documentn5

Tensile performance improvement of low nanoparticleslled-polypropylene composites

Chun Lei Wua, Ming Qiu Zhangb,*, Min Zhi Rongb, Klaus Friedrichc

aKey Laboratory for Polymeric Composite and Functional Materials of Ministry of Education, Zhongshan University,

Guangzhou 510275, PR ChinabMaterials Science Institute, Zhongshan University, Guangzhou 510275, PR China

cInstitute for Composite Materials (IVW), University of Kaiserslautern, D-67663 Kaiserslautern, Germany

Received 13 November 2001; received in revised form 19 March 2002; accepted 5 April 2002

Abstract

It was found beforehand that low nanoparticles loaded polymer composites with improved mechanical performance can be pre-pared by conventional compounding technique in which the nanoparticles are pre-grafted by some polymers using irradiation. Toexamine the applicability of the approach, a tougher polypropylene (PP) was compounded with nano-silica by industrial-scale twinscrew extruder and injection molding machine in the present work. The results of tensile tests indicated that the nanoparticles can

simultaneously provide PP with stiening, strengthening and toughening eects at a rather low ller content (typically 0.5% byvolume). The presence of grafting polymers on the nanoparticles improves the tailorability of the composites. Due to the visco-elastic nature of the matrix and the grafting polymers, the tensile performance of the composites lled with untreated and treated

nanoparticles is highly dependent on loading rate. With increasing the crosshead speed for the tensile tests, the dominant failuremode changed from plastic yielding of the matrix to brittle cleavage.# 2002 Elsevier Science Ltd. All rights reserved.

Keywords: A. Particle-reinforced composites; B. Mechanical properties; B. Surface treatments; Nanoparticles

1. Introduction

Mineral llers are added to polymers in commercialproduction primarily for the reasons of cost reductionand stiness improvement [1,2]. Although most studiesdealing with the modication of semi-crystalline poly-mers with inorganic particulates reported embrittlingeects by comparing ultimate elongation and impactstrength of composite materials with those of unlledresins [35], some researchers showed the enhancementof toughness in rigid particles lled polypropylene [6,7]and polyethylene [8,9].It is worth noting that in the case of micrometer-sized

particulates, high ller content (typically higher than20% by volume) is generally required to bring theabove-stated positive eects of the llers into play. Thiswould detrimentally aect some important properties ofthe matrix polymers such as processability, appearance,

density and ageing performance. Therefore, a compositewith improved performance and low particle con-centration is highly desired. With regard to this, thenewly developed nanocomposites would be competitivecandidates.The extremely high surface area is one of the most

attractive characteristics of nanoparticles because itfacilitates creating a great amount of interphase in acomposite. Introduction of nanoparticles into a polymerchanges the intermolecular interaction of the matrix[10]. As estimated by Reynaud et al. [11], an interphase1 nm thick represents roughly 0.3% of the total volumeof polymer in the case of microparticle lled composites,whereas it can reach 30% of the total volume in the caseof nanocomposites. That is, the non-negligible con-tribution made by the interphase provides diverse pos-sibilities of performance tailoring, and is able toinuence the properties of matrices to a much greaterextent under a rather low nano-ller loading. The cruxof the matter lies in that how to well distribute nano-particles over a polymer matrix and how to improvenanoparticles/matrix interaction.

0266-3538/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved.PI I : S0266-3538(02 )00079-9

Composites Science and Technology 62 (2002) 13271340

www.elsevier.com/locate/compscitech

* Corresponding author. Tel.: +86-20-8403-6576; fax: +86-20-

8403-6564.

E-mail address: [email protected] (M.Q. Zhang).

From a practical point of view, dispersive mixing inpreparing polymer based particulate composites hasimportant technical meaning. However, a homogenousdispersion of nanoparticles in a polymer is very dicultby using the existing compounding techniques due tothe strong tendency of the ne particles to agglomerateand the high melt viscosity of the matrix. In many cases,the so-called nanoparticles lled polymers contain anumber of loosened clusters of nanoparticles, asdemonstrated by the work of Jana and Jain [12] dealingwith untreated nanosilica/polyethersulphone composites.When the composites are subjected to force, the nano-particle agglomerates can be split easily and a prematurefailure of the materials would thus take place [1215].To overcome this dilemma and to give full play to

nanoparticles, we used irradiation grafting polymeri-zation method to modify nanoparticles rst, and thenthe treated particles were mechanically mixed with apolymer as usual [16]. Owing to the low molecularweight nature, the grafting monomers can penetrateinto the agglomerated nanoparticles easily and reactwith the activated sites of the nanoparticles inside aswell as outside the agglomerates [17]. As a result, thefollowing eects can be obtained [18]: (i) Hydro-phobicity of the nanoparticles is increased, facilitatingthe ller/matrix miscibility; (ii) ller/matrix interactionis enhanced through the entanglement between thegrafting polymer and the polymer matrix; (iii) thenanoparticle agglomerates become stronger becausethey turn into a nanocomposite microstructure com-prising the nanoparticles and the grafted, homo-polymerized polymer; (iv) the interfacial characteristicsbetween the treated nanoparticles and the matrix poly-mer can be tailored by changing the species of thegrafting monomers and the grafting conditions. In thiscontext, a uniform dispersion of nanoparticles in thematrix might no longer be critical.Mechanical testing of polypropylene lled with nano-

