lypctnc

,*, M

Minis

n Un

sity of

d fore 5 N

sults of a series of thermoplastics based composites nanoparticles of the agglomerates as restricted by thelarger molecules, is helpless to well disperse the nano-particles. In some cases, as a result, the composites withthe addition of nanoparticles would exhibit propertiesworse than microcomposites. To bring the eect of the

.

* Correspondence author. Tel.: +86 208 411 2283/403 6576; fax:+86 208 403 6576.

E-mail address: ceszmq@zsu.edu.cn (M.Q. Zhang).

COMPOSITESSCIENCE AND

Composites Science and Technolog0266-3538/$ - see front matter 2004 Elsevier Ltd. All rights reserved1. Introduction

In recent years, inorganic nanoparticles lled polymercomposites have received increasing research interests ofmaterials scientists because the ller/matrix interface inthese composites might constitute a much greater areaand hence inuence the composites properties to amuch greater extent at rather low ller concentrationas compared to conventional micro-particulate compos-ites. Considering the versatility of production facilitiesand raw materials, dispersive mixing of ready-madenanoparticles and polymers is still one of the main man-ufacturing methods to make nanocomposites. The re-

prepared in this way have been reported, like nano-SiO2/polypropylene (PP) [1,2], nano-SiO2/high-densitypolyethylene (HDPE) [3], nano-CaCO3/PP [46], nano-SiO2/poly(ethylene terephthalate) (PET) [7], nano-TiO2/polystyrene (PS) [8], nano-SiO2/acrylic latex [9],and nano-SiO2/polyethersulfone (PES) [10].

It is worth noting that the market available nanopar-ticles generally take the form of agglomerates, which arehard to be broken apart during compounding due to thestrong interaction among the nanoparticles, the limitedshear force provided by the mixing device and the highmelt viscosity of polymer melts. Modication with cou-pling agents, which can only react with the exteriorAbstract

The current paper is a continuation of the authors work on mechanical performance of nano-silica/polypropylene (PP) compos-ites. Unlike the fumed nano-silica used in the previous studies, precipitated nano-silica is employed in the present investigation. Theresults indicate that graft polymerization onto the precipitated nano-silica (that has been successfully applied to the surface mod-ication of fumed nano-silica) is still an eective method to pre-treat the particles, which leads to an overall improvement of thecomposites properties. In addition to the grafting polymers covalently attached to the nanoparticles, matrix ductility and nanopar-ticles size are important factors that inuence the extent of performance enhancement of the composites. In the case of suitable com-bination of these factors, the positive eect of the nanoparticles can be maximized. 2004 Elsevier Ltd. All rights reserved.

Keywords: A. Particle-reinforced composites; B. Mechanical properties; B. Surface treatments; NanoparticlesSilica nanoparticles lled posurface treatment, matrix du

mechanical performa

Chun Lei Wu a, Ming Qiu Zhang b

a Key Laboratory for Polymeric Composite and Functional Materials ofb Materials Science Institute, Zhongsha

c Institute for Composite Materials (IVW), Univer

Received 9 November 2003; received in reviseAvailable onlindoi:10.1016/j.compscitech.2004.09.004ropylene: eects of particleility and particle species one of the composites

in Zhi Rong b, Klaus Friedrich c

try of Education, Zhongshan University, Guangzhou 510275, PR China

iversity, Guangzhou 510275, PR China

Kaiserslautern, D-67663 Kaiserslautern, Germany

m 14 July 2004; accepted 17 September 2004ovember 2004

www.elsevier.com/locate/compscitech

TECHNOLOGYy 65 (2005) 635645

the following steps called gaseous graft polymerization,which is dierent from the method used previously [1].

ce annanoparticles into play, graft polymerization onto theparticulates surface was developed by the authors as apre-treatment technique [13,5,6]. The low molecularweights of the grafting monomers allow them to pene-trate into the agglomerated nanoparticles and react withthe particles both inside and outside the agglomerates.Taking the advantage of this, the following benetscan be gained: (i) the hydrophilic particles surfaces areconverted into hydrophobic in favor of improving misci-bility between the components; (ii) the loosened nano-particles agglomerates are turned into compactnanocomposite structure consisting of the particles, thegrafting polymers and the homopolymers generated inthe course of graft polymerization; and (iii) the interfa-cial interaction between the ller particles and the sur-rounding matrix is enhanced through entanglement ofthe grafting polymers attached to the nanoparticles withthe matrix molecules. Therefore, stress can be trans-ferred to all the nanoparticles when the composites aresubjected to applied force, while stiening, reinforcingand toughening eects are observed at very low nano-particle content. It is believed that a double percolationof eective stress volumes takes the responsibility for theoverall enhancement of the composites [11].

