Highly ductile polypropylene-based nanocomposites by dispersing monodisperse silica nanospheres in functionalized polypropylene containing hydroxyl groups

Polymer 99 (2016) 63e71

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Highly ductile polypropylene-based nanocomposites by dispersing monodisperse silica nanospheres in functionalized polypropylene containing hydroxyl groups Ryota Watanabe, Masao Kunioka, Junji Mizukado, Hiroyuki Suda, Hideaki Hagihara* Research Institute for Sustainable Chemistry, National Institute of Advanced Industrial Science and Technology (AIST), 1-1-1 Higashi, Tsukuba 305-8565, Japan

a r t i c l e i n f o

a b s t r a c t

Article history: Received 19 April 2016 Received in revised form 14 June 2016 Accepted 21 June 2016 Available online 21 June 2016

We fabricated nanocomposites by dispersing uniform-sized silica nanospheres (SNSs, 49e282 nm, 0 e30 wt%) into polypropylene (PP) and two types of functionalized PP to clarify the effects of matrix-filler interfacial adhesion forces on the mechanical properties in PP-based nanocomposites. Since there is good affinity between silica surfaces and hydroxyl group in PP, poly(5-hexen-1-ol-co-propylene) with 1.3 and 6.4 mol% of hydroxyl groups were employed as functionalized PP (PPOH). Tensile tests revealed that some mechanical properties of the composites with PPOH were higher than those with PP. In particular, composites with PPOH maintained high ductility, whereas the ductility of composites with PP were significantly reduced by dispersing SNSs. Fourier-transform infrared (FTIR) and X-ray photoelectron spectroscopy (XPS) studies indicated the formation of hydrogen bonds between PPOH and the silica surface, which are attributed to the high ductility. © 2016 Elsevier Ltd. All rights reserved.

Keywords: Functionalized polypropylene Silica nanosphere Ductility of nanocomposites

1. Introduction Polypropylene (PP) based nanocomposites exhibit markedly improved mechanical strength [1e9], flame retardancy [10,11], gas barrier properties [12] and thermal conductivity [13] over original PP. In particular, extensive work has been carried out in reinforcement of PP by dispersing nanofillers into the material. In most cases, the addition of nanofillers to PP leads to significantly diminishing ductility [14e16]. Therefore, further advances are required to realize improved ductility over existing PP-based nanocomposites. This lower ductility can be attributed in part to void development at interfaces between the PP matrix and the filler, due to their low affinity for each other [17,18]. Hence, preventing detachment at the matrix-filler interface by enhancement of the interfacial adhesion is a promising approach to overcome the poor ductility in PP-based nanocomposites. Phua et al. demonstrated a highly ductile polyurethane/dopamine-modified clay nanocomposite by formation of high matrix-filler interfacial adhesion via hydrogen bonding interactions, resulting in improved ductility [19]. In the

* Corresponding author. E-mail address: [email protected] (H. Hagihara). http://dx.doi.org/10.1016/j.polymer.2016.06.054 0032-3861/© 2016 Elsevier Ltd. All rights reserved.

case of PP-based nanocomposites, on the other hand, interactions between the matrix and filler are difficult to take place, due to the inertness of PP (which has no functional groups). To improve the affinity of PP for polar materials, maleic anhydride-grafted PP (MAPP) is extensively used as an additive in PP-based composites. Several PP-based nanocomposites in which the fillers were highly dispersed have been reported, such as PP/ octadecylamine-functionalized clay [4], PP/silica [20] and PP/ titania [21]. In general, the modulus and mechanical strength of a nanocomposite was also improved by the addition of MAPP, whereas the ductility of the PP-based nanocomposite was drastically decreased. The polar group content of commercially available MAPP is usually under 1 mol% per monomer units, and the average molecular weight (Mn) of that is under 10,000; this indicates that there is room for improvement. Refinement of the additive reagent to further enhance the affinity between matrix and fillers, namely interfacial adhesion force, is a possible path to improved ductility. We previously developed functionalized PP with hydroxyl groups by copolymerization of propylene and a-olefin containing hydroxyl groups, denoted as PPOH [22e25]. The amount of hydroxyl groups in PPOH can be arbitrarily controlled without reduction in molecular weight by varying the propylene/comonomer ratio during polymerization. The hydroxyl groups in PPOH