SiO2 [18] and nano-CaCO3 [19] demonstrated the feasi-bility of the above approach. Only a small amount ofmodied nanoparticles (typically less than 3% byvolume) can eectively improve modulus, strength,toughness and thermal deformation temperature of thematrix polymer. Such an improvement in overall prop-erties of polymers can scarcely be observed in conven-tional microparticulate composites. It was found thatthe deformation habit but not the crystallization char-acteristics of the matrix polymer remarkably changewith the addition of the treated nanoparticles. Toexplain the specic inuence generated by the nano-particles at low-ller loading regime, a double percola-tion of stress volumes, characterized by the appearanceof connected shear yielded networks throughout thecomposite, was proposed [20].Considering that the polypropylene used in our pre-

vious works [1820] is a brittle type and the composites

were prepared with a lab-scale single screw extruder andcompression molding, the results might not have ade-quate applicability. Therefore, a commercial poly-propylene with higher toughness was compounded withnano-silica by means of industrial-scale twin screwextruder and injection molding machine in the currentwork. Tensile performance and fractured surfaces areanalyzed as a function of particulate treatment, llercontent and crosshead speed to reveal the structureproperty relationships of the composites, and themechanical role of the nanoparticles as well.

2. Experimental

2.1. Materials

Isotactic polypropylene (PP) homopolymer T30S1

was supplied by Qilu Petrochemical Industrial Co.,China. It has a melt ow index of 3.2 g/10 min (2.16 kgat 230 C). Fumed silica with an average primary parti-cle size of 15 nm and a specic surface area of 374 m2/gwas produced by Shenyang Chemical Engineering Ltd.,China. Commercial monomers, styrene and ethyl acry-late, were used as grafting monomers without furtherpurication.

2.2. Irradiation grafting of nano-SiO2

Modication of nano-silica proceeded according tothe following steps. The nanoparticles were pretreatedat 140 C under vacuum for 6 h to eliminate possibleabsorbed water on the surface of the particles. Then amixture of nanoparticles/monomer (100/20 by weight)and a certain amount of n-hexane was irradiated by60Co g-ray under atmosphere at room temperature.After exposure to a dose of 4 Mrad, the solvent wasrecovered, and the dried residual powder was availablefor the subsequent compounding.

2.3. Characterization of the irradiation grafted products

To evaluate the results of grafting and to characterizethe grafted nanoparticles, the grafting polymer and thehomopolymer, which were generated during the irra-diation polymerization of the monomers, should beseparated. For this purpose, a certain amount of theirradiation products were extracted by benzene in aSoxhlet apparatus for 36 h. In this way the homo-polymer was isolated. The residual material was thendried in vacuum at 80 C until a constant weight wasreached. By using a Shimadzu TA-50 thermogravimetre(TG) and a Bruker Equinox 55 Fourier transforminfrared spectroscope (FTIR), the weight of the gratingpolymer and the chemical structure of the modiednanoparticles were characterized, respectively. To further

1328 C.L. Wu et al. / Composites Science and Technology 62 (2002) 13271340

separate the grafting polymer from the treated nano-particles, nano-silica accompanied with the unex-tractable grafting polymer was immersed in 20% HFsolution for 72 h to remove the inorganic particles. Themolecular weights of the grafting and the homo-polymerized polymers were determined by a Waters 991gel permeation chromatography (GPC), with tetra-hydrofuran as the solvent.To observe the morphologies of the nanoparticles,

untreated SiO2 and grafted SiO2 (without homo-polymer) were added into ethanol and toluene to pre-pare 0.001 g/ml solutions, respectively. With the aid ofsonication for 30 min, the solutions were transferred tocopper gauzes by droppers. After the evaporation of thesolvents, a JEM-100CXZ transmission electron micro-scopy (TEM) was used to examine the appearance ofthe particles.