In our previous studies [13,6], fumed nano-silicawas employed as the predominant ller particles andthe feasibility of graft pre-treatment of these nanopar-ticles for acquiring mechanical properties improvementhas been investigated by using lab-scale and industrialscale compounding machines, respectively. Besides, sur-face morphologies of the nanoparticles before and aftergrafting were also characterized [12]. For conductingsystematic researches, precipitated nano-silica, whichis synthesized by a process dierent from the one formaking fumed silica, is used in the present work. Onthe basis of this ller selection, some other importantfactors that have not yet been reported in the litera-ture, like the eects of matrix ductility and particulatesize on the composites mechanical performance, arestudied hereinafter. To maintain the continuity of ourwork on this subject, PP acts as the matrix polymeronce more.

Precipitated silica is manufactured by a wet proce-dure by treating silicates with mineral acids to obtainne hydrated silica particles in the course of precipita-tion. The reaction and drying conditions determine theporosity, surface area, surface chemistry and the degreeof impurities in the precipitated silica. In general, precip-itated silicas are cheap and have a particle size higherthan 10 lm [13]. However, they can be made as tiny asnano-scale under specic circumstances [14]. Fumed sil-icas are manufactured by high-temperature hydrolysis ofsilicon tetrachloride in a ame. Silanol and siloxanegroups are created on the silica surface, leading tohydrophilic nature of the particles. The use of fumed sil-

636 C.L. Wu et al. / Composites Scienica as llers in thermoplastics has been well documentedThe nanoparticles were pre-treated at 140 C under vac-uum for 5 h to eliminate possible absorbed water on thesurface of the particles. Then they were lled into a glassvessel and absorb certain amount of monomer vaporunder vacuum. Afterwards, the sealed vessel containingnot only by our own works stated above but also byother groups [7,9,10]. It is thus expected that the appli-cability of nano-silica in thermoplastics modicationwould be further broadened if precipitated silicas proveto be as useful as fumed ones.

2. Experimental

2.1. Materials

The precipitated nano-SiO2 with an average primaryparticle size of 10 nm and a specic surface area of640 m2/g was supplied by Zhoushan NanomaterialsCo., China. For purposes of comparative study, fumednano-SiO2 with an average primary particle size of15 nm and a specic surface area of 374 m2/g waspurchased from Shenyang Chemical Engineering Ltd.,China. The two types of silicas are denoted by p-SiO2and f-SiO2 for the convenience of discussion in the fol-lowing text, respectively.

Isotactic polypropylene (PP) homopolymer T30S,supplied by Qilu Petrochemical Industrial Co., China,was used as the matrix polymer. It has a melt ow index(MFI) of 3.2 g/10 min (2.16 kg at 230 C). To reveal theeect of matrix ductility, however, other two types of PPwere also introduced. They are isotactic PP homopoly-mer PP700 (MFI = 13 g/10 min) produced by Guangz-hou Petroleum Chemical Co., China, and blockcopolymerized PP EPS30R (consisting of the segmentsof ethylene and propylene, MFI = 1.9 g/10 min) by QiluPetrochemical Industrial Co., China.

For carrying out graft polymerization onto thenano-silica particles, various commercial monomers:styrene, methyl methacrylate, ethyl acrylate and butylacrylate, were used as grafting monomers without fur-ther purication, respectively. It is known that themacromolecular chains constructed by these monomershave rigidities ranking in the order of their appearanceas written above. The later three monomers have thesame backbones but dierent lengths of the side chains,which might help to understand the interfacial eecteasily.

2.2. Pre-treatment of the nanoparticles through graft

polymerization and the related analysis

Modication of nano-silica proceeded according to

d Technology 65 (2005) 635645the nanoparticles was irradiated by 60Co c-ray under

atmosphere at room temperature. After exposure to adose of 4 Mrad, the powder was available for the subse-quent compounding.

To evaluate the results of grafting and to characterizethe grafted nanoparticles, the grafting polymer and thehomopolymer, which were generated during the irradia-tion polymerization of the monomers, should be sepa-rated. For this purpose, a certain amount of theirradiation products were extracted by benzene in aSoxhlet apparatus for 36 h. In this way the homopoly-mer was isolated. The residual material was then driedin vacuum at 80 C until a constant weight was reached.By using a Shimadzu TA-50 thermogravimetre (TG),the weight of the grafting polymer on the modiednanoparticles was determined and the percent graftingcan be calculated accordingly. To further separate the

Charpy impact bars (GB/T 1043-93) with a CJ150MZinjection-molding machine at 215 C.

Room temperature tensile testing of the compositeswas conducted on a Hounseld-5KN universal testingmachine. Unless otherwise specied, the crosshead speedwas set at 50 mm/min. Five samples were tested for eachcase. A Hitachi S-520 scanning electron microscope(SEM) was used to observe the fractured surfaces. Un-notched Charpy impact strengths of the compositeswere measured by an XJJ-50 impact tester. Eight sam-ples were tested for each case.

