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Tensile performance improvement of low nanoparticles filled-polypropylene composites
Composites Science and Technology 62 (2002) 1327–1340 www.elsevier.com/locate/compscitech
Tensile performance improvement of low nanoparticles filled-polypropylene composites Chun Lei Wua, Ming Qiu Zhangb,*, Min Zhi Rongb, Klaus Friedrichc a
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 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 prepared by conventional compounding technique in which the nanoparticles are pre-grafted by some polymers using irradiation. To examine the applicability of the approach, a tougher polypropylene (PP) was compounded with nano-silica by industrial-scale twin screw extruder and injection molding machine in the present work. The results of tensile tests indicated that the nanoparticles can simultaneously provide PP with stiffening, strengthening and toughening effects at a rather low filler content (typically 0.5% by volume). The presence of grafting polymers on the nanoparticles improves the tailorability of the composites. Due to the viscoelastic nature of the matrix and the grafting polymers, the tensile performance of the composites filled with untreated and treated nanoparticles is highly dependent on loading rate. With increasing the crosshead speed for the tensile tests, the dominant failure mode 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 fillers are added to polymers in commercial production primarily for the reasons of cost reduction and stiffness improvement [1,2]. Although most studies dealing with the modification of semi-crystalline polymers with inorganic particulates reported embrittling effects by comparing ultimate elongation and impact strength of composite materials with those of unfilled resins [3–5], some researchers showed the enhancement of toughness in rigid particles filled polypropylene [6,7] and polyethylene [8,9]. It is worth noting that in the case of micrometer-sized particulates, high filler content (typically higher than 20% by volume) is generally required to bring the above-stated positive effects of the fillers into play. This would detrimentally affect some important properties of the matrix polymers such as processability, appearance, * Corresponding author. Tel.: +86-20-8403-6576; fax: +86-208403-6564. E-mail address: [email protected] (M.Q. Zhang).
density and ageing performance. Therefore, a composite with improved performance and low particle concentration is highly desired. With regard to this, the newly developed nanocomposites would be competitive candidates. The extremely high surface area is one of the most attractive characteristics of nanoparticles because it facilitates creating a great amount of interphase in a composite. Introduction of nanoparticles into a polymer changes the intermolecular interaction of the matrix [10]. As estimated by Reynaud et al. [11], an interphase 1 nm thick represents roughly 0.3% of the total volume of polymer in the case of microparticle filled composites, whereas it can reach 30% of the total volume in the case of nanocomposites. That is, the non-negligible contribution made by the interphase provides diverse possibilities of performance tailoring, and is able to influence the properties of matrices to a much greater extent under a rather low nano-filler loading. The crux of the matter lies in that how to well distribute nanoparticles over a polymer matrix and how to improve nanoparticles/matrix interaction.
0266-3538/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(02)00079-9
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From a practical point of view, dispersive mixing in preparing polymer based particulate composites has important technical meaning. However, a homogenous dispersion of nanoparticles in a polymer is very difficult by using the existing compounding techniques due to the strong tendency of the fine particles to agglomerate and the high melt viscosity of the matrix. In many cases, the so-called nanoparticles filled polymers contain a number of loosened clusters of nanoparticles, as demonstrated by the work of Jana and Jain [12] dealing with untreated nanosilica/polyethersulphone composites. When the composites are subjected to force, the nanoparticle agglomerates can be split easily and a premature failure of the materials would thus take place [12–15]. To overcome this dilemma and to give full play to nanoparticles, we used irradiation grafting polymerization method to modify nanoparticles first, and then the treated particles were mechanically mixed with a polymer as usual [16]. Owing to the low molecular weight nature, the grafting monomers can penetrate into the agglomerated nanoparticles easily and react with the activated sites of the nanoparticles inside as well as outside the agglomerates [17]. As a result, the following effects can be obtained [18]: (i) Hydrophobicity of the nanoparticles is increased, facilitating the filler/matrix miscibility; (ii) filler/matrix interaction is enhanced through the entanglement between the grafting polymer and the polymer matrix; (iii) the nanoparticle agglomerates become stronger because they turn into a nanocomposite microstructure comprising the nanoparticles and the grafted, homopolymerized polymer; (iv) the interfacial characteristics between the treated nanoparticles and the matrix polymer can be tailored by changing the species of the grafting monomers and the grafting conditions. In this context, a uniform dispersion of nanoparticles in the matrix might no longer be critical. Mechanical testing of polypropylene filled with nanoSiO2 [18] and nano-CaCO3 [19] demonstrated the feasibility of the above approach. Only a small amount of modified nanoparticles (typically less than 3% by volume) can effectively improve modulus, strength, toughness and thermal deformation temperature of the matrix polymer. Such an improvement in overall properties of polymers can scarcely be observed in conventional microparticulate composites. It was found that the deformation habit but not the crystallization characteristics of the matrix polymer remarkably change with the addition of the treated nanoparticles. To explain the specific influence generated by the nanoparticles at low-filler loading regime, a double percolation of stress volumes, characterized by the appearance of connected shear yielded networks throughout the composite, was proposed [20]. Considering that the polypropylene used in our previous works [18–20] is a brittle type and the composites
were prepared with a lab-scale single screw extruder and compression molding, the results might not have adequate applicability. Therefore, a commercial polypropylene with higher toughness was compounded with nano-silica by means of industrial-scale twin screw extruder and injection molding machine in the current work. Tensile performance and fractured surfaces are analyzed as a function of particulate treatment, filler content and crosshead speed to reveal the structure– property relationships of the composites, and the mechanical 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 flow index of 3.2 g/10 min (2.16 kg at 230 C). Fumed silica with an average primary particle size of 15 nm and a specific surface area of 374 m2/g was produced by Shenyang Chemical Engineering Ltd., China. Commercial monomers, styrene and ethyl acrylate, were used as grafting monomers without further purification. 2.2. Irradiation grafting of nano-SiO2 Modification of nano-silica proceeded according to the following steps. The nanoparticles were pretreated at 140 C under vacuum for 6 h to eliminate possible absorbed water on the surface of the particles. Then a mixture of nanoparticles/monomer (100/20 by weight) and a certain amount of n-hexane was irradiated by 60 Co g-ray under atmosphere at room temperature. After exposure to a dose of 4 Mrad, the solvent was recovered, and the dried residual powder was available for the subsequent compounding. 2.3. Characterization of the irradiation grafted products 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) and a Bruker Equinox 55 Fourier transform infrared spectroscope (FTIR), the weight of the grating polymer and the chemical structure of the modified nanoparticles were characterized, respectively. To further
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separate the grafting polymer from the treated nanoparticles, nano-silica accompanied with the unextractable grafting polymer was immersed in 20% HF solution for 72 h to remove the inorganic particles. The molecular weights of the grafting and the homopolymerized polymers were determined by a Waters 991 gel permeation chromatography (GPC), with tetrahydrofuran as the solvent. To observe the morphologies of the nanoparticles, untreated SiO2 and grafted SiO2 (without homopolymer) were added into ethanol and toluene to prepare 0.001 g/ml solutions, respectively. With the aid of sonication for 30 min, the solutions were transferred to copper gauzes by droppers. After the evaporation of the solvents, a JEM-100CXZ transmission electron microscopy (TEM) was used to examine the appearance of the particles. 2.4. Composites preparation and characterization 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) 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 at crosshead speeds of 10, 30, 50, 100 and 500 mm/min, respectively. Five samples were tested for each case. The fractured surfaces of the samples were observed with a Hitachi S-520 scanning electron microscope (SEM) at an accelerating voltage of 20 kV.
