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Polymer 55 (2014) 3758e3769

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Polymer

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b-nucleated propyleneeethylene random copolymer filledwith multi-walled carbon nanotubes: Mechanical, thermaland rheological properties

Dimitrios G. Papageorgiou a, **, Lazaros Tzounis b, George Z. Papageorgiou c,Dimitrios N. Bikiaris c, Konstantinos Chrissafis a, *

a Solid State Physics Section, Physics Department, Aristotle University of Thessaloniki, 541 24 Thessaloniki, Greeceb Leibniz-Institut für Polymerforschung Dresden, Hohe Straße 6, 01069 Dresden, Germanyc Laboratory of Polymer Chemistry and Technology, Department of Chemistry, Aristotle University of Thessaloniki, 541 24 Thessaloniki, Greece

a r t i c l e i n f o

Article history:Received 16 April 2014Received in revised form2 June 2014Accepted 5 June 2014Available online 14 June 2014

Keywords:b-Polypropylene random copolymerPolymer nanocompositesMulti-walled carbon nanotubes

* Corresponding author. Tel.: þ30 2310 998188.** Corresponding author. Tel.: þ30 2310 998053.

E-mail addresses: dpapage@auth.gr (D.G. Papageogr (K. Chrissafis).

http://dx.doi.org/10.1016/j.polymer.2014.06.0120032-3861/© 2014 Elsevier Ltd. All rights reserved.

a b s t r a c t

Nanocomposites of a b-nucleated propyleneeethylene random copolymer filled with different amountsof multi-walled carbon nanotubes (МWCNTs) were prepared by a melt mixing process. The complexcrystalline structure that the samples exhibited was attributed to the strong a-nucleating effect ofMWCNTs and the presence of calcium pimelate, as it was revealed by X-ray diffraction. It was found thatthe presence of MWCNTs at concentrations higher than 2.5 wt.% practically diminishes the effect of theb-nucleating agent, and this was also reflected on the mechanical properties of the samples. The qualityof MWCNT dispersion in the polymer matrix was evaluated by transmission electron microscopy. Me-chanical and thermal properties of the nanocomposites were studied as a function of the crystallinestructure and the filler content by tensile and impact testing, thermogravimetric analysis and differentialscanning calorimetry. A crystallization kinetics study was also performed in order to export conclusionsregarding the rates and parameters of crystallization.

© 2014 Elsevier Ltd. All rights reserved.

1. Introduction

Nanocomposite materials based on multi-walled carbon nano-tubes have attracted the interest of the scientific community overthe last decades. The main reason behind that is the importantenhancement of the physicochemical properties of the compositesamples, which are attributed mainly to the presence of MWCNTsand their unique mechanical, thermal and electrical properties.Ajayan et al. first reported an epoxy matrix reinforced withMWCNTs [1] and since then, a vast amount of papers has beendedicated to the processing and resulting mechanical, thermal andelectrical properties of polymer nanocomposites based onMWCNTs. Most of the efforts have been concentrated on the suc-cessful dispersion of MWCNTs into the polymeric matrices, sincethe mechanical properties are most usually hampered due to theformation of bundles and curling of MWCNTs [2].

rgiou), hrisafis@physics.auth.

One of the most popular polymeric matrices is polypropylene,which is a semicrystalline polymer, existing in various crystallineforms, such as the monoclinic a-form and the hexagonal b-form.The different crystalline forms attribute different characteristics tothe polymeric matrix and the nanocomposite samples [3,4], whilethe most efficient way of altering the crystalline phase of poly-propylene from a (the dominant phase due to the higher thermo-dynamic stability) to b form, is the inclusion of a b-nucleating agent[5e7]. In addition, the copolymerization of polypropylene withother olefins such as ethylene, evokes a more balanced behavior ofthe polymeric matrix in terms of rigidity and toughness. Therefore,the physicochemical behavior of a b-nucleated propyleneeethylenerandom copolymer matrix, presents very interesting aspects, asdemonstrated in various literature reports [8e12].

The a-form of polypropylene, which is most commonly useddue to its higher thermodynamic stability, has been extensivelyapplied in film, rigid packaging, sanitary and heating applicationsover the last decades. Its low cost and good physicochemicalproperties have ensured its wide commercial penetration. Oneefficient way to improve even more the impact strength, theelongation at break and heat distortion properties of a

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e3769 3759

propyleneeethylene random copolymer (PPR) matrix, is theincorporation of b-crystalline phase. The differences in the physi-cochemical properties between the two crystalline forms of poly-propylene can be attributed to the higher ductility that b-spherulites exhibit, when compared with the cross-hatchedmorphology of a-lamellae. Additionally, b-phase is very stablemechanically in the initial range of elastic deformation up to theyielding point, and the energy dissipation process during yield,improves even more the mechanical behavior of such systems [13].For the above reasons, the mechanical and rheological behavior ofb-nucleated propyleneeethylene random copolymers presentsseveral interesting aspects, while the addition of fillers which canefficiently reinforce their properties, can lead to the production of amaterial effective enough to withstand extreme conditions inadvanced applications.

Several methods have been proposed for the fabrication ofpolymer nanocomposites [14,15]. Most of the protocols have usedcommon methods such as solution mixing/casting [16], in situpolymerization [17] and melt mixing [18e20], while some re-searchers have employed the most advanced electrospinning [21],coagulation [22] and latex technology techniques [23]. Melt-mixingmethod was selected in this work, since it combines speed,simplicity and low cost. Furthermore, in melt-mixing no solventswhich can alter the chemical composition of the samples, are used.Prior to the melt-mixing, a ball-milling process was performed inorder to ensure good miscibility between the matrix and MWCNTs,which would play an important role on the final dispersion of thefiller in the nanocomposite samples.

