ð8Þ
where, the contribution of POE was not taken into consid-eration due to a negligible yield strength comparing withthat of PPE. Eq. (8) could describe the experimental dataof the conventional blends well, except for the date witha POE content of about 52 vol.%, which deviated from thetheoretic curve and moved close to the upper bound. Itwas attributed to the increase of parallel branch in the con-ventional blended PPE/POE (48/52, vol/vol), resulting fromthe change of phase morphology with the increasing POEcontent.
For the stratified composites, almost all the data distrib-uted around the upper bound. The experimental accor-dance with theoretic prediction depicted the parallelmodel was a valid tool to simulate the yield strengths ofmultilayered composites till the number of layers in-creased to 64.
According to EBM, the yield strength of a multilayeredmaterial was determined by the yield behavior and the le-vel of phase continuity of each component, mainly higheryield strength and content component. In order to under-stand intensively the nature of the increase in yieldstrengths of multilayered composites, the yield behaviorand continuity of PPE layer were, respectively, studied.
A lot of reporters indicated that the yield behavior ofsemi-crystalline polymers was induced by the slip of crys-talline blocks each other, which was related to the degreeof crystallinity [30,31], lamellae thickness [20,32], compo-sition distribution [33], tensile rate [34], etc. It was gener-ally considered that the crystallinity and lamellae
thickness was related to the enthalpy value and meltingpoint, respectively, [35]. So, the effects of morphologyand composition on these values were investigated in ourwork. The DSC results shown in Fig. 6 and Table 2 indicatedthat there was little change of melting point and enthalpywith the addition of different POE content and the varia-tion of phase morphologies, which was consistent withthe study of Yamaguchi et al. [36]. It was suggested thatthere might be a similar effect of crystalline region on yieldstrengths of blended and multilayered samples because oftheir same crystalline properties. Therefore, the yieldbehavior and mechanism of PPE phase might not be influ-enced by the different phase morphologies.
Further more, we have studied the effect of strain rates(c) on the yield strengths (ry) of blended and stratifiedsamples. As seen in Fig. 7, there was a similar trend ofthe linear dependence of ry upon the log10(c) for all sam-ples. According to Eyring’s theory [37], no change occurredin yielding mechanism at the testing strain rates. There-fore, it could be deduced that, in the range of these strainrates, the distinctions of the yield strengths betweenblended and multilayer samples might be resulted fromother reasons.
1010.10
10
20
30
Blended samples
64-layer samplesY
ield
str
engt
h, σ
C (
MP
a)
Strain rate, ε (min-1)
Pure PPE
2-layer samples
Fig. 7. Effect of strain rate on the yield strengths of PPE, blended and laminating samples with 55% POE.
J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281 3275
As discussed above, the yield behavior of the compos-ites was also related to the level of phase continuity inthe direction of the acting force. In the case of conventionalPPE/POE blends, POE domains might primarily surroundthe crystalline regions of PPE rather than dispersed in thedefective regions between crystalline blocks, as assumedby Plaza et al. [38]. Therefore, the addition of elastomericcomponent might reduce the capability to transmit the ap-plied stress to the crystalline regions, which might affectthe slip of crystalline blocks each other in PPE. As for themultilayered samples, PPE and POE components were
0.00 0.250.0
0.5
1.0
1.5
2.0
Con
tinui
ty d
egre
e of
PP
E, C
PP
E
Volume fractio
CPPE
=[(VPPE
-0.16)/(1-0.16)]1.8/VP
CPPE
=1
Fig. 8. The effect of volume fraction of PO
alternately aligned and continuous along the direction ofextrusion. These two-components might not disrupt oneanother without concerning the interfacial interdiffusionand processing limitation, so that the load was able to betransmitted along each layer without being baffled.According to Eq. (7), the continuity of PPE (CPPE) of multi-layered and conventionally blended PPE/POE with differentPOE content could be quantitatively calculated. Fig. 8 illus-trated that in the case of blended samples, the values ofCPPE decrease gradually with increasing the volume frac-tion of POE (VPOE), indicating that the addition of elasto-
Fig. 9. The effect of volume fraction of POE on the yield strengths of 128-layer samples: dash line (– –) and short dash line (- - -) are theoretical predictionby additive principle and anti-additive principle.