2.4. Composites preparation and characterization

The nanoparticles were rstly compounded with PP(1:2 by weight) using an X(S)R-160 two-roll mill at195 C to produce composite masterbatch. Then, themasterbatch was mixed with neat PP to dilute the llerloading to desired values through an SHJN-25 twin-screw extruder at 210230 C. The rotation speed of theextruder was set to 180 rpm. Finally, the resultant pel-lets were molded into dog-bone-shaped tensile bars(ASTM D63897 Type IV specimen) with a CJ150MZinjection-molding machine at 215 C.Room temperature tensile testing of the composites

was conducted on a Hounseld-5KN universal testingmachine at crosshead speeds of 10, 30, 50, 100 and 500mm/min, respectively. Five samples were tested for eachcase. The fractured surfaces of the samples wereobserved with a Hitachi S-520 scanning electron micro-scope (SEM) at an accelerating voltage of 20 kV.

3. Results and discussion

3.1. Eect of irradiation grafting polymerization onnano-SiO2

Since the present work aims to study the eect ofmodied nano-silica on the mechanical behavior of PPcomposites, variation in the chemical structure of theparticles should be known at the very beginning of thediscussion. FTIR spectra of untreated and treated nano-silica are shown in Fig. 1. To eliminate the inuence ofhomopolymers, both polystyrene-grafted nano-SiO2(SiO2-g-PS) and polyethyl acrylate-grafted nano-SiO2(SiO2-g-PEA) used for the FTIR examinations wereseparated from the homopolymers in advance. In com-parison with the spectrum of SiO2 as-received, theadsorptions at 690, 1460 and 2960 cm1 appearing in

the spectrum of SiO2-g-PS represent the bending modeof CH in benzene rings and the stretching modes ofCC and CH, respectively. In addition, the band at1725 cm1 in the spectrum of SiO2-g-PEA indicates theexistence of carbonyl groups. These prove that poly-styrene (PS) and polyethyl acrylate (PEA) have beenchemically connected to the surface of nano-silica duringthe irradiation polymerization processes as expected.On the basis of above qualitative analysis, the quan-

titative results of the grafting polymerization on nano-silica are given in Table 1. It is seen that when otherconditions being equal, the percentage grafting and thegrafting eciency of styrene onto the nanoparticles aremuch higher than those of ethyl acrylate, while themonomer conversion of styrene is lower than that ofethyl acrylate. This reects the dierence in reactivefeature between the monomers. In general, room tem-perature irradiation grating polymerization onto inor-ganic particles is controlled by the mechanism of freeradical polymerization. Under the same irradiationdose, acrylic monomers would generate much moreradicals than styrene, and moreover, ethyl acrylateradicals have higher activity than styrene radicals [21].Therefore, the conversion of ethyl acrylate is superior tothat of styrene, leading to higher homopolymer fractionof PEA than PS. On the other hand, partial surface ofnanosilica might be connected with ethyl acrylate in thesolvent through hydrogen bonding prior to the irradia-tion processing. When the composite system is exposedto the irradiation, radicals can be formed on both thenanoparticles and the ethyl acrylate molecules con-nected to the particles. The latter would result in eithergrafting polymerization or homopolymerization. Thatis, the amount of grafted PEA has to be lower as com-pared with the styrene/silica mixture where no chemicalconnection between the monomers and the particles isestablished before the irradiation. It should be respon-sible for the dierence in percentage grafting betweenethyl acrylate/silica and styrene/silica.

Fig. 1. FTIR spectra of untreated SiO2 and grafted SiO2.

C.L. Wu et al. / Composites Science and Technology 62 (2002) 13271340 1329

By further examining the data listed in Table 1, it canbe found that the grafting polymers attached to thenanoparticles possess higher molecular weights andbroader molecular weight distributions than the homo-polymers. Similar phenomena were also reported byFukano and Kageyama [22] when they studied radiationgrafting of styrene onto silica-gel. These can be attri-buted to the characteristics of the grafting reaction.That is, the reaction was a typical heterogeneous onebecause the activated sites were created by irradiationon the nanoparticles and the chain growth of the graft-ing polymers had to proceed in solidliquid state. Themobility of the growing chains was thus worse as com-pared with a homogeneous reaction in liquid whereboth ends of the macromolecular chains can movefreely. As a result, the probability of chain terminationbetween the radicals became relatively dicult, leadingto higher molecular weight of the grafting polymers. Inaddition, because the radicals are not simultaneouslyproduced on the nanoparticles, the grafting polymersformed at earlier stage would wrap the surface of thenanoparticles and reduce the probability of collisionbetween the monomers and the radicals formed on theparticles at latter stage. This accounts for the highermolecular weight polydispersity indexes of the gratingpolymers.Morphologies of the nanoparticles before and after

grafting polymerization are illustrated in Fig. 2. Achain-like branched structure of the agglomerated SiO2particles [17] can be observed in the solutions. Thesmallest perceivable units are approximately 15 nm indiameter in the case of untreated nanoparticles[Fig. 2(a)]. When grafting polymers are introduced ontothe particles, the sizes of the agglomerates become lar-ger and the edges are no longer clearly discernible[Fig. 2(b, c)]. Such a change demonstrates the role of thegrafting polymer, i.e. separating and connecting thenanoparticles. To estimate the thickness of the polymerlayer adhered to the particles, the data of SiO2-g-PS isused as an example, i.e. percent grafting=4.64%, density