In the current manuscript, a detailed physicochemical charac-terization of the complex system of a b-nucleated propylene/ethylene random copolymer filled with MWCNTs was carried out.The competitive crystallization tendency of the nanocompositesamples due to the strong a-nucleating effect from MWCNTs andthe b-nucleating effect from the addition of calcium pimelate,forced thematerials to crystallize at different crystalline forms. Thishad as a result the appearance of behavioral turning point of themechanical behavior of the samples, after a specific filler content.Theoretical models were also applied for the prediction of me-chanical properties based on the characteristics of the matrix andthe filler and the simultaneous estimation of the quality of adhe-sion and stress transfer mechanisms between the filler and thematrix. The effect of the preparation procedure on the dispersionquality was evaluated by examining the microstructure of thenanocomposite by TEM, while the crystalline structure was evalu-ated by XRD. Finally, the thermally-dependent properties of thesamples, such as DMA, TGA and DSC were tested in relation to thefiller content and the crystalline structure of the nanocomposites. Acrystallization kinetics study was additionally performed in orderto export conclusions on the kinetic parameters of the nano-composites and the nucleation activity of the fillers. It is importantto note that in order to eliminate any possible structural changesduring the preparation of the samples, which will cause a differentbehavior during the characterization procedure, all nano-composites were prepared carefully in the same way (quenchingafter hot pressing) for the use during the experimental part.

2. Experimental

2.1. Materials and methods

Polypropylene random copolymer (PP-R) was supplied by BasellPolyolefines with the commercial name Hostalen PP H5416Random Copolymer containing 7 wt.% ethylene content and meltflow rate 0.29 g/10 min (463 K/5 kg). Multi-walled carbon

nanotubes were purchased from Chendu Organic Chemicals Co. Ltd.Their density was 2.1 g/cm3 and their surface area about 500 m2/g.

2.2. Synthesis of calcium pimelate nucleator

The calcium pimelate nucleator was synthesized by doubledecomposition [24]. In detail, calcium pimelate was preparedthrough the reaction between equimolar CaCl2 and sodium pime-late aqueous solution with stirring. The precipitation was filtered,washed twice with distilled water and ethanol, respectively, anddried under vacuum at room temperature. The calcium pimelatewas used at concentration of 0.1 wt.% as b-nucleator for PP-R. Theprepared b-nucleated polymeric matrix will be referred onwards asPPR.

2.3. Nanocomposites preparation

PPR/MWCNT nanocomposites containing 0.5, 1, 2.5, 3.5 and5 wt.% of multi-walled carbon nanotubes were prepared by melt-mixing in a HaakeeBuchler Rheomixer (model 600) at 200 �Cand 30 rpm for 15 min. Prior to melt-mixing, the nanoparticleswere dried by heating in a vacuum oven at 120 �C for 24 h. Thecomponents were physically premixed before being fed into therheomixer. In order to achieve a better dispersion of the MWCNTsin PPR, a RETSCH planetary ball mill model S100 was used. The PPRflakes along with the proper amount of MWCNTs were fed into a ‘C’type stainless steel grinding jar with a capacity of 25 ml. Five steelspheres were also added as a grinding medium. The milling was setat 500 rpm for a period of 1 h. The physicochemical properties ofthematerials prior and after themilling process were tested, and nosignificant differences were observed. The nanocomposite samplesthat were prepared and studied will be referred onwards as PPR/MWCNTs 0.5 wt.%, PPR/MWCNTs 1 wt.%, PPR/MWCNTs 2.5 wt.%,PPR/MWCNTs 3.5 wt.%, PPR/MWCNTs 5 wt.%. The preparation of allsamples for the characterization procedure involved hot pressing ofthe extrudates in order to take the form of films of various thick-nesses and sizes (depending on the characterization technique) andafterwards quenching into cold water. This way, any possible dif-ferences between the samples due to the processing conditions,were eliminated.

2.4. Transmission electron microscopy (TEM)

The TEM investigations were performed using the Libra 200(HR-TEM, Carl Zeiss AG, Germany) high resolution transmissionelectron microscope operating at an acceleration voltage of 200 KV.Thin sections of approximately 70 nm thickness were prepared bymeans of cryo-ultramicrotomy using the Leica Ultramicrotome(Reichert Ultracut S). Diamond knives (Diatome) for cryo-temperatures were utilized for the trimming (model cryotrim45�) as well as the thin-section cutting process (model cryo 35�).The cutting was carried out at temperatures between �120and �160 �C. The thin sections were placed on copper grids(300 mesh Cu, Agar) and investigated with TEM. Bright field TEMimages were recorded, while the energy filtering and contrast ap-ertures were inserted in order to enhance the contrast of theimages.

2.5. X-ray diffraction (XRD)

XRD analysis was performed on the polymeric matrix and thenanocomposites using a Rigaku Ultima þ diffractometer using CuKa radiation, a step size of 0.05� and a step time of 3 s, operating at40 kV and 30 mA.

Fig. 1. TEM images of the PPR/MWCNT composites showing the state of MWCNT nano-dispersion at 0.5, 2.5 and 3.5 wt.% loadings.

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e37693760

2.6. Mechanical properties

Measurements of tensile mechanical properties of the preparednanocomposites were performed on an Instron 3344 dynamom-eter, in accordance with ASTM D638, using a crosshead speed of50 mm/min. Thin films of about 350 ± 25 mmwere prepared usingan Otto Weber, Type PW 30 hydraulic press connected with anOmron E5AX Temperature Controller, at a temperature of175 ± 5 �C. The molds were rapidly cooled by immersing them inwater at 20 �C. From these films, dumb-bell-shaped tensile testspecimens (central portions 5 � 0.5 mm thick, 22 mm gaugelength) were cut in aWallace cutting press and conditioned at 25 �Cand 55e60% relative humidity for 48 h. From stressestrain curves,the values of tensile strength at yield point and at break, elongationat break and Young's modulus were determined. At least fivespecimens were tested for each sample and the average values,together with the standard deviations, are reported. Izod impacttests were performed using a Tinius Olsen apparatus in accordancewith ASTM D256 method. Five measurements were conducted foreach sample, and the results were averaged to obtain a mean value.