The isothermal crystallization of PP and its compos-ites was conducted on a PerkinElmer dierential scan-ning calorimetry (DSC-7) instrument. The samples wereheated from room temperature to 210 C at a rate of10 C/min in N , and kept at 210 C for 5 min. Then,

on the particles surface is not discussed in detail.Instead, the investigation results in this aspect [15] are

ipitate

g poly

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635645 637grafting polymer from the treated nanoparticles, nano-silica accompanied with the unextractable graftingpolymer was immersed in 20% HF solution for 72 h toremove the inorganic particles. Table 1 lists the param-eters quantifying the grafting reaction on the nanoparti-cles for reference.

A Micromeritics ASAP 2100 surface area analyzerwas employed to measure the micropore volumes andspecic surface areas of the nano-silica before and afterthe graft treatment. Prior to the tests, the samples weredried in vacuum under 150 C for 24 h.

2.3. 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 ofthe extruder was set to 180 rpm. Finally, the resultantpellets were molded into dog-bone-shaped tensile bars(ASTM D638-97 Type IV specimen) and rectangular

Table 1Irradiation induced graft polymerization of vinyl monomers onto prec

Samples Percent graftingb (%)

p-SiO2g-PSe 11.4

p-SiO2g-PMMAf 12.7

p-SiO2g-PEAg 15.3

p-SiO2g-PBAh 13.4

a Irradiation dose: 4 Mrad; monomer/silica = 20 wt%.b Percent grafting = weight of grafting polymer/weight of silica.c Grafting eciency = weight of grafting polymer/weight of graftind Monomer conversion = weight of polymer/weight of monomer.e PS: polystyrene.f PMMA: polymethyl methacrylate.g PEA: polyethyl acrylate.

h PBA: polybutyl acrylate.summarized below, providing knowledge basis for thecomposites analysis. According to the infrared andX-ray photoelectron spectroscopy studies, it is provedthat after irradiation the grafting polymers are chemi-cally attached to silica surface through SiOC andSiC bonds as expected. The weight average molecular

d nano-silica particlesa

Grafting eciencyc (%) Monomer conversiond (%)

57.1 10063.3 10076.5 10066.9 100

mer and homopolymer.2

the samples were cooled down to 130 C at a rate of80 C/min and started to isothermally crystallize. Inthe meantime, the exothermic curves were collected.When the crystallization is completed, the samples wereheated again to 210 C at a rate of 10 C/min and themelting behaviors were recorded.

3. Results and discussion

3.1. Eect of surface grafting onto the nanoparticles

As the current paper is focused on the mechanicalperformance of the PP composites lled with nano-silica, the chemistry related to the graft polymerization

weights of the grafting polymers are about 105 and those

nano-silica. In addition, the pore diameters correspond-ing to the maximum pore volumes of the treated nano-particles shift from 0.55 to 0.71 nm. It seems that the

0.0 0.2 0.4 0.6 0.8 1.00

70

140

210

280

350

3

2

1

Ads

orbe

d vo

lum

e[cm3

/g]

Relative pressure

1: p-SiO2 as-received2: p-SiO2-g-PS3: p-SiO2-g-PEA

Adsorption curve Desorption curve

Fig. 1. Nitrogen absorption isotherms of nano-silica particles.

p-SiO2 as-received p-SiO2-g-PS p-SiO2-g-PEA

0.0 0.5 1.0 1.5 2.0 2.50.00

0.03

0.06

0.09

0.12

0.15

Diff

eren

tial p

ore

volu

me

[cm3 /g

]

Diff

eren

tial p

ore

volu

me

[cm3 /g

]

Pore diameter [nm]0.000

0.003

0.006

0.009

0.012

0.015

Fig. 2. Micropore size distribution of nano-silica particles.

638 C.L. Wu et al. / Composites Science and Technology 65 (2005) 635645of the homopolymers are 104.To have an image of the surface structure variation,

nitrogen absorption isotherms of the nanoparticles weremeasured (Fig. 1). The curve proles are something be-tween type II and IV isotherms as viewed from Brunauerclassication system [16]. Although hysteresis loops canbe seen on the isotherms of the nanoparticles with andwithout the grafting polymers, those are not closed inthe case of p-SiO2g-PS and p-SiO2g-PEA. It meansthat the capillary condensation occurring under higherrelative pressure is severer in the grafted nano-silica.In comparison with p-SiO2 as-received, both p-SiO2g-PS and p-SiO2g-PEA have signicantly reducedabsorbance. Evidently, it can be attributed to the factthat the micropores and macropores of the nanoparticleagglomerates are lled by the grafting polymers, whichin turn demonstrates that the grafting polymers havebeen planted onto the nanoparticles. The plots showingthe micropore size distribution in Fig. 2 also providesupporting evidence for the estimation. After graftingtreatment, the pore volumes of the nanoparticles areabout ten times smaller than that of the untreatedFig. 3. SEM microphotos of tensile fracture surfaces of nano-silica/PP compvol%); (b) p-SiO2g-PS/PP (nano-silica content = 2.75 vol%).graft polymerization was initiated at the smallest micro-pores of the nanoparticles. The porous structure of theparticles has been changed accordingly.