Table 1

Results of irradiation grafting polymerization onto nano-silica

Monomer Monomer

conversiona

(%)

Percent

graftingb

(%)

Grafting

eciencyc

(%)

MWd of

grafting

polymer

(104)

MW of

homopolymer

(104)

Styrene 52.8 4.64 43.9 1.3 (de=1.70) 1.1 (d=1.46)

Ethyl acrylate 93.5 1.56 10.5 7.1 (d=2.13) 1.1 (d=1.33)

a Monomer conversion=weight of polymer/weight of monomer.b Percent grafting=weight of grafting polymer/weight of nano-SiO2.c Grafting eciency=weight of grafting polymer/weights of grafting polymer and homopolymer.d MW=weight average molecular weight.e d=Molecular weight polydispersity index.

Fig. 2. TEM microphotos (magnication=105) of (a) SiO2 as-

received, (b) SiO2-g-PS, and (c) SiO2-g-PEA in solvents. In the latter

two specimens, homopolymers were removed in advance.


of PS=1.05 g/cm3, specic surface area of the silica=374m2/g. Supposing a complete coverage on each silica par-ticles, the thickness of the grafting PS should equal to0.0464/(374 m2/g1.05 g/cm3)0.12 nm. Evidently, thisis a reasonable value as evidenced by the TEM photosin Fig. 2. Due to the low percent grafting and the thingrafting polymer layer, there are still many unreactedhydroxyl groups on the surface of nano-SiO2, which isresponsible for the appearance of the larger agglomeratesof the grafted particles in the solvent.

3.2. Tensile properties of the composites

Fig. 3 shows the results of tensile testing of PP rein-forced by nano-silica as a function of ller content,determined at a moderate crosshead speed of 50 mm/min. Although both the treated and the untreatednanoparticles can impart the high stiness of the llersto the matrix polymer as expected, the compositesincorporated with the modied particles exhibit lowermodulus over the whole range of ller loading of inter-ests [Fig. 3(a)]. Usually the capability of compositeinterface to transfer elastic deformation depends to agreat extent upon the interfacial stiness and staticadhesion strength [23,24]. A high interfacial stinesscorresponds to a high composite modulus. Since thegrafting polymers and the homopolymers introducedonto the nanoparticles form a relatively compliantinterlayer at the particles/matrix interface, the high

stiness of the particles has to be masked under the lowstress level [9] and the composites have to show lowermodulus as compared with the case of untreated SiO2composites. With a rise in ller content, the increasedamount of grafting polymers further increases the mod-ulus mismatching of the ller and the matrix, and redu-ces the stiening eect of SiO2, leading to the drop inYoungs moduli of the composites at high ller regime.In addition, due to the higher rigidity of PS moleculesthan PEA, the interfacial elastic stress is less ecientlytransferred in SiO2-g-PEA/PP composites than in SiO2-g-PS/PP, especially when the fraction of the modiednanoparticles is high.In contrast, an approximately linear composition

dependence of tensile modulus is perceived in the com-posites lled with SiO2 as-received [Fig. 3(a)]. This canbe interpreted as the absence of a soft interphase andthe appearance of larger agglomerates of the nano-particles in the matrix as a result of poor ller/matrixmiscibility. Under the same ller volume fraction [25],the latter eect would provide higher load carryingability within small strain range. The above results fac-tually reect the contradiction between the stieningeect of the rigid particles and the weakening eect ofthe soft interlayer.Fig. 3(b) gives the tensile strengths of the composites

versus SiO2 content. Clearly, both untreated and treatednanoparticles exhibit the strengthening ability. It is wellknown that the tensile strength of a particulate composite

Fig. 3. Tensile properties of PP composites as a function of nano-SiO2 volume fraction: (a) Youngs modulus, (b) tensile strength, (c) elongation to

break, and (d) area under stressstrain curve. Crosshead speed=50 mm/min.


Fig. 5. Youngs modulus of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.

