Designing Tough and Fracture Resistant Polypropylene-multi Wall Carbon Nanotubes Nanocomposites by Controlling Stereo-complexity and Dispersion Morphology

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Designing tough and fracture resistant polypropylene/multi wall

carbon nanotubes nanocomposites by controlling stereo-complexity

and dispersion morphology

Dibyendu Das, Bhabani K. Satapathy

Centre for Polymer Science and Engineering, Indian Institute of Technology Delhi, Hauz Khas, New Delhi 110016, India

a r t i c l e i n f o

Article history:

Received 7 June 2013Accepted 18 August 2013Available online 29 August 2013

Keywords:

PolypropyleneMulti wall carbon nanotubesTacticityToughnessFractureCrack propagation

a b s t r a c t

A remarkable toughness enhancement (>330%) of multi wall carbon nanotubes (MWCNT) filled stereo-complex polypropylene (PP) matrix i.e. blend ofisotactic-PP andsyndiotactic-PP (70:30) with differencesin stereo-regularity has been observed. The enhancement has been correlated to quantifiable morpholog-ical parameters such as free-space lengths concerning dispersion and relatively greater reduction in crys-tallite size/lamellar thickness. Systematic analysis of glass transition data and estimation of multi wallcarbon nanotubes induced reduction in interfacial polymer chain immobilization reiterates susceptibilityof polymer segments to ready-mobility. The extent of toughening has quantitatively been analyzed byfracture-energy partitioning, essential work of fracture (EWF), approach enabling the detection of asemi-ductile-to-tough-to-quasi-brittle transition in the MWCNT filled stereo-complex polypropylene.Real-time fracture kinetics analysis revealed toughening mechanism to be primarily blunting-assisted;an aspect also corroborated by extensive plastic flow without much energy dissipation in the inner frac-ture process zone. Thus the study establishes a new pathway of tacticity-defined matrix modification totoughen nanocomposites.

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1. Introduction

Difference in stereo-regularity (tacticity) induced matrix modi-fication enabling the manipulation of the polymer nanocompositesstructure favorably towards enhanced energy dissipation (tough-ening mode) is a theme not so outrightly attempted. This is be-cause molecular level interaction involving topologicalconstraints, van der Waals forces, Lennard-Jones potential and Lon-don dispersion forces may be very complex to explain the precisenature of adhesive/cohesive interactions in understanding tough-ening mechanisms. The topological attributes to the mechanismof toughening has the significance in the sense that the conformalproximity of the zig-zagarrangements and the wrapping-up phe-nomenon via the helical conformation of the polymer chain config-urations facilitate in effectively enhancing the stress transfermechanism by increasing the interfacial interaction. The size scaleadvantage in this sense plays a crucial role since smaller dimensionof the hollow cylinders extend a larger amount of interfacial areafor molecular level interactions. This is because the effective vol-ume fraction of the polymer chains that potentially can be ad-sorbed per unit surface area of the reinforcing second phase is

much lower and hence may lead to promoting efficient polymer-nanotube interaction. Such interactions involving less of looselyentangled polymer chains renders greater extent of chain immobi-lization enabling the interfacial polymer-nanotube network to actas a composite phase and hence larger energy dissipation andtoughness[13]. The pioneering work of Lordi and Yao[4] postu-lating the role of helical conformations in increasing the effectivebinding of the polymer-nanotube interface and subsequently thecounter approach view point in evolving the theory of a strongpolymer-nanotube interface for enhanced toughening by Jiangand Penn [5] reiterates the above understanding pertaining tonanoscale reinforcement especially that of carbon nanotubes.

Wong et al.[6]have reported the determining role of polymer-nanotube interfacial characteristics using molecular mechanicssimulations and elasticity calculations. Relatively higher interfacialshear stresses in case of nanotubes than their micron sized coun-terparts have reportedly been attributed to intimate solid phasecontact between polymer and nanotube at the molecular level.Lee et al. in another fundamental work concerning the role of dis-persion and exfoliation of single wall carbon nanotubes (SWCNT)and multi wall carbon nanotubes (MWCNT) in polypropylene(PP) matrix have postulated on the temperature dependence ofinterfacial strength across polymer/nanotube interface and thebuilding-up of a three dimensional percolation network[7,8]. Cole-mann et al.[9]in their pioneering work emphasized the role of an

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Corresponding author. Tel.: +91 11 26596043; fax: +91 011 26591421.

E-mail address:[email protected](B.K. Satapathy).

Materials and Design 54 (2014) 712726

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ordered interfacial stiff polymer region in enhancing the reinforce-ment efficiency, particularly with respect to enhancements inmodulus. Schaefer and Justice[10]while commenting on the elu-sive nano effects have put forth an explanation based on smallangle X-ray scattering (SAXS), light scattering and electron imagingstudies for the discrepancy pertaining to mechanical property in-crease in terms of ubiquitous large scale disorder of the filler dis-

persion morphology. Lin et al. [11] have reported regarding thenanoplastic flow of glassy polymer chains interacting withMWCNT in polymer nanocomposites where intrinsic differencesin the fundamental behavior of the entangled polymer chains withrespect to crazing and shear yielding mechanism determining theoverall extensibility/ductility of the nanocomposites were high-lighted. Interestingly, the entangled network responses in polymernanocomposites have been proved not only to be responsible forthe mode of deformation but also their interaction with individualnanotubes. Mu and Winey[12]have inferred the possible relation-ship between enhanced load-transfer effectiveness and increasedpolymer molecular weight with their molecular level implicationsto the larger size of the polymer chain (radius of gyration-Rg) rela-tive to that of the nanofiller. It was elaborated to the extent thatthe criteria 2RgD for stronger interfacial interaction between Polymer/carbon nanotube were semi-empirically proposed to qualitativelyexplain the nature of interface.

Comprehensive investigations by Satapathy et al. [13] and Gansset al.[14]on the morphology and fracture behavior relationship inisotactic polypropylene (i-PP)/MWCNT nanocomposites have re-vealed a ductile-to-semi-ductile transition in the nanotube loadingrange of 0.51.5 wt.%. Further the details of such a transition werecorroborated by fracture kinetics and localization of strain contoursanalysis. Qualitatively, such fracture transitions have been experi-mentally characterized by a transition from stable non-steady tosteady state crack tip opening displacement rate and thereby caus-ing a delayed yielding phenomenon. Extending these observations,further, to understand the MWCNT induced toughening effects in

i-PP/MWCNT nanocomposites the temperature dependence ofcreep response was also intensively discussed by Ganss et al. [15].Bao and Tjong[16]have assessed the strain rate and temperaturesensitivity of i-PP/MWCNT nanocomposites where the increase inheat deflection temperature and glass transitiontemperatures havebeen reported to increase with MWCNT content and the yieldingphenomenon could be explained by a time defined Eyrings equa-tion. In consequence to such observations by Bao and Tjong [16]the reinforcing effects were observed to enhance at elevated tem-peratures. Karevan et al. [17]have reported the dominant role ofinterphase in defining the stress-transfer efficiency whilevalidatingtheexperimentallyobtainedtensilemodulitothatofthepredictionsfrom HalpinTsai model. The interrelationship of fiber strength,interfacestrengthand critical lengthof thecarbonnanotubes as pro-

posed in the KellyTyson model (arguably valid for nanocompos-ites) further indicates that toughness as a bulk nanomaterialresponse is substantially dependent on interface strength assumingnanotube related parameters remain unchanged [18]. Prashanthaet al. [19] have reported on the enhanced impact resistance fornotched samples where nanotubes effectively limits the crack prop-agation of PP/MWCNT nanocomposites prepared by masterbatchdilutionroute.Thiswas accompanied with an increase inyieldstressand lesser extent of reduction in ductility when compared to theclassical carbon fiber reinforced PP composites.