2.7. Thermogravimetric analysis (TGA)

Thermogravimetric analysis was carried out with a SETARAMSETSYS TG-DTA 16/18. Samples (5.0 ± 0.2 mg) were placed inalumina crucibles. An empty alumina crucible was used as refer-ence. All materials were heated from ambient temperature to550 �C in a 50 ml/min flow of nitrogen at heating rate of 20 �C/min.Continuous recordings of sample temperature, weight and heatflow were performed.

2.8. Differential scanning calorimetry (DSC)

A PerkineElmer, Pyris Diamond differential scanning calorim-eter (DSC), calibrated with indium and zinc standards, was used forthe study of crystallization and melting of the prepared nano-composites. Samples of 5± 0.1mgwere used in the tests. Theyweresealed in aluminum pans and heated at a heating rate of 20 �C/minto 200 �C which is about 50 �C above the melting point and quiteabove the equilibrium melting point of the studied samples. Thesamples were held at 200 �C for 2 min in order to erase any thermalhistory. For isothermal crystallizations the sample was rapidlycooled to 20 �C above the crystallization temperature and finallycooled to that at a rate 50 �C/min, to achieve equilibration of theinstrument. The crystallization exothermic peakwas then recorded.Crystallization temperatures were in the range between 100 and125 �C.

For non-isothermal crystallization, the same melting proceduredescribed above was applied before the cooling scan. Cooling scanswere performed at rates 2.5, 5, 7.5, 10, 15 and 20 �C/min. A freshsample was used in each run. Heating rate was in most tests 20 �C/min if not otherwise stated.

3. Results and discussion

3.1. Microstructure of nanocomposites

Transmission electronmicroscopy (TEM)was employed in orderto study the nanoscale dispersion of MWCNTs in the PPR matrix.Fig. 1 depicts the TEM images of PPR/MWCNT nanocomposites fil-led with 0.5, 2.5 and 3.5 wt.% MWCNTs, respectively. The images

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e3769 3761

indicate that the polymer melt has efficiently infiltrated theMWCNT primary agglomerates, especially at low filler loadings,therefore single MWCNTs were found dispersed in the polymermatrix. However, at higher filler loadings, aggregates are inevitablyformed, which has as a result the deterioration of some of thephysicochemical properties of the nanocomposite samples [2]. Thequality of the MWCNTs dispersion can be considered quite satis-factory, and it could be attributed to the high energy ball millingprocedure that the MWCNTs were subjected prior to the melt-mixing and the optimization of the melt-mixing conditions.

3.2. X-ray diffraction

The X-ray diffraction profiles of PPR and PPR/MWCNTs nano-composites are shown in Fig. 2. The peak at 16.1� is attributed to thereflections from the (300) plane of b-crystals, while the peaks at14.1�, 16.9�, 18.5� correspond to the principal reflections from the(110), (040) and (130) planes of a-crystals. As it can be seen, thesamples which contain small concentrations of MWCNTs (up to2.5% wt.), include the peak which corresponds to the reflectionsfrom b-crystals, while for higher concentrations of MWCNTs, thepeak disappears due to the strong a-nucleating effect of MWCNTs[25]. Obviously, the competitive crystallization tendency due to a-nucleation fromMWCNTs and the b-nucleation from the nucleatingagent is dependent on the filler concentration, which eventuallydetermines the crystalline structure of the nanocomposite sample.The results are in good agreement with the previously reportedones from the melting behavior of the nanocomposites observedduring the DSC experiments [26]. The b-crystals content in the PPRand the nanocomposites was calculated according to the equationproposed by Turner and Jones [7]:

Kb ¼ Ið300ÞbIð300Þb þ Ið110Þa þ Ið040Þa þ Ið130Þa

(1)

The overall crystallinity Xc was determined by:

Xc ¼ Ac

Ac þ Aa(2)

Fig. 2. X-ray diffraction profiles for the PPR/MWCNTs set of samples. The arrow pointstowards the peak which is attributed to the b-crystals (2q ¼ 16.1�).

where Ac and Aa are the areas under the crystalline peaks andamorphous halo, respectively. In order to proceed to the abovecalculations, a deconvolution of the X-ray diffractograms was per-formed in order to separate the amorphous and the crystallinecontent. The results are presented in Table 1. It can be seen that thedegree of crystallinity decreases slightly with the incorporation ofMWCNTS. Thus, the regions between lamellar stacks exhibit higheramorphous material content, due to the presence of MWCNTs andthe nucleating effect which is associated to them [27]. Anotherinteresting fact is that the b-crystals content at low filler content(up to 2.5 wt.% MWCNTs) presents a small increase, being in goodagreement with the findings of Zhang et al. [28].

3.3. Mechanical properties

In the present study, MWCNTs were used simultaneously with ab-nucleating agent in order to elaborate their competitive or syn-ergistic effect on mechanical properties. The final strength of thepolymer/MWCNT nanocomposites generally depends on two fac-tors. Firstly, the quality of adhesion between the polymeric matrixand the nanotubes should be high, because the load should betransferred effectively between the matrix and the nanotubes [29].If this does not happen, the nanotubes behave as flaws and inducelocal stress concentrations which during the tensile testing act asfailure points [2,5,30]. Secondly, a fine dispersion of the nanotubesshould be secured, since poor dispersionwill induce the presence ofbundles of nanotubes, which are expected to reduce the strength ofthe nanocomposite. As it was observed during the characterizationof the nanocomposites with electronmicroscopy, the samples filledwith MWCNTs present a satisfying dispersion quality, which isexpected to attribute an enhancement on the physicochemicalproperties of the final material.