Figs. 3 and 4 show the typical fracture surfaces andmechanical properties of PP composites with nano-sil-ica, respectively. The untreated nanoparticles are se-verely agglomerated in PP matrix (Fig. 3(a)), while thetreated ones are well separated into tiny aggregates(Fig. 3(b)). The SEM observations evidence the authorsexpectation of graft treatment of the nanoparticles. Itmeans that the composites containing the untreatednanoparticles are provided with heterogeneous micro-structure and would exhibit worse reinforcing andtoughening eects as compared to those having morehomogeneous appearances due to graft pre-treamentof the particles.

As for the details of the composites mechanical per-formance, it is seen that both treated and untreatednano-silicas are able to stien the matrix, as reectedby the proportional relationship between Youngs mod-ulus and the ller loading (Fig. 4(a)). Comparatively, thestiness of the composites with grafted nano-silica isosites (PP: T30S). (a) p-SiO2 as-received/PP (nano-silica content = 2.74

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635645 6391.2

1.4

1.6

1.8

2.0(a)Yo

ung'

s m

odul

us [G

Pa]

p-SiO2 as-received/PPp-SiO2-g-PS/PPp-SiO2-g-PMMA/PPp-SiO-g-PEA/PP

2p-SiO2-g-PBA/PPlower than that of the untreated nanoparticles compos-ites. As interfacial stress transfer eciency depends onthe stiness of the interphase, higher interfacial stinessfavors improvement of the composites modulus [17].The grafting polymer adhering to the nano-silica as wellas the surrounding homopolymer establish a compliantinterlayer between the particles and the matrix, andhence decrease the stiening eect of the particles. Onthe other hand, it is interesting to note that the moduliof the composites with grafted nano-silica at identicalparticle content are arranged in the following order:p-SiO2g-PS/PP > p-SiO2g-PMMA/PP > p-SiO2g-

0.0 0.5 1.0 1.5 2.0 2.5

SiO2 [vol%]

0.0 0.5 1.0 1.5 2.0 2.5

SiO2 [vol%]

0.0 0.5 1.0

SiO2 [

p-SiO2 as-received/PPp-SiO2-g-PS/PPp-SiO2-g-PMMA/PPp-SiO-g-PEA/PP

2p-SiO2-g-PBA/PP

35

36

37

38

39

40(b)

(e)

Tens

ile s

treng

th [M

Pa]

0

25

50

75

100

125

Impa

ct s

treng

th [k

J/m2 ]

Fig. 4. Mechanical properties of nano-silica/PP composites as a function ofelongation at break; (d) area under tensile stressstrain curve; and (e) impacp-SiO2 as-received/PPp-SiO2-g-PS/PPp-SiO2-g-PMMA/PPp-SiO-g-PEA/PP

2p-SiO2-g-PBA/PP

(c)

0

100

200

300

400

Elon

gatio

n to

bre

ak [%

]PEA/PP > p-SiO2g-PBA/PP. It exactly coincides withthe exibilities of the grafting polymer chains. That is,PBA possesses the highest exibility and consequentlymasks the high stiness of the particles in the compositesto the greatest extent. The phenomenon again proves theabove analysis concerning the dependence of compositesmoduli on interfacial stiness.

With respect to the composites under higher nano-sil-ica content, it is found that the moduli increment of thecomposites with treated nano-silica can no longer keepthe linearity. For p-SiO2g-PEA/PP and p-SiO2g-PBA/PP composites, their moduli even decline. This

1.5 2.0 2.5

vol%]

0.0 0.5 1.0 1.5 2.0 2.5

SiO2 [vol%]

0.0 0.5 1.0 1.5 2.0 2.5

SiO2 [vol%]

p-SiO2 as-received/PPp-SiO2-g-PS/PPp-SiO2-g-PMMA/PPp-SiO-g-PEA/PP

2p-SiO2-g-PBA/PP

(d)

0

30

60

90

120

Are

a un

der s

tress

-st

rain

cur

ve [M

Pa]

p-SiO2 as-received/PPp-SiO2-g-PS/PPp-SiO2-g-PMMA/PPp-SiO-g-PEA/PP

2p-SiO2-g-PBA/PP

silica content (PP: T30S): (a) Youngs modulus; (b) tensile strength; (c)t strength.

ce anshould be the result of the increasing amount of the softinterphase comprised of the grafting polymers andhomoploymers with a rise in the content of the graftednanoparticles.

