Silica nanoparticles filled polypropylene: effects of particle surface treatment, matrix ductility and particle species on mechanical performance of the composites

COMPOSITES SCIENCE AND TECHNOLOGY Composites Science and Technology 65 (2005) 635–645 www.elsevier.com/locate/compscitech

Silica nanoparticles filled polypropylene: effects of particle surface treatment, matrix ductility and particle species on mechanical performance of the composites Chun Lei Wu a, Ming Qiu Zhang a

b,*

, Min Zhi Rong b, Klaus Friedrich

c

Key Laboratory for Polymeric Composite and Functional Materials of Ministry of Education, Zhongshan University, Guangzhou 510275, PR China b Materials Science Institute, Zhongshan University, Guangzhou 510275, PR China c Institute for Composite Materials (IVW), University of Kaiserslautern, D-67663 Kaiserslautern, Germany Received 9 November 2003; received in revised form 14 July 2004; accepted 17 September 2004 Available online 5 November 2004

Abstract The current paper is a continuation of the authorsÕ work on mechanical performance of nano-silica/polypropylene (PP) composites. Unlike the fumed nano-silica used in the previous studies, precipitated nano-silica is employed in the present investigation. The results indicate that graft polymerization onto the precipitated nano-silica (that has been successfully applied to the surface modification of fumed nano-silica) is still an effective method to pre-treat the particles, which leads to an overall improvement of the composites properties. In addition to the grafting polymers covalently attached to the nanoparticles, matrix ductility and nanoparticles size are important factors that influence the extent of performance enhancement of the composites. In the case of suitable combination of these factors, the positive effect of the nanoparticles can be maximized.
2004 Elsevier Ltd. All rights reserved. Keywords: A. Particle-reinforced composites; B. Mechanical properties; B. Surface treatments; Nanoparticles

1. Introduction In recent years, inorganic nanoparticles filled polymer composites have received increasing research interests of materials scientists because the filler/matrix interface in these composites might constitute a much greater area and hence influence the compositesÕ properties to a much greater extent at rather low filler concentration as compared to conventional micro-particulate composites. Considering the versatility of production facilities and raw materials, dispersive mixing of ready-made nanoparticles and polymers is still one of the main manufacturing methods to make nanocomposites. The results of a series of thermoplastics based composites *

Correspondence author. Tel.: +86 208 411 2283/403 6576; fax: +86 208 403 6576. E-mail address: [email protected] (M.Q. Zhang). 0266-3538/$ - see front matter
2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.compscitech.2004.09.004

prepared in this way have been reported, like nanoSiO2/polypropylene (PP) [1,2], nano-SiO2/high-density polyethylene (HDPE) [3], nano-CaCO3/PP [4–6], nanoSiO2/poly(ethylene terephthalate) (PET) [7], nanoTiO2/polystyrene (PS) [8], nano-SiO2/acrylic latex [9], and nano-SiO2/polyethersulfone (PES) [10]. It is worth noting that the market available nanoparticles generally take the form of agglomerates, which are hard to be broken apart during compounding due to the strong interaction among the nanoparticles, the limited shear force provided by the mixing device and the high melt viscosity of polymer melts. Modification with coupling agents, which can only react with the exterior nanoparticles of the agglomerates as restricted by the larger molecules, is helpless to well disperse the nanoparticles. In some cases, as a result, the composites with the addition of nanoparticles would exhibit properties worse than microcomposites. To bring the effect of the

636

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

nanoparticles into play, graft polymerization onto the particulates surface was developed by the authors as a pre-treatment technique [1–3,5,6]. The low molecular weights of the grafting monomers allow them to penetrate into the agglomerated nanoparticles and react with the particles both inside and outside the agglomerates. Taking the advantage of this, the following benefits can be gained: (i) the hydrophilic particles surfaces are converted into hydrophobic in favor of improving miscibility between the components; (ii) the loosened nanoparticles agglomerates are turned into compact nanocomposite structure consisting of the particles, the grafting polymers and the homopolymers generated in the course of graft polymerization; and (iii) the interfacial interaction between the filler particles and the surrounding matrix is enhanced through entanglement of the grafting polymers attached to the nanoparticles with the matrix molecules. Therefore, stress can be transferred to all the nanoparticles when the composites are subjected to applied force, while stiffening, reinforcing and toughening effects are observed at very low nanoparticle content. It is believed that a double percolation of effective stress volumes takes the responsibility for the overall enhancement of the composites [11]. In our previous studies [1–3,6], fumed nano-silica was employed as the predominant filler particles and the feasibility of graft pre-treatment of these nanoparticles for acquiring mechanical properties improvement has been investigated by using lab-scale and industrial scale compounding machines, respectively. Besides, surface morphologies of the nanoparticles before and after grafting were also characterized [12]. For conducting systematic researches, precipitated nano-silica, which is synthesized by a process different from the one for making fumed silica, is used in the present work. On the basis of this filler selection, some other important factors that have not yet been reported in the literature, like the effects of matrix ductility and particulate size on the composites mechanical performance, are studied hereinafter. To maintain the continuity of our work on this subject, PP acts as the matrix polymer once more. Precipitated silica is manufactured by a wet procedure by treating silicates with mineral acids to obtain fine hydrated silica particles in the course of precipitation. The reaction and drying conditions determine the porosity, surface area, surface chemistry and the degree of impurities in the precipitated silica. In general, precipitated silicas are cheap and have a particle size higher than 10 lm [13]. However, they can be made as tiny as nano-scale under specific circumstances [14]. Fumed silicas are manufactured by high-temperature hydrolysis of silicon tetrachloride in a flame. Silanol and siloxane groups are created on the silica surface, leading to hydrophilic nature of the particles. The use of fumed silica as fillers in thermoplastics has been well documented