In an effort to understand the mechanism of increase instrength in polymercarbon nanotube nanocomposites interfacialcharacteristics including restrained relaxation of polymer chainsin the interphase, tailored interfaces via surface modification of

MWCNT, role of nanoparticle size, loading and shape and evolutionof interface across polymer-carbon nanotubes have also been

discussed in the literature [2022]. Kovalchuk et al. [23] haveinvestigated on the purified and alkyl functionalized MWCNT filledi-PP and syndiotactic polypropylene (s-PP) nanocomposites wheresignificant improvement in nanocomposite plasticity and modulusof s-PP was achieved. However, the possible role of varying the tac-tility/stereo regularity of the polymer matrix as blend of the twostereo-forms of PP on the mechanical performance of nanocompos

ites has not been attempted yet, despite the theoretical possibilitythat syndiotacticity of PP may enhance the toughness of PPMWCNT nanocomposites by interfering with the nature of theinterface because of the differences in the helical conformation oisotactic and syndiotactic polypropylenes. Therefore, the presentpaper deals with the assessment of the morphological and fracturetoughness properties of PP/MWCNT nanocomposites where theconventional matrix (i-PP) based system is modified and the ma-trix is replaced by a 70:30 blend of i-PP to s-PP giving rise to a ma-trix with asymmetric stereo-complexity.

2. Experimental details

2.1. Materials and composites preparation

The details of the raw materials selected and the processing con-ditions are given inTable 1andTable 2. The polymer nanocomposites viz. IPNC (i-PP/PP-g-MA/MWCNT) and ISPNC (i-PP/s-PP(70:30)PP-g-MA/MWCNT) were prepared by melt mixingof i-PP, s-PP, polypropylene grafted maleicanhydride (PP-g-MA) with the commerciamasterbatchPlasticylPP 2001containing20 wt.% MWCNTin a corotating type twin screw extruder (Steer Omega 20)at a screw speedof 250 rpm. The continuous strands obtained from the twin screwextruder were later chopped in a granulator and dried in a vacuumoven at 80 C for 2 h before the injection molding on an L&T Demaginjection molding machine (Model PFY 40-LNC4P) to obtain80 mm 80 mm square plates of1.5 mm thickness. The process-ing conditions with temperature profiles in twin screw extruder

and injection molding are given in Table 2. The composites designation and compositions are given inTable 3.

2.2. Morphology of the nanocomposites

Transmission electron microscopy (TEM) was carried out oncryo-microtomed specimen sections of thickness of about 7090 nm. The samples were cut from the central portion of the injec-tion molded sample using a Leica Ultracut EM UC6/EM FC6 ultra-microtome (Model Leica Mikrosysteme GmbH, Wien, Austria)The ultramicrotome is equipped with a diamond knife with a cutangle of 35. The TEM used is an EM 912 (LEO, Oberkochen, Ger-many) operated at 120 kV.

2.3. Structural characterization by 2D wide angle X-ray (WAXD)diffraction

WAXD measurements were done with an X-ray diffractometemodel number PW304060 Xpert PRO (Netherland) with 40 kVvoltage, 30 mA current and Cu Ka = 1.54 , in a 2h range from550, to evaluate the crystallinity of the nanocomposites. Thecrystallinity was calculated by applying the peak-area integrationmethod in the range of 2h= 540(as typical for i-PP). The amor-phous scattering curve was drawn by an approximation based onstandard experimental and theoretical principles.

2.4. Light microscopic measurement

The spherulitic morphology of i-PP, s-PP, IPNC and ISPNC arestudied using a Instec HCS-302 (Meiji Techno-Japan) microscope

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The sample was placed between a microscope glass slide and a

cover slip. The sample was heated up to 200

C and then kept for2 min to ensure complete melting of the samples. Then samplesare cooled down to 130 C at a cooling rate of 5C per minute. Aftercomplete crystallization the images were captured, using a2 mega-pixel camera attached to the microscope, with a resolutionof 20-fold magnification.

2.5. Thermal characterization

2.5.1. Differential scanning calorimetry (DSC)/Thermo gravimetric

analysis (TGA)

Differential scanning calorimetry (DSC) measurements werecarried out (in TA instruments, Q200) at a heating rate of 10 C/min to obtain the crystallinity of the composites. The samples were

first heated up to 220 C then allowed to stand still at that temper-ature for 1 min (to relieve its thermal history) followed by cooling

down to 20 C at a cooling rate 10C/min and then subsequently

heated up to 220

C at a heating rate 10

C/min. Apparent enthal-pies of fusion have been calculated from the area under the exo-thermic peaks and the crystallinity (%) of the i-PP, i-PP/s-PPblends and its MWCNT filled nanocomposites have been deter-mined using following equations:

Xc;iPP DHm;iPP=DH0m;iPP wt:%iPP 1

Xc;sPP DHm;sPP=DH0m;sPP wt:%sPP 2

Xc;blends Xc;iPP wt:%iPPXc;sPP wt:%sPP 3

whereDHmis the heat of fusion of i-PP, s-PP in the MWCNT nano-composites and DHom is the heat of fusion of 100% crystalline i-PPi.e., taken as 209 J/g and that of s-PP as 196 J/g[24]. The crystallite

sizes have been determined following Scherrers equation [7,25].The referred equation for the crystallite size may be given as,

Table 1

Data sheet of raw materials and their characteristics.

Raw materials Grade Supplier Characteristics

Isotactic polypropylene homopolymer REPOLH033MG

RelianceIndustriesLimited

Melt flow index of 3.30 g/10 min. at 230 C and 2.26 kg load; Tm= 165C

Syndiotactic polypropylenehomopolymer

452149 SigmaAldrich Melt flow index of 2.20 g/10 min. at 230C and 2.26 kg load; Tm= 125C

PP-g-MA OPTIM P-405

PLUSS Polymer Density (q) = 0.91 g/ml; Tm= 163C; melt flow index of 20 g/10 min at 190 C and 2.16 kgload; Maleic anhydride content range 1.62.5%

Isotactic polypropylene/MWCNTcommercial master batch

PlastiCylPP 2001

Nanocyl 20 wt.% MWCNT in master batch, average diameter = 9.5 nm; average length = 1.5lm andpurity > 90%

Table 2

Processing parameters used in extrusion and injection molding machine.

Extrusion temperature profileZones Barrel 2 Barrel 3 Barrel 4 Barrel 5 Barrel 6 Die adaptor Die headTemperature C 180 190 200 210 220 225 230

Temperature profile set in injection molding machine

Feed (C) Z-I (C) Z-II (C) Z-III (C) Nozzle (C)

40 190 210 220 230Processing conditions used in injection molding machine

Process parameter ValueInjection pressure 1059 barInjection time 4 sCooling time 25 s

Table 3

Details of the compositions and designation of i-PP/PP-g-MA/MWCNT nanocomposites (IPNC) and i-PP/s-PP/PP-g-MA/MWCNT nanocomposites (ISPNC).