Besides stiening eect, the precipitated nano-silicacan also provide PP with reinforcing eect at ratherlow ller concentration (Fig. 4(b)), which is similar tothe situation of fumed nano-silica [1,2]. In the case ofp-SiO2g-PS/PP and p-SiO2g-PBA/PP, for example,tensile strengths of the composites are higher than thatof neat PP. This is dierent from what is observed inconventional micrometer particles/polymer composites,i.e., tensile strength of the composites remarkably de-creases with the addition of the particulate llers dueto the poor bonding at the interface [18,19]. Jancaret al. [20] suggested that a strong ller/matrix adhesionwould lead to enhanced strength of particulate compos-ites. According to their consideration, it is known thatthe improvement of tensile strength of the compositesexhibited in Fig. 4(b) should also be interpreted as theimprovement of the interfacial interaction. Especiallywhen PS- and PBA-grafted nano-silica is incorporated,the chain entanglement between the grafting polymersand the matrix polymer guarantees eective interfacialbonding over the whole ller content range of interests.However, grafting treatment does not always take eectas revealed by the composition dependence of tensilestrength of p-SiO2g-PMMA/PP and p-SiO2g-PEA/PP. This contradicts the results of PP composites withfumed nano-silica grafted by the same species of poly-mers [1]. It means that in addition to the nature of thegrafting polymers on the nanoparticles, species of nano-particles themselves and other unknown factors alsogreatly inuence the reinforcing eect of the graftednanoparticles. More detailed study in this direction is re-quired to have a reasonable conclusion.

From engineering point of view, elongation-to-breakis an important parameter describing the rupture behav-ior of composite materials. The addition of mineral par-ticulates into polymers used to lower it, even though thematrix has high impact toughness [21]. Fig. 4(c) clearlyindicates that this is not the case when nano-silica isused. Either untreated or treated nano-silica is able toincrease elongation-to-break of PP. Compared to theuntreated and PS-grafted ones, acrylic polymers-graftednano-silicas are more eective, in particular whenp-SiO2g-PMMA is concerned. It implies that in princi-ple compliant grafting polymers can induce more matrixpolymer to be involved in plastic deformation. However,chain exibility of the grafting polymers on thenanoparticles is not the only prerequisite, otherwise p-SiO2g-PBA/PP should have the highest value of elon-gation-to-break. In contrast to tensile strength of thecomposites that needs strong ller/matrix bonding, themicro-deformation mechanism involved in elongation-

640 C.L. Wu et al. / Composites Sciento-break of the composites depends upon extensionalityof the nanoparticles agglomerates in the composites [22].Balanced viscoelasticity of the grafting polymers on thenanoparticles might thus be necessary for the improve-ment of elongation-to-break of the composites.

Fig. 4(d) shows the areas under tensile stressstraincurves of the composites, another parameter character-izing the static toughness. The magnitude order of thevalues at given ller content is similar to that in Fig.4(c), suggesting that the deformation features of thecomposites still govern their areas under tensile stressstrain curves. Grafted nano-silica again performs wellthan the untreated nanoparticles. The maximum valueof p-SiO2g-PMMA/PP is about 3.8 times higher thanthat of neat PP. The results demonstrate the role ofgrafting polymers, i.e., interconnecting the nanoparticlesthrough chemical bonding and correlating the graftednanoparticles with the matrix by chain entanglement.Under the applied force, plastic deformation of largeamount of the matrix polymer beside the grafted nano-particles is induced, leading to signicantly high elonga-tion-to-break and areas under tensile stressstrain curveof the composites. In the case of untreated nanoparticleslled PP, voiding and disintegration of the nanoparticleagglomerates are the main ways of energy dissipationdue to the lack of extensionality.

Unnotched impact strengths of the composites are gi-ven in Fig. 4(e) as a function of silica fraction. It is clearthat after grafting treatment the toughening ability ofthe nanoparticles is greatly increased. At the silica con-tent of 0.45 vol%, the impact ductility of p-SiO2g-PBA/PP is over 3 times higher than that of unlled PP, and allthe impact strengths of the composites at this ller load-ing are arranged in the order contrary to what observedin Fig. 4(a) concerning Youngs modulus: p-SiO2/PP < p-SiO2g-PS/PP < p-SiO2g-PMMA/PP < p-SiO2>g-PEA/PP < p-SiO2g-PBA/PP. Since unnotched impactstrength reects the energy consumed by the plastic defor-mation prior to crack initiation, the above results manifestthat the exible macromolecular chains grafted onto thenanoparticles surfaces must have made contribution to thispart of energy. In comparison to the data shown inFig. 4(c) and (d), it is known that the dynamic toughnessof the composites (i.e., impact strength) is more sensitiveto the dispersion status of the grafted nanoparticles thanthe static toughness (i.e., elongation-to-break and area un-der tensile stressstrain curve). With a rise in the ller con-tent, homogeneity of the particles distribution might beworse as a result of the increased viscosity of the compositesystems. Under this circumstance, the grafting polymerscould not response in step with the impact load and hencedecline of impact strength is observed when the content ofnano-silica exceeds 0.45 vol%.