not only by our own works stated above but also by other groups [7,9,10]. It is thus expected that the applicability of nano-silica in thermoplastics modification would be further broadened if precipitated silicas prove to be as useful as fumed ones.

2. Experimental 2.1. Materials The precipitated nano-SiO2 with an average primary particle size of 10 nm and a specific surface area of 640 m2/g was supplied by Zhoushan Nanomaterials Co., China. For purposes of comparative study, fumed nano-SiO2 with an average primary particle size of 15 nm and a specific surface area of 374 m2/g was purchased from Shenyang Chemical Engineering Ltd., China. The two types of silicas are denoted by p-SiO2 and f-SiO2 for the convenience of discussion in the following text, respectively. Isotactic polypropylene (PP) homopolymer T30S, supplied by Qilu Petrochemical Industrial Co., China, was used as the matrix polymer. It has a melt flow index (MFI) of 3.2 g/10 min (2.16 kg at 230 C). To reveal the effect of matrix ductility, however, other two types of PP were also introduced. They are isotactic PP homopolymer PP700 (MFI = 13 g/10 min) produced by Guangzhou Petroleum Chemical Co., China, and block copolymerized PP EPS30R (consisting of the segments of ethylene and propylene, MFI = 1.9 g/10 min) by Qilu Petrochemical Industrial Co., China. For carrying out graft polymerization onto the nano-silica particles, various commercial monomers: styrene, methyl methacrylate, ethyl acrylate and butyl acrylate, were used as grafting monomers without further purification, respectively. It is known that the macromolecular chains constructed by these monomers have rigidities ranking in the order of their appearance as written above. The later three monomers have the same backbones but different lengths of the side chains, which might help to understand the interfacial effect easily. 2.2. Pre-treatment of the nanoparticles through graft polymerization and the related analysis Modification of nano-silica proceeded according to the following steps called gaseous graft polymerization, which is different from the method used previously [1]. The nanoparticles were pre-treated at 140 C under vacuum for 5 h to eliminate possible absorbed water on the surface of the particles. Then they were filled into a glass vessel and absorb certain amount of monomer vapor under vacuum. Afterwards, the sealed vessel containing the nanoparticles was irradiated by 60Co c-ray under

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

637

atmosphere at room temperature. After exposure to a dose of 4 Mrad, the powder was available for the subsequent compounding. To evaluate the results of grafting and to characterize the grafted nanoparticles, the grafting polymer and the homopolymer, which were generated during the irradiation polymerization of the monomers, should be separated. For this purpose, a certain amount of the irradiation products were extracted by benzene in a Soxhlet apparatus for 36 h. In this way the homopolymer was isolated. The residual material was then dried in 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 modified nanoparticles was determined and the percent grafting can be calculated accordingly. To further separate the grafting polymer from the treated nanoparticles, nanosilica accompanied with the unextractable grafting polymer was immersed in 20% HF solution for 72 h to remove the inorganic particles. Table 1 lists the parameters quantifying the grafting reaction on the nanoparticles for reference. A Micromeritics ASAP 2100 surface area analyzer was employed to measure the micropore volumes and specific surface areas of the nano-silica before and after the graft treatment. Prior to the tests, the samples were dried in vacuum under 150 C for 24 h.

Charpy impact bars (GB/T 1043-93) with a CJ150MZ injection-molding machine at 215 C. Room temperature tensile testing of the composites was conducted on a Hounsfield-5KN universal testing machine. Unless otherwise specified, the crosshead speed was set at 50 mm/min. Five samples were tested for each case. A Hitachi S-520 scanning electron microscope (SEM) was used to observe the fractured surfaces. Unnotched Charpy impact strengths of the composites were measured by an XJJ-50 impact tester. Eight samples were tested for each case. The isothermal crystallization of PP and its composites was conducted on a Perkin–Elmer differential scanning calorimetry (DSC-7) instrument. The samples were heated from room temperature to 210 C at a rate of 10 C/min in N2, and kept at 210 C for 5 min. Then, the samples were cooled down to 130 C at a rate of 80 C/min and started to isothermally crystallize. In the meantime, the exothermic curves were collected. When the crystallization is completed, the samples were heated again to 210 C at a rate of 10 C/min and the melting behaviors were recorded.