3276 J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281
meric component disrupted the continuity of PPE in thedirection of the acting force and the yield behavior wouldbe influenced by the interfacial adhesion between phases.However, in the case of multilayered samples, the valuesof CPPE were generally distributed around unity irrespec-tive of the VPOE. Thus, the higher CPPE of multilayered com-posites compared with the conventional blends was anessential reason causing a higher yield strength.
On the base of above studies, we increased the numberof layers of multilayered PPE/POE and studied the effect ofVPOE on the yield strengths of 128-layer samples. Thecurves presented in Fig. 9 revealed that all experimentaldata of 128-layer composites were intervenient betweenthe predictive lines acquired by additive and anti-additiveprinciple, respectively, like the conventional blends asshown in Fig. 5. It indicated the existence of series branchin 128-layer composites. We assumed that the discrepancyof 128-layer and lower-layer composites might be relatedto the interfacial layers between PPE and POE layers. ThePOE was partly compatible with the PPE, because the 1-oc-tene in the former could dissolve in the amorphous regionof the latter, as studied by Yamaguchi [39]. This led to theexistence of interfacial layers in multilayered composites,where the chains of PPE and POE were tangled each other.Therefore, these interfacial layers could be treated as theconventional blends including the series structure. Clearly,when the layer number was lower, the contribution of theinterfacial layers to the yield strength could be ignored, sothat the yield strength could be calculated according toparallel model (Eq. (3)). But for 128-layer composites, theseries branch could not be ignored due to the more inter-facial layers. So, their yield strengths must be simulatedby Eq. (5) and was reasonably lower than that obtainedby parallel model.
The existence of interfacial layers could be proved byexamining the tensile-induced morphologies of samples.
Fig. 10 showed SEM micrographs of neat PPE, 2-, 64- and128-layer composites which were subjected to a deforma-tion of 200% and then relaxed. As shown in Fig. 10a, the ori-entation of pure PPE was quite conspicuous because of thegood plastic property of PPE. Fig. 10b showed the morphol-ogy of PPE layer in 2-layer composites was similar to thatof pure PPE. This indicated that high elastic POE couldnot compel PPE to relax because of only one interfaciallayer in 2-layer composites. Fig. 10c and d showed that,with the increase of the number of layers and interfaces,POE caused the deformation of PPE layer to decreasethrough the interfacial layers. When the layer number in-creased to 128, with 127 interfacial layers, the deformationof PPE layers almost totally relaxed.
4.3. Dynamic tensile properties
The distinctive layered structure and interface of multi-layered composites also led to some different dynamicmechanical behaviors from the conventional blends.Fig. 11a described the temperature dependence of storagemoduli (E0) of conventional blended and 64-layer sampleswith about 13% POE. According to the prediction of EBMdiscussed above, the E0 of multilayered PPE/POE were rea-sonably higher than that of conventional blends with thesame volume fraction of POE when the temperature waslower than the glass transition temperature (Tg) of PPE,i.e., 15 �C. In the conventional blends, when alternatingstress acted on the discontinuous phases, the modulusmight be greatly influenced by the component in the seriesstructure which could be deformed more easily and theinterfacial adhesion between phases. Therefore, the exis-tence of lower-modulus POE phase in the series structurecaused a lower-modulus of conventional PPE/POE blends.However, the PPE and POE layers of 64-layer compositesgave simultaneously the response to the outer oscillating
Fig. 10. SEM micrographs of liquid N2 fractured surface of vertical section for PPE, 2-, 64- and 128-layer samples.
J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281 3277
action due to their parallel structure. When flexible POEdeformed, there still were rigid continuous PPE phases tomaintain a higher modulus. Once the temperature washigher than Tg of PPE, meaning the defrosting of PPE chainsegments, the E0 of PPE decreased greatly. This caused thesame E0 of multilayered composites as that of conventionalblends.
But there was some difference between PPE/POE sam-ples with POE content of 13 and 55% comparing Fig. 11aand b. It could be found that, when the temperature waslower than �55 �C, E0 values of samples with higher POEcontent were almost same. It might be attributed to thehigher POE content, which led to the phase transition fromsea-island to co-continuous morphology and then themore parallel branch, and the higher E0 of POE with theglass state.