Since PP is a semi-crystalline polymer and itsmechanical properties would change with the crystallinestructure and crystallinity, the inuence of the addition

d Technology 65 (2005) 635645of the nanoparticles should be known. As listed in

Table 2, the untreated nano-silica remarkably acceler-ates the crystallization of PP matrix as revealed by thevalues of t1/2, tf and k. This nucleating activity can be ex-plained by the thermodynamic model proposed by

Table 2Kinetic parameters of isothermal crystallization of PP and its composites at

Samples t1/2a (min)

PP 8.68p-SiO2 as-received/PP (nano-silica content = 0.86 vol%) 4.69p-SiO2g-PS/PP (nano-silica content = 1.066 vol%) 6.17p-SiO2g-PEA/PP (nano-silica content = 0.82 vol%) 9.79

a t1/2: half-crystallization time.b tf: the time at which the crystallization is completed.c n: Avrami index.d k: rate constant of crystallization.e DH: enthalpy of crystallization.

C.L. Wu et al. / Composites Science anEbengou [23]. That is, when PP chains were absorbedon the silica surface, the congurational entropy of theentire chain decreased, forming a nucleus of a certainvolume within the adsorbed chains costs less energy.In the case of grafted nanoparticles, the nucleation ef-fects are less profound because the grafting polymersshielded the nanoparticles from the direct contacts withPP, which coincides with the results of PP/elastomer/particles composites characterized by core-shell micro-structures [24]. Fig. 5 further shows the melting behav-iors of PP and its composites. In comparison with neatPP, the endothermic peak proles and temperatures ofthe composites are almost the same, suggesting the crys-tal microstructure of PP has not been changed. Rela-tively, the crystallinity of PP in the composites isslightly reduced. On the whole, however, it cannot beconcluded that the aforesaid mechanical performancevariations of PP composites with the incorporation ofsilica nanoparticles results from the variations in PPcrystalline structure and crystallinity. Although un-

Samples Tm (oC) X

c (%)

1 166.5 46.62 166.4 45.73 166.5 45.4140 150 160 170 180 190 200

4 166.1 44.5432

1

Endo

>

Temperature [oC]Fig. 5. DSC heating traces of PP and its composites having beenisothermally crystallized at 130 C (PP: T30S). (1) Neat PP; (2) p-SiO2as-received/PP (nano-silica content = 0.86 vol%); (3) p-SiO2g-PS/PP(nano-silica content = 1.06 vol%); (4) p-SiO2g-PEA/PP (nano-silicacontent = 0.82 vol%). Tm: peak melting temperature; Xc: crystallinity.treated nano-silica has signicant nucleating ability, itfails to improve PP properties equivalently. Similarly,the grafted nanoparticles only lead to marginal decreaseof PP crystallinity, which is also out of proportion to theperformance enhancement.

3.2. Eects of matrix ductility and nanoparticle species

Mechanical performance of composite materials is afunction of ller and matrix characteristics. When stud-ying nano-CaCO3/PP composites, for example, Renet al. [25] showed the importance of the matrix tough-ness. A much more remarkable increase in impact resist-ance was observed in the composites with a PPcopolymer possessing higher ductility as matrix, whilethe same nanoparticles did not result in a similarimprovement in a PP homopolymer with lower ductility.As a result, it was suggested that the polymer to betoughened by nanoparticles should possess at least a cer-tain toughness, which is dierent from the case whenelastomer acts as toughener.

To nd out whether nano-silica/PP composites followthe same law, three types of PP: EPS30R, T30S andPP700, are used in the current work. Relatively,EPS30R has the highest toughness and PP700 the low-est. For making comparative tests, the crosshead speedof the tensile tests was raised to 200 mm/min to matchthe large failure strain of EPR30S. As illustrated byFig. 6(a) and (b), the addition of nano-silica into either

130 C

tfb (min) nc kd ( 103minn) DHce (J/g)

17.8 2.59 2.57 97.18.89 3.04 6.31 97.011.39 3.04 2.75 95.017.7 2.92 1.08 97.0

d Technology 65 (2005) 635645 641EPR30S or PP700 results in reduced values of elonga-tion-to-break and area under tensile stressstrain curve,no matter the particles have been treated or not. Itresembles the behavior of micro-particles lled polymercomposites. However, nano-silica lled T30S has ac-quired signicantly high static toughness. The untreatednanoparticles have already certain ability to increaseelongation-to-break and area under tensile stressstraincurve of T30S PP. It means that under tensile loading,considerable matrix yielding of nano-silica lled PPoccurs only in the case of matrix polymer possessingmoderate toughness. The synergetic eect brought aboutby the grafted nanoparticles depends on the matchingof matrix toughness and exibility of the grafting

ce an642 C.L. Wu et al. / Composites Scienpolymers. For the composites based on EPR30S andPP700, the nanoparticles are ineective to induce local-ized matrix drawing.