2.3. Composites preparation and characterization

As the current paper is focused on the mechanical performance of the PP composites filled with nanosilica, the chemistry related to the graft polymerization on the particles surface is not discussed in detail. Instead, the investigation results in this aspect [15] are summarized below, providing knowledge basis for the composites analysis. According to the infrared and X-ray photoelectron spectroscopy studies, it is proved that after irradiation the grafting polymers are chemically attached to silica surface through Si–O–C and Si–C bonds as expected. The weight average molecular

The nanoparticles were firstly compounded with PP (1:2 by weight) using an X(S)R-160 two-roll mill at 195 C to produce composite masterbatch. Then, the masterbatch was mixed with neat PP to dilute the filler loading to desired values through an SHJN-25 twinscrew extruder at 210–230 C. The rotation speed of the extruder was set to 180 rpm. Finally, the resultant pellets were molded into dog-bone-shaped tensile bars (ASTM D638-97 Type IV specimen) and rectangular

3. Results and discussion 3.1. Effect of surface grafting onto the nanoparticles

Table 1 Irradiation induced graft polymerization of vinyl monomers onto precipitated nano-silica particlesa Samples

Percent graftingb (%)

Grafting efficiencyc (%)

Monomer conversiond (%)

p-SiO2–g-PSe p-SiO2–g-PMMAf p-SiO2–g-PEAg p-SiO2–g-PBAh

11.4 12.7 15.3 13.4

57.1 63.3 76.5 66.9

100 100 100 100

a b c d e f g h

Irradiation dose: 4 Mrad; monomer/silica = 20 wt%. Percent grafting = weight of grafting polymer/weight of silica. Grafting efficiency = weight of grafting polymer/weight of grafting polymer and homopolymer. Monomer conversion = weight of polymer/weight of monomer. PS: polystyrene. PMMA: polymethyl methacrylate. PEA: polyethyl acrylate. PBA: polybutyl acrylate.

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

280

210

1: p-SiO2 as-received 2: p-SiO2-g-PS 3: p-SiO2-g-PEA Adsorption curve Desorption curve

Differential pore volume [cm3/g]

Adsorbed volume[cm3/g]

350

1

140

2

70 3

0 0.0

0.2

0.4

0.6

0.8

1.0

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

weights of the grafting polymers are about 105 and those of the homopolymers are 104. To have an image of the surface structure variation, nitrogen absorption isotherms of the nanoparticles were measured (Fig. 1). The curve profiles are something between type II and IV isotherms as viewed from Brunauer classification system [16]. Although hysteresis loops can be seen on the isotherms of the nanoparticles with and without the grafting polymers, those are not closed in the case of p-SiO2–g-PS and p-SiO2–g-PEA. It means that the capillary condensation occurring under higher relative pressure is severer in the grafted nano-silica. In comparison with p-SiO2 as-received, both p-SiO2–gPS and p-SiO2–g-PEA have significantly reduced absorbance. Evidently, it can be attributed to the fact that the micropores and macropores of the nanoparticle agglomerates are filled by the grafting polymers, which in turn demonstrates that the grafting polymers have been planted onto the nanoparticles. The plots showing the micropore size distribution in Fig. 2 also provide supporting evidence for the estimation. After grafting treatment, the pore volumes of the nanoparticles are about ten times smaller than that of the untreated

0.015

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

0.12

0.012

0.09

0.009

0.06

0.006

0.03

0.003

0.00 0.0

0.5

1.0

1.5

2.0

0.000 2.5

3 Differential pore volume [cm /g]

638

Pore diameter [nm]

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

nano-silica. In addition, the pore diameters corresponding to the maximum pore volumes of the treated nanoparticles shift from 0.55 to 0.71 nm. It seems that the graft polymerization was initiated at the smallest micropores of the nanoparticles. The porous structure of the particles has been changed accordingly. Figs. 3 and 4 show the typical fracture surfaces and mechanical properties of PP composites with nano-silica, respectively. The untreated nanoparticles are severely agglomerated in PP matrix (Fig. 3(a)), while the treated ones are well separated into tiny aggregates (Fig. 3(b)). The SEM observations evidence the authorsÕ expectation of graft treatment of the nanoparticles. It means that the composites containing the untreated nanoparticles are provided with heterogeneous microstructure and would exhibit worse reinforcing and toughening effects as compared to those having more homogeneous appearances due to graft pre-treament of the particles. As for the details of the composites mechanical performance, it is seen that both treated and untreated nano-silicas are able to stiffen the matrix, as reflected by the proportional relationship between YoungÕs modulus and the filler loading (Fig. 4(a)). Comparatively, the stiffness of the composites with grafted nano-silica is

Fig. 3. SEM microphotos of tensile fracture surfaces of nano-silica/PP composites (PP: T30S). (a) p-SiO2 as-received/PP (nano-silica content = 2.74 vol%); (b) p-SiO2–g-PS/PP (nano-silica content = 2.75 vol%).