Designation Compositions

i-PP (wt.%) i-PP/s-PP a (70:30) (wt.%) PP-g-MA (wt.%) MWCNT (wt.%)

i-PP 100.00 0.00 0.0IPNC-0(iPP/PP-g-MA) 95.00 5.00 0.0IPNC-0.5(iPP/PP-g-MA/0.5 wt.% MWCNT) 94.53 4.97 0.5IPNC-1.0(iPP/PP-g-MA/1.0 wt.% MWCNT 94.05 4.95 1.0IPNC-1.5(iPP/PP-g-MA/1.5 wt.% MWCNT) 93.58 4.92 1.5IPNC-2.0(iPP/PP-g-MA/2.0 wt.% MWCNT) 93.10 4.90 2.0IPNC-3.0(iPP/PP-g-MA/3.0 wt.% MWCNT) 92.15 4.85 3.0ISPNC-0(i-PP/s-PP/PP-g-MA) 95.00 5.00 0.0ISPNC-0.5(iPP/s-PP/PP-g-MA/0.5 wt.% MWCNT) 94.53 4.97 0.5IPSNC-1.0(iPP/s-PP/PP-g-MA/1.0 wt.% MWCNT) 94.05 4.95 1.0IPSNC-1.5(iPP/s-PP/PP-g-MA/1.5 wt.% MWCNT) 93.58 4.92 1.5IPSNC-2.0(iPP/s-PP/PP-g-MA/2.0 wt.% MWCNT) 93.10 4.90 2.0IPSPN-3.0(iPP/s-PP/PP-g-MA/3.0 wt.% MWCNT) 92.15 4.85 3.0

a Stereo-complex Polypropylene matrix (70:30-Asymmetric composition).

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Lcrystallite Kk

B2h cos h 4

whereL is the crystallite size, k is wave length of radiation,B is fullwidth half maxima (FWHM) and Kis proportionality constant (0.94for FWHM of spherical crystals with cubic symmetry).

Furthermore the non-oxidative thermal stability of the nano-composites has been measured using a Pyris 6 TGA instrument,

Perkin-Elmer. The thermo gravimetric analysis (TGA) were carriedout in nitrogen atmosphere at a heating rate 10 C/min in the tem-perature range 30700 C.

2.6. Dynamic mechanical analysis (DMA)

Nanocomposite specimens with dimensions of 20 8 1.5 mm3

were subjected to dynamic mechanical analysis (DMA) measure-ments in single cantilever mode. The measurements were done inthe temperature range of30 C to 165 C at a frequency of 1.0 rad/s and heating rate of 5 K/min on an Q800 machine (TA Instruments,USA)to characterize loss tangent (tand) for thequalitativeestimationof damping/energy dissipation characteristics and for quantitativelyascertaining shift (if any) in glass transition temperature (Tg) of the

MWCNT filled virgin isotactic (IPNC systems) and stereo-complex(ISPNC system) polypropylene matrices.

2.7. Determination of fracture behavior

The fracture behavior of the nanocomposites has been mea-sured following the plane stress essential work of fracture (EWF)approach [26,27]. Rectangular injection molded bars of80 mm 20 mm 1.5 mm dimension were used as specimensfor the preparation of double edge notch tension (DENT) specimensfor the measurement of fracture parameters, EWF and Non- essen-tial work of fracture (N-EWF). The specimens were prepared insuch a manner that the injection molding direction and the direc-tion of the application of uni-axial tensile force onto the bars re-

mained identical. The samples were pre-notched with differentligament lengths varying from 210 mm. The fracture of thesepre-notched specimens was carried out using a universal tensiletesting machine (Zwick Z250) under constant extension speed(2 mm/min) to obtain individual loaddisplacement curves. Theplane-stress essential work of fracture (EWF) method has beenused since conceptually it enables the distinguishing betweentwo terms representing the resistance to crack initiation (EWF: we)and resistance to crack propagation (N-EWF: bwp) correspondingto inner fracture process zone (IFPZ) and outer plastic deformationzone (OPDZ) respectively. The precondition for the validity of EWFapproach has been demonstrated by the self-similar nature of theloaddisplacement diagrams of these nanocomposites. In thisstudy, the fracture mechanical tests (test speed: 2 mm/min, i.e.

e9

= 0.033 mm/s, room temperature) were performed on double-edge-notched-tension (DENT) specimen by a universal testing ma-chine with mechanical grips. The clamp distance was 40 mm. Fornotching, a special device with fresh razor blades (notch tip radiusof 0.20 lm) was used to realize that both notches are similarsized. For each material, at least 8 specimens with different liga-ment lengths were tested.

The total work of fracture W (kJ/m2) dissipated in a notchedspecimen under plane-stress conditions can be partitioned intotwo components, We and Wp characterizing the inner fractureprocess zone (IFPZ) and outer plastic deformation zone (OPDZ)respectively, as schematically shown inFig. 1. Therefore,

W WeWp weBlbwpB:l2

5

whereB,l andbare specimen thickness, ligament length and shapefactor of the plastic zone respectively. The specific work of fracture

w may be obtained on normalizing (dividing) Wby the ligament(notched) area i.e., B l. The relationship (with the quantities wand bwp are given in N/mm and N/mm

2 units respectively) maybe represented as,

w webwpl 6

Based on the fact that the intrinsic fracture process takes placein the inner fracture process zone (IFPZ), the term EWF, the essen-tial work of fracture, is experimentally determined by extrapola-tion of w as a function of l to zero ligament length. For thequantitative determination of these fracture parameters (we andbwp), several similar sized specimens with different ligamenlengths are monotonically loaded to obtain several data points inthe plot ofw versus l , the intercept and the slope of which gives

rise to we (EWF: resistance to crack initiation) and bwp (N-EWFresistance to crack propagation) respectively[28].

2.8. Fractured surface morphology

The post-yield fracture surface morphologies of the DENT spec-imens of MWCNT filled isotactic and stereo-complex (iso-syndioblend) polypropylene matrices based nanocomposites have beeninvestigated using scanning electron microscopy (SEM) on a ZeissEVO-50 electron microscope to analyze the associated fracture mi-cro mechanisms involving micro-deformation, micro-fibrillationshear yielding and layer peeling off characteristics. The failednanocomposite surfaces were gold sputter-coated prior to exami-nation to make the surfaces conductive.

3. Results and discussion

3.1. Morphology of the nanocomposites

The distribution and dispersion of the MWCNT in i-PP and instereo-complex PP matrices (i-PP/s-PP: 70/30) are shown inFig. 2. It was observed that the MWCNT formed partially agglomer-ated domains at the distribution level that are uniformly spaced inthe PP matrix. The state of dispersion involving entangled, curvedand partially looped MWCNT that are randomly spaced apart fromeach other could be observed. The free space length (lf) in the ma-trix are estimated by resorting to the principles proposed by Burris[29]with the exception that the sides of all the maximum square

spaces are averaged linearly. Subsequently the infiltrated free-space lengths (linf) have been estimated by locating the maximum

Fig. 1. Double edge notch tension (DENT) specimen showing inner fracture processzone (IFPZ) and outer fracture process zone (OPDZ).

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possible free spaces in between the individual MWCNTs whileexcluding the squares that have sides less than or equal to thediameter of the incorporated MWCNT. The diameter of theMWCNT used, in the present study is in the range of 810 nm.The ratio l inf/lfis taken as a quantitative measure to ascertain theextent of dispersion. The higher the ratio, the better is the disper-sion. From theTable 4, it could be well observed that the dispersionof the MWCNT in the stereo-complex PP matrix is fairly uniformwhere infiltrated PP rich region could be observed in the inter-tubespacing spanning with a dimension more than that of the MWCNT.The increase in MWCNT content has led to the appearance of mas-

terbatch rich regions indicating inefficient infiltration of PP chains.This may be attributed to a phenomenon resembling agglomera-tion of MWCNT (as seen in ISPNC-3.0 and is discussed in theSEM micrographs). Interestingly the distributive mixing of IPNChas been observed to be better than that of ISPNC since isomorph-icity of the matrix PP may facilitate the easy distribution of the