Fig. 4. Tensile properties of PP composites as a function of crosshead speed: (a) Youngs modulus, (b) tensile strength, (c) elongation to break, and

(d) area under stressstrain curve.


is usually reduced with ller content following a powerlaw in the case of a poor ller/matrix bonding [26,27].That means, the strength of the composite cannot begreater than that of the unlled version because the llerparticles do not bear any fraction of the external load.This contradicts the results shown in Fig. 3(b). In fact,when bonding between llers and matrix is strongenough, the tensile yield strength of a particulate com-posite can be higher than that of the matrix polymer[28,29]. Although these models were developed based onthe cases of microsized particulate composites, they arestill valid for the explanation of the composites lledwith nanoparticles [20]. Therefore, the extremely eec-tive improvement of tensile strength of the compositeswith grafted nano-SiO2 should result from chain inter-diusion and entanglement between the macro-molecules of the grafting polymers and the matrix. It isworth noting that when the amount of the graftednanoparticles is increased (e.g. >0.5 vol%), the contentof compliant PEA chains is raised accordingly in theinterlayer and the interfacial stress transfer eciencyhas to be decreased. This accounts for the relatively lowstrength of SiO2-g-PEA/PP composites at higher parti-culate loading.

By comparing Fig. 3(b) with the data of Ref. [18], themost distinct dierence lies in the results of untreatedSiO2/PP. That is, for the composites prepared by a lab-scale single screw extruder and compression molding,the addition of untreated SiO2 lowers the tensilestrength of PP in the lower loading region but then leadsto a slight increase in strength when the particle fractionreaches 4.68 vol% [18]. With respect to the SiO2/PPcomposites of the present work, which were manu-factured through twin screw extrusion and injectionmolding, a continuous increase in the strength with SiO2content is detected [Fig. 3(b)]. So far as we know, cor-otating twin screw extruders are able to provide moresucient homogenization in comparison with singlescrew extruders. The above dierence manifests that theimportant role of even distribution of untreated nano-particles in the composites. For SiO2 as-received, themore particles are exposed to the matrix polymer, themore possibly the interaction between the particles andthe matrix can be enhanced. So, an improved homo-geneity of the untreated SiO2/PP composites would cer-tainly be benecial to the stress transfer.This again shows the signicance of grafting modi-

cation of the nanoparticles, which reduces the sensitivity

Fig. 6. Tensile strength of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.


of composites strength performance to the dispersionstate of the particles. As the grafted nanoparticleagglomerates turn into a nonocomposite microstructureconsisting of the particles and the grafted, homo-polymerized secondary polymer [18], they are broughtinto play as an integral when the composites are sub-jected to the applied force [20]. The situation is com-pletely dierent from the composites with untreatednanoparticles, in which agglomerated particles have tobe deagglomerated as much as possible to reduce theprobability of premature failure. As a result, a uniformdispersion of the nanoparticles in the matrix is abso-lutely necessary for obtaining the reinforcing eect incase untreated nanoparticles are used.Failure strain can partially assess the rupture beha-

vior of a composite material. The incorporation of par-ticulate llers usually results in a decrease in thisparameter regardless of the interfacial adhesion [26]. Itis true even in the system exhibiting impressive impacttoughness improvement with the addition of mineralllers [8]. However, the plots shown in Fig. 3(c)demonstrate that the values of elongation-to-break of

PP can be signicantly increased by using nano-SiO2,,implying a failure mechanism dierent from thoseinvolved in conventional composites. Comparatively,SiO2-g-PEA is able to provide a stable improvementover the entire ller content range of interests. Thereduction in elongation-to-break of the composite lledwith untreated SiO2 suggests that the llers cause areduction in matrix deformation due to an introductionof mechanical restrains. In contrast, the improvement ofelongation-to-break with the incorporation of the graf-ted nanoparticles is a result of interfacial viscoelasticdeformation and matrix yielding. Evidently, the graftedPEA makes more eective contribution.The area under the tensile stressstrain curve can

more reasonably characterize the toughness potential ofthe composites than elongation-to-break under statictensile loading conditions [30]. As conrmed byFig. 3(d), the grafted nanoparticles indeed improve theductility of PP at a silica content as low as 0.5 vol%. Theresults are somewhat opposite to those observed in Ref.[18], which reports a deteriorated eect of grafting PEAon the tensile behavior of the composites. Considering

Fig. 7. Elongation to break of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.