3. Results and discussion 3.1. Effect of irradiation grafting polymerization on nano-SiO2 Since the present work aims to study the effect of modified nano-silica on the mechanical behavior of PP composites, variation in the chemical structure of the particles should be known at the very beginning of the discussion. FTIR spectra of untreated and treated nanosilica are shown in Fig. 1. To eliminate the influence of homopolymers, both polystyrene-grafted nano-SiO2 (SiO2-g-PS) and polyethyl acrylate-grafted nano-SiO2 (SiO2-g-PEA) used for the FTIR examinations were separated from the homopolymers in advance. In comparison with the spectrum of SiO2 as-received, the adsorptions at 690, 1460 and 2960 cm 1 appearing in
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Fig. 1. FTIR spectra of untreated SiO2 and grafted SiO2.
the spectrum of SiO2-g-PS represent the bending mode of C–H in benzene rings and the stretching modes of C–C and C–H, respectively. In addition, the band at 1725 cm 1 in the spectrum of SiO2-g-PEA indicates the existence of carbonyl groups. These prove that polystyrene (PS) and polyethyl acrylate (PEA) have been chemically connected to the surface of nano-silica during the irradiation polymerization processes as expected. On the basis of above qualitative analysis, the quantitative results of the grafting polymerization on nanosilica are given in Table 1. It is seen that when other conditions being equal, the percentage grafting and the grafting efficiency of styrene onto the nanoparticles are much higher than those of ethyl acrylate, while the monomer conversion of styrene is lower than that of ethyl acrylate. This reflects the difference in reactive feature between the monomers. In general, room temperature irradiation grating polymerization onto inorganic particles is controlled by the mechanism of free radical polymerization. Under the same irradiation dose, acrylic monomers would generate much more radicals than styrene, and moreover, ethyl acrylate radicals have higher activity than styrene radicals [21]. Therefore, the conversion of ethyl acrylate is superior to that of styrene, leading to higher homopolymer fraction of PEA than PS. On the other hand, partial surface of nanosilica might be connected with ethyl acrylate in the solvent through hydrogen bonding prior to the irradiation processing. When the composite system is exposed to the irradiation, radicals can be formed on both the nanoparticles and the ethyl acrylate molecules connected to the particles. The latter would result in either grafting polymerization or homopolymerization. That is, the amount of grafted PEA has to be lower as compared with the styrene/silica mixture where no chemical connection between the monomers and the particles is established before the irradiation. It should be responsible for the difference in percentage grafting between ethyl acrylate/silica and styrene/silica.
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Table 1 Results of irradiation grafting polymerization onto nano-silica Monomer
Monomer conversiona (%)
Percent graftingb (%)
Grafting efficiencyc (%)
MWd of grafting polymer (104)
MW of homopolymer (104)
Styrene Ethyl acrylate
52.8 93.5
4.64 1.56
43.9 10.5
1.3 (de=1.70) 7.1 (d=2.13)
1.1 (d=1.46) 1.1 (d=1.33)
a b c d e
Monomer conversion=weight of polymer/weight of monomer. Percent grafting=weight of grafting polymer/weight of nano-SiO2. Grafting efficiency=weight of grafting polymer/weights of grafting polymer and homopolymer. MW=weight average molecular weight. d=Molecular weight polydispersity index.
By further examining the data listed in Table 1, it can be found that the grafting polymers attached to the nanoparticles possess higher molecular weights and broader molecular weight distributions than the homopolymers. Similar phenomena were also reported by Fukano and Kageyama [22] when they studied radiation grafting of styrene onto silica-gel. These can be attributed to the characteristics of the grafting reaction. That is, the reaction was a typical heterogeneous one because the activated sites were created by irradiation on the nanoparticles and the chain growth of the grafting polymers had to proceed in solid–liquid state. The mobility of the growing chains was thus worse as compared with a homogeneous reaction in liquid where both ends of the macromolecular chains can move freely. As a result, the probability of chain termination between the radicals became relatively difficult, leading to higher molecular weight of the grafting polymers. In addition, because the radicals are not simultaneously produced on the nanoparticles, the grafting polymers formed at earlier stage would wrap the surface of the nanoparticles and reduce the probability of collision between the monomers and the radicals formed on the particles at latter stage. This accounts for the higher molecular weight polydispersity indexes of the grating polymers. Morphologies of the nanoparticles before and after grafting polymerization are illustrated in Fig. 2. A chain-like branched structure of the agglomerated SiO2 particles [17] can be observed in the solutions. The smallest perceivable units are approximately 15 nm in diameter in the case of untreated nanoparticles [Fig. 2(a)]. When grafting polymers are introduced onto the particles, the sizes of the agglomerates become larger and the edges are no longer clearly discernible [Fig. 2(b, c)]. Such a change demonstrates the role of the grafting polymer, i.e. separating and connecting the nanoparticles. To estimate the thickness of the polymer layer adhered to the particles, the data of SiO2-g-PS is used as an example, i.e. percent grafting=4.64%, density
Fig. 2. TEM microphotos (magnification=105) of (a) SiO2 asreceived, (b) SiO2-g-PS, and (c) SiO2-g-PEA in solvents. In the latter two specimens, homopolymers were removed in advance.
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of PS=1.05 g/cm3, specific surface area of the silica=374 m2/g. Supposing a complete coverage on each silica particles, the thickness of the grafting PS should equal to 0.0464/(374 m2/g1.05 g/cm3)0.12 nm. Evidently, this is a reasonable value as evidenced by the TEM photos in Fig. 2. Due to the low percent grafting and the thin grafting polymer layer, there are still many unreacted hydroxyl groups on the surface of nano-SiO2, which is responsible for the appearance of the larger agglomerates of the grafted particles in the solvent. 3.2. Tensile properties of the composites Fig. 3 shows the results of tensile testing of PP reinforced by nano-silica as a function of filler content, determined at a moderate crosshead speed of 50 mm/ min. Although both the treated and the untreated nanoparticles can impart the high stiffness of the fillers to the matrix polymer as expected, the composites incorporated with the modified particles exhibit lower modulus over the whole range of filler loading of interests [Fig. 3(a)]. Usually the capability of composite interface to transfer elastic deformation depends to a great extent upon the interfacial stiffness and static adhesion strength [23,24]. A high interfacial stiffness corresponds to a high composite modulus. Since the grafting polymers and the homopolymers introduced onto the nanoparticles form a relatively compliant interlayer at the particles/matrix interface, the high
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stiffness of the particles has to be masked under the low stress level [9] and the composites have to show lower modulus as compared with the case of untreated SiO2 composites. With a rise in filler content, the increased amount of grafting polymers further increases the modulus mismatching of the filler and the matrix, and reduces the stiffening effect of SiO2, leading to the drop in Young’s moduli of the composites at high filler regime. In addition, due to the higher rigidity of PS molecules than PEA, the interfacial elastic stress is less efficiently transferred in SiO2-g-PEA/PP composites than in SiO2g-PS/PP, especially when the fraction of the modified nanoparticles is high. In contrast, an approximately linear composition dependence of tensile modulus is perceived in the composites filled with SiO2 as-received [Fig. 3(a)]. This can be interpreted as the absence of a soft interphase and the appearance of larger agglomerates of the nanoparticles in the matrix as a result of poor filler/matrix miscibility. Under the same filler volume fraction [25], the latter effect would provide higher load carrying ability within small strain range. The above results factually reflect the contradiction between the stiffening effect of the rigid particles and the weakening effect of the soft interlayer. Fig. 3(b) gives the tensile strengths of the composites versus SiO2 content. Clearly, both untreated and treated nanoparticles exhibit the strengthening ability. It is well known that the tensile strength of a particulate composite
Fig. 3. Tensile properties of PP composites as a function of nano-SiO2 volume fraction: (a) Young’s modulus, (b) tensile strength, (c) elongation to break, and (d) area under stress–strain curve. Crosshead speed=50 mm/min.
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Fig. 4. Tensile properties of PP composites as a function of crosshead speed: (a) Young’s modulus, (b) tensile strength, (c) elongation to break, and (d) area under stress–strain curve.
Fig. 5. Young’s modulus of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.
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is usually reduced with filler content following a power law in the case of a poor filler/matrix bonding [26,27]. That means, the strength of the composite cannot be greater than that of the unfilled version because the filler particles do not bear any fraction of the external load. This contradicts the results shown in Fig. 3(b). In fact, when bonding between fillers and matrix is strong enough, the tensile yield strength of a particulate composite can be higher than that of the matrix polymer [28,29]. Although these models were developed based on the cases of microsized particulate composites, they are still valid for the explanation of the composites filled with nanoparticles [20]. Therefore, the extremely effective improvement of tensile strength of the composites with grafted nano-SiO2 should result from chain interdiffusion and entanglement between the macromolecules of the grafting polymers and the matrix. It is worth noting that when the amount of the grafted nanoparticles is increased (e.g. > 0.5 vol%), the content of compliant PEA chains is raised accordingly in the interlayer and the interfacial stress transfer efficiency has to be decreased. This accounts for the relatively low strength of SiO2-g-PEA/PP composites at higher particulate loading.