Generally, the presence of MWCNTs is expected to increase themodulus of elasticity, the strength at yield and break point but alsoto decrease the strain at break, indicating a decrease in the com-posite's toughness and flexibility. This particular behavior has beenobserved also in the case of the prepared nanocomposites, indi-cating the effective stress transfer mechanism between the pro-pyleneeethylene random copolymer matrix and the carbonnanotubes (Table 2).

The nanocomposite samples were additionally tested in order toobserve their impact strength. The results are presented in Table 2.

It is obvious that the results are quite satisfactory, and thebehavior of the nanocomposites can be attributed to several fac-tors. Firstly, since it was shown from the XRD results that thepresence of MWCNTs promotes the formation of a-crystallites,which are known to enhance the modulus of elasticity and thetensile strength at yield point, the higher values of those prop-erties (especially at high MWCNTs concentrations) are anotherindication of the behavioral turning point of the nanocompositesafter the concentration of 2.5 wt.% MWCNTs. It can be seen thatthe yield values of the nanocomposites remain quite stable up to2.5 wt.%, while from that point on, the increase that is recorded,can be attributed to the presence of a-crystalline phase [31].Regarding the modulus of elasticity, the degree of crystallinity isknown to influence also significantly the Young's modulus, since

Table 1XRD results from the nanocomposite samples.

MWCNTs content (wt.%) 0 0.5 1 2.5 3.5 5

Crystallinity Xc (%) 53.9 53.2 53.7 53.2 55.5 57.6b-content (%) 9 9.2 9.4 11.1 e e

Table 2Mechanical properties of the nanocomposites.

MWCNTs content (wt.%) Tensile strength at yield (MPa) Tensile strength at break (MPa) Elongation at break (%) Young's modulus (MPa) Impact strength (J/m)

0 21.3 ± 0.6 25.1 ± 1 813 ± 50 575 ± 61 8.3 ± 0.40.5 22.3 ± 0.6 27.3 ± 1.1 506 ± 51 637 ± 55 8.6 ± 0.81 22.8 ± 1.4 28.7 ± 1 423 ± 49 656 ± 65 8.5 ± 0.72.5 23.2 ± 1.3 29.2 ± 1.5 405 ± 45 686 ± 59 9.5 ± 0.73.5 27.4 ± 1.2 27.4 ± 1.3 360 ± 39 800 ± 52 7.6 ± 0.95 28.6 ± 1.4 26.6 ± 1.8 302 ± 42 877 ± 45 7.2 ± 0.4

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e37693762

the crystalline regions are stiffer than the amorphous phase andthe results from XRD indicate an increase of Xc at higher fillercontents. The tensile stress at break presents an increase for theinitial MWCNTs contents, where the b-crystalline structure ismaintained, while at higher filler content, presents a drop due tothe formation of aggregates and the alteration of the crystallinephase. Furthermore, the higher packing density that the a-crystalsexhibit along with the stress concentration effect of MWCNTs dueto agglomeration and the higher crystalline content of thosesamples, lower significantly the elongation at break properties ofthe nanocomposites [32]. Generally, the good dispersion of thenanotubes due to the mechanochemical ball milling procedureperformed prior to the melt-mixing process, allowed a moreuniform distribution of the applied load, which acted in favor ofmost of the mechanical properties of the materials. Regarding theimpact strength results, it can be seen that the strength of thenanocomposites is quite stable at lower concentrations of filler,increases slightly at the 2.5 wt.% MWCNTs sample, due to thesmall increase in the b-crystalline content as was concluded byXRD experiments, while at higher concentrations a significantdrop can be observed. The reason for this behavior is the crys-talline structure of the nanocomposite samples. It has been shownthat a-polypropylene presents lower impact strength values dueto its brittle nature, thus this drop was expected for the specificsamples.

3.3.1. Modeling of the mechanical propertiesVarious theoretical models have been proposed for the pre-

diction of the mechanical properties of composite materials,which have attempted to take into account important aspectswhich affect the mechanical properties such as aggregates con-centration, stiffening effects, interfacial adhesion and fillerdispersion among others [33,34]. The main problem of thosemicromechanical theories is that most of them have beendeveloped for the estimation of the performance of micro-composite systems and not nanocomposites, as in our case.Therefore, some of them are expected to fit quite satisfying theexperimental values and some others not. In addition, some ofthe parameters or approaches introduced by those theories arepartly empirical, thus the results should be exported withcaution and should not be considered absolute.

Initially, the well-known Einstein equation was tested, underthe premise that it is only valid at low concentration of filler [35].The Einstein equation is described by the following formula:

EcEm

¼�1þ BVf

�(3)

where Ec and Em are the Young's moduli of the composite and thematrix respectively, B is a parameter related to the filler geometryand interface adhesion between the filler and the polymeric ma-trix and Vf is the volume of the filler. The B parameter takes valuesfrom 1 to 2.5, which indicate that the adhesion between the filler

and the matrix is poor (B ¼ 1) or when the adhesion is excellent(B > 2.5).

Guth and Smallwoodmodified the Einstein equation by adding asecond power term [36]. The equation obtained was as follows:EcEm

¼�1þ 2:5Vf þ 14:1V2

f

�(4)

where the first term relates to the stiffening effect from individualnanotubes, and the second power term corresponds to the contri-bution from the interaction between the nanotubes and the matrix.

Kerner has developed another equation in his attempt todescribe the modulus of elasticity of expansible polymers whichinclude fillers [37]. The equation which describes his theory is:

EcEm

¼ 1þ 15ð1� vÞ

8� 10v$

Vf1� Vf

!(5)

where n is the Poisson ratio, which is approximately assumed to be0.35 [38].