Fig. 12 gave the comparison between experimentallymeasured E0 values of 2- and 64-layer composites and the-oretically calculated E0 values by additive principle. It couldbe found from Fig. 12a that, the experimental and theoret-ical curves of 2-layered composites with different POE con-tents almost overlapped at any temperature, which wasconsistent with the results acquired by Y.J. Kim [40]. Itindicated the validity of parallel model in the case of low-er-layer composites. But for 64-layer composites, Fig. 12bshowed that the storage modulus accorded with the addi-tive principle until �35 �C, close to the glass transitiontemperature of POE. In this temperature range, both chainsegments of PPE and POE were nearly frozen and their
strains could quickly recover after the discharge of outerstress. Accordingly, their tensile behaviors could be de-scribed by Hooke’s law, so that the additive principle wasreasonably fit for predicting the storage modulus in this re-gion. However, once the temperature was higher than�35 �C, it was found that the experimental data did notagree with the theoretical curve. We considered that itmight be related to the interfacial layers between PPEand POE layers. As discussed in 4.2, flexible POE layerscould compel PPE layers to relax through the interfaciallayers. Therefore, the movement ability of PPE moleculeswas enhanced by defrosted POE molecules. This caused adecrease of storage modulus of 64-layer composites com-pared with the theoretical value, which did not take theinteraction at interfaces into consideration.
The loss modulus (E00), resulted from the energy lost tofriction and internal motions in macromolecules during asinusoidal oscillation, was regarded as a characteristic ofviscous behavior, and the temperature corresponding tothe maximum value of E00 in DMA spectrum could be de-fined as the glass transition temperature (Tg) [41]. In con-ventional PPE/POE blended system, there generally weretwo separated loss peaks around the Tg of each phase, asshown in Fig. 13a, where the peak at about �40 �C corre-sponded to the Tg of POE (Tg, POE) and that in the vicinityof 16 �C corresponded to the Tg of PPE (Tg, PPE). Fig. 13bshowed the temperature dependence of loss moduli ofmultilayered samples. 2-layer composites showed twoabsorbing peaks of Tg, POE and Tg, PPE, like the conventional
0
1000
2000
3000
4000
5000
Sto
rage
Mod
ulus
, E' (
MP
a)
Temperature
Neat PPE Blended sample 64-layer sample
Tg,PPE
(a) 13% POE
-90 -60 -30 0 30 60
-90 -60 -30 0 30 60
0
500
1000
1500
2000
2500
Neat POE Blended sample 64-layer sample
Sto
rage
Mod
ulus
, E' (
MP
a)
Temperature
Tg,POE
(b) 55% POE
Fig. 11. Comparison of storage moduli between blended and 64-layer samples.
3278 J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281
blends. However, with the increase of number of layers, thethird peak appeared in DMA spectrum of 32-, 64- and 128-layer samples. This new peak was located between Tg, POEand Tg, PPE, as indicated by arrows in Fig. 13b. This mightbe related to the plane interfaces between POE and PPEphases parallel to the tensile direction, different from thespherical those in the conventional blends. Some research-ers [2] studied the effects of fillers on dynamic mechanicalproperties of polymers and found that laminar fillers suchas graphite and mica were effective in increasing thedamping of some polymers. They ascribed this increaseto the friction between thin laminas or sheets of the filler
as they slipped over one another with the deformation ofpolymeric composites. In the 32-, 64- and 128-layer PPE/POE composites, there were some multilayered interfaceslike laminar fillers. According to the conclusions of refer-ence [38], we considered that the peculiar intermediatepeak might be caused by the shearing friction betweenPOE and PPE layers through interfacial layers, due to theirdifferent movement abilities at the temperatures rangedfrom Tg, POE to Tg, PPE. In addition, it was found that whenthe number of layers increased to 128, the intermediatepeak was more obvious, which might also support theassumption described above.
10
100
1000
Sto
rage
mod
ulus
, E' (
MP
a)
Temperature
neat PPE
2-3
2-5
neat POE
(a)
-80 -60 -40 -20 0 20 40 60
-75 -50 -25 0 25 50
10
100
1000
Sto
rage
mod
ulus
, E' (
MP
a)
Temperature
neat PPE 64-1 sample 64-3 sample 64-4 sample neat POE
(b)
Fig. 12. Temperature dependence of storage moduli of multilayered samples: (a) solid line (—) and dash line (– –) were theoretical curves of E0 of 2–3 and2–5 samples by additive principle, respectively. (b) Solid line (—), dash line (– –), and short dash line (- - -) were theoretical curves of E0 of 64–1, 64–3, and64–4 samples by additive principle, respectively.