Fig. 6(c) shows the impact strength of the composites.Although EPS30R and its composites are too tough tobe broken by the impact load under the current testingconditions, some useful hints can still be yielded by com-paring the performance of T30S and PP700 based com-posites. It is interesting to see that the nanoparticlesexert toughening eect on both types of PP. However,the relative increments of the impact strengths of T30Sbased composites are much higher than those of

Fig. 6. Mechanical properties of nano-silica/PP composites withdierent PP as matrices: (a) elongation at break; (b) area under tensilestressstrain curve; and (c) impact strength. The percentage numeralsquantify the relative variation of the mechanical property withreference to the value of the corresponding neat PP. Nano-silicacontent = 0.5 vol%, crosshead speed of the tensile tests = 200 mm/min.PP700 based composites. Relating the results ofFig. 6(a)(c) to each other, it is clear that the matrix duc-tility or the capability of the matrix to plastically deformis a key factor inuencing the toughening eect of thenanoparticles. In the case of suitable matrix ductility likeT30S, large scale of matrix polymer is successfully in-volved in plastic deformation as induced by the nano-particles, leading to high impact toughness of thecomposites.

On the other hand, it has been known that the inter-face between the ller particles and the matrix in a pol-ymer nanocomposite constitutes a much greater areawithin the bulk material as compared with conventionalcomposites containing micrometer-sized particles, andhence inuences the composites properties to a muchgreater extent. However, there is no information aboutthe eect of nanoparticles having dierent sizes. In thecurrent work, fumed nano-silica (15 nm) acts as a refer-ence for the precipitated nano-silica (10 nm). To bringthe positive eect of the nanoparticles into full play,both types of the particles were grafted with PBA (thepercent grafting of f-SiO2g-PBA is 8.26%, close to thatof p-SiO2g-PBA: 13.4%) and compounded with PP(T30S) under the same conditions.

First Youngs moduli of the composites are comparedin Fig. 7(a). The ller content dependences of modulusof the two types of composites are almost the same,but the precipitated nano-silica leads to higher compos-ites stiness. It must result from the increased interfacialarea in the composites with ner nano-silica, which pro-motes the stress transfer eciency within small strainrange. The reduction of Youngs modulus at highersilica content manifests the softening eect of the com-pliant PBA interlayer, which becomes more and moreevident with a rise in content of the graftednanoparticles.

With respect to tensile strength of the composites, onthe contrary, the larger particles (fumed silica) seem tohave more remarkable reinforcing ability (Fig. 7(b)).Generally, composites strength depends on the ller/ma-trix bonding under large strain condition. Since the pre-sent nanoparticles are grafted with the same polymer(i.e., PBA) and the amounts of grafting polymer are sim-ilar in the two types of nano-silica, the interfacial adhe-sion in the composites should be almost identical atgiven ller concentration [26]. The signicant dierencein the tensile strengths of the composites might be attrib-uted to the load bearing ability of the nano-silica itself.That is, the fumed nano-silica particles might be stron-ger than the precipitated ones. To the authors knowl-edge, the strength of nano-silica particles has not yetbeen reported because of the experimental diculties,but one might nd some traces on the basis of their syn-thesis processes. Precipitated silica particles are obtainedin low-temperature wet process and might form soft

d Technology 65 (2005) 635645agglomerates (more physical bonding) during drying.

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635645 6431.3

1.4

1.5

1.6

1.7

1.8Yo

ung'

s m

odu

lus

[G

Pa] p-SiO2-g-PBA/PP

f-SiO2-g-PBA/PP

(a)As for fumed silica particles, they are produced in high-temperature gaseous process, which would made themform hard agglomerates (more chemical bonding). Dur-ing grafting pre-treatment, the former particle agglomer-ates are easier to be broken apart so that the graftedprecipated nano-silicas can be well dispersed than thefumed ones. When the composites are subjected to ap-plied stress, the intrinsic high strength of the fumed silicaaggregates oers higher strength for the composites.