Mooney inserted a crowding factor (S) to Einstein's equation, forthe ratio of the apparent volume occupied by the nanofillers totheir true volume, and its value lies between 1.0 and 2.0 [39]. Theequation that describes his theory is:

EcEm

¼ exp

2:5Vf1� SVf

!(6)

The equations that were developed in the early 90s from Hal-pineTsai for the prediction of the modulus of elasticity of com-posite materials demonstrated efficiently the effect of the fillergeometry and orientation of the filler to the final properties of thematerial [40]. The initially developed equations are:

EcEm

¼ 1þ chVf1� hVf

; n ¼Ef.Em � 1

Ef.Em þ c

(7)

where c ¼ 2(l/d) is a constant shape factor related to the aspectratio of the reinforcement length l and diameter d. According to thesupplier's data sheet, the length of MWCNTs was 25 mm and thediameter was 10 nm. The effective Young's modulus of MWCNTswas adopted from earlier literature reports equal to 1 TPa [8].

The primary form of the HalpineTsai equations was developedfor composites with unidirectional reinforcement and for thisreason, Thostenson and Chou [41] modified this approach whiletaking into account that the outer shell of the nanofiller wouldwithstand the load which will be applied, as a logical assumption ofthe relatively low bonding with inner layers, especially when thesurface of the nanofiller has not been chemically modified. Thus,the effective MWCNTs elastic modulus was evaluated by inserting aterm to the HalpineTsai equation, which considers that the appli-cation of all loads is only to the outer section of the filler (Eeff ¼ 4t/dEf). The Thostenson equation is:

E11Em

¼241þ 2

�l

d

�0@�Ef.Em���df.4t�

�Ef.Em�þ�lf

.2t�1AVf

35$241�

0@�Ef.Em���df.4t�

�Ef.Em�þ�lf

.2t�1AVf

35�1

(8)

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e3769 3763

where E11 is the composite modulus of elasticity in the principalmaterial direction, Vf is the volume fraction of the multi-wallednanotubes, Ef, Em are the moduli of the nanotubes and the matrixrespectively, lf is the length of the nanotubes, df is the thickness ofnanotubes and t is the layer thickness of the nanotube (¼0.34 nm).The results from the application of all the above theories plottedtogether with our experimental observations can be seen in Fig. 3.

From the results of the application of the theories, some con-clusions can be reached. Initially, the fact that Einstein's theory ismore accurately represented by B values as high as 5, is an indi-cation of the good interfacial adhesion between the polymericmatrix and the filler. The theory proposed by Guth and Smallwoodis also relatively close to the experimental value, due to itsexponential shape, implying that the interaction between theMWCNTs and the matrix is quite adequate. Kerner's equation isonly valid and close to the experimental values at low fillerloadings, while Mooney's theory is relatively closer to the

Fig. 3. Experimental Young's modulus values fitted by different theorie

experimental values to its exponential form. However, the factthat those two theories were initially developed for sphericalfillers makes the results rather unreliable. The results from theapplication of the original HalpineTsai equation are ratherconfusing, since this theory is based on perfect adhesion betweenthe matrix and the filler, while in our case the experimental valuesare higher than the ones predicted theoretically. This assumptionis not valid in the case of our nanocomposites since polypropyleneis non-polar compared to the MWCNTs, thus it can be concludedthat the HalpineTsai equation omits some contributions fromMWCNTs that play an important role on the mechanical proper-ties of the nanocomposites. Finally, the Thostenson and Chouequation simulates quite successfully the behavior of the modulusof elasticity of the nanocomposite samples. Therefore, their orig-inal assumption that the load is practically applied only to theouter surface of MWCNTs along with the insertion of the layerthickness of the nanotube was correct.

s. The line on the experimental values is only a guide to the eye.

Fig. 4. Dynamic mechanical relaxation behavior of the prepared PPR/MWCNT nano-composites: (a) storage modulus (E0) and (b) tan d.

Fig. 5. a) Mass loss (TG%) curves of PPR and nanocomposites, b) differential mass loss(DTG) of PPR and nanocomposites under nitrogen atmosphere.

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e37693764

3.4. Dynamic mechanical analysis (DMA)

Dynamic mechanical analysis was also employed in order toinvestigate the macromolecular chain conformations of the nano-composites during heating. The results from the DMA of PPR/MWCNTs nanocomposites can be seen in Fig. 4. A clearly increasingtrend of the storage modulus (E0) can be seen for the nano-composites with increasing filler content, with the sample filledwith 3.5 wt.% MWCNTs presenting the higher stiffness between thenanocomposites. The sample filled with 5 wt.% MWCNTs does notpresent the highest storage modulus, most probably due to theformation of large aggregates. The stiffening effect of MWCNTs,even at low filler contents has been reported by several groups[25,42e47] and the enhanced interactions between the filler andthe matrix as they were reported earlier with TEM and tensile/impact testing reduce the mobility of the chains due to the increaseof stiffness. Additionally, the transition zone which is formed be-tween the matrix and the filler and it is described several times inliterature for polymer/MWCNTs nanocomposites, exhibits highmodulus, since the adsorption of the polymeric chains on theMWCNTs surface, reduces the mobility of the macromolecularsegment and increases the stiffness of the samples [43,48,49].

The b-relaxation peak as it is presented in Fig. 4b represents theglass transition temperature and it is also increasing withincreasing filler content due to the previously mentioned enhancedpolymerefiller interactions. It is also known that the glass transi-tion temperature depends mainly on the mobility of the chainsegments of the macromolecular chains of the polymeric matrix.The motion of the chain segment proceeds more slowly and chainmovement is more difficult at the Tg when the molecular chain isrestricted. Therefore, this phenomenon can be attributed to thenanoconfinement effect on the macromolecular chains fromMWCNTs, which increases the glass transition temperature, asdescribed by Chen et al. [50].