J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281 3279
4.4. Brittleness
In this study, static and dynamic tensile tests of PPE/POE multilayered materials have been performed, respec-tively. Hence, the brittleness of polymeric materials (B)might be calculated by combining the strain at break (eb)obtained from stress–strain curves with the storage modu-lus (E0) obtained from dynamic mechanical spectra, as fol-lowing equation formulated by Witold Brostow et al.[42,43].
B ¼ 1eb � E0
The brittleness values of some PPE/POE blended and 64-layer samples are listed in Table 3. Going from B1 to B5,one had a fivefold increase in the concentration of the elas-tomeric component. The result was a nearly threefoldincrease in the tensile elongation at break. At the sametime the dynamic storage modulus E0 had decreased, butat a rate slower than the elongation at break increase. Thisled to a decrease in brittleness, not very large but clear. Forthe 64-layer samples, the elongation at break remainedapproximately constant with increasing the concentrationof the elastomer, which was attributed to parallel align-ment of POE layers along the direction of external force.E0 decreases with increasing POE content because of less
Loss
mod
ulus
Temperature
Neat POE
Neat PPE
B1
B5
(a)
-75 -50 -25 0 25 50
-75 -50 -25 0 25 50
Loss
mod
ulus
Temperature
Neat POE
Neat PPE
2 layers
32 layers
64 layers
128 layers
(b)
Fig. 13. Temperature dependence of loss moduli of conventionally blended and multilayered samples: (a) Conventional blends with POE content of 10%(B1) and 52% (B5). (b) 2-, 32-, 64- and 128-layer composites with POE content of 45%, 21%, 14% and 14%, respectively.
3280 J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281
adhesion between layers; consequently the brittleness Bincreases.
5. Conclusions
The experimental results showed that the equivalentbox model (EBM) was a valid tool to predict the mechani-cal properties of laminar polymer materials when theinterfaces between two different layers could be ignored.The yield strengths of multilayered composites were high-er than those of conventional blends with the same POEcontent according to EBM. The DSC and tensile tests con-firmed that the yield behavior and mechanism of each
component did not change in both multilayered and con-ventionally blended composites. Therefore, the higherphase continuity became an essential factor causing thehigher yield strengths of multilayered composites com-pared with conventional blends.
The interfaces between PPE and POE layers, where thechains of PPE and POE were tangled each other, could betreated as the conventional blends including series branch.With increasing number of layers of multilayered compos-ites, the content of interfaces increased. These interfaciallayers led to the deviation of mechanical properties of mul-tilayered composites from the theoretical curves predictedby EBM. For example, the yield strength of 128-layer com-
Table 3The comparison of elongation at break (eb), storage modulus (E0) andbrittleness (B) of conventional blends and layered composites.
J. Shen et al. / European Polymer Journal 45 (2009) 3269–3281 3281
posite was lowered compared with that calculated by theadditive principle. The storage moduli of 64-layer compos-ites were lower than theoretical values predicted accordingto EBM. At the same time, the existence of interfacial layersled to a new absorbing peak in the DMA spectrum of32- and 64-layer composites due to the shearing frictionbetween POE and PPE layers through interfaces. Addition-ally, unlike that in blended samples, the brittleness in64-layer composites increased with increasing the POEcontent. Increased B is a consequence of poorer adhesionbetween layers at higher POE contents.
Acknowledgements
The authors are grateful to the Special Funds for MajorState Basic Research Projects of China (2005CB623800),National Natural Science Foundation of China (50603016,50773047), and Funds for Doctoral Disciplines of the Min-istry of Education of China (20050610028) for financialsupport of this work.
References
[1] Veenstra Harm, Verkooijen Paul CJ, van Lent Barbara JJ, van DamJaap, de Boer Albertus P, Nijhof Abe Posthuma HJ. Polymer2000;41:1817.
[2] Nielsen LE, Lawrence E. Mechanical properties of polymers. London:Reinhold Publishing Corporation; 1962.