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By comparing Fig. 3(b) with the data of Ref. [18], the most distinct difference lies in the results of untreated SiO2/PP. That is, for the composites prepared by a labscale single screw extruder and compression molding, the addition of untreated SiO2 lowers the tensile strength of PP in the lower loading region but then leads to a slight increase in strength when the particle fraction reaches 4.68 vol% [18]. With respect to the SiO2/PP composites of the present work, which were manufactured through twin screw extrusion and injection molding, a continuous increase in the strength with SiO2 content is detected [Fig. 3(b)]. So far as we know, corotating twin screw extruders are able to provide more sufficient homogenization in comparison with single screw extruders. The above difference manifests that the important role of even distribution of untreated nanoparticles in the composites. For SiO2 as-received, the more particles are exposed to the matrix polymer, the more possibly the interaction between the particles and the matrix can be enhanced. So, an improved homogeneity of the untreated SiO2/PP composites would certainly be beneficial to the stress transfer. This again shows the significance of grafting modification 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.
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of composites strength performance to the dispersion state of the particles. As the grafted nanoparticle agglomerates turn into a nonocomposite microstructure consisting of the particles and the grafted, homopolymerized secondary polymer [18], they are brought into play as an integral when the composites are subjected to the applied force [20]. The situation is completely different from the composites with untreated nanoparticles, in which agglomerated particles have to be deagglomerated as much as possible to reduce the probability of premature failure. As a result, a uniform dispersion of the nanoparticles in the matrix is absolutely necessary for obtaining the reinforcing effect in case untreated nanoparticles are used. Failure strain can partially assess the rupture behavior of a composite material. The incorporation of particulate fillers usually results in a decrease in this parameter regardless of the interfacial adhesion [26]. It is true even in the system exhibiting impressive impact toughness improvement with the addition of mineral fillers [8]. However, the plots shown in Fig. 3(c) demonstrate that the values of elongation-to-break of
PP can be significantly increased by using nano-SiO2,, implying a failure mechanism different from those involved in conventional composites. Comparatively, SiO2-g-PEA is able to provide a stable improvement over the entire filler content range of interests. The reduction in elongation-to-break of the composite filled with untreated SiO2 suggests that the fillers cause a reduction in matrix deformation due to an introduction of mechanical restrains. In contrast, the improvement of elongation-to-break with the incorporation of the grafted nanoparticles is a result of interfacial viscoelastic deformation and matrix yielding. Evidently, the grafted PEA makes more effective contribution. The area under the tensile stress–strain curve can more reasonably characterize the toughness potential of the composites than elongation-to-break under static tensile loading conditions [30]. As confirmed by Fig. 3(d), the grafted nanoparticles indeed improve the ductility of PP at a silica content as low as 0.5 vol%. The results are somewhat opposite to those observed in Ref. [18], which reports a deteriorated effect of grafting PEA on 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.
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that SiO2-g-PEA employed in the present work possesses almost the same percent grafting and grafting efficiency as that in Ref. [18], it can be concluded that the molecular entanglement between the grafting PEA and matrix PP was not sufficiently formed during melt compounding when the single screw extruder was used. As a result, localized plastic deformation or matrix drawing cannot be efficiently induced by SiO2-g-PEA as in 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 at higher filler content regime are lower than that of the neat PP. It implies that the short range interaction at SiO2/PP interface is not good at inducing plastic deformation of the matrix polymer. Due to the encounter of the propagating neck with a larger agglomerate, particularly in the case of higher filler content, final failure of the composites might be initiated easily [8]. Obviously, the resultant embrittling effect can be prevented by the application of grafted nanoparticles. Since polymer composites maintain the viscoelasticity of polymers, the dependence of tensile properties on crosshead speed should be known for engineering purposes. As
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can be seen from Fig. 4(a), a linear increase in Young’s modulus with increasing crosshead speed is valid for all the specimens. The presence of silica leads to the values of modulus being higher than that for the unfilled PP and being slightly less speed dependent. These observations 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 strengths increase in principle, there is a significant transition in the slope. Accordingly, the first derivation of the strength in Fig. 4(b) with respect to crosshead speed would yield a peak between 50 and 100 mm/min, which clearly corresponds to an energy-activated process in tensile fracture of the materials. Further efforts should be made to understand whether the Eyring’s theory that considers the effect of an applied stress is to reduce the height of a potential energy barrier [31,32] is applicable to the present systems. Both elongation-to-break and area under the tensile stress-strain curve have similar dependence on crosshead speed [Fig. 4(c) and (d)]. With a rise in the speed, a drastic decrease of the two parameters is followed by a
Fig. 8. Area under stress–strain curve of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.
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gradual reduction. SiO2-g-PEA/PP is able to keep its ductility superior to other systems when crosshead speed is slower than 100 mm/min. As compared with the plots’ profiles in Fig. 4(b), it can be deduced that different failure mechanisms take effects when crosshead speed is faster or slower than 100 mm/min. Vu-Khanh and Denault found that the dynamic fracture toughness of glass-flake/PP composites sharply decreases with impact speed [33]. It was speculated that due to the low thermal conductivity of the composite, the relaxation process in the matrix led to an increase in temperature at the crack tip. The temperature rise caused a decrease in the fracture toughness with loading rate. It seems their analysis can also explain the experimental data shown in Fig. 4(c) and (d). To have a more comprehensive understanding of the interdependence of tensile properties of the composites on filler content and crosshead speed, three-dimensional diagrams are drawn in Figs. 5–8. In the case of a crosshead speed of 10 mm/min, Young’s moduli of SiO2-g-PS/ PP and SiO2-g-PEA/PP composites increase with filler content and then decrease (Fig. 5). When crosshead speed is raised, the aforesaid decreasing trend of moduli is gradually replaced by a slight increase or a plateau.
This is different from the performance of SiO2/PP, which exhibits a continuous increase in the stiffness with silica content at each crosshead speed investigated. Evidently, the viscoelastic nature of the interphase due to the appearance of grating polymers in SiO2-g-PS/PP and SiO2-g-PEA/PP composites should be responsible for the distinct behavior. The most obvious characteristics of Fig. 6 is that the strengths measured at 100 and 500 mm/min are much higher than those obtained at a slower crosshead speed. As suggested previously, it should be indicative of a change in failure modes due to the different viscoelastic responses as found in conventional polymers. For SiO2g-PEA/PP composites, the addition of the grafted nanoparticles used to slightly decrease the strength of PP at crosshead speeds of 100 and 500 mm/min [Fig. 6(c)]. It means that the grafting PEA molecules become less efficient to transfer stress under high strain rate. In contrast, 0.5 vol% of SiO2-g-PS can still provide the reinforcing effect at 500 mm/min [Fig. 6(b)]. Figs. 7 and 8 illustrate the elongation-to-break and the area under tensile stress–strain curve of the composites 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 content=1.06 vol%); (e) SiO2-g-PEA/PP (SiO2 content=0.82 vol%). Crosshead speed=50 mm/min.
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different from the grafted SiO2 composites especially in the case of higher particle content. At a silica volume fraction of about 2.7 vol%, for example, the elongation-to-break and the area under tensile stress-strain curve of untreated SiO2/PP are rather small and nearly independent of crosshead speed in comparison with the treated SiO2/PP composites. This should be interpreted as that splitting of the large nanoparticle agglomerates accumulated in
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the matrix due to the fact that the increased particle content governs the failure process. By examining the curves of SiO2-g-PS/PP and SiO2-gPEA/PP determined at different crosshead speeds (Figs. 7 and 8), it can be found that the toughening effect exerted by the modified particles becomes remarkable only at moderate speeds. The improved elongation-to-break and area under tensile stress–strain
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.