MWCNT-rich regions of PP/MWCNT masterbatch via physicalentanglements of the loose PP chains at melt temperatures. How-ever, the incorporation of PP with a difference in stereo-regularityfacilitates the dispersion of the MWCNT through the inherent ten-dency of the stereo-complex PP blend tending to demix due to ahigher repulsive interaction parameter operating across the twostereo regular forms of the PP matrices. The enthalpy of demixingof s-PP and i-PP may well be construed from their differences insolubility parameters (d) and their repulsive interaction parame-ters (v). The solubility parameters of i-PP and s-PP at a tempera-ture 453 K are reported to be 5.64 (cal/cm3)1/2 and 5.76 (cal/

cm3

)1/2

respectively [30]. The solubility parameter of i-PP/s-PP(stereo-complex) blend may be calculated, by assuming that thesegmental volume change is negligible due to mixing, using thefollowing equation:

diPP=sPP /1diPP 1 /1dsPP 7

The critical interaction parameter may theoretically be esti-mated, while assuming monodispersity of the two components,by the following equation[31]:

v 0:5N0:51 N0:52

2; 8

whereN1andN2are the degrees of polymerization of i-PP and s-PPrespectively. Thev value so obtained is 10.02 104 which is in

striking contrast to reported value ofv of3.54 104

correspond-ing to 50:50 blend composition[30]. It must also be noted that the

Fig. 2. Dispersion and distribution of MWCNT: TEM images of the nanocomposites (a) IPNC-0.5 at low magnification, (b) IPNC-0.5 at high magnification, (c) IPNC-1.0 at highmagnification, (d) ISPNC-0.5 at low magnification, (e) ISPNC-1.0 at low magnification, and (f) ISPNC-1.0 at higher magnification.

Table 4

The free space length in the matrix and infiltrated free-space lengths of

nanocomposites.

lf(nm) linf(nm) linf/lf

IPNC-0.5 911.01 36.68 0.040IPNC-1.0 797.75 29.51 0.037ISPNC-0.5 774.22 55.83 0.072ISPNC-1.0 483.07 77.89 0.161



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segmental dynamics of s-PP is reported to be 1.7 times lower thanthat of i-PP at 500 K[32]. These estimations corroborate the factthat the two stereo-regular phases of PP retain the differences intheir dynamics of miscibility and that may lead to a difference inthe stereo specific relaxation response causing an enhancement inthe extent of dispersion of MWCNT.

3.2. Structural characterization by 2D wide angle X-ray diffraction(WAXD)

The structural characteristics of the IPNC (i-PP/PP-g-MA/MWCNT nanocomposites) and ISPNC (i-PP/s-PP/PP-g-MA/MWCNTnanocomposites) have been obtained from WAXD of the injectionmolded specimens. The X-ray diffraction plots in terms of intensity(I) versus 2hmeasured in the 2hrange of 550are shown inFig. 3.Four distinct peaks at 2hof 14.08, 16.83, 18.45and 21.7corre-sponding to the (110), (040), (130) and (111)/(041) net planeshave been observed in both, i-PP and its nanocomposites. Thesepeaks crystallographically correspond to the monoclinic a-crystal-line phase with complete absence of theb-crystalline phase, whichshows two strong peaks at 2hof 16.2and 21.2[33,34]. But in caseof ISPNC two additional peaks at 2hof 12.23and 24.70have beenobserved which indicate the characteristic peaks of (200), (310/400) plane of helical form I of s-PP[35,36]. The crystallinity resultsobtained from WAXD and DSC (2nd heating curve) are given inTa-ble 5. It was observed that the percentage crystallinity of ISPNC-0and its nanocomposites are much lower than IPNC-0 and itsnanocomposites.

3.3. Light optical microscopy

The optical micrographs of the virgin polymer matrix and nano-composite melts cooled to a temperature of 130 C are shown inFig. 4. It has been observed that incorporation of MWCNT into i-PP matrix led to a distinct reduction in the spherulites radius, i.e.from 24.48 lm in IPNC-0 to 14.36 lm in IPNC-0.5. On increas-ing the MWCNT content to IPNC-1.0 and IPNC-3.0 the spherulitessizes tended to reduce further. In contrast the incorporation ofMWCNT into the stereo-complex matrix i.e. ISPNC-0 has led to sig-nificant changes in the spherulitic dimensions, where the spheru-lites have appeared to be much finer with worm-like structuresdistributed uniformly in the relatively darker amorphous PP phase.Such alterations in the bulk spherulitic morphology may be attrib-uted to the relatively slower crystallization process of the syndio-tactic phase of the stereo-complex PP matrix of ISPNC basedsystems, when compared to the virgin i-PP matrix of IPNC. Thesubstantial reduction in spherulitic dimensions observed in ISPNCmay be construed to have an analogy with relatively stiffer crystal-line phase of PP dispersed in a softer amorphous phase of PP. Such

a scenario leads to the understanding that stiffer crystalline micro-domains dispersed in a randomly entangled amorphous PP matrixmay lead to a morphology that is conducive to toughening.

3.4. Thermal and mechanical properties of the nanocomposites

3.4.1. Differential scanning calorimetry (DSC)

The effects of MWCNT on the melting behavior of i-PP, in IPNCand ISPNC are shown inFig. 5. It has been observed that in IPNC themelt temperature and crystallinity remains nearly unaffected withthe addition of MWCNT indicating the absence of MWCNT inducedchanges in crystalline packing of the base matrix, i.e. i-PP. Interest-ingly the nature of the melting peaks corresponding to IPNC andISPNC showed characteristic difference in their correspondingendotherms. In IPNC a single melting peak at 165 C is observedwithout any appreciable MWCNT-induced shift in the Tm. On theother hand ISPNC showed two distinct endothermic transitions at120125 C and 155165 C corresponding to stable helical-1(2:1)[37]and double helix (3:1) conformations of s-PP and i-PPrespectively. Furthermore on increasing the MWCNT contenabove 0.5 wt.% MWCNT, a shift of the melting endothermic peakto a lower temperature of 155 C could be observed. Such a de-crease by 10 C in theTmmay be attributed to faster rate of crys-tallization indicating possible role of MWCNT as a-nucleatingagents. Typically faster nucleation process facilitates the reductionof the crystallite size as has already been observed from the crystallite sizes determined by using Scherrers equation based onWAXD data for ISPNC nanocomposites [38]. Such a reductionmay eventually lead to closer crystalline packing that may be con-strued as a consequence of stereo-complexity. Since such a phe-nomenon ofa-nucleation could not be observed in IPNC unlikeISPNC, the possibility of higher extent of dispersion of MWCNT inISPNC may be conceptually presumed. Theoretically, when twoisomorphic phases are melt blended with differences in their melt-ing points (Tm), the possibility of a nanoscopic inclusion gettingdispersed becomes a diffusion controlled process. Such a diffusion

controlled process may facilitate by the partial melting of s-PPphase around 120 C and thereby may improve the dispersionof MWCNT. Thus these calorimetric observations reiterate the rel-atively better state of dispersion morphology in ISPNC as comparedto IPNC systems that are already discussed (Fig. 1).

3.4.2. Dynamic mechanical analysis of the nanocomposites

3.4.2.1. Loss-tangent (tand) response and Tg. The results from solidstate dynamic mechanical analysis in terms of variation of loss tan-gent (tand) with temperature of IPNC and ISPNC are shown inFig. 6a and b respectively. From the figure it could be clearly ob-served that in case of both IPNC and ISPNC transitions occur ataround 20 C to 25 C, which correspond to the glass-rubber tran-

5 10 15 20 25 30 35 40 45 50

IPNC-3.0

IPNC-2.0

IPNC-1.5

IPNC-1.0

IPNC-0.5Intensity

(a.u.)

2

i-PP

IPNC-0

5 10 15 20 25 30 35 40 45 50

ISPNC-3.0

ISPNC-2.0

ISPNC-1.5

ISPNC-1.0Intensity

(a.u.)

ISPNC-0.5

ISPNC-0

2

(a) (b)

Fig. 3. Wide angle X-ray diffractogram (WAXD) measurement of nanocomposites: variation of intensity with 2h(a) IPNC (b) ISPNC.