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

1.8

p-SiO2 as-received/PP p-SiO2-g-PS/PP p-SiO2-g-PMMA/PP p-SiO-g-PEA/PP 2 p-SiO2-g-PBA/PP

(c) 400 Elongation to break [%]

Young's modulus [GPa]

(a) 2.0

1.6

1.4

1.2 0.0

0.5

1.0

1.5

2.0

300

200 p-SiO2 as-received/PP p-SiO2-g-PS/PP p-SiO2-g-PMMA/PP p-SiO-g-PEA/PP 2 p-SiO2-g-PBA/PP

100

0 0.0

2.5

0.5

39

38

37

36

0.5

1.0

1.5

2.0

(d) Area under stressstrain curve [MPa]

Tensile strength [MPa]

p-SiO2 as-received/PP p-SiO2-g-PS/PP p-SiO2-g-PMMA/PP p-SiO-g-PEA/PP 2 p-SiO2-g-PBA/PP

40

35 0.0

2.0

2.5

90

60 p-SiO2 as-received/PP p-SiO2-g-PS/PP p-SiO2-g-PMMA/PP p-SiO-g-PEA/PP 2 p-SiO2-g-PBA/PP

30

0 0.0

2.5

0.5

1.0

1.5

2.0

2.5

SiO2 [vol%] p-SiO2 as-received/PP p-SiO2-g-PS/PP p-SiO2-g-PMMA/PP p-SiO-g-PEA/PP 2 p-SiO2-g-PBA/PP

125

2

Impact strength [kJ/m ]

1.5

120

SiO2 [vol%] (e)

1.0

SiO2 [vol%]

SiO2 [vol%]

(b)

639

100

75

50

25

0 0.0

0.5

1.0

1.5

2.0

2.5

SiO2 [vol%] Fig. 4. Mechanical properties of nano-silica/PP composites as a function of silica content (PP: T30S): (a) YoungÕs modulus; (b) tensile strength; (c) elongation at break; (d) area under tensile stress–strain curve; and (e) impact strength.

lower than that of the untreated nanoparticles composites. As interfacial stress transfer efficiency depends on the stiffness of the interphase, higher interfacial stiffness favors improvement of the composites modulus [17]. The grafting polymer adhering to the nano-silica as well as the surrounding homopolymer establish a compliant interlayer between the particles and the matrix, and hence decrease the stiffening effect of the particles. On the other hand, it is interesting to note that the moduli of the composites with grafted nano-silica at identical particle content are arranged in the following order: p-SiO 2 –g-PS/PP > p-SiO 2 –g-PMMA/PP > p-SiO 2 –g-

PEA/PP > p-SiO2–g-PBA/PP. It exactly coincides with the flexibilities of the grafting polymer chains. That is, PBA possesses the highest flexibility and consequently masks the high stiffness of the particles in the composites to the greatest extent. The phenomenon again proves the above analysis concerning the dependence of composites moduli on interfacial stiffness. With respect to the composites under higher nano-silica content, it is found that the moduli increment of the composites with treated nano-silica can no longer keep the linearity. For p-SiO2–g-PEA/PP and p-SiO2– g-PBA/PP composites, their moduli even decline. This

640

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

should be the result of the increasing amount of the soft interphase comprised of the grafting polymers and homoploymers with a rise in the content of the grafted nanoparticles. Besides stiffening effect, the precipitated nano-silica can also provide PP with reinforcing effect at rather low filler concentration (Fig. 4(b)), which is similar to the situation of fumed nano-silica [1,2]. In the case of p-SiO2–g-PS/PP and p-SiO2–g-PBA/PP, for example, tensile strengths of the composites are higher than that of neat PP. This is different from what is observed in conventional micrometer particles/polymer composites, i.e., tensile strength of the composites remarkably decreases with the addition of the particulate fillers due to the poor bonding at the interface [18,19]. Jancar et al. [20] suggested that a strong filler/matrix adhesion would lead to enhanced strength of particulate composites. According to their consideration, it is known that the improvement of tensile strength of the composites exhibited in Fig. 4(b) should also be interpreted as the improvement of the interfacial interaction. Especially when PS- and PBA-grafted nano-silica is incorporated, the chain entanglement between the grafting polymers and the matrix polymer guarantees effective interfacial bonding over the whole filler content range of interests. However, grafting treatment does not always take effect as revealed by the composition dependence of tensile strength of p-SiO2–g-PMMA/PP and p-SiO2–g-PEA/ PP. This contradicts the results of PP composites with fumed nano-silica grafted by the same species of polymers [1]. It means that in addition to the nature of the grafting polymers on the nanoparticles, species of nanoparticles themselves and other unknown factors also greatly influence the reinforcing effect of the grafted nanoparticles. More detailed study in this direction is required to have a reasonable conclusion. From engineering point of view, elongation-to-break is an important parameter describing the rupture behavior of composite materials. The addition of mineral particulates into polymers used to lower it, even though the matrix has high impact toughness [21]. Fig. 4(c) clearly indicates that this is not the case when nano-silica is used. Either untreated or treated nano-silica is able to increase elongation-to-break of PP. Compared to the untreated and PS-grafted ones, acrylic polymers-grafted nano-silicas are more effective, in particular when p-SiO2–g-PMMA is concerned. It implies that in principle compliant grafting polymers can induce more matrix polymer to be involved in plastic deformation. However, chain flexibility of the grafting polymers on the nanoparticles is not the only prerequisite, otherwise pSiO2–g-PBA/PP should have the highest value of elongation-to-break. In contrast to tensile strength of the composites that needs strong filler/matrix bonding, the micro-deformation mechanism involved in elongationto-break of the composites depends upon extensionality