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potentially interact with surface hydroxyl groups of inorganic oxide fillers (e.g., silica, alumina and titania) via hydrogen bonding. Therefore, PPOH can be regarded as a promising material for not only improving the mechanical strength of PP-based nanocomposites, but also maintaining the ductility of PP-based nanocomposites. Silica nanoparticles are one of the most widely-used fillers for polymer nanocomposites [15,20,26]. Particle size is one factor influencing the mechanical properties of polymer composites with silica nanoparticles. However, particle size effects have not been widely studied because conventionally used silica nanoparticles have broad particle size distributions and non-uniform shapes. Recently, uniform silica nanospheres (SNSs) with a controllable size ranging from 10 to 550 nm were successfully synthesized [27e30]. SNSs are model nanofillers for PP-based nanocomposites for the purpose of determining the effects of particle size on the matrixfiller interfacial adhesion and the mechanical properties of the nanocomposites, because of their high monodispersibility, controllable particle size on the nanoscale, and absence of anisotropy, unlike commercial silica nanoparticles. In this article, we developed new nanocomposites by dispersing uniform-sized SNSs with sizes ranging from 49 to 282 nm into PP and PPOH containing 1.3 and 6.4 mol% hydroxyl groups. In order to clarify the effects of the matrix-filler interfacial adhesion forces on the overall mechanical properties, tensile and impact tensile tests were performed (Scheme 1). The interfacial adhesion force between the matrix and SNSs was evaluated using theoretical methods, such as Einstein’s equation [5,20]. Fourier-transform infrared spectroscopy (FTIR) and X-ray photoelectron spectroscopy (XPS) were used for characterization of chemical structures at the matrix-filler interface. The objective of this work is to better understand the role of matrix-filler interfacial adhesion on the mechanical properties of nanocomposites.

Mw ¼ 146 000, Mw/Mn ¼ 1.95, mm ¼ 98.4%) were used, and their chemical structures are shown in Scheme 2. Densities of PP, PPOH1.3 and PPOH6.4 are 0.9 g/cm3. 2.2. Synthesis of SNSs Silica nanospheres (SNSs) were synthesized by a modified regrowth method [30]. First, 0.174 g of L-arginine (Arg, SigmaAldrich, >98%) and 174 g of ionized water were added to a 250 mL PP bottle and stirred at room temperature. Tetraethyl orthosilicate (TEOS, 10.41 g, Tokyo Chemical Industry, >96%) was rapidly added to the Arg-H2O solution with stirring at 70  C for 10 h, yielding monodisperse SNS 14 nm in size to be used as a seed solution. Further expansion of the SNS size was attained by regrowth of the 14 nm SNS. An appropriate amount of the seed solution was added to a mixture of ionized water and ethanol (EtOH, Wako, >99.5) containing Arg with compositions shown in Table 1. TEOS was then added to the solution, and the resulting mixture was placed in an oven at 70  C for 10 h and a white powder was collected by evaporation of the solvent at 100  C for 24 h. The final product was obtained after calcination at 600  C for 5 h. Synthetic conditions and physical properties are summarized in Table 1. SNSs are denoted together with their particle sizes as SNS49, 103, 160 and 282. 2.3. Fabrication of nanocomposites The nanocomposites were prepared by melt-mixing of PP, PPOH1.3 or PPOH6.4 with SNSs using a Labo-Plastmill kneading

2. Experimental 2.1. Polymer matrices PP pellets (Novatec® MA3, weight-average molecular weight Mw ¼ 397,000, molecular weight distribution Mw/Mn ¼ 3.57, isotactic triad (mm) ¼ 93.7%) were purchased from Japan Polypropylene Corporation. PPOH was synthesized using previously reported methods [23]. Poly(5-hexen-1-ol-co-propylene) containing 1.3 mol% of 5-hexen-1-ol comonomer (PPOH1.3; Mw ¼ 503,000, Mw/Mn ¼ 1.90, mm ¼ 98.0%) and poly(5-hexen-1ol-co-propylene) containing 6.4 mol% of comonomer (PPOH6.4;

Scheme 2. The chemical structures of the polymers.