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curve represent the improved deformation ability of the composites in relation to plastic stretching of the matrix polymer induced by the grafted nanoparticles. In addition, the entanglement between the grafting polymers and the matrix polymer is also viscoelastic in nature. 3.3. Microscopic observation of fractured surface To have clear images of the failure patterns of the composites under tension, SEM fractographs of the specimens with different filler contents tested at different crosshead speeds are discussed hereinafter. Fig. 9 shows the tensile fractured surfaces of neat PP and the composites with relatively low silica fraction generated at a crosshead speed of 50 mm/min. The unfilled PP has a relatively smooth fractured surface in association with terraced markings [Fig. 9(a)], indicating weak resistance to crack propagation. In the case of SiO2/PP (Vf=0.86 vol%), the fractured surface becomes rougher but the traces of plastic deformation are still less [Fig. 9(b)]. Many silica agglomerates (41 mm in size) are dispersed in the matrix without clear signs of stretching of the surrounding matrix. Occasionally, elongated matrix polymer can be found around silica agglomerates inside a large cavity on the composites’ surface, as highlighted by an arrow in Fig. 9(c). These demonstrate not only the insufficiently interfacial interaction between the particles and the matrix, but also the poor toughening capability of the composites. For SiO2-g-PS/PP and SiO2-g-PEA/ PP composites, the fractured surfaces are full of extensive matrix fibrils [Fig. 9(d) and (e)]. Therefore, it can be evidenced that the grafting polymers on the nanoparticles enhance the interfacial interaction and the dissipation of energy through matrix stretching. When the content of nanosilica approaches around 2.7 vol%, the morphologies of the composites’ fractured surfaces become somewhat different (Fig. 10). A number of cavities appear on the surface of SiO2/PP composites [Fig. 10(a)]. In fact, they are produced due to the debonding of the untreated particles, as illustrated by a magnified view [Fig. 10(b)]. In general, an increased content of untreated SiO2 would lead to larger agglomerates and hence greater probability of debonding due to the poor interfacial adhesion. As there is not enough time for inducing matrix yielding after the extensive particles debonding, the matrix beside the cavities seems to be rather flat [Fig. 10(a)]. This coincides with the reduction of toughness of SiO2/PP at high SiO2 loading [Fig. 3(c) and (d)]. When SiO2-g-PS is incorporated [Fig. 10(c)], concentric matrix-fibrillated circles around nanoparticle agglomerates (as indicated by the upper arrow) and voids left as a result of agglomerated particles detachment (as indicated by the lower arrow) can be found on the fractured surface. As suggested in Ref. [18], the appearance of the fibrillated matrix circles are probably the result of a successive debonding of the modified
nanoparticles from the matrix accompanied by an unconstrained plastic stretching of the interparticulate matrix ligaments [Fig. 10(d)]. Such a deformation process would certainly consume more energy than that dominated only by debonding as shown in Fig. 10(a) and (b). In the case of SiO2-g-PEA/PP, the concentric fibrillar circles around nanoparticle agglomerates emerge next to each other [Fig. 10(e)]. Besides, the matrix surrounding the agglomerated nanoparticles has turn into plastically drawn fibrils [indicated by the arrow in Fig. 10(f)]. Evidently, the inherent flexibility of PEA has made important contribution. In accordance with the model describing the double percolation of yielded zones [20], these should result from the superposition of stress volumes around the agglomerates and the nanoparticles. It explains the cause for that SiO2-g-PEA/PP is still able to maintain higher static ductility at a relatively high nanoparticle concentration [Fig. 3(c) and (d)]. Fig. 10(g) exhibits the SEM observation result of the side surface of the SiO2-g-PEA/PP tensile specimen. Elongated cavities can be seen around partially debonded nanoparticle agglomerates (indicated by the arrows). Again, plastic flow of the bridging matrix is clear visible. On the other hand, Fig. 10(c), (e), (f) and (g) confirm the estimation that the grafted nanoparticle agglomerates have 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 content=1.06 vol%); (d) SiO2-g-PEA/PP (SiO2 content=0.82 vol%). Crosshead speed=10 mm/min.
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By summarizing the data shown in Fig. 4 and the fracture morphology of Figs. 11 and 12, it can be known that the tensile performance of the composites is close to that of neat matrix polymer when the specimens are tested under low or high tensile speeds. Neither the untreated nor the treated particles can take effects under these circumstances. This is particularly true in the case of low silica fraction.
4. Conclusions
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 content=1.06 vol%); (d) SiO2-g-PEA/PP (SiO2 content=0.82 vol%). Crosshead speed=500 mm/min.
are brought into play as an integral. In one sense, the grafted nanoparticle agglomerates can be taken as the folded polymer chains configuration because they are provided with the capability of deforming and releasing locally concentrated stress instead of simply splitting. Fig. 11 gives the fractographs of neat PP and the composites with low SiO2 content tested at a crosshead speed of 10 mm/min. They all have similar appearances characterized by ductile failure except some fine microfibrils 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 observed probably because of the shields of the highly elongated matrix and the low filler content as well. It is worth noting that the fracture modes of the same materials can be changed when a higher testing speed is applied (Fig. 12). The neat PP shows cleavage fracture feature under the crosshead speed of 500 mm/min [Fig. 12(a)]. The striation structure resulting from the joining of different fractured planes on the surface of SiO2/PP composites demonstrates that the particles have little resistance to the crack propagation [Fig. 12(b)]. Similarly, the grafted nanoparticles agglomerates cannot induce effective matrix yielding on a large scale [Fig. 12(c) and (d)]. The mild profiles of the deformation circles [as indicated by the arrows in Fig. 12(c) and (d)] suggest a low plastic deformation level.
Based on the above results and discussions, the following statements can be drawn. (1) The addition of nanoparticles into PP can bring in both reinforcing and toughening effects at filler content as low as 0.5 vol%. Such a simultaneous improvement in modulus, strength and elongation-to-break is hard to be observed in conventional microsized particulate composites. (2) Modification of nanosilica by means of grafting polymerization helps to provide the composites with balanced performance. In addition, different species of the grafting monomers result in different interfacial interactions and different ultimate properties of the composites. (3) With respect to the manufacturing aspect, dispersion homogeneity of the composites filled with untreated nanoparticles is critical, while it is not necessarily realized in the case of grafted nanoparticles. (4) As compared with single screw extruder, twin screw extruder can further decrease the amount of the nanoparticles needed for the composites performance enhancement. (5) The relative increment of the areas under the tensile stress-strain curves of the current composites is similar to the values reported by Ref. [18], although there is a significant difference in ductility between the PP used in the two works. This somewhat contradicts the results of Ren and co-workers, who found that a tougher PP would gain more remarkable improvement of fracture toughness with the addition of nanoparticles [34]. It means that continuous effort should be paid to understand the role of the matrix toughness in nanoparticles composites. (6) Owing to the viscoelastic nature of the grafting polymers, the influence of the modified nanoparticles on the tensile properties of PP is also a function of loading speed.
Acknowledgements The authors are grateful to the support of the Deutsche Forschungsgemeinschaft (DFG FR675/40–1) for the cooperation between the German and Chinese institutes on the topic of nanocomposites. Further thanks
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are due to the National Natural Science Foundation of China (Grant: 50133020), the Key Program of the Ministry of Education of China (Grant: 99198), the Team Project of the Natural Science Foundation of Guangdong, China, the Natural Science Foundation of Guangdong, China (Grant: 990277), and the Key Program of the Science and Technology Department of Guangdong, China (Grant: A10172).