that SiO2-g-PEA employed in the present work pos-sesses almost the same percent grafting and graftingeciency as that in Ref. [18], it can be concluded thatthe molecular entanglement between the grafting PEAand matrix PP was not suciently formed during meltcompounding when the single screw extruder was used.As a result, localized plastic deformation or matrixdrawing cannot be eciently induced by SiO2-g-PEA asin the composites prepared by a twin screw extruder.In the case of untreated nano-SiO2, the areas under

the tensile stress-strain curves of the composites athigher ller content regime are lower than that of theneat PP. It implies that the short range interaction atSiO2/PP interface is not good at inducing plastic defor-mation of the matrix polymer. Due to the encounter ofthe propagating neck with a larger agglomerate, parti-cularly in the case of higher ller content, nal failure ofthe composites might be initiated easily [8]. Obviously,the resultant embrittling eect can be prevented by theapplication of grafted nanoparticles.Since polymer composites maintain the viscoelasticity of

polymers, the dependence of tensile properties on cross-head speed should be known for engineering purposes. As

can be seen from Fig. 4(a), a linear increase in Youngsmodulus with increasing crosshead speed is valid for allthe specimens. The presence of silica leads to the valuesof modulus being higher than that for the unlled PPand being slightly less speed dependent. These observa-tions are exactly as expected.The values of tensile strength of the materials are

plotted as a function of crosshead speed in Fig. 4(b).With a rise in crosshead speed, although the strengthsincrease in principle, there is a signicant transition inthe slope. Accordingly, the rst derivation of thestrength in Fig. 4(b) with respect to crosshead speedwould yield a peak between 50 and 100 mm/min, whichclearly corresponds to an energy-activated process intensile fracture of the materials. Further eorts shouldbe made to understand whether the Eyrings theory thatconsiders the eect of an applied stress is to reduce theheight of a potential energy barrier [31,32] is applicableto the present systems.Both elongation-to-break and area under the tensile

stress-strain curve have similar dependence on crossheadspeed [Fig. 4(c) and (d)]. With a rise in the speed, adrastic decrease of the two parameters is followed by a

Fig. 8. Area under stressstrain curve of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.


gradual reduction. SiO2-g-PEA/PP is able to keep itsductility superior to other systems when crossheadspeed is slower than 100 mm/min. As compared with theplots proles in Fig. 4(b), it can be deduced that dier-ent failure mechanisms take eects when crossheadspeed is faster or slower than 100 mm/min. Vu-Khanhand Denault found that the dynamic fracture toughnessof glass-ake/PP composites sharply decreases withimpact speed [33]. It was speculated that due to the lowthermal conductivity of the composite, the relaxationprocess in the matrix led to an increase in temperatureat the crack tip. The temperature rise caused a decreasein the fracture toughness with loading rate. It seemstheir analysis can also explain the experimental datashown in Fig. 4(c) and (d).To have a more comprehensive understanding of the

interdependence of tensile properties of the compositeson ller content and crosshead speed, three-dimensionaldiagrams are drawn in Figs. 58. In the case of a cross-head speed of 10 mm/min, Youngs moduli of SiO2-g-PS/PP and SiO2-g-PEA/PP composites increase with llercontent and then decrease (Fig. 5). When crossheadspeed is raised, the aforesaid decreasing trend of moduliis gradually replaced by a slight increase or a plateau.

This is dierent from the performance of SiO2/PP,which exhibits a continuous increase in the stiness withsilica content at each crosshead speed investigated. Evi-dently, the viscoelastic nature of the interphase due tothe appearance of grating polymers in SiO2-g-PS/PPand SiO2-g-PEA/PP composites should be responsiblefor the distinct behavior.The most obvious characteristics of Fig. 6 is that the

strengths measured at 100 and 500 mm/min are muchhigher than those obtained at a slower crosshead speed.As suggested previously, it should be indicative of achange in failure modes due to the dierent viscoelasticresponses as found in conventional polymers. For SiO2-g-PEA/PP composites, the addition of the graftednanoparticles used to slightly decrease the strength ofPP at crosshead speeds of 100 and 500 mm/min[Fig. 6(c)]. It means that the grafting PEA moleculesbecome less ecient to transfer stress under high strainrate. In contrast, 0.5 vol% of SiO2-g-PS can still providethe reinforcing eect at 500 mm/min [Fig. 6(b)].Figs. 7 and 8 illustrate the elongation-to-break and

the area under tensile stressstrain curve of the compo-sites as a function of silica fraction and crosshead speed.The untreated SiO2/PP composites exhibit performance

Fig. 9. SEM graphs of tensile fractured surface of: (a) neat PP; (b) and (c) SiO2/PP (SiO2 content=0.86 vol%); (d) SiO2-g-PS/PP (SiO2 con-

tent=1.06 vol%); (e) SiO2-g-PEA/PP (SiO2 content=0.82 vol%). Crosshead speed=50 mm/min.


dierent from the grafted SiO2 composites especially in thecase of higher particle content. At a silica volume fractionof about 2.7 vol%, for example, the elongation-to-breakand the area under tensile stress-strain curve of untreatedSiO2/PP are rather small and nearly independent of cross-head speed in comparison with the treated SiO2/PPcomposites. This should be interpreted as that splittingof the large nanoparticle agglomerates accumulated in

the matrix due to the fact that the increased particlecontent governs the failure process.By examining the curves of SiO2-g-PS/PP and SiO2-g-