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sition of the i-PP. The magnitude of tan d increases significantly

with the incorporation of 0.5 wt.% MWCNT in the IPNC-0. But withthe further loading of MWCNTs the tandvalues decreases substan-tially. In the entire composition range the tan dvalues remained inbetween that of IPNC-0.5 and i-PP with an exception of IPNC-3.0.This increment qualitatively attributed to a better dispersion ofMWCNT up to 0.5 wt.% loading. Similarly in case of ISPNC the mag-nitude of tand is highest in case of 1.0 wt.% MWCNT content(ISPNC-1.0) and it decreases with the further loading of MWCNTs.Qualitatively this indicated that in case of ISPNC dispersion ofMWCNT is better up to a higher percentage loading. Interestinglyit could be also observed from Fig. 6 that in case of IPNC theglass-rubber transition i.e. glass transition temperature (Tg) re-mained almost unaffected whereas in case of ISPNC theTgis shiftedslightly to relatively lower values. Theoretically, the Tgof the ma-

trix should be expected to increase with the incorporation of thenanofillers, since the nanofillers mostly infiltrate into the amor-phous region imposing restrictions to the chain segmental mobil-ity. Similar results for reduction in Tg based on dynamicmechanical properties in case of polycarbonate/MWCNT nanocom-posites has been reported and attributed to processing inducedthermal degradation of polymer chain[39]. But in our case the on-set degradation temperature obtained from the TGA (listed in Ta-ble 5) is increased by 8 C for IPNC and by 20 C for ISPNC.These results however do not support the assumption of process-ing induced thermal degradation of i-PP matrix.

3.4.2.2. Influence of interface and MWCNT content on Tg. The promi-nent reduction in Tgof the ISPNC systems relative to the IPNC

nanocomposites may be attributed to an overall reduction in thepercentage crystallinity estimated from WAXD and supported bythe crystallinity data obtained by calorimetric (DSC) measure-ments. To further analyze the phenomenon two complementaryapproaches were adopted by resorting to the assumptions that Tg(of the nanocomposites) may get affected by the correspondingthermal responses (a) of the nanoinclusion with respect to matrixand (b) of the immobilized matrix polymer chains by dispersedMWCNT. In the former approach the reduction in Tg equivalentdue to the incorporated MWCNT (Tg2), that is responsible for themanifestation of the overall decrease in Tgof the nanocomposite,may theoretically be estimated by Fox equation[40]. The Fox equa-tion may be stated as below

1

Tg

w1

Tg1

w2

Tg2 9

whereTg,Tg1andTg2correspond to the glass transition temperature

of the nanocomposites, matrix without MWCNT, reduced Tg-equiv-alent due to incorporated MWCNT respectively, w1and w2are theweight fraction of the matrix and nanotube respectively.

However, in the later approach the immobilized volume frac-tion (T) may be estimated from Kerners equation [41]. The Tmay be construed as responsible for storage modulus ( E0) enhance-ment. The T is taken into account in order to quantify approxi-mately the interfacial polymer chain immobilization influence onthe overall Tg of the nanocomposites. The rearranged Kernersequation for composite to estimate the immobilized volume frac-tion of polymer chains (T) may be given as below.

/T EcT EmT

EmT aEcT 10

whereEm(T) andEc(T) are the storage moduli of the matrix (IPNC-0and ISPNC-0 matrices) and their nanocomposites respectively at areference temperature (i.e. at T= 25 C) and

a24 5tm7 5tm

11

wheremmis Poissons ratio of the composites and is taken as 0.30 inthe present case. For the estimation ofTgof the immobilized phasein IPNC-0.5 and IPNC-1.0, IPNC-0 has been taken as the matrix,whereas for the same compositions of ISPNC series ISPNC-0 hasbeen taken as the matrix. Since theoretically the unaffected Tg ina nanocomposite system indicates the unaffected bulk relaxation

phenomenon of the polymer chains, hence the nanocompositeswith Tg identical to that of the unfilled matrix, may be assumedto be the system with virtual absence of any immobilization/mobi-lization effects. The Tgestimates obtained by adhering to the abovetwo approaches are given in Table.5. It could be clearly observedthat theTgof the immobilized PP phase in IPNC in the compositionrange of IPNC-0.5 to IPNC-1.0 remained at 10C. Such observa-tions indicated the formation of an ordered interphase that isdynamically (unaltered) resembling i-PP chains. In contrast inISPNC the Tgof the immobilized phase remained in the range of12 C to +4.7 C. These indicate that the incorporation ofMWCNT into ISPNC matrix readily enhances the segmental mobil-ity. This fundamentally reiterates the possibility of MWCNT inducednano-structural reorganization facilitating an increase in the overall

amorphous free volume space (fraction); which in turn, microme-chanically, may lead to corresponding ductile toughening effects.

Table 5

Nanotube induced immobilized volume fraction at the interface, thermal property and crystalline morphological properties.

T@ 25 C Tg(C) Tg

# (C) Tg## (C) TGA Crystallite size$ (nm) Crystallinity from

WAXD (%)Crystallinity fromDSC (%)

Llamellar(nm)Tonset(C) (040)i-PP

IPNC-0 0 23 443.6 3.85 55.10 53.78 1.997IPNC-0.5 0.103 25 2 10 449.7 3.80 56.03 54.06 1.983IPNC-1.0 0.076 21 2 10 451.5 3.83 56.84 55.52 1.991IPNC-1.5 0.056 23 450.8 3.70 56.83 55.21 1.993

IPNC-2.0 0.102 23 449.6 3.79 56.94 56.63 1.994IPNC-3.0 0.166 23 450.3 3.77 58.46 58.02 1.990ISPNC-0 0 24 416.1 3.66 35.10 30.67 1.987ISPNC-0.5 0.026 26 2 12 440.9 3.47 34.61 25.57 1.927ISPNC-1.0 0.049 20 1 5 443.5 3.12 34.64 30.20 1.930ISPNC-1.5 0.053 22 3 9 425.2 3.49 35.83 29.20 1.933ISPNC-2.0 0.071 24 422.2 4.25 35.94 29.61 1.924ISPNC-3.0 0.113 23 10 17 422.2 3.90 35.46 31.10 1.925

T= Volume fraction of immobilized polymer chains at 25 C;Tg= glass transition temperature of the composites, Tg# = Reduced equivalent glass transition temperature due

to incorporation of MWCNT;Tg## = glass transition temperature of immobilized volume fraction of polymer chain (immobilized volume fraction of polymer are converted to

corresponding weight fraction of the polymer to apply Fox equation), $ = crystallite size estimated from Scherrers equation, Llamellar= lamellar thickness deduced from theThomson Gibbs equation.



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3.5. Fracture behavior of the nanocomposites

3.5.1. Load displacement diagram and the validity of the EWF

approach

The loaddisplacement diagram of i-PP, IPNC-0 (i-PP/PP-g-MA),ISPNC-0 (i-PP/s-PP/PP-g-MA) and its MWCNT filled nanocompos-ites are shown inFig. 7. All the composites viz. IPNC and ISPNCand their neat components have not only shown thermoplasticbehavior but also have shown self-similar nature of the forcedisplacement diagrams, indicating pre-conditional validity ofEWF approach. It is also clear from the loaddisplacement diagramthat displacement becomes higher in case of i-PP/PP-g-MA ascompared to i-PP although maximum load required for the stable

crack propagation remain almost unaffected. In case of IPNC-0.5(i-PP/PP-g-MA/0.5 wt.% MWCNT) the displacement is higher as

compared to the neat component, but further addition oMWCNT affect the displacement as well as the maximum loadfor the stable crack propagation. Interestingly, in the stereo-com-plex matrix system, ISPNC-0 the displacement becomes nearlydouble as compared to its neat counterpart (i.e. IPNC-0) and re-mains unaffected up to ISPNC-1.5. The plane stress criteria forthe applicability of post-yield fracture mechanics (PYFM) conceptis ensured by the Hills analysis [42]as shown inFig. 8. The anal-ysis revealed that the net section stress (rn) remained indepen-dent of the ligament length (l). The full yielding of the entireligament length region in the DENT specimens occurred at max-imum load (Fmax) and prior to the resumption of crack propaga-tion, which is visually ensured. The total works of fractures (W

for various nanocomposites were obtained by integration of thetotal area under the loaddisplacement diagrams. After normali-

Fig. 4. Light optical micrographs of (a) i-PP, (b) s-PP (c) IPNC-0, (d) ISPNC-0, (e) IPNC-0.5, (f) ISPNC-0.5, (g) IPNC-1.0, (h) ISPNC-1.0.