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[15] Sumita M, Tsukumo Y, Miyasaka K, Ishikawa K. Tensile yield stress of polypropylene composites filled with ultrafine particles. J Mater Sci 1983;18:1758–64. [16] Rong MZ, Zhang MQ, Zheng YX, Zeng HM. Chinese patent application no.: CN99116017, 1999. [17] 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. [18] 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. [19] 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. [20] 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. [21] Ma RD. Techniques of radiation processing. Sichuan Science and Technology Press; 1984. p. 32–67 [in Chinese]. [22] Fukano K, Kageyama E. Study on radiation-induced polymerization of vinyl monomers adsorbed on inorganic substances I. Radiation-induced polymerization of styrene adsorbed on several inorganic substances. J Polym Sci Polym Chem 1975;13:1309–24. [23] Kerner EH. Electrical conductivity of composites media. Proc Phys Soc 1956;B69:802–7. [24] Zhang M, Zeng H, Zhang L, Lin G, Li RKY. Fracture characteristics of discontinuous carbon fibre-reinforced PPS and PESC composites. Polym Polym Compos 1993;1:357–65. [25] Ahmed A, Jones FR. A review of particulate reinforcement theories for polymer composites. J Mater Sci 1990;25:4933–42. [26] Nielsen LE. Simple theory of stress-strain properties of filled polymer. J Appl Polym Sci 1966;10:97–103. [27] Nicolais L, Narkis M. Stress-strain behavior of styrene-acrylonitrile/glass bead composites in the glassy region. Polym Eng Sci 1971;11:194–9. [28] Jancar J, Dianselmo A, Dibenedetto AT. The yield strength of particulate reinforced thermoplastic composites. Polym Eng Sci 1992;32:1394–9. [29] Fekete E, Molnar SZ, Kim GM, Michler GH, Pukanszky B. Aggregation, fracture initiation, and strength of PP/CaCO3 composites. J Macromol Sci Phys 1999;B38:885–99. [30] Friedrich K. Microstructural efficiency and fracture toughness of short fiber/thermoplastic matrix composites. Compos Sci Technol 1985;22:43–74. [31] Eyring H. Viscosity, plasticity and diffusion as examples of absolute reaction rates. J Chem Phys 1936;4:283–95. [32] Kinloch AJ, Young RJ. Fracture behaviour of polymers. Applied Science 1983. [33] Vu-Khanh T, Denault J. Toughness of reinforced ductile thermoplastics. J Compos Mater 1992;26:2262–77. [34] Ren X, Bai L, Wang G. Reinforcement and toughening of polypropylene composites by nanoparticle CaCO3. Chem World 2000;41(2):83–7 (in Chinese).
Tensile performance improvement of low nanoparticles filled-polypropylene composites Chun Lei Wua, Ming Qiu Zhangb,*, Min Zhi Rongb, Klaus Friedrichc a
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 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 prepared by conventional compounding technique in which the nanoparticles are pre-grafted by some polymers using irradiation. To examine the applicability of the approach, a tougher polypropylene (PP) was compounded with nano-silica by industrial-scale twin screw extruder and injection molding machine in the present work. The results of tensile tests indicated that the nanoparticles can simultaneously provide PP with stiffening, strengthening and toughening effects at a rather low filler content (typically 0.5% by volume). The presence of grafting polymers on the nanoparticles improves the tailorability of the composites. Due to the viscoelastic nature of the matrix and the grafting polymers, the tensile performance of the composites filled with untreated and treated nanoparticles is highly dependent on loading rate. With increasing the crosshead speed for the tensile tests, the dominant failure mode 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 fillers are added to polymers in commercial production primarily for the reasons of cost reduction and stiffness improvement [1,2]. Although most studies dealing with the modification of semi-crystalline polymers with inorganic particulates reported embrittling effects by comparing ultimate elongation and impact strength of composite materials with those of unfilled resins [3–5], some researchers showed the enhancement of toughness in rigid particles filled polypropylene [6,7] and polyethylene [8,9]. It is worth noting that in the case of micrometer-sized particulates, high filler content (typically higher than 20% by volume) is generally required to bring the above-stated positive effects of the fillers into play. This would detrimentally affect some important properties of the matrix polymers such as processability, appearance, * Corresponding author. Tel.: +86-20-8403-6576; fax: +86-208403-6564. E-mail address: [email protected] (M.Q. Zhang).
density and ageing performance. Therefore, a composite with improved performance and low particle concentration is highly desired. With regard to this, the newly developed nanocomposites would be competitive candidates. The extremely high surface area is one of the most attractive characteristics of nanoparticles because it facilitates creating a great amount of interphase in a composite. Introduction of nanoparticles into a polymer changes the intermolecular interaction of the matrix [10]. As estimated by Reynaud et al. [11], an interphase 1 nm thick represents roughly 0.3% of the total volume of polymer in the case of microparticle filled composites, whereas it can reach 30% of the total volume in the case of nanocomposites. That is, the non-negligible contribution made by the interphase provides diverse possibilities of performance tailoring, and is able to influence the properties of matrices to a much greater extent under a rather low nano-filler loading. The crux of the matter lies in that how to well distribute nanoparticles over a polymer matrix and how to improve nanoparticles/matrix interaction.
0266-3538/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved. PII: S0266-3538(02)00079-9
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From a practical point of view, dispersive mixing in preparing polymer based particulate composites has important technical meaning. However, a homogenous dispersion of nanoparticles in a polymer is very difficult by using the existing compounding techniques due to the strong tendency of the fine particles to agglomerate and the high melt viscosity of the matrix. In many cases, the so-called nanoparticles filled polymers contain a number of loosened clusters of nanoparticles, as demonstrated by the work of Jana and Jain [12] dealing with untreated nanosilica/polyethersulphone composites. When the composites are subjected to force, the nanoparticle agglomerates can be split easily and a premature failure of the materials would thus take place [12–15]. To overcome this dilemma and to give full play to nanoparticles, we used irradiation grafting polymerization method to modify nanoparticles first, and then the treated particles were mechanically mixed with a polymer as usual [16]. Owing to the low molecular weight nature, the grafting monomers can penetrate into the agglomerated nanoparticles easily and react with the activated sites of the nanoparticles inside as well as outside the agglomerates [17]. As a result, the following effects can be obtained [18]: (i) Hydrophobicity of the nanoparticles is increased, facilitating the filler/matrix miscibility; (ii) filler/matrix interaction is enhanced through the entanglement between the grafting polymer and the polymer matrix; (iii) the nanoparticle agglomerates become stronger because they turn into a nanocomposite microstructure comprising the nanoparticles and the grafted, homopolymerized polymer; (iv) the interfacial characteristics between the treated nanoparticles and the matrix polymer can be tailored by changing the species of the grafting monomers and the grafting conditions. In this context, a uniform dispersion of nanoparticles in the matrix might no longer be critical. Mechanical testing of polypropylene filled with nanoSiO2 [18] and nano-CaCO3 [19] demonstrated the feasibility of the above approach. Only a small amount of modified nanoparticles (typically less than 3% by volume) can effectively improve modulus, strength, toughness and thermal deformation temperature of the matrix polymer. Such an improvement in overall properties of polymers can scarcely be observed in conventional microparticulate composites. It was found that the deformation habit but not the crystallization characteristics of the matrix polymer remarkably change with the addition of the treated nanoparticles. To explain the specific influence generated by the nanoparticles at low-filler loading regime, a double percolation of stress volumes, characterized by the appearance of connected shear yielded networks throughout the composite, was proposed [20]. Considering that the polypropylene used in our previous works [18–20] is a brittle type and the composites
were prepared with a lab-scale single screw extruder and compression molding, the results might not have adequate applicability. Therefore, a commercial polypropylene with higher toughness was compounded with nano-silica by means of industrial-scale twin screw extruder and injection molding machine in the current work. Tensile performance and fractured surfaces are analyzed as a function of particulate treatment, filler content and crosshead speed to reveal the structure– property relationships of the composites, and the mechanical 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 flow index of 3.2 g/10 min (2.16 kg at 230 C). Fumed silica with an average primary particle size of 15 nm and a specific surface area of 374 m2/g was produced by Shenyang Chemical Engineering Ltd., China. Commercial monomers, styrene and ethyl acrylate, were used as grafting monomers without further purification. 2.2. Irradiation grafting of nano-SiO2 Modification of nano-silica proceeded according to the following steps. The nanoparticles were pretreated at 140 C under vacuum for 6 h to eliminate possible absorbed water on the surface of the particles. Then a mixture of nanoparticles/monomer (100/20 by weight) and a certain amount of n-hexane was irradiated by 60 Co g-ray under atmosphere at room temperature. After exposure to a dose of 4 Mrad, the solvent was recovered, and the dried residual powder was available for the subsequent compounding. 2.3. Characterization of the irradiation grafted products 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) and a Bruker Equinox 55 Fourier transform infrared spectroscope (FTIR), the weight of the grating polymer and the chemical structure of the modified nanoparticles were characterized, respectively. To further
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separate the grafting polymer from the treated nanoparticles, nano-silica accompanied with the unextractable grafting polymer was immersed in 20% HF solution for 72 h to remove the inorganic particles. The molecular weights of the grafting and the homopolymerized polymers were determined by a Waters 991 gel permeation chromatography (GPC), with tetrahydrofuran as the solvent. To observe the morphologies of the nanoparticles, untreated SiO2 and grafted SiO2 (without homopolymer) were added into ethanol and toluene to prepare 0.001 g/ml solutions, respectively. With the aid of sonication for 30 min, the solutions were transferred to copper gauzes by droppers. After the evaporation of the solvents, a JEM-100CXZ transmission electron microscopy (TEM) was used to examine the appearance of the particles. 2.4. Composites preparation and characterization 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) 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 at crosshead speeds of 10, 30, 50, 100 and 500 mm/min, respectively. Five samples were tested for each case. The fractured surfaces of the samples were observed with a Hitachi S-520 scanning electron microscope (SEM) at an accelerating voltage of 20 kV.