PEA/PP determined at dierent crosshead speeds(Figs. 7 and 8), it can be found that the tougheningeect exerted by the modied particles becomesremarkable only at moderate speeds. The improvedelongation-to-break and area under tensile stressstrain

Fig. 10. SEM graphs of tensile fractured surface of: (a) and (b) SiO2/PP (SiO2 content=2.74 vol%); (c) and (d) SiO2-g-PS/PP (SiO2 content=2.75

vol%); (e), (f) and (g) SiO2-g-PEA/PP (SiO2 content=2.75 vol%). (g) was taken from the side face of the specimen within the stress-whitened neck

zone. Crosshead speed=50 mm/min.


curve represent the improved deformation ability of thecomposites in relation to plastic stretching of the matrixpolymer induced by the grafted nanoparticles. In addi-tion, the entanglement between the grafting polymersand the matrix polymer is also viscoelastic in nature.

3.3. Microscopic observation of fractured surface

To have clear images of the failure patterns of thecomposites under tension, SEM fractographs of thespecimens with dierent ller contents tested at dierentcrosshead speeds are discussed hereinafter. Fig. 9 showsthe tensile fractured surfaces of neat PP and the com-posites with relatively low silica fraction generated at acrosshead speed of 50 mm/min. The unlled PP has arelatively smooth fractured surface in association withterraced markings [Fig. 9(a)], indicating weak resistanceto crack propagation. In the case of SiO2/PP (Vf=0.86vol%), the fractured surface becomes rougher but thetraces of plastic deformation are still less [Fig. 9(b)].Many silica agglomerates (41 mm in size) are dispersedin the matrix without clear signs of stretching of thesurrounding matrix. Occasionally, elongated matrixpolymer can be found around silica agglomerates insidea large cavity on the composites surface, as highlightedby an arrow in Fig. 9(c). These demonstrate not only theinsuciently interfacial interaction between the particlesand the matrix, but also the poor toughening capabilityof the composites. For SiO2-g-PS/PP and SiO2-g-PEA/PP composites, the fractured surfaces are full of exten-sive matrix brils [Fig. 9(d) and (e)]. Therefore, it can beevidenced that the grafting polymers on the nano-particles enhance the interfacial interaction and the dis-sipation of energy through matrix stretching.When the content of nanosilica approaches around 2.7

vol%, the morphologies of the composites fractured sur-faces become somewhat dierent (Fig. 10). A number ofcavities appear on the surface of SiO2/PP composites[Fig. 10(a)]. In fact, they are produced due to the debond-ing of the untreated particles, as illustrated by a magniedview [Fig. 10(b)]. In general, an increased content ofuntreated SiO2 would lead to larger agglomerates andhence greater probability of debonding due to the poorinterfacial adhesion. As there is not enough time for indu-cing matrix yielding after the extensive particles debond-ing, the matrix beside the cavities seems to be rather at[Fig. 10(a)]. This coincides with the reduction of toughnessof SiO2/PP at high SiO2 loading [Fig. 3(c) and (d)].When SiO2-g-PS is incorporated [Fig. 10(c)], con-

centric matrix-brillated circles around nanoparticleagglomerates (as indicated by the upper arrow) andvoids left as a result of agglomerated particles detach-ment (as indicated by the lower arrow) can be found onthe fractured surface. As suggested in Ref. [18], theappearance of the brillated matrix circles are probablythe result of a successive debonding of the modied

nanoparticles from the matrix accompanied by anunconstrained plastic stretching of the interparticulatematrix ligaments [Fig. 10(d)]. Such a deformation pro-cess would certainly consume more energy than thatdominated only by debonding as shown in Fig. 10(a)and (b). In the case of SiO2-g-PEA/PP, the concentricbrillar circles around nanoparticle agglomeratesemerge next to each other [Fig. 10(e)]. Besides, thematrix surrounding the agglomerated nanoparticles hasturn into plastically drawn brils [indicated by thearrow in Fig. 10(f)]. Evidently, the inherent exibility ofPEA has made important contribution. In accordancewith the model describing the double percolation of yiel-ded zones [20], these should result from the superpositionof stress volumes around the agglomerates and the nano-particles. It explains the cause for that SiO2-g-PEA/PP isstill able to maintain higher static ductility at a relativelyhigh nanoparticle concentration [Fig. 3(c) and (d)].Fig. 10(g) exhibits the SEM observation result of the sidesurface of the SiO2-g-PEA/PP tensile specimen. Elon-gated cavities can be seen around partially debondednanoparticle agglomerates (indicated by the arrows).Again, plastic ow of the bridging matrix is clear visible.On the other hand, Fig. 10(c), (e), (f) and (g) conrm the

estimation that the grafted nanoparticle agglomerateshave turned into a nonocomposite microstructure and

Fig. 11. SEM graphs of tensile fractured surface of: (a) neat PP;

(b) SiO2/PP (SiO2 content=0.86 vol%); (c) SiO2-g-PS/PP (SiO2 con-

tent=1.06 vol%); (d) SiO2-g-PEA/PP (SiO2 content=0.82 vol%).