Scheme 1. Schematic of the relationship between ductility and matrix-filler interfaces of nanocomposites composed of functionalized polypropylene containing hydroxyl group nanocomposites (PPOH) and silica nanospheres (SNSs).

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Table 1 Synthetic conditions of uniform sized silica nanospheres (SNSs). Sample

SNS49 SNS103 SNS160 SNS282 a b c d e f g

Reactants’ compositionsa [g] TEOSb

Seed solutionc

Argd

H2O

EtOH

5.09 10.41 10.41 10.41

2.02 0.52 0.12 0.02

0.17 0.17 0.26 0.26

47 41 41 41

125 133 133 133

Particle sizee [m2/g]

CVf [%]

BET surface areag [m2/g]

49 103 160 282

4.0 4.2 5.2 4.3

154 100 63 39

Mixture were placed in an oven at 70  C for 10 h. The products were obtained by evaporation of the solvent at 100  C followed by calcination in an oven at 600  C for 5 h. Tetraethyl orthosilicate. Seed solution is aquatic SNS14 solution (ca. 5.7 wt%) in ionized water. L-Arginine. Mean particle sizes estimated by FE-SEM images. Coefficient of variation of particle size of SNS, defined as the ratio of standard deviation and mean particle size. Determined by N2 adsorption measurement.

machine (Toyo Seiki Seisaku-sho, Japan) equipped with a KF6 twin rotary mixer (5 mL in volume) at 180  C and 100 rpm for 30 min. Original polymer materials without additional SNSs were also melt-mixed to determine a coherent thermal history. The total amount of polymer material and 0e30 wt% of SNSs was 3 g during each process. These melt-mixing processes were repeated to obtain sufficient amounts of nanocomposites. Sample sheets (100  100  0.5 mm) for structural analysis and mechanical tests were prepared using 4 g of melt-mixed samples by hot pressing at 180  C under 5 MPa for 3 min, and then under 15 MPa for 10 min using a Naflon® sheet (Nichias, Japan), stainless window flame (0.5 mm in thickness) and stainless plates, and rapidly quenched to room temperature. 2.4. Analyses The mean particle size of SNSs and the dispersion of SNSs in the nanocomposites were evaluated using field-emission scanning electron microscopy (FE-SEM, S-4800, Hitachi High-Tech Science Corporation, Japan) operated at 1 kV. Mean particle sizes were determined by measuring 100 nanoparticles in each FE-SEM image. The densities of original polymers were measured using an electronic densimeter (MDS-300, Alfa Mirage, Japan). SNSs and freezefracture surfaces of the nanocomposites were observed in secondary electron and back-scattered electron modes, respectively, without any metal coating. Differential scanning calorimetry (DSC) measurements were performed using an EXSTAR DSC 7020 (Hitachi High-Tech Science Corporation, Japan) at a heating rate of 10  C/min under N2. The sample crystallinity was determined with a melting endotherm in the first heat cycle, where the samples were heated to 200  C at 10  C/min. Fourier-transform infrared spectroscopy (FTIR) spectra were recorded using KBr pellets (SNS49) and films 0.2 mm in thickness (original polymers and nanocomposites) in transmission mode with 64 scans using a FT-IR 6100 spectrometer (Jasco, Japan), equipped with a TGS detector. Difference FTIR spectra were obtained by subtracting the spectra of original polymers (as references) from the corresponding nanocomposites by considering the peak intensity at both 1375 and 1458 cm1, which are assigned to eCH3 symmetrical and asymmetrical bending vibration, respectively [31]. X-ray photoelectron spectroscopy (XPS) spectra were recorded on a VG Theta Probe (Thermo Fisher Scientific, USA) with an Al anode Xray source. The shifts in binding energy due to relative surface charging were corrected using the C1s level at 284.6 eV as an internal standard. The Brunauer-Emmet-Teller (BET) specific surface areas were calculated from the N2 adsorption data measured by autosorb-iQ (Quantachrome instruments, USA) at 196  C at a relative pressure of 0.04e0.2 of N2 gas [32]. Tensile properties were measured by a multi-purpose stretching tester (EZ-L, Shimadzu, Japan) using a dumbbell-shaped specimen (0.5 mm in thickness,