3. Results and discussion 3.1. Effect of irradiation grafting polymerization on nano-SiO2 Since the present work aims to study the effect of modified nano-silica on the mechanical behavior of PP composites, variation in the chemical structure of the particles should be known at the very beginning of the discussion. FTIR spectra of untreated and treated nanosilica are shown in Fig. 1. To eliminate the influence of homopolymers, both polystyrene-grafted nano-SiO2 (SiO2-g-PS) and polyethyl acrylate-grafted nano-SiO2 (SiO2-g-PEA) used for the FTIR examinations were separated from the homopolymers in advance. In comparison with the spectrum of SiO2 as-received, the adsorptions at 690, 1460 and 2960 cm 1 appearing in
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Fig. 1. FTIR spectra of untreated SiO2 and grafted SiO2.
the spectrum of SiO2-g-PS represent the bending mode of C–H in benzene rings and the stretching modes of C–C and C–H, respectively. In addition, the band at 1725 cm 1 in the spectrum of SiO2-g-PEA indicates the existence of carbonyl groups. These prove that polystyrene (PS) and polyethyl acrylate (PEA) have been chemically connected to the surface of nano-silica during the irradiation polymerization processes as expected. On the basis of above qualitative analysis, the quantitative results of the grafting polymerization on nanosilica are given in Table 1. It is seen that when other conditions being equal, the percentage grafting and the grafting efficiency of styrene onto the nanoparticles are much higher than those of ethyl acrylate, while the monomer conversion of styrene is lower than that of ethyl acrylate. This reflects the difference in reactive feature between the monomers. In general, room temperature irradiation grating polymerization onto inorganic particles is controlled by the mechanism of free radical polymerization. Under the same irradiation dose, acrylic monomers would generate much more radicals than styrene, and moreover, ethyl acrylate radicals have higher activity than styrene radicals [21]. Therefore, the conversion of ethyl acrylate is superior to that of styrene, leading to higher homopolymer fraction of PEA than PS. On the other hand, partial surface of nanosilica might be connected with ethyl acrylate in the solvent through hydrogen bonding prior to the irradiation processing. When the composite system is exposed to the irradiation, radicals can be formed on both the nanoparticles and the ethyl acrylate molecules connected to the particles. The latter would result in either grafting polymerization or homopolymerization. That is, the amount of grafted PEA has to be lower as compared with the styrene/silica mixture where no chemical connection between the monomers and the particles is established before the irradiation. It should be responsible for the difference in percentage grafting between ethyl acrylate/silica and styrene/silica.
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Table 1 Results of irradiation grafting polymerization onto nano-silica Monomer
Monomer conversiona (%)
Percent graftingb (%)
Grafting efficiencyc (%)
MWd of grafting polymer (104)
MW of homopolymer (104)
Styrene Ethyl acrylate
52.8 93.5
4.64 1.56
43.9 10.5
1.3 (de=1.70) 7.1 (d=2.13)
1.1 (d=1.46) 1.1 (d=1.33)
a b c d e
Monomer conversion=weight of polymer/weight of monomer. Percent grafting=weight of grafting polymer/weight of nano-SiO2. Grafting efficiency=weight of grafting polymer/weights of grafting polymer and homopolymer. MW=weight average molecular weight. d=Molecular weight polydispersity index.
By further examining the data listed in Table 1, it can be found that the grafting polymers attached to the nanoparticles possess higher molecular weights and broader molecular weight distributions than the homopolymers. Similar phenomena were also reported by Fukano and Kageyama [22] when they studied radiation grafting of styrene onto silica-gel. These can be attributed to the characteristics of the grafting reaction. That is, the reaction was a typical heterogeneous one because the activated sites were created by irradiation on the nanoparticles and the chain growth of the grafting polymers had to proceed in solid–liquid state. The mobility of the growing chains was thus worse as compared with a homogeneous reaction in liquid where both ends of the macromolecular chains can move freely. As a result, the probability of chain termination between the radicals became relatively difficult, leading to higher molecular weight of the grafting polymers. In addition, because the radicals are not simultaneously produced on the nanoparticles, the grafting polymers formed at earlier stage would wrap the surface of the nanoparticles and reduce the probability of collision between the monomers and the radicals formed on the particles at latter stage. This accounts for the higher molecular weight polydispersity indexes of the grating polymers. Morphologies of the nanoparticles before and after grafting polymerization are illustrated in Fig. 2. A chain-like branched structure of the agglomerated SiO2 particles [17] can be observed in the solutions. The smallest perceivable units are approximately 15 nm in diameter in the case of untreated nanoparticles [Fig. 2(a)]. When grafting polymers are introduced onto the particles, the sizes of the agglomerates become larger and the edges are no longer clearly discernible [Fig. 2(b, c)]. Such a change demonstrates the role of the grafting polymer, i.e. separating and connecting the nanoparticles. To estimate the thickness of the polymer layer adhered to the particles, the data of SiO2-g-PS is used as an example, i.e. percent grafting=4.64%, density
Fig. 2. TEM microphotos (magnification=105) of (a) SiO2 asreceived, (b) SiO2-g-PS, and (c) SiO2-g-PEA in solvents. In the latter two specimens, homopolymers were removed in advance.
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of PS=1.05 g/cm3, specific surface area of the silica=374 m2/g. Supposing a complete coverage on each silica particles, the thickness of the grafting PS should equal to 0.0464/(374 m2/g1.05 g/cm3)0.12 nm. Evidently, this is a reasonable value as evidenced by the TEM photos in Fig. 2. Due to the low percent grafting and the thin grafting polymer layer, there are still many unreacted hydroxyl groups on the surface of nano-SiO2, which is responsible for the appearance of the larger agglomerates of the grafted particles in the solvent. 3.2. Tensile properties of the composites Fig. 3 shows the results of tensile testing of PP reinforced by nano-silica as a function of filler content, determined at a moderate crosshead speed of 50 mm/ min. Although both the treated and the untreated nanoparticles can impart the high stiffness of the fillers to the matrix polymer as expected, the composites incorporated with the modified particles exhibit lower modulus over the whole range of filler loading of interests [Fig. 3(a)]. Usually the capability of composite interface to transfer elastic deformation depends to a great extent upon the interfacial stiffness and static adhesion strength [23,24]. A high interfacial stiffness corresponds to a high composite modulus. Since the grafting polymers and the homopolymers introduced onto the nanoparticles form a relatively compliant interlayer at the particles/matrix interface, the high
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stiffness of the particles has to be masked under the low stress level [9] and the composites have to show lower modulus as compared with the case of untreated SiO2 composites. With a rise in filler content, the increased amount of grafting polymers further increases the modulus mismatching of the filler and the matrix, and reduces the stiffening effect of SiO2, leading to the drop in Young’s moduli of the composites at high filler regime. In addition, due to the higher rigidity of PS molecules than PEA, the interfacial elastic stress is less efficiently transferred in SiO2-g-PEA/PP composites than in SiO2g-PS/PP, especially when the fraction of the modified nanoparticles is high. In contrast, an approximately linear composition dependence of tensile modulus is perceived in the composites filled with SiO2 as-received [Fig. 3(a)]. This can be interpreted as the absence of a soft interphase and the appearance of larger agglomerates of the nanoparticles in the matrix as a result of poor filler/matrix miscibility. Under the same filler volume fraction [25], the latter effect would provide higher load carrying ability within small strain range. The above results factually reflect the contradiction between the stiffening effect of the rigid particles and the weakening effect of the soft interlayer. Fig. 3(b) gives the tensile strengths of the composites versus SiO2 content. Clearly, both untreated and treated nanoparticles exhibit the strengthening ability. It is well known that the tensile strength of a particulate composite
Fig. 3. Tensile properties of PP composites as a function of nano-SiO2 volume fraction: (a) Young’s modulus, (b) tensile strength, (c) elongation to break, and (d) area under stress–strain curve. Crosshead speed=50 mm/min.