Crosshead speed=10 mm/min.


are brought into play as an integral. In one sense, thegrafted nanoparticle agglomerates can be taken as thefolded polymer chains conguration because they areprovided with the capability of deforming and releasinglocally concentrated stress instead of simply splitting.Fig. 11 gives the fractographs of neat PP and the

composites with low SiO2 content tested at a crossheadspeed of 10 mm/min. They all have similar appearancescharacterized by ductile failure except some ne micro-brils on the surface of SiO2-g-PS/PP and SiO2-g-PEA/PP. In comparison with the images taken at 50 mm/min(Fig. 9), no nanoparticles agglomerates are observedprobably because of the shields of the highly elongatedmatrix and the low ller content as well.It is worth noting that the fracture modes of the same

materials can be changed when a higher testing speed isapplied (Fig. 12). The neat PP shows cleavage fracturefeature under the crosshead speed of 500 mm/min[Fig. 12(a)]. The striation structure resulting from thejoining of dierent fractured planes on the surface ofSiO2/PP composites demonstrates that the particleshave little resistance to the crack propagation[Fig. 12(b)]. Similarly, the grafted nanoparticlesagglomerates cannot induce eective matrix yielding on alarge scale [Fig. 12(c) and (d)]. The mild proles of thedeformation circles [as indicated by the arrows inFig. 12(c) and (d)] suggest a low plastic deformation level.

By summarizing the data shown in Fig. 4 and thefracture morphology of Figs. 11 and 12, it can beknown that the tensile performance of the composites isclose to that of neat matrix polymer when the specimensare tested under low or high tensile speeds. Neither theuntreated nor the treated particles can take eects underthese circumstances. This is particularly true in the caseof low silica fraction.

4. Conclusions

Based on the above results and discussions, the fol-lowing statements can be drawn.(1) The addition of nanoparticles into PP can bring in

both reinforcing and toughening eects at ller contentas low as 0.5 vol%. Such a simultaneous improvementin modulus, strength and elongation-to-break is hard tobe observed in conventional microsized particulatecomposites.(2) Modication of nanosilica by means of grafting

polymerization helps to provide the composites withbalanced performance. In addition, dierent species ofthe grafting monomers result in dierent interfacialinteractions and dierent ultimate properties of thecomposites.(3) With respect to the manufacturing aspect, disper-

sion homogeneity of the composites lled withuntreated nanoparticles is critical, while it is not neces-sarily realized in the case of grafted nanoparticles.(4) As compared with single screw extruder, twin

screw extruder can further decrease the amount of thenanoparticles needed for the composites performanceenhancement.(5) The relative increment of the areas under the ten-

sile stress-strain curves of the current composites issimilar to the values reported by Ref. [18], althoughthere is a signicant dierence in ductility between thePP used in the two works. This somewhat contradictsthe results of Ren and co-workers, who found that atougher PP would gain more remarkable improvementof fracture toughness with the addition of nanoparticles[34]. It means that continuous eort should be paid tounderstand the role of the matrix toughness in nano-particles composites.(6) Owing to the viscoelastic nature of the grafting poly-

mers, the inuence of the modied nanoparticles on thetensile properties of PP is also a function of loading speed.

Acknowledgements

The authors are grateful to the support of theDeutsche Forschungsgemeinschaft (DFG FR675/401)for the cooperation between the German and Chineseinstitutes on the topic of nanocomposites. Further thanks

Fig. 12. SEM graphs of tensile fractured surface of: (a) neat PP;

(b) SiO2/PP (SiO2 content=0.86 vol%); (c) SiO2-g-PS/PP (SiO2 con-

tent=1.06 vol%); (d) SiO2-g-PEA/PP (SiO2 content=0.82 vol%).

Crosshead speed=500 mm/min.


are due to the National Natural Science Foundation ofChina (Grant: 50133020), the Key Program of the Minis-try of Education of China (Grant: 99198), the Team Pro-ject of the Natural Science Foundation of Guangdong,China, the Natural Science Foundation of Guangdong,China (Grant: 990277), and the Key Program of theScience and Technology Department of Guangdong,China (Grant: A10172).

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documentn5

Documents

introduction of nanoparticles

low particle

ageing performance

low ller content

presence of grafting

zhongshan university

university of kaiserslautern

polymer composites