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zation by ligament area (B. l) the specific work of fracture (w) isobtained.

3.5.2. Work of fracture-composition relationship and transition in

fracture modes

The specific work of fracture as a function of ligament length ( l)for i-PP, IPNC and ISPNC are shown in Fig. 9. The linear fit of thedata points across the valid range ofl leads to the determinationof slope as non-essential work of fracture (N-EWF;bwp) and inter-cept as essential work of fracture (EWF; we) for each composition.The variation of we and bwp as a function of MWCNT content isshown inFig. 10.It has been observed that with the incorporation

of 5 wt.% PP-g-MA into i-PP as the matrix, i.e. in IPNC-0, we in-creased by 125% with respect to i-PP which is followed by a sharpdrop of 70% in we value in the nanocomposites containing0.5 wt.% MWCNT (IPNC-0.5). On further addition MWCNT intothe nanocomposites, wevalues are found to remain broadly unaf-fected till 1.5 wt.% of MWCNT followed by a linear increase tillIPNC-3. In contrast blending of i-PP matrix with 30 wt.% of s-PP(ISPNC-0) lead to an enhancement in we by 135%, whereas onincorporation of MWCNT into ISPNC as the matrix till ISPNC-1.5a consistent reduction in wecould be observed. However, despitethe reduction in weof ISPNC the gross magnitude of the same re-mained well above their corresponding IPNC counterparts. Theseinevitably indicate an intrinsic enhancement in the resistance tocrack initiation of the nanocomposites based on stereo-complex

PP-matrix, i.e. ISPNC when compared to i-PP based IPNC systems.Mechanistically, these observations indicate the non-linear depen-

dency of the crack growth prior to failure. The resistance to crackpropagation (bwp) increased by 9% and 77% for IPNC-0 andIPNC-0.5 respectively, when compared to i-PP matrix (Fig. 10b).Such a remarkable toughening may be attributed to an improvedstate of dispersion of MWCNT in the polymer as evident from theTEM micrographs (Fig. 1). These findings are in striking contrastto the earlier reports where an increase in bwp by 15% was re-ported in case of i-PP/MWCNT nanocomposites in the absence ofPP-g-MA [13]. The role of PP-g-MA in improving the dispersionof MWCNT in polymer matrix was reported independently by Pra-santha et al. [43] and Yang et al. [44] while investigating themechanical and toughness properties of nylon and PP nanocom-

posites. However upon further increasing the amount of MWCNTup to 1.5 wt.% as in IPNC-1.5, the bwpdecreases by 57% comparedto IPNC-0.5. Following such a reduction, upon further addition ofMWCNT the magnitude ofbwpremained nearly unaffected.

Interestingly in ISPNC-0.5 and ISPNC-1.0 the bwp increased by123% and 339% respectively when compared to the ISPNC-0matrix. The maxima in bwpin ISPNC-1.0 indicate the compositionwith maximum resistance to crack propagation that essentially ac-counts for the energy dissipation in the outer-plastic deformationzone (OPDZ). However, on further increasing the MWCNT contentto ISPNC-1.5 and ISPNC-2.0 a sharp reduction in bwpby 60% wasobserved. The magnitude of we showed a nearly linear decreasewith the increase in MWCNT content from ISPNC-0.5 to ISPNC-1.5 indicating a systematic decrease in the resistance to crack ini-

tiation of the nanocomposites. Theoretically resistance to crack ini-tiation has a qualitative correspondence to the energy dissipated in

80 90 100 110 120 130 140 150 160 170 180

IPNC-3.0

IPNC-2.0

IPNC-1.5

IPNC-1.0

IPNC-0.5

IPNC-0

i-PP

Temperature (C)

Hea

tflow

(en

do

)[

H(J/g)]

110 120 130 140 150 160 170

Temperature (C)

Helical form I

melting of s-PP

Shift of Tm

a- ISPNC-0

b- ISPNC-0.5

c- ISPNC-1.0

d- ISPNC-1.5

e- ISPNC-2.0

f- ISPNC-3.0

a

b

c

d

e

f

Hea

tflow

(en

do

)[H(J/g)]

(a) (b)

Fig. 5. Differential scanning calorimetry (DSC) heat scans for the nanocomposites (a) IPNC (b) ISPNC.

-25 0 25 50 75 100 125 150

0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

0.16

0.18

Temperature (C) Temperature (C)

i-PP

IPNC-0

IPNC-0.5

IPNC-1.0

IPNC-1.5

IPNC-2.0

IPNC-3.0

tan

-25 0 25 50 75 100 125 150

0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

0.16

0.18

0.20

0.22

0.24

tan

ISPNC-0

ISPNC-0.5

ISPNC-1.0ISPNC-1.5

ISPNC-2.0

ISPNC-3.0

(a) (b)

Fig. 6. Variation of loss tangent (tan d) with temperature (T) of the nanocomposites (a) IPNC (b) ISPNC.



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0

0

50

100

150

200250

300

350

400

450

500

550

600

650

700

750

LOAD(N)

DISPLACEMENT (mm)

3.89 mm

4.67 mm

6.49 mm

7.46 mm

9.56 mm

i-PP

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 0 1 2 3 4 5 6 7 8 9

0

50

100

150

200250

300

350

400

450

500

550

600

650

700

DISPLACEMENT (mm)

LOAD(N)

3.65 mm

4.28 mm

5.91 mm

6.74 mm

7.78 mm

9.39 mm

IPNC-0

1 2 3 4 5 6 7 8 9 10 110

100

200

300

400

500

600

DISPLACEMENT (mm)

LOAD(N)

3.81 mm

4.86 mm

5.74 mm

6.79 mm

7.32 mm

9.48 mm

IPNC-0.5

0 1 2 3 4 5 60

75

150

225

300

375

450

525

600

675

L

OAD(N)

DISPLACEMENT (mm)

3.89 mm

4.99 mm

5.98 mm

7.34 mm

9.69 mm

IPNC-3.0

0 4 8 12 16 20 24

0

75

150

225

300

375

450

DISPLACEMENT (mm)

LOAD(N)

1.87 mm

3.66 mm

4.67 mm

5.94 mm

6.61 mm

7.38 mm

ISPNC-0

0 2 4 6 8 10 12 14 16 18 200

50

100

150

200

250

300

350

400

450

500

550

Lo

ad(N)

Displacement (mm)

2.35 mm

3.87 mm

4.83 mm

6.65 mm

8.34 mm

ISPNC-1.0

Fig. 7. Self-similarity of loaddisplacement diagrams.