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Fig. 4. Tensile properties of PP composites as a function of crosshead speed: (a) Young’s modulus, (b) tensile strength, (c) elongation to break, and (d) area under stress–strain curve.
Fig. 5. Young’s modulus of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.
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is usually reduced with filler content following a power law in the case of a poor filler/matrix bonding [26,27]. That means, the strength of the composite cannot be greater than that of the unfilled version because the filler particles do not bear any fraction of the external load. This contradicts the results shown in Fig. 3(b). In fact, when bonding between fillers and matrix is strong enough, the tensile yield strength of a particulate composite can be higher than that of the matrix polymer [28,29]. Although these models were developed based on the cases of microsized particulate composites, they are still valid for the explanation of the composites filled with nanoparticles [20]. Therefore, the extremely effective improvement of tensile strength of the composites with grafted nano-SiO2 should result from chain interdiffusion and entanglement between the macromolecules of the grafting polymers and the matrix. It is worth noting that when the amount of the grafted nanoparticles is increased (e.g. > 0.5 vol%), the content of compliant PEA chains is raised accordingly in the interlayer and the interfacial stress transfer efficiency has to be decreased. This accounts for the relatively low strength of SiO2-g-PEA/PP composites at higher particulate loading.
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By comparing Fig. 3(b) with the data of Ref. [18], the most distinct difference lies in the results of untreated SiO2/PP. That is, for the composites prepared by a labscale single screw extruder and compression molding, the addition of untreated SiO2 lowers the tensile strength of PP in the lower loading region but then leads to a slight increase in strength when the particle fraction reaches 4.68 vol% [18]. With respect to the SiO2/PP composites of the present work, which were manufactured through twin screw extrusion and injection molding, a continuous increase in the strength with SiO2 content is detected [Fig. 3(b)]. So far as we know, corotating twin screw extruders are able to provide more sufficient homogenization in comparison with single screw extruders. The above difference manifests that the important role of even distribution of untreated nanoparticles in the composites. For SiO2 as-received, the more particles are exposed to the matrix polymer, the more possibly the interaction between the particles and the matrix can be enhanced. So, an improved homogeneity of the untreated SiO2/PP composites would certainly be beneficial to the stress transfer. This again shows the significance of grafting modification 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.
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of composites strength performance to the dispersion state of the particles. As the grafted nanoparticle agglomerates turn into a nonocomposite microstructure consisting of the particles and the grafted, homopolymerized secondary polymer [18], they are brought into play as an integral when the composites are subjected to the applied force [20]. The situation is completely different from the composites with untreated nanoparticles, in which agglomerated particles have to be deagglomerated as much as possible to reduce the probability of premature failure. As a result, a uniform dispersion of the nanoparticles in the matrix is absolutely necessary for obtaining the reinforcing effect in case untreated nanoparticles are used. Failure strain can partially assess the rupture behavior of a composite material. The incorporation of particulate fillers usually results in a decrease in this parameter regardless of the interfacial adhesion [26]. It is true even in the system exhibiting impressive impact toughness improvement with the addition of mineral fillers [8]. However, the plots shown in Fig. 3(c) demonstrate that the values of elongation-to-break of
PP can be significantly increased by using nano-SiO2,, implying a failure mechanism different from those involved in conventional composites. Comparatively, SiO2-g-PEA is able to provide a stable improvement over the entire filler content range of interests. The reduction in elongation-to-break of the composite filled with untreated SiO2 suggests that the fillers cause a reduction in matrix deformation due to an introduction of mechanical restrains. In contrast, the improvement of elongation-to-break with the incorporation of the grafted nanoparticles is a result of interfacial viscoelastic deformation and matrix yielding. Evidently, the grafted PEA makes more effective contribution. The area under the tensile stress–strain curve can more reasonably characterize the toughness potential of the composites than elongation-to-break under static tensile loading conditions [30]. As confirmed by Fig. 3(d), the grafted nanoparticles indeed improve the ductility of PP at a silica content as low as 0.5 vol%. The results are somewhat opposite to those observed in Ref. [18], which reports a deteriorated effect of grafting PEA on 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.
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that SiO2-g-PEA employed in the present work possesses almost the same percent grafting and grafting efficiency as that in Ref. [18], it can be concluded that the molecular entanglement between the grafting PEA and matrix PP was not sufficiently formed during melt compounding when the single screw extruder was used. As a result, localized plastic deformation or matrix drawing cannot be efficiently induced by SiO2-g-PEA as in 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 at higher filler content regime are lower than that of the neat PP. It implies that the short range interaction at SiO2/PP interface is not good at inducing plastic deformation of the matrix polymer. Due to the encounter of the propagating neck with a larger agglomerate, particularly in the case of higher filler content, final failure of the composites might be initiated easily [8]. Obviously, the resultant embrittling effect can be prevented by the application of grafted nanoparticles. Since polymer composites maintain the viscoelasticity of polymers, the dependence of tensile properties on crosshead speed should be known for engineering purposes. As
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can be seen from Fig. 4(a), a linear increase in Young’s modulus with increasing crosshead speed is valid for all the specimens. The presence of silica leads to the values of modulus being higher than that for the unfilled PP and being slightly less speed dependent. These observations 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 strengths increase in principle, there is a significant transition in the slope. Accordingly, the first derivation of the strength in Fig. 4(b) with respect to crosshead speed would yield a peak between 50 and 100 mm/min, which clearly corresponds to an energy-activated process in tensile fracture of the materials. Further efforts should be made to understand whether the Eyring’s theory that considers the effect of an applied stress is to reduce the height of a potential energy barrier [31,32] is applicable to the present systems. Both elongation-to-break and area under the tensile stress-strain curve have similar dependence on crosshead speed [Fig. 4(c) and (d)]. With a rise in the speed, a drastic decrease of the two parameters is followed by a
Fig. 8. Area under stress–strain curve of PP composites as a function of nano-SiO2 volume fraction and crosshead speed.
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gradual reduction. SiO2-g-PEA/PP is able to keep its ductility superior to other systems when crosshead speed is slower than 100 mm/min. As compared with the plots’ profiles in Fig. 4(b), it can be deduced that different failure mechanisms take effects when crosshead speed is faster or slower than 100 mm/min. Vu-Khanh and Denault found that the dynamic fracture toughness of glass-flake/PP composites sharply decreases with impact speed [33]. It was speculated that due to the low thermal conductivity of the composite, the relaxation process in the matrix led to an increase in temperature at the crack tip. The temperature rise caused a decrease in the fracture toughness with loading rate. It seems their analysis can also explain the experimental data shown in Fig. 4(c) and (d). To have a more comprehensive understanding of the interdependence of tensile properties of the composites on filler content and crosshead speed, three-dimensional diagrams are drawn in Figs. 5–8. In the case of a crosshead speed of 10 mm/min, Young’s moduli of SiO2-g-PS/ PP and SiO2-g-PEA/PP composites increase with filler content and then decrease (Fig. 5). When crosshead speed is raised, the aforesaid decreasing trend of moduli is gradually replaced by a slight increase or a plateau.
This is different from the performance of SiO2/PP, which exhibits a continuous increase in the stiffness with silica content at each crosshead speed investigated. Evidently, the viscoelastic nature of the interphase due to the appearance of grating polymers in SiO2-g-PS/PP and SiO2-g-PEA/PP composites should be responsible for the distinct behavior. The most obvious characteristics of Fig. 6 is that the strengths measured at 100 and 500 mm/min are much higher than those obtained at a slower crosshead speed. As suggested previously, it should be indicative of a change in failure modes due to the different viscoelastic responses as found in conventional polymers. For SiO2g-PEA/PP composites, the addition of the grafted nanoparticles used to slightly decrease the strength of PP at crosshead speeds of 100 and 500 mm/min [Fig. 6(c)]. It means that the grafting PEA molecules become less efficient to transfer stress under high strain rate. In contrast, 0.5 vol% of SiO2-g-PS can still provide the reinforcing effect at 500 mm/min [Fig. 6(b)]. Figs. 7 and 8 illustrate the elongation-to-break and the area under tensile stress–strain curve of the composites 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 content=1.06 vol%); (e) SiO2-g-PEA/PP (SiO2 content=0.82 vol%). Crosshead speed=50 mm/min.