3.5. Thermogravimetric analysis

The thermal stability of the polymeric matrix and the nano-composites was evaluated by thermogravimetric analysis tech-nique. The samples were heated under Nitrogen atmosphere up to550 �C and the heating rate that was applied was 20 �C/min. Themass loss and the derivative of the mass loss thermograms arepresented in Fig. 5.

From the results presented in Fig. 5, it is obvious that theMWCNTs have induced a significant enhancement of the thermalstability of the nanocomposites, indicating the formation of astrong interface between the two components, as was determined

Fig. 6. (a) Evolution of relative crystallinity as a function of crystallization time for PPRat various temperatures. The geometrical points represent experimental data while thecontinuous line represents the results of the theoretical Avrami model, (b) reciprocalcrystallization half-time of nanocomposite samples.

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e3769 3765

also from the mechanical testing. Kashiwaghi et al. [51] observedalso this effect and came into the conclusion that MWCNTs form aprotective layer, which acts as a thermal insulation layer and also asa barrier for the gasification of the degradation products. However,if we take into consideration that the filler concentration corre-sponds to a very small fraction of the initial polymeric matrix, theresults cannot be securely interpreted, even if we suppose a perfectdispersion of the filler. Chen et al. [50] introduced the nano-confinement concept, according to which, the filler particles cancreate locally confined areas where the macromolecular chains ofthe polymeric matrix are more or less confined, thus their

Table 3Results of the Avrami analysis for isothermal crystallization of PPR/CNT nanocomposites

Tc (�C) PPR PPR/MWCNTs 0.5 wt.% PPR/MWCNTs 1 wt.% P

n Zc n Zc n Zc n

100 5.4 1.9 5.8 2.21 5.6 2.34 5105 5.3 1.51 5.5 1.93 5.1 1.94 5110 4.9 1.03 3.7 1.26 4.1 1.28 4115 2.8 0.37 2.9 0.65 2.2 0.53 3120 1.9 0.1 2.0 0.2 2.8 0.16 2125 1.8 0.02 3.0 0.06 1.9 0.04 2

movement is restricted, their regular coil conformation is notmaintained and the thermal degradation is enhanced. Furthermore,the nanoconfinement effect leads to lower chemical reactivity ofthe chains and higher activation energy for the nanocompositesamples. This theory can be considered valid for the prepared set ofsamples, since the thermal stability of the nanocomposites in-creases with increasing filler content, indicating that the higher theMWCNTs content, the higher is the confinement of the macromo-lecular chains due to the presence of MWCNTs. Additionally, thisconcept is in agreement with the results exported from DMA,where the increase of the glass transition temperature for thenanocomposites (Fig. 4b), indicates the restricted movement of themacromolecular chains, induced by MWCNTs.

3.6. Differential scanning calorimetry

The nucleating effect of MWCNTs and the effect of the simul-taneous presence of calcium pimelate during the isothermal andnon-isothermal crystallization of the nanocomposites were studiedby differential scanning calorimetry. For the isothermal crystalli-zation study, the samples were crystallized isothermally at varioustemperatures from 100 to 125 �C. Also, in order to proceed to thenon-isothermal crystallization kinetics study, several cooling rateswere applied, namely 2.5, 5, 7.5, 10, 15 and 20 �C/min.

3.6.1. Isothermal crystallization kineticsCrystallization of polymer melt is usually accompanied by sig-

nificant heat release which can be measured by differential scan-ning calorimetry. The isothermal crystallization behavior of the b-nucleated polymeric matrix and the effect of MWCNTs contentwere investigated by means of DSC over a temperature range be-tween 100 and 125 �C. The relative degree of crystallinity X(t) canbe obtained if we make the assumption that the evolution ofcrystallinity is linearly proportional to the evolution of heatreleased during the crystallization and it can be formulated as:

XðtÞ ¼

Z t

0ðdHc=dtÞdtZ ∞

0ðdHc=dtÞdt

(9)

where dHc denotes the measured enthalpy of crystallization duringan infinitesimal time interval dt. The limits t and∞ on the integralsare used to denote the elapsed time during the course of crystal-lization and at the end of the crystallization process, respectively. Inthe DSC traces it was obvious that as the supercooling, i.e., thedifference between the melting and crystallization temperature,decreases, the crystallization rate becomes slower and theexothermic peak becomes broader. Thus, the time to reach the peakincreases. The relative crystallinity as a function of time wasdetermined and it is presented in Fig. 6a only for PPR for brevityreasons. The half time of crystallization (t1/2) is an indication of thecrystallization rate and it can be obtained from the curves depictingthe degree of crystallinity versus time. According to the results

.

PR/MWCNTs 2.5 wt.% PPR/MWCNTs 3.5% PPR/MWCNTs 5 wt.%

Zc n Zc n Zc

.8 2.22 5.7 2.21 5.6 2.24

.5 1.92 5.3 1.91 5.1 2.1

.4 1.39 4.4 1.4 4.6 1.67

.3 0.77 3.2 0.77 3.8 0.99

.6 0.29 2.8 0.31 3.5 0.45

.6 0.09 2.8 0.1 2.7 0.14

Fig. 7. Evolution of the crystallization peak temperature as a function of the coolingrate for PPR and nanocomposites.

Fig. 8. Plots of log[-ln(1�X(t))] vs. log t for non-isothermal crystallization of PPR (a)and PPR/MWCNTs 2.5% (b).

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e37693766

which are presented in Fig. 6b, the reciprocal of the half time ofcrystallization decreases almost exponentially with increasingcrystallization temperature. This is an indication that t1/2 increaseswith increasing Tc and decreases with increasing MWCNTs content,as a result of the strong nucleating effect of MWCNTs.