In contrast to Youngs modulus and tensile strength,elongation-to-break and area under tensile stressstraincurve of the composites shown in Fig. 7(c) and (d) indi-

0.0 0.5 1.0 1.5 2.0 2.5

SiO2 [vol%]

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gatio

n at

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ak [%

]

SiO2 [vol%]

p-SiO2-g-PBA/PPf-SiO2-g-PBA/PP

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70

105

140

Impa

ct st

ren

gth

[kJ/m

2 ]

SiO

p-SiO2-gf-SiO2-g-

(e)

(c)

Fig. 7. Mechanical properties of PP composites lled with dierent nano-silictensile strength; (c) elongation at break; (d) area under tensile stressstrain c36

37

38

39

40

Tens

ile s

tren

gth

[MPa

]

p-SiO2-g-PBA/PP f-SiO2-g-PBA/PP

(b)cate that the species of the nanoparticles nearly havenothing to do with the two parameters when the llercontent 60.84 vol%. The characteristics of the graftingpolymer play the leading role in this case. This is reason-able because under tensile loading the extent of matrixstretching is mainly related to the chain entanglementbetween the grafting polymers and the matrix molecules.For the composites with nano-silica content higher than0.84 vol%, the decrease in elongation-to-break and areaunder tensile stressstrain curve of f-SiO2g-PB/PPshould be due to the poor dispersion of the nanoparti-cles. Compared with precipitated nano-silica, fumed

0.0 0.5 1.0 1.5 2.0 2.5

SiO2 [vol%]

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30

45

60

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Are

a un

der s

tress

-st

rain

cur

ve [M

Pa]

SiO2 [vol%]

p-SiO2-g-PBA/PPf-SiO2-g-PBA/PP

1.5 2.0 2.5

2 [vol%]

-PBA/PPPBA/PP

(d)

as as a function of silica content (PP: T30S): (a) Youngs modulus; (b)urve; and (e) impact strength.

When the composites specimens are subjected to im-pact load, however, dierent nanoparticles give dierent

cles become more ecient to improve the strength andtoughness of the composites. The nature of the grafting

[18] Nielsen LE. Simple theory of stressstrain properties of lledpolymer. J Appl Polym Sci 1966;10:97103.

[19] Nicolais L, Narkis M. Stressstrain behavior of styreneacrylo-

ce and Technology 65 (2005) 635645polymer chains plays an important role in the propertiesenhancement.

Ductility of the matrix PP determines the toughen-ing eect of the nanoparticles. Only in the case ofmoderate matrix ductility, the composites can receivethe highest extent of toughness increase. Besides, thesize and surface area of the nanoparticles are alsoimportant inuencing factors. The smaller nanoparti-cles lead to higher Youngs modulus and impactstrength of the composites, and reduce the sensitivityof the static toughness to the status of llerdistribution.

Acknowledgements

The authors are grateful to the support of the Deut-sche Forschungsgemeinschaft (DFG FR675/40-4) forthe cooperation between the German and Chinese insti-tutes on the topic of nanocomposites. Further thanksare due to the National Natural Science Foundationof China (Grant No. 50133020), the Team Project ofresponses (Fig. 7(e)). Similar to the results in Fig. 7(a),precipitated nano-silica lled PP has higher impactstrength than fumed nano-silica/PP in most cases. Thisimplies that besides the interphase formed by the graft-ing polymers and matrix, the interfacial area betweenthe nanoparticles and the surrounding polymers is alsoimportant to the toughening eect. The latter factormight facilitate the generation of crazes during the im-pact test, absorbing certain amount of the input energyadditionally. Therefore, the composites lled with nano-silica with smaller particle size exhibit greater resistanceto impact loading.

4. Conclusions

Precipitated nano-silica is able to provide PP withstiening, reinforcing and toughening eects at ratherlow ller concentration as fumed nano-silica. Havingbeen grafted with dierent polymers onto the surfacesin terms of gaseous graft polymerization, the nanoparti-nano-silica has much lower packing density. As a result,homogeneous distribution of fumed nano-silica duringcompounding is more dicult to be achieved than pre-cipitated nano-silica especially at higher ller content.It explains that the two parameters reecting statictoughness maintain almost unchanged over the wholeller content range of interests in the case of p-SiO2g-PB/PP (Fig. 7(c) and (d)).

644 C.L. Wu et al. / Composites Scienthe Natural Science Foundation of Guangdong, Chinanitrile/glass bead composites in the glassy region. Polym Eng Sci1971;11:1949.

[20] Jancar J, Dianselmo A, Dibenedetto AT. The yield strength ofparticulate reinforced thermoplastic composites. Polym Eng Sci(Grant No. 20003038), and the Key Program of the Sci-ence and Technology Department of Guangdong, China(Grant No. A10172).

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C.L. Wu et al. / Composites Science and Technology 65 (2005) 635645 645

Silica nanoparticles filled polypropylene: effects of particle surface treatment, matrix ductility and particle species on mechanical performance of the compositesIntroductionExperimentalMaterialsPre-treatment of the nanoparticles through graft polymerization and the related analysisComposites preparation and characterization

Results and discussionEffect of surface grafting onto the nanoparticlesEffects of matrix ductility and nanoparticle species

ConclusionsAcknowledgementsReferences