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different from the grafted SiO2 composites especially in the case of higher particle content. At a silica volume fraction of about 2.7 vol%, for example, the elongation-to-break and the area under tensile stress-strain curve of untreated SiO2/PP are rather small and nearly independent of crosshead speed in comparison with the treated SiO2/PP composites. This should be interpreted as that splitting of the large nanoparticle agglomerates accumulated in
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the matrix due to the fact that the increased particle content governs the failure process. By examining the curves of SiO2-g-PS/PP and SiO2-gPEA/PP determined at different crosshead speeds (Figs. 7 and 8), it can be found that the toughening effect exerted by the modified particles becomes remarkable only at moderate speeds. The improved elongation-to-break and area under tensile stress–strain
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.
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curve represent the improved deformation ability of the composites in relation to plastic stretching of the matrix polymer induced by the grafted nanoparticles. In addition, the entanglement between the grafting polymers and the matrix polymer is also viscoelastic in nature. 3.3. Microscopic observation of fractured surface To have clear images of the failure patterns of the composites under tension, SEM fractographs of the specimens with different filler contents tested at different crosshead speeds are discussed hereinafter. Fig. 9 shows the tensile fractured surfaces of neat PP and the composites with relatively low silica fraction generated at a crosshead speed of 50 mm/min. The unfilled PP has a relatively smooth fractured surface in association with terraced markings [Fig. 9(a)], indicating weak resistance to crack propagation. In the case of SiO2/PP (Vf=0.86 vol%), the fractured surface becomes rougher but the traces of plastic deformation are still less [Fig. 9(b)]. Many silica agglomerates (41 mm in size) are dispersed in the matrix without clear signs of stretching of the surrounding matrix. Occasionally, elongated matrix polymer can be found around silica agglomerates inside a large cavity on the composites’ surface, as highlighted by an arrow in Fig. 9(c). These demonstrate not only the insufficiently interfacial interaction between the particles and the matrix, but also the poor toughening capability of the composites. For SiO2-g-PS/PP and SiO2-g-PEA/ PP composites, the fractured surfaces are full of extensive matrix fibrils [Fig. 9(d) and (e)]. Therefore, it can be evidenced that the grafting polymers on the nanoparticles enhance the interfacial interaction and the dissipation of energy through matrix stretching. When the content of nanosilica approaches around 2.7 vol%, the morphologies of the composites’ fractured surfaces become somewhat different (Fig. 10). A number of cavities appear on the surface of SiO2/PP composites [Fig. 10(a)]. In fact, they are produced due to the debonding of the untreated particles, as illustrated by a magnified view [Fig. 10(b)]. In general, an increased content of untreated SiO2 would lead to larger agglomerates and hence greater probability of debonding due to the poor interfacial adhesion. As there is not enough time for inducing matrix yielding after the extensive particles debonding, the matrix beside the cavities seems to be rather flat [Fig. 10(a)]. This coincides with the reduction of toughness of SiO2/PP at high SiO2 loading [Fig. 3(c) and (d)]. When SiO2-g-PS is incorporated [Fig. 10(c)], concentric matrix-fibrillated circles around nanoparticle agglomerates (as indicated by the upper arrow) and voids left as a result of agglomerated particles detachment (as indicated by the lower arrow) can be found on the fractured surface. As suggested in Ref. [18], the appearance of the fibrillated matrix circles are probably the result of a successive debonding of the modified
nanoparticles from the matrix accompanied by an unconstrained plastic stretching of the interparticulate matrix ligaments [Fig. 10(d)]. Such a deformation process would certainly consume more energy than that dominated only by debonding as shown in Fig. 10(a) and (b). In the case of SiO2-g-PEA/PP, the concentric fibrillar circles around nanoparticle agglomerates emerge next to each other [Fig. 10(e)]. Besides, the matrix surrounding the agglomerated nanoparticles has turn into plastically drawn fibrils [indicated by the arrow in Fig. 10(f)]. Evidently, the inherent flexibility of PEA has made important contribution. In accordance with the model describing the double percolation of yielded zones [20], these should result from the superposition of stress volumes around the agglomerates and the nanoparticles. It explains the cause for that SiO2-g-PEA/PP is still able to maintain higher static ductility at a relatively high nanoparticle concentration [Fig. 3(c) and (d)]. Fig. 10(g) exhibits the SEM observation result of the side surface of the SiO2-g-PEA/PP tensile specimen. Elongated cavities can be seen around partially debonded nanoparticle agglomerates (indicated by the arrows). Again, plastic flow of the bridging matrix is clear visible. On the other hand, Fig. 10(c), (e), (f) and (g) confirm the estimation that the grafted nanoparticle agglomerates have 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 content=1.06 vol%); (d) SiO2-g-PEA/PP (SiO2 content=0.82 vol%). Crosshead speed=10 mm/min.
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By summarizing the data shown in Fig. 4 and the fracture morphology of Figs. 11 and 12, it can be known that the tensile performance of the composites is close to that of neat matrix polymer when the specimens are tested under low or high tensile speeds. Neither the untreated nor the treated particles can take effects under these circumstances. This is particularly true in the case of low silica fraction.
4. Conclusions
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 content=1.06 vol%); (d) SiO2-g-PEA/PP (SiO2 content=0.82 vol%). Crosshead speed=500 mm/min.
are brought into play as an integral. In one sense, the grafted nanoparticle agglomerates can be taken as the folded polymer chains configuration because they are provided with the capability of deforming and releasing locally concentrated stress instead of simply splitting. Fig. 11 gives the fractographs of neat PP and the composites with low SiO2 content tested at a crosshead speed of 10 mm/min. They all have similar appearances characterized by ductile failure except some fine microfibrils 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 observed probably because of the shields of the highly elongated matrix and the low filler content as well. It is worth noting that the fracture modes of the same materials can be changed when a higher testing speed is applied (Fig. 12). The neat PP shows cleavage fracture feature under the crosshead speed of 500 mm/min [Fig. 12(a)]. The striation structure resulting from the joining of different fractured planes on the surface of SiO2/PP composites demonstrates that the particles have little resistance to the crack propagation [Fig. 12(b)]. Similarly, the grafted nanoparticles agglomerates cannot induce effective matrix yielding on a large scale [Fig. 12(c) and (d)]. The mild profiles of the deformation circles [as indicated by the arrows in Fig. 12(c) and (d)] suggest a low plastic deformation level.
Based on the above results and discussions, the following statements can be drawn. (1) The addition of nanoparticles into PP can bring in both reinforcing and toughening effects at filler content as low as 0.5 vol%. Such a simultaneous improvement in modulus, strength and elongation-to-break is hard to be observed in conventional microsized particulate composites. (2) Modification of nanosilica by means of grafting polymerization helps to provide the composites with balanced performance. In addition, different species of the grafting monomers result in different interfacial interactions and different ultimate properties of the composites. (3) With respect to the manufacturing aspect, dispersion homogeneity of the composites filled with untreated nanoparticles is critical, while it is not necessarily realized in the case of grafted nanoparticles. (4) As compared with single screw extruder, twin screw extruder can further decrease the amount of the nanoparticles needed for the composites performance enhancement. (5) The relative increment of the areas under the tensile stress-strain curves of the current composites is similar to the values reported by Ref. [18], although there is a significant difference in ductility between the PP used in the two works. This somewhat contradicts the results of Ren and co-workers, who found that a tougher PP would gain more remarkable improvement of fracture toughness with the addition of nanoparticles [34]. It means that continuous effort should be paid to understand the role of the matrix toughness in nanoparticles composites. (6) Owing to the viscoelastic nature of the grafting polymers, the influence of the modified nanoparticles on the tensile properties of PP is also a function of loading speed.
Acknowledgements The authors are grateful to the support of the Deutsche Forschungsgemeinschaft (DFG FR675/40–1) for the cooperation between the German and Chinese institutes on the topic of nanocomposites. Further thanks
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are due to the National Natural Science Foundation of China (Grant: 50133020), the Key Program of the Ministry of Education of China (Grant: 99198), the Team Project of the Natural Science Foundation of Guangdong, China, the Natural Science Foundation of Guangdong, China (Grant: 990277), and the Key Program of the Science and Technology Department of Guangdong, China (Grant: A10172).
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