0

20

40

60

80

100

i-PP

IPNC-0

IPNC-0IPNC-1.0

IPNC-1.5

IPNC-2.0

IPNC-3.0

Ne

tsec

tions

tress

(N/mm2

)

Ligament length (mm)

(a)

3 4 5 6 7 8 9 10 1 2 3 4 5 6 7 8 9 10 110

10

20

30

40

50

60

70

80

90

100


Ne

tsec

tions

tress

(N/mm

2

)

ISPNC-0

ISPNC-0.5

ISPNC-1.0

ISPNC-1.5

ISPNC-2.0

ISPNC-3.0

(b)

5.

Fig. 8. Hills analysis plot: variation of net section stress (rn) with ligament length (l) (a) IPNC (b) ISPNC.

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the inner fracture process zone (IFPZ). On a critical look into thedata, it was observed that ISPNC-0 and ISPNC-2.0 showed we-max-ima andwe-minima respectively. In contrast, ISPNC-1.0 showed anintermediate value ofwethat is 60% lower than that of ISPNC-0.5and 40% higher than that of ISPNC-1.5. These observations imper-atively indicate the existence of a semi-ductile-to-tough-to-quasi-brittle transition in ISPNC in the composition range of 0.51.5 wt.% of MWCNT, unlike the failure transition involving defor-mation with plastic-flow to deformation without plastic-flow inIPNC. The semi-ductile, tough and quasi-brittle natures of therespective ISPNC compositions are characterized by the high-we/low-bwp, high-we/high-bwp and low-we/low-bwp combina-tions of the fracture mechanics parameters. Comprehensively such

a nature of the failure transition in ISPNC underlines the role of tac-ticity-defined modification of the polymer matrix on the fracturetoughness of MWCNT filled polymer nanocomposites. Furthermorethe fracture surface morphology of these compositions by SEMhave revealed a systematic transition in the nature of their failurecharacteristics and are discussed in a subsequent section.

3.6. Morphology and fracture toughness correlation

The observed significant and moderate enhancements in tough-ness due to tacticity differences in the stereo-complex PP matrix(ISPNC) and in the conventional (single form crystalline matrix)i-PP (IPNC) may be correlated to morphological parameters; qual-itatively with extent of dispersion and quantitatively with the

lamellar thickness (Llamellar) of the PP chains. Several authors haveattributed toughness and stiffness enhancements to dispersion and

distribution uniformity of nanotubes[32,45,46]. The length-scale(radius of gyration of polymer chains and nanotube dimensions)and time-scale (segmental relaxation of polymer and mobility ofnanotube within the framework of entangled polymer chains) cor-respondences in nanocomposites are theoretically well demon-strated and was reported that a fundamentally differenttoughening mechanism operates in these class of nano-structuredmaterials where the singular mobility of the nano-inclusion in thepolymer matrix under deformation defines the energy dissipationcapability of the material [47]. Interestingly the preceding argu-ment seems to be in agreement to the fact that the singular mobil-ity of MWCNT may get accentuated by the increase in the overallamorphous nature of the matrix. This is well in agreement to our

observations manifested in the form of significant reduction in Tgand reduction in crystallite size (as obtained from Scherrers equa-tion) upon blending of s-PP, i.e. upon rendering the nature of ma-trix-polymer stereo-complex. To further correlate themorphological parameters, composition specific lamellar thick-nesses (Llamellar) was estimated following the ThomsonGibbsequation and was analyzed vis--vis the fracture parameters weand bwpvalues. The ThomsonGibbs equation may be given as,

Llamellar 2rT0m

DHmT0mTm

12

whereris lamellar surface free energy (0.122 J m2), DHvis melt-ing enthalpy of lamellar with infinite thickness (192.28 106 -

J m3

), T0m is equilibrium melting temperature [48,49]. The

calculated lamellar thickness for the investigated IPNC and ISPNC

3 4 5 6 7 8 9 10

25

50

75

100

125

150

175

200

225

250

275

300

i-PP

IPNC-0

IPNC-0.5

IPNC-1.0

IPNC-1.5

IPNC-2.0

IPNC-3.0

Specificwork

offracture


(a)

1 2 3 4 5 6 7 8 9 10

50

100

150

200

250300

350

400

450

500

550

600

650


Specificworko

ffracture(N/mm)

ISPNC-0

ISPNC-0.5

ISPNC-1.0

ISPNC-1.5

ISPNC-2.0

ISPNC-3.0

(b)

(N/mm)

Fig. 9. Variation of specific work of fracture (w) with ligament length for (a) IPNC (b) ISPNC.

Fig. 10. Variation of plane stress fracture parameters with MWCNT content (a) essential work of fracture (EWF; we) and (b). non-essential work of fracture (N-EWF; bwp).



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compositions are shown inTable 5. It could be clearly observed thatthe we-essential work of fracture has a direct correspondence tolamellar thickness indicating that the resistance to stable crack ini-tiation has a direct correlation to lamellar thickness and crystallitearrangement[50]. However, the inverse correlation tobwpmay the-oretically be construed as lamellar break-up/destruction phenome-non leading to an increase in resistance to stable crack propagation.

The scenario may be simplistically explained as the availability ofmore of random polymer chains segments in the larger frame ofthe material architecture facilitating energy dissipation mechanismby promoting the singular mobility of the embedded MWCNT in theentangled polymer chain networks with substantially lesserhindrance.

3.7. Kinetics of crack propagation

The investigated nanocomposites showed ductile yielding,though the nature of the ductile response remained dependenton MWCNT content and therefore this observation encouragedthe authors to probe into the nature of such transitions in failuredynamics in real-time framework following the kinetics of crack

growth approach. In both type of nanocomposites the matrix-dom-inant ductile-yielding behavior has been observed to get alteredwith the state of dispersion/distribution of nanoscopic inclusionsand hence with their consequent effects on morphological changes.The determination of kinetics parameters of crack growth has beendone by using an optical video monitoring system in combinationwith digital image correlation (DIC) techniques. The real-timecrack growth has further been monitored with periodic recordingof crack extension (Da), crack tip opening displacement (CTOD,d) and CTOD rate (dd/dt) data at various deformation stages asshown inFig. 11. The systems i-PP (IPNC-0) and IPNC-2.0/IPNC-3.0 have undergone much faster crack extension than the otherIPNC systems while the slowest crack extension was registeredfor IPNC-0.5. In contrast a drastic retardation (>5-fold) of the crack

extension could be observed in ISPNC when compared to i-PP(IPNC-0), an observation solely attributed to tacticity induced ste-reo-complexity of the matrix. On incorporation of MWCNT toISPNC-0 matrix the crack extension rates have been observed toget further retarded, especially at larger time scales, with ISPNC-0.5 exhibiting the slowest crack extension rate. InterestinglyISPNC-2.0/ISPNC-3.0 has also been observed to undergo failure atmuch faster rate than ISPNC-0 matrix, an observation resemblingIPNC-2.0/IPNC-3.0. These observations corroborate the intrinsiccorrespondence between poor dispersion and faster crack exten-sion leading to catastrophic failure. The CTOD data, indicating theinherent blunting efficiency of material, of IPNC revealed a similartrend to that ofDashowing the lowest CTOD-rate of IPNC-0.5. Incontrast the CTOD-rate of ISPNC showed nearly identical values

irrespective of the extent of MWCNT incorporation. Since isolatedanalysis of the kinetic parameters may lead to discrepancies quan-titatively, the analysis of the fracture parameters has been carriedout in a cooperative fashion by analyzing the plots ofdd/dtversusDa.