[3] Meijer Han EH, Govaert Leon E. Prog Polym Sci 2005;30:915.[4] Willemse RC, Speijer A, Langeraar A, Posthuma de Boer AE. Polymer
[6] Xie ZM, Sheng J, Wan ZM. J Macromol Sci Phys 2001;40:251.[7] Stachurski ZH. Polymer 2003;44:6067.[8] Li YY, Hu SW, Sheng J. Eur Polym J 2007;43:561.[9] Joseph S, Thomas S. J Polym Sci B Polym Phys 2002;40:755.
[10] Mustafa O, Mehmet. J Appl Polym Sci 2005;98:1445.[11] Machiels AGC, Denys KFJ, van Dam J, Posthuma de Boer A. Polym Eng
Sci 1996;36:2451.[12] Harm V, Verkooijen Paul CJ, Van Lent Barbara JJ, van Dam Jaap, de
Boer Albertus P, Nijhof Abe Posthuma HJ. Polymer 2000;41:1817.[13] Kolarik J. Polymer 1996;37:887.[14] Kopczynska A, Ehrenstein GW. J Mater Educ 2007;29:325.[15] Kolarik J. Polym Eng Sci 1996;36:2518.[16] Halpin JC, Kardos JL. Polym Eng Sci 1976;16:344.[17] Takayanagi M, Minami S, Uemara S. J Polym Sci 1964;C5:113.[18] Boyd RH. J Polym Sci B Polym Phys 1983;21:493.[19] Mueller CD, Nazarenko S, Ebeling T, Schuman TL, Hiltner A, Baer E.
Polym Eng Sci 1997;37:355.[20] Jarus D, Hiltner A, Baer E. Polymer 2000;43:2401.[21] Sekelik DJ, Stepanov EV, Nazarenko S, Schiraldi D, Hiltner A, Baer E. J
Polym Sci B Polym Phys 1999;37:847.[22] Radford JA, Alfrey T, Schrenk WJ. Polym Eng Sci 1973;13:216.[23] Tangirala R, Baer E, Hiltner A, Weder C. Adv Funct Mater
2004;14:595.[24] Xu Shuangxi, Wen Ming, Li Jiang, Guo Shaoyun, Wang Ming, Du Qin,
et al. Polymer 2008;49:4861.[25] Gregory BL, Siegmann A, Im J, Hiltner A, Baer E. J Mater Sci
1987;22:532.[26] Ma M, Im J, Hiltner A, Baer E. J Appl Polym Sci 1990;40:669.[27] Ebeling T, Hiltner A, Baer E. Polymer 1999;40:1985.[28] Kern J, Hsieh A, Hiltner A, Baer E. J Appl Polym Sci 2000;77:1545.[29] de Gennes Pierre-Gilles. J Phys Lett (Paris) 1976;37:L-1.[30] Brooks NW, Ghazali M, Duckett RA, Unwin AP, Ward IM. Polymer
1999;40:821.[31] Popli R, Mandelkern L. J Polym Sci B Polym Phys 1987;25:441.[32] Kazmierczak T, Galeski A, Argon AS. Polymer 2005;46:8926.[33] Xu XR, Xu JT, Feng LX. Polym Int 2002;51:458.[34] Viana JC. Polymer 2005;46:11773.[35] Cheng SZD, Zhu L, Li CY, Honigfort PS, Keller A. Thermochim Acta
1999;332:105.[36] Yamaguchi M, Miyata H, Nitta KH. J Appl Polym Sci 1996;62:87.[37] Ward IM. Mechanical properties of solid polymers. New
York: Wiley-Interscience; 1983.[38] Plaza AR, Ramos E, Manzur A, Olayo R, Escobar A. J Mater Sci
1997;32:549.[39] Yamaguchi M, Miyata H, Nitta K. J Appl Polym Sci 1996;62:87.[40] Kim YJ, Han CD, Song BK, Kouassi E. J Appl Polym Sci 1984;29:2359.[41] Menard KP. Dynamic mechanical analysis – an introduction. New
York: CRC Press; 2008.[42] Brostow W, Hagg Lobland HE. J Mater Res 2006;21:2422.[43] Brostow W, Hagg Lobland HE. Polym Eng Sci 2008;48:1982.