The development of relative crystallinity was analyzed using theAvrami equation. A time-dependent relative volumetric crystal-linity X(t) for an isothermal crystallization process can be expressedas:

XðtÞ ¼ 1� expð�ktnÞ0XðtÞ ¼ 1� exp��ðKtÞn� (10)

where,n is the Avrami exponent, which is a function of the nucle-ation process, and k is the growth function, which is dependent onnucleation and crystal growth. Since the units of k are a function ofn, Equation (10) can be written in the composite Avrami form usingK instead of k (where k ¼ Kn). The Avrami equation in the simpleform represents unimpeded spherical crystal growth. In order totake into consideration the whole range of data that correspond tothe relative crystallinity of the studied sample, the non-linearfitting procedure with Eq. (10) was employed, based on the Mar-quardteLevenberg algorithm. This algorithm seeks the parametervalues that minimize the sum of the squared differences betweenthe observed and predicted values of the dependent variables by aniterative process, until convergence with a predefined tolerance isobtained. The theoretical lines that were produced are plotted inFig. 6a and compared to the experimental data. As it can be seen,the quality of the fit is very good (R2 > 0.99), meaning that theAvrami equation is capable of describing the isothermal crystalli-zation kinetics of PPR/MWCNTS nanocomposites. Some minor de-viations appear only for data that correspond to very low (<5%) orvery high (>95%) relative degrees of crystallinity. The Avrami pa-rameters after the fitting procedure were calculated and presentedin Table 3. The Avrami exponent n depends on both the mechanismof nucleation and the morphology of crystal growth and ideally nwould be an integer whose values would represent different typesof growth. The n values that were calculated were in the vicinity of3e4 and they should possibly be related with three-dimensionalgrowth. In addition, the Avrami parameters k and K are indicativeof the crystallization rate and they both decrease with increasingtemperature.

3.6.2. Non-isothermal crystallization kineticsDuring the non-isothermal crystallization kinetics study, the

samples were cooled down from the melt at several heating rates,namely 2.5, 5, 7.5, 10, 15 and 20 �C/min and the cooling curves wererecorded. The variation of the peak temperature of non-isothermalcrystallization for all samples is visible in Fig. 7, where an increaseof the crystallization temperature occurs simultaneously with theincrease of MWCNTs content; thus higher rates of crystallizationare achieved on the nanocomposite samples. This was expected,since more nucleation sites are offered to the samples during thecrystallization due to the simultaneous presence of MWCNTs andthe b-nucleating agent. Their presence facilitates practically theheterogeneous nucleation as both fillers seem to work synergisti-cally this time for the faster crystallization of the nanocomposites.Regarding the contribution to crystallization from each filler, in ourprevious study it was shown that the insertion of the b-nucleatingagent into the PPRmatrix, increases the crystallization temperaturefor almost 10 �C [27], while the presence of the highest MWCNTsconcentration (5 wt.%) increases the crystallization temperature (ofthe b-nucleated matrix) almost 4e5 �C.

Table 4Non-isothermal crystallization kinetic parameters of PPR and PPR/MWCNT nanocomposite samples analyzed by the Avrami theory.

Filler content (wt.%) 0 0.5 1 2.5 3.5 5

b (oC/min) n Zc n Zc n Zc n Zc n Zc n Zc

Primary stage2.5 4.0 0.18 3.4 0.7 3.9 0.65 4.6 0.65 5.1 0.62 4.7 0.675 4.3 0.61 3.4 0.91 4.3 0.88 4.7 0.88 5.5 0.87 4.3 0.937.5 4.1 0.90 3.6 0.97 4.2 0.94 4.5 0.95 5.3 0.95 4 1.0110 4.5 1.00 3.7 0.99 4.1 0.98 4.1 0.99 4.4 0.99 4.2 1.0215 4.6 1.03 4 1.1 3.9 1.06 4.4 1.1 4.5 1.08 4.4 1.0920 4.7 1.05 4.1 1.07 4.1 1.09 4.4 1.08 4.4 1.1 4.1 1.17Secondary stage2.5 2.4 0.40 2.5 0.43 2.6 0.44 3.3 0.46 3.4 0.41 3.9 0.455 2.0 0.75 2.4 0.87 2.7 0.93 2.8 0.93 3.3 0.95 3.2 0.947.5 1.9 0.95 2.3 0.95 2.8 0.98 2.3 0.92 3.4 0.98 3 1.0510 1.8 1.00 2.4 1.02 2.3 1.09 2.2 1.06 2.7 1.09 3.5 1.1715 1.6 1.02 2 1.05 1.9 1.03 2.1 1.03 2.4 1.07 3.2 1.1120 1.6 1.04 2.3 1.04 2.1 1.04 2.4 1.05 2.3 1.05 2.9 1.09

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e3769 3767

3.6.2.1. Avrami analysis of non-isothermal crystallization.Several theories have been proposed in order to describe the non-isothermal crystallization kinetics of polymers. One of the mostused is the Avrami theory [52e54], modified by Jerziorny [55]:

1� XðtÞ ¼ expð�ZttnÞ (11)

where X(t) is the relative crystallinity, Zt is the crystallization rateconstant and n is the Avrami exponent, which is dependent on thenucleation mechanism and the growth dimension. Although thereis no direct physical meaning of n and Zt for the study of non-isothermal crystallization, their study can provide further insightto the crystallization phenomena. According to Jerziorny, the rate ofthe non-isothermal crystallization should be adequately correctedsince it depends on the cooling rate. Thus, the corresponding rateconstant at unit of cooling rate Zc can be obtained by: logZc ¼ logZt/b. In order to take into consideration the whole range of data thatcorrespond to the relative crystallinity of the studied sample, thelinear fitting procedure was employed. The results for the poly-meric matrix and the sample filled with 2.5 %wt.% MWCNTs ispresented in Fig. 8. As it can be seen, there is a deviation from

Fig. 9. (left graph): Dobreva plots for the evaluation of nucleation activity of PPR and PPR/

linearity for both samples and for all cooling rates, in contrast withthe Avrami plots of a non-nucleated polypropylene, which werecompletely linear and they were presented in an earlier report [56].This is an indication of the secondary crystallization and the het-erogeneous nucleation process due to the presence of the nucle-ating agent and the MWCNTs [57,58]. According to the resultsreported in Table 4, the Avrami exponent presents variations whichare indicative of the complication of the nucleationmechanism dueto the simultaneous presence of fillers.