of the nanoparticles agglomerates in the composites [22]. Balanced viscoelasticity of the grafting polymers on the nanoparticles might thus be necessary for the improvement of elongation-to-break of the composites. Fig. 4(d) shows the areas under tensile stress–strain curves of the composites, another parameter characterizing the static toughness. The magnitude order of the values at given filler content is similar to that in Fig. 4(c), suggesting that the deformation features of the composites still govern their areas under tensile stress– strain curves. Grafted nano-silica again performs well than the untreated nanoparticles. The maximum value of p-SiO2–g-PMMA/PP is about 3.8 times higher than that of neat PP. The results demonstrate the role of grafting polymers, i.e., interconnecting the nanoparticles through chemical bonding and correlating the grafted nanoparticles with the matrix by chain entanglement. Under the applied force, plastic deformation of large amount of the matrix polymer beside the grafted nanoparticles is induced, leading to significantly high elongation-to-break and areas under tensile stress–strain curve of the composites. In the case of untreated nanoparticles filled PP, voiding and disintegration of the nanoparticle agglomerates are the main ways of energy dissipation due to the lack of extensionality. Unnotched impact strengths of the composites are given in Fig. 4(e) as a function of silica fraction. It is clear that after grafting treatment the toughening ability of the nanoparticles is greatly increased. At the silica content of 0.45 vol%, the impact ductility of p-SiO2–g-PBA/ PP is over 3 times higher than that of unfilled PP, and all the impact strengths of the composites at this filler loading are arranged in the order contrary to what observed in Fig. 4(a) concerning YoungÕs modulus: p-SiO2/ PP < p-SiO2–g-PS/PP < p-SiO2–g-PMMA/PP < p-SiO2>– g-PEA/PP < p-SiO2–g-PBA/PP. Since unnotched impact strength reflects the energy consumed by the plastic deformation prior to crack initiation, the above results manifest that the flexible macromolecular chains grafted onto the nanoparticles surfaces must have made contribution to this part of energy. In comparison to the data shown in Fig. 4(c) and (d), it is known that the dynamic toughness of the composites (i.e., impact strength) is more sensitive to the dispersion status of the grafted nanoparticles than the static toughness (i.e., elongation-to-break and area under tensile stress–strain curve). With a rise in the filler content, homogeneity of the particles distribution might be worse as a result of the increased viscosity of the composite systems. Under this circumstance, the grafting polymers could not response in step with the impact load and hence decline of impact strength is observed when the content of nano-silica exceeds 0.45 vol%. Since PP is a semi-crystalline polymer and its mechanical properties would change with the crystalline structure and crystallinity, the influence of the addition of the nanoparticles should be known. As listed in

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

641

Table 2 Kinetic parameters of isothermal crystallization of PP and its composites at 130 C Samples

t1/2a (min)

tfb (min)

nc

kd (· 10 3min n)

DHce (J/g)

PP p-SiO2 as-received/PP (nano-silica content = 0.86 vol%) p-SiO2–g-PS/PP (nano-silica content = 1.066 vol%) p-SiO2–g-PEA/PP (nano-silica content = 0.82 vol%)

8.68 4.69 6.17 9.79

17.8 8.89 11.39 17.7

2.59 3.04 3.04 2.92

2.57 6.31 2.75 1.08

97.1 97.0 95.0 97.0

a b c d e

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

Table 2, the untreated nano-silica remarkably accelerates the crystallization of PP matrix as revealed by the values of t1/2, tf and k. This nucleating activity can be explained by the thermodynamic model proposed by Ebengou [23]. That is, when PP chains were absorbed on the silica surface, the configurational entropy of the entire chain decreased, forming a nucleus of a certain volume within the adsorbed chains costs less energy. In the case of grafted nanoparticles, the nucleation effects are less profound because the grafting polymers shielded the nanoparticles from the direct contacts with PP, which coincides with the results of PP/elastomer/ particles composites characterized by core-shell microstructures [24]. Fig. 5 further shows the melting behaviors of PP and its composites. In comparison with neat PP, the endothermic peak profiles and temperatures of the composites are almost the same, suggesting the crystal microstructure of PP has not been changed. Relatively, the crystallinity of PP in the composites is slightly reduced. On the whole, however, it cannot be concluded that the aforesaid mechanical performance variations of PP composites with the incorporation of silica nanoparticles results from the variations in PP crystalline structure and crystallinity. Although un-