The variations of CTOD-rate (dd/dt) with crack extension (Dafor IPNC and ISPNC are shown inFig. 11. It has been observed tha

barring i-PP and IPNC-0 all the investigated nanocomposites failedwithin Da 2.0 mm, i.e. at one third of the ligament lengthl 6 mm. Interestingly such a range ofDawhen compared to thedisplacement measured in the loaddisplacement diagrams fol 6 mm remains well beyond the load-maxima in the failure ofDENT specimens. However, the maximum crack extension for i-PP and IPNC-0 exceeds the generally observed extent ofDato fail-ure for the nanocomposites irrespective of tacticity modificationi.e. stereo-complexity. Further the dd/dthas been observed to beleveled off within dd/dt 0.04 mm/s. It must be mentioned thatthe leveling off limit ofdd/dtis little (1.5 times) higher than thestrain rate (e9 = 0.033 mm/s) at which the fracture experimentshave been conducted. Based on these observations, the kineticallydefined failure map of the nanocomposites may be divided intofour failure-regimes as indicated inFig. 11. The four macro-failureregimes and their associated deformation attributes may be sum-marized as given inTable 6.

It was observed that in IPNC and ISPNC systems the failure mapremained mostly confined in regime-II indicating the possibility oMWCNT induced toughening, though the extent of toughening andtheir underlying mechanistic attributes may not be easy to quantify. However, from the plots it is amply evident that in ISPNCthe CTOD-rate has been found to be more uniform as comparedto IPNC systems. Additionally the nature of clustering of the mea-sured data points in regime-II essentially indicates the inherentdependence of the macroscopic kinetic parameters on the morpho-logical attributes of the systems. For ISPNC the less random clus-tering of the data points may imply a better state of dispersionunlike their IPNC counterparts. A more comprehensive analysis

based on qualitative estimation of J-integral and tear modulus re-sponses may lead to further details of the influence of stereo-com-plexity (tacticity-defined blending) of matrix on fracture toughnessresponse of these nanocomposites, an aspect which will be fol-lowed up in future.

3.8. Fracture surface morphology by SEM

The fracture surface morphology responsible for the associatedmicro-deformation and crack propagation mechanisms of the i-PPand the stereo-complex PP blend (i-PP/s-PP: 70/30) i.e. ISPNC-0 asthe matrices and their MWCNT filled nanocomposites have beenevaluated by scanning electron microscopy and is shown in

0.00

0.02

0.04

0.06

0.08

0.10

0.12

Regime-IRegime-IV

Regime-IIIRegime-II

Steady state d

/dt for i-PP/ PP-g-MA

Steady state d

/dt for nanocomposites

CTODra

ted

/dt[mm

/s]

Crack extension a (mm)

i-PP

IPNC-0

IPNC-0.5

IPNC-1.0

IPNC-1.5

IPNC-2.0

IPNC-3.0Steady state d/dt for i-PP

0 1 2 3 4 5 6 0 1 2 3 4 5 6

0.00

0.02

0.04

0.06

0.08

0.10

0.12

Regime-IRegime-IV

Regime-III

Regime-II

Steady state d

/dt for

nanocomposites

Steady state d/dt for i-PP/ PP-g-MA

CTODra

ted

/dt[mm

/s]

Crack extension

a (mm)

IPNC-0

ISPNC-0

ISPNC-0.5

ISPNC-1.0

ISPNC-1.5

ISPNC-2.0

ISPNC-3.0

(a) (b)

Fig. 11. Kinetics of crack propagation of the nanocomposites crack tip opening displacement (CTOD) rate dependence on crack growth (a) IPNC nanocomposites (b) ISPNCnanocomposites.

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Fig. 10. It was clearly observed that IPNC-0 showed a homogeneoussurfacemorphologywhereasIPNC-0.5showeddistinctsignsofpeel-ing-off tendencies of fracture surface layers indicating the fractureresistance to be controlled more by crackpropagatingalongthe frac-tureplane.Theobservationisalsocorroboratedbythefactthatthereis a small decrease in the we(essential work of fracture process) ofIPNC-0.5 as compared to IPNC-0 as shown inFig. 12. Interestinglythe failure surface morphology of the stereo-complex blend matrixi.e., ISPNC-0, the surface showed signs of alternately patternedpeeled-off strips indicating the presence of a quasi-co-continuousmorphologydue to asymmetricstereo complexity.On incorporating0.5 wt.% of MWCNT i.e., the nanocomposite ISPNC-0.5 revealedintensive fibrillation process involving thin fibrils and fibriler stripsto be involved in the course of fracture under uni-axial tension. Onfurther increasing the MWCNT content to 1.0 wt.% i.e., in ISPNC-1.0, the nanocomposite showed relatively distinct and dense popu-

lationof fibrillated structuresspanningacross the crackalso referredas crack-bridging, highly stretched fibrils (0.10.3 lm) indicatingstrong interfacial interaction between PP and MWCNT due to ste-reo-complexity and very long fibrillar strips (210 lm) undergo-ing fragmentation to lead to smaller strips (24 lm) andfibrils(0.81.0 lm) at their tips prior to failure. On increasing theMWCNT content to 3.0 wt.%, as in ISPNC-3.0, the fracture surface re-vealed i-PP/MWCNT (80:20) masterbatch rich domains as the re-gions responsible for failure initiation. Such a failure mechanismmay be attributed to highly localized morphology with improperdispersion and distribution since the MWCNT rich region does notget infiltrated effectively by the excessively added i-PP or s-PP soas to lead to a state of real dilution. Hence the fracture surface inthese regions showed the relative absence of any signs of plastic-deformation/shear-yielding prior to failure causing quasi-brittle/ductility-arrested failure characteristics.

Table 6

The macro-failure regimes and their associated deformation.

Regimes dd/dt Da Deformation attributes

Regime-I High (>0.04 mm/s) Low (


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4. Conclusions

A new phenomenon of matrix stereo-complexity enhancing thefracture toughness by >330% and improvement of the extent of dis-persion of MWCNT in polypropylene has been observed. Dispersionmorphology was quantitatively explained in terms of the ratio offree space length to infiltrated free space lengths. Such morpholog-

ical consequences leading to a decrease in Tgof MWCNT filled ste-reo-complex or tacticity modified blends of polypropylene relativeto the nanocomposites based on isotactic-only matrix by 28 Chas been observed. The fracture toughness performances of thesenanocomposites have been observed to be significantly affectedas estimated by essential work of fracture approach. A semi-duc-tile-to-tough-to-quasi-brittle transition in ISPNC in the composi-tion range of 0.51.5 wt.% of MWCNT has been observed thatwas primarily attributed to a corresponding transition in the nat-ure of deformation occurring with (in ISPNC) and without (in IPNC)plastic flows respectively. A set of distinctly evolved and novel cri-teria for the semi-ductile, tough and quasi-brittle natures ISPNCcompositions have emerged that are characterized by high-we/low-bwp, high-we/high-bwp and low-we/low-bwp combina-tions of the fracture mechanics parameters respectively. A fivefoldretardation of the crack extension in tacticity modified relative tothe unmodified virgin isotactic PP matrix could be observed. Theenhanced dispersion in stereocomplex PP matrix based nanocom-posite have shown a direct bearing on the CTOD rate as reflectedin the blunting response. Fracture surface morphology revealedsignificant crack-bridging, fibrils stretching across a growing crackaccompanied by extensive fibrillation and formation of microstripsgetting terminally fragmented into micro-fibrils as the possibletoughening mechanism in stereocomplex PP matrix based MWCNTfilled nanocomposites.

Our study demonstrates a new pathway for toughening of car-bon nanotube based thermoplastic polymer nanocomposites byblending of two different stereo-regular forms of the chemicallyidentical polymer matrix. The conceptual validation has success-fully been established with regard to PP/MWCNT nanocomposites,where remarkable toughness enhancement has been corroboratedby significant (nearly fivefold) retardation of the crack-growthphenomenon assisted by efficient crack-blunting mechanism.

Acknowledgements

This paper is dedicated to late Prof. Roland Weidisch. Authorsgratefully acknowledge the financial support, vide Grant No.SR/S3/ME/0010/2008 by Department of Science and Technology NewDelhi, India.

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