From the results presented in the table, the crystallizationconstant increases with increasing cooling rate, while the exponentn does not follow any trend. This is reasonable since Zc measuresthe crystallization rate which gets faster with supercooling. Thevalues of Zc become constant only at high cooling rates (b > 10 �C/min). During the secondary stage of nucleation of the samples, thecrystalline perfection of the samples takes place, thus the decreaseof n2 with increasing cooling rate corresponds to a faster and moreeffective crystallization stage. This is mainly caused by the imper-fect and metastable crystals grown during the primary crystalli-zation stage. Additionally, since the growth rate is mostlycontrolled by the rate of secondary nucleation, the increasing filler

MWCNT nanocomposites, (right graph): calculated B*/B ratio from Dobreva's method.

D.G. Papageorgiou et al. / Polymer 55 (2014) 3758e37693768

content provides more nucleating sites, favoring in this way thesecondary nucleation.

3.6.2.2. Nucleation activity. During the study of the crystallizationbehavior of the PPR/MWCNT nanocomposites, the strong nucle-ating effect of MWCNTs was examined. Dobreva and Gutzow pro-posed a relatively simple method for the estimation of nucleationactivity by fillers [59,60]. If the reinforcement is extremely active,the nucleation activity will be close to zero, while for inert rein-forcement, it will be close to unity. According to Dobreva's method,for homogeneous and heterogeneous nucleation from the polymermelt, the nucleation activity can be calculated by the ratio: 4¼ В*/Вwhere B is a parameter which can be calculated by the followingequation: B ¼ us3V2

m=3nkT0mDS2m where u is a geometric factor, s isa specific energy, Vm is the molar volume of the crystallizing sub-stance, n is the Avrami exponent, DSm is the entropy of melting andT0m the infinite crystal melting temperature. T0

m was assumed to beequal to 151 �C for PPR, according to previous calculations [27].

In addition, B can be experimentally determined from the slopeof Eq. 13 obtained by plotting ln(b) versus the inverse squareddegree of supercooling 1/(DTp)2 (DTp ¼ Tm�Tp):

ln b ¼ A� B2:303DT2P

(12)

ln b ¼ A� B*

2:303DT2P(13)

where A is constant, B and B* are the constants related to homo-geneous and heterogeneous nucleation and DΤp is the degree ofsupercooling and equals to Τm�Tc, where Tc is the temperaturecorresponding to the peak temperature of DSC crystallization andTm is the melting temperature. Plots of ln(b) versus 1/(DTp)2 areshown in Fig. 9. From the slopes of PPR and PPR/MWCNT nano-composites, the values of В and B* were calculated and the ratio ofB*/B (which represents the nucleation activity) versus MWCNTscontent, is plotted in the right graph of Fig. 9. It should be notedagain that the lowest values of B*/B ratio can be interpreted as thehighest nucleation activity. Therefore from Fig. 9, the nucleationactivity increases with increasing MWCNT content, with maximumactivity being reached by the samples with 5 wt.% MWCNTs (lowestB*/B ratio). It is worth to notice that in this case, the presence of theb-nucleating agent which is included into the polymeric matrix,makes it more difficult to increase significantly the nucleation ac-tivity of the nanocomposites with the addition of MWCNTs, sincethe nucleation activity of the matrix is already high due to thenucleation sites offered by the b-nucleating agent.

4. Conclusions

b-nucleated propyleneeethylene random copolymer/MWCNTnanocomposites have been prepared by melt-mixing. The TEMimages revealed that the preparation procedure induced a quitesatisfactory dispersion of the nanotubes, a fact which was reflectedon the physicochemical properties of the samples. The competitiveaction between the a-nucleating effect of MWCNTs and b-nucle-ating effect of calcium pimelate was verified by XRD, since after acertain filler content (2.5 wt.% MWCNTs) the samples crystallizedsolely at a-form, a fact which afterwards affected the mechanicalbehavior of the nanocomposite samples. Additionally, the physicalpresence of MWCNTs ensured a general improvement on the me-chanical and rheological properties, while the application of severalmicromechanical models revealed that the adhesion quality be-tween the filler and the matrix is satisfying. Furthermore, the

thermal stability and the glass transition temperature of thenanocomposites increased with increasing filler content, due to thenanoconfinement effect on the macromolecular chains byMWCNTs. The nanoconfinement led to the restricted movement ofthe macromolecular chains and the difficulty of the matrix tomaintain the regular coil conformation of the chains. Finally, thesimultaneous presence of the b-nucleating agent and MWCNTsprovided a significant number of nucleating sites in order to pre-sent an increase in the crystallization temperature with increasingMWCNT concentration. The isothermal and non-isothermal crys-tallization kinetics revealed the three-dimensional growth of thenanocomposites and the effect of the fillers on the growth rateparameters, while the nucleating efficiency was improved due tothe presence of MWCNTs.

Acknowledgment

This research has been co-financed by the European Union(European Social Funde ESF) and Greek national funds through theOperational Program “Education and Lifelong Learning” of theNational Strategic Reference Framework (NSRF) e Research Fund-ing Program: Heracleitus II. Investing in knowledge society throughthe European Social Fund.

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