Endo >

Samples 1 2 3 4

o

Tm ( C) 166.5 166.4 166.5 166.1

Xc (%) 46.6 45.7 45.4 44.5

4 3 2

1

140

150

160

170

180

190

200

Temperature [oC] Fig. 5. DSC heating traces of PP and its composites having been isothermally crystallized at 130 C (PP: T30S). (1) Neat PP; (2) p-SiO2 as-received/PP (nano-silica content = 0.86 vol%); (3) p-SiO2–g-PS/PP (nano-silica content = 1.06 vol%); (4) p-SiO2–g-PEA/PP (nano-silica content = 0.82 vol%). Tm: peak melting temperature; Xc: crystallinity.

treated nano-silica has significant nucleating ability, it fails to improve PP properties equivalently. Similarly, the grafted nanoparticles only lead to marginal decrease of PP crystallinity, which is also out of proportion to the performance enhancement. 3.2. Effects of matrix ductility and nanoparticle species Mechanical performance of composite materials is a function of filler and matrix characteristics. When studying nano-CaCO3/PP composites, for example, Ren et al. [25] showed the importance of the matrix toughness. A much more remarkable increase in impact resistance was observed in the composites with a PP copolymer possessing higher ductility as matrix, while the same nanoparticles did not result in a similar improvement in a PP homopolymer with lower ductility. As a result, it was suggested that the polymer to be toughened by nanoparticles should possess at least a certain toughness, which is different from the case when elastomer acts as toughener. To find out whether nano-silica/PP composites follow the same law, three types of PP: EPS30R, T30S and PP700, are used in the current work. Relatively, EPS30R has the highest toughness and PP700 the lowest. For making comparative tests, the crosshead speed of the tensile tests was raised to 200 mm/min to match the large failure strain of EPR30S. As illustrated by Fig. 6(a) and (b), the addition of nano-silica into either EPR30S or PP700 results in reduced values of elongation-to-break and area under tensile stress–strain curve, no matter the particles have been treated or not. It resembles the behavior of micro-particles filled polymer composites. However, nano-silica filled T30S has acquired significantly high static toughness. The untreated nanoparticles have already certain ability to increase elongation-to-break and area under tensile stress–strain curve of T30S PP. It means that under tensile loading, considerable matrix yielding of nano-silica filled PP occurs only in the case of matrix polymer possessing moderate toughness. The synergetic effect brought about by the grafted nanoparticles depends on the matching of matrix toughness and flexibility of the grafting

642

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

Fig. 6. Mechanical properties of nano-silica/PP composites with different PP as matrices: (a) elongation at break; (b) area under tensile stress–strain curve; and (c) impact strength. The percentage numerals quantify the relative variation of the mechanical property with reference to the value of the corresponding neat PP. Nano-silica content = 0.5 vol%, crosshead speed of the tensile tests = 200 mm/min.

polymers. For the composites based on EPR30S and PP700, the nanoparticles are ineffective to induce localized matrix drawing. Fig. 6(c) shows the impact strength of the composites. Although EPS30R and its composites are too tough to be broken by the impact load under the current testing conditions, some useful hints can still be yielded by comparing the performance of T30S and PP700 based composites. It is interesting to see that the nanoparticles exert toughening effect on both types of PP. However, the relative increments of the impact strengths of T30S based composites are much higher than those of

PP700 based composites. Relating the results of Fig. 6(a)–(c) to each other, it is clear that the matrix ductility or the capability of the matrix to plastically deform is a key factor influencing the toughening effect of the nanoparticles. In the case of suitable matrix ductility like T30S, large scale of matrix polymer is successfully involved in plastic deformation as induced by the nanoparticles, leading to high impact toughness of the composites. On the other hand, it has been known that the interface between the filler particles and the matrix in a polymer nanocomposite constitutes a much greater area within the bulk material as compared with conventional composites containing micrometer-sized particles, and hence influences the compositeÕs properties to a much greater extent. However, there is no information about the effect of nanoparticles having different sizes. In the current work, fumed nano-silica (15 nm) acts as a reference for the precipitated nano-silica (10 nm). To bring the positive effect of the nanoparticles into full play, both types of the particles were grafted with PBA (the percent grafting of f-SiO2–g-PBA is 8.26%, close to that of p-SiO2–g-PBA: 13.4%) and compounded with PP (T30S) under the same conditions. First YoungÕs moduli of the composites are compared in Fig. 7(a). The filler content dependences of modulus of the two types of composites are almost the same, but the precipitated nano-silica leads to higher compositesÕ stiffness. It must result from the increased interfacial area in the composites with finer nano-silica, which promotes the stress transfer efficiency within small strain range. The reduction of YoungÕs modulus at higher silica content manifests the softening effect of the compliant PBA interlayer, which becomes more and more evident with a rise in content of the grafted nanoparticles. With respect to tensile strength of the composites, on the contrary, the larger particles (fumed silica) seem to have more remarkable reinforcing ability (Fig. 7(b)). Generally, composites strength depends on the filler/matrix bonding under large strain condition. Since the present nanoparticles are grafted with the same polymer (i.e., PBA) and the amounts of grafting polymer are similar in the two types of nano-silica, the interfacial adhesion in the composites should be almost identical at given filler concentration [26]. The significant difference in the tensile strengths of the composites might be attributed to the load bearing ability of the nano-silica itself. That is, the fumed nano-silica particles might be stronger than the precipitated ones. To the authorsÕ knowledge, the strength of nano-silica particles has not yet been reported because of the experimental difficulties, but one might find some traces on the basis of their synthesis processes. Precipitated silica particles are obtained in low-temperature wet process and might form soft agglomerates (more physical bonding) during drying.

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

1.7

(b) 40 p-SiO2-g-PBA/PP f-SiO2-g-PBA/PP

Tensile strength [MPa]

Young's modulus [GPa]

(a) 1.8

1.6

1.5

1.4

1.3 0.0

0.5

1.0

1.5

2.0

39

38

37

36 0.0

2.5

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

0.5

(c) 300

1.5

2.0

2.5

(d) 90

Area under stressstrain curve [MPa]

Elongation at break [%]

1.0

SiO2 [vol%]

SiO2 [vol%]

250

200

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

100

50 0.0

643

0.5

1.0

75

60

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

30

1.5

2.0

15 0.0

2.5

0.5

SiO2 [vol%]

1.0

1.5

2.0

2.5

SiO2 [vol%]

2

Impact strength [kJ/m ]

(e) 140

105

70

35

0 0.0

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

0.5

1.0

1.5

2.0

2.5

SiO2 [vol%] Fig. 7. Mechanical properties of PP composites filled with different nano-silicas as a function of silica content (PP: T30S): (a) YoungÕs modulus; (b) tensile strength; (c) elongation at break; (d) area under tensile stress–strain curve; and (e) impact strength.

As for fumed silica particles, they are produced in hightemperature gaseous process, which would made them form hard agglomerates (more chemical bonding). During grafting pre-treatment, the former particle agglomerates are easier to be broken apart so that the grafted precipated nano-silicas can be well dispersed than the fumed ones. When the composites are subjected to applied stress, the intrinsic high strength of the fumed silica aggregates offers higher strength for the composites. In contrast to YoungÕs modulus and tensile strength, elongation-to-break and area under tensile stress–strain curve of the composites shown in Fig. 7(c) and (d) indi-

cate that the species of the nanoparticles nearly have nothing to do with the two parameters when the filler content 60.84 vol%. The characteristics of the grafting polymer play the leading role in this case. This is reasonable because under tensile loading the extent of matrix stretching is mainly related to the chain entanglement between the grafting polymers and the matrix molecules. For the composites with nano-silica content higher than 0.84 vol%, the decrease in elongation-to-break and area under tensile stress–strain curve of f-SiO2–g-PB/PP should be due to the poor dispersion of the nanoparticles. Compared with precipitated nano-silica, fumed

644

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645

nano-silica has much lower packing density. As a result, homogeneous distribution of fumed nano-silica during compounding is more difficult to be achieved than precipitated nano-silica especially at higher filler content. It explains that the two parameters reflecting static toughness maintain almost unchanged over the whole filler content range of interests in the case of p-SiO2– g-PB/PP (Fig. 7(c) and (d)). When the composites specimens are subjected to impact load, however, different nanoparticles give different responses (Fig. 7(e)). Similar to the results in Fig. 7(a), precipitated nano-silica filled PP has higher impact strength than fumed nano-silica/PP in most cases. This implies that besides the interphase formed by the grafting polymers and matrix, the interfacial area between the nanoparticles and the surrounding polymers is also important to the toughening effect. The latter factor might facilitate the generation of crazes during the impact test, absorbing certain amount of the input energy additionally. Therefore, the composites filled with nanosilica with smaller particle size exhibit greater resistance to impact loading.

4. Conclusions Precipitated nano-silica is able to provide PP with stiffening, reinforcing and toughening effects at rather low filler concentration as fumed nano-silica. Having been grafted with different polymers onto the surfaces in terms of gaseous graft polymerization, the nanoparticles become more efficient to improve the strength and toughness of the composites. The nature of the grafting polymer chains plays an important role in the properties enhancement. Ductility of the matrix PP determines the toughening effect of the nanoparticles. Only in the case of moderate matrix ductility, the composites can receive the highest extent of toughness increase. Besides, the size and surface area of the nanoparticles are also important influencing factors. The smaller nanoparticles lead to higher YoungÕs modulus and impact strength of the composites, and reduce the sensitivity of the static toughness to the status of filler distribution.

Acknowledgements The authors are grateful to the support of the Deutsche Forschungsgemeinschaft (DFG FR675/40-4) for the cooperation between the German and Chinese institutes on the topic of nanocomposites. Further thanks are due to the National Natural Science Foundation of China (Grant No. 50133020), the Team Project of the Natural Science Foundation of Guangdong, China

(Grant No. 20003038), and the Key Program of the Science and Technology Department of Guangdong, China (Grant No. A10172).

References [1] Rong MZ, Zhang MQ, Zheng YX, Zeng HM, Walter R, Friedrich K. Structure–property relationships of irradiation grafted nano-inorganic particle filled polypropylene composites. Polymer 2001;42:167–83. [2] Wu CL, Zhang MQ, Rong MZ, Friedrich K. Tensile performance improvement of low nanoparticles filled-polypropylene composites. Compos Sci Technol 2002;62:1327–40. [3] Zhang MQ, Rong MZ, Zhang HB, Friedrich K. Mechanical properties of low nano-silica filled high-density polyethylene composites. Polym Eng Sci 2003;43:490–500. [4] Wang G, Chen XY, Huang R, Zhang L. Nano-CaCO3/polypropylene composites made with ultra-high-speed mixer. J Mater Sci Lett 2002;21:985–6. [5] Zhang MQ, Rong MZ, Pan SL, Friedrich K. Tensile properties of polypropylene filled with nanoscale calcium carbonate particles. Adv Compos Lett 2002;11:293–8. [6] Rong MZ, Zhang MQ, Zheng YX, Zeng HM, Walter R, Friedrich K. Irradiation graft polymerization on nano-inorganic particles: an effective means to design polymer based nanocomposites. J Mater Sci Lett 2000;19:1159–61. [7] Chung SC, Hahm WG, Im SS, Oh SG. Poly(ethylene terephthalate) (PET) nanocomposites filled with fumed silicas by melt compounding. Macromol Res 2002;10:221–9. [8] Zhang J, Wang X, Lu L, Li D, Yang X. Preparation and performance of high-impact polystyrene (HIPS)/nano-TiO2 nanocomposites. J Appl Polym Sci 2003;87:381–5. [9] Xiong MN, Wu LM, Zhou SX, You B. Preparation and characterization of acrylic latex/nano-SiO2 composites. Polym Int 2002;51:693–8. [10] Jana SC, Jain S. Dispersion of nanofillers in high-performance polymers using reactive solvents as processing aids. Polymer 2001;42:6897–905. [11] Rong MZ, Zhang MQ, Zheng YX, Zeng HM, Friedrich K. Improvement of tensile properties of nano-SiO2/PP composites in relation to percolation mechanism. Polymer 2001; 42:3301–4. [12] Zhang MQ, Rong MZ, Zeng HM, Schmitt S, Wetzel B, Friedrich K. An atomic force microscopy study on structure and properties of irradiation grafted silica particles in polypropylene based nanocomposites. J Appl Polym Sci 2001;80:2218–27. [13] Torro´-Palau AM, Ferna´ndez-Garcı´a JC, Orgile´s-Barcelo´ AC, Martı´n-Martı´nez JM. Characterization of polyurethanes containing different silicas. Int J Adhes Adhes 2001;21:1–9. [14] Wang G, Zhang L, He G. Chinese Patent (CN1046921); 1996. [15] Wu CL. PhD thesis, Zhongshan University; 2002. [16] Brunauer S, Deming L, Deming W, Teller E. On a theory of the van der Waals adsorption of gases. J Am Chem Soc 1940;62:1723–32. [17] Kerner EH. Electrical conductivity of composites media. Proc Phys Soc B 1956;69:802–7. [18] Nielsen LE. Simple theory of stress–strain properties of filled polymer. J Appl Polym Sci 1966;10:97–103. [19] Nicolais L, Narkis M. Stress–strain behavior of styrene–acrylonitrile/glass bead composites in the glassy region. Polym Eng Sci 1971;11:194–9. [20] Jancar J, Dianselmo A, Dibenedetto AT. The yield strength of particulate reinforced thermoplastic composites. Polym Eng Sci 1992;32:1394–9.

C.L. Wu et al. / Composites Science and Technology 65 (2005) 635–645 [21] Manson JA, Sperling LH. Polymer blends and composites Plenum Press. New York; 1976. [22] Wu CL, Zhang MQ, Rong MZ, Lehmann B, Friedrich K. Deformation characteristics of nano-SiO2 filled polypropylene composites. Polym Polym Compos 2003;11:559–62. [23] Ebengou RH. Adsorption as a mechanism for nucleating activity: a thermodynamic explanation. J Polym Sci Polym Phys 1997;35:1333–8.

645

[24] Long Y, Shanks RA. PP/elastomer/filler hybrids. II. Morphologies and fracture. J Appl Polym Sci 1996;62:639–46. [25] Ren X, Bai L, Wang G. Reinforcement and toughening of polypropylene composites by nanoparticle CaCO3. Chem World 2000;41(2):83–7. [in Chinese]. [26] Rong MZ, Zhang MQ, Pan SL, Lehmann B, Friedrich K. Analysis of the interfacial interaction in polypropylene/silica nanocomposites. Polym Int 2004;53:176–83.