4 mm in width and 15 mm of length of parallel part) at a crosshead speed of 10 mm/min at room temperature. Three specimens were tested for each sample. Elastic modulus was determined from data in the strain range between 0.1 mm and 0.2 mm according to JIS K 7161 (equivalent to ISO527-1). Impact tensile strength was measured by a digital impact tester (DG-UB, Toyo Seiki Seisaku-sho, Japan) using a type-L dumbbell shaped specimen (0.5 mm in thickness and 25 mm of length of parallel part). Five specimens were tested for each sample. Tests were performed with 1 J of hammer capacity and a 150 lifting angle. 3. Results and discussion 3.1. Fabrication of PP and PPOH nanocomposites melt-mixed with SNSs In order to characterize the matrix-filler interfacial adhesion, nanocomposites consisting of (i) uniform-sized SNSs in different size ranges and (ii) PP and PPOH containing 1.3 and 6.4 mol% of comonomer with hydroxyl groups (PPOH1.3 and PPOH6.4) were fabricated. Fig. 1 shows FE-SEM images of SNSs and nanocomposites melt-mixed with SNSs at 180  C for 30 min. The formation of highly uniform-sized SNSs (49e282 nm) with particle size distributions ranging from 4.0 to 5.2% CV was confirmed (Fig. 1aed and Table 1). The FE-SEM images of the nanocomposites showed that SNS49 and SNS103 formed compact aggregates in PP (Fig. 1e and f), whereas the fillers were well dispersed in PPOH1.3 and PPOH6.4 (Fig. 1i, j, m and n). These results suggest that the hydroxyl groups in PPOH improve the affinity between the matrix and SNS49. On the other hand, the larger-sized SNS160 and 282 were finely dispersed in both PP and PPOH. The difference in dispersibility of SNSs in polymers may be due to the difference in cohesive energy depending on the particle size, because the surface properties were almost unchanged in SNSs with different sizes. According to Eq (1), the cohesive energy density E between nanoparticles usually increases with decreasing particle size as [33].

E∝

nA 24R2 d

(1)

where n is the number of nanoparticles interacting in the elementary mesh, A is the Hamaker constant, R is the radius of the nanoparticles, and d is the distance between two particles. To disperse cohesive nanoparticles by melt mixing, shear energy higher than the cohesive energy is needed. The poor dispersibility of SNS49 and 103 in PP may result from an insufficient shear strength for dispersion of SNS49 and 103, because the cohesive energies of small sized SNS49 and 103 were higher than those of the large-sized SNS160 and 282.

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Fig. 1. FE-SEM images of (aed) SNSs, (eef) PP/SNS composites, (iel) PPOH1.3/SNS composites, (mep) PPOH6.4/SNS composites, sorted according to size of SNS. The contents of SNSs in present composites are 20 wt%.

3.2. Mechanical properties of PP and PPOH nanocomposites meltmixed with SNS49 At first, in order to investigate the effect of filler content on mechanical properties, the composites containing various amounts of SNS49 were employed for testing. Tensile tests were performed for the nanocomposites of PP, PPOH1.3 and PPOH6.4 melt-mixed with 0e30 wt% of SNS49 [34]. In Fig. 2, the representative stressstrain curves of the nanocomposites melt-mixed with 10 wt% of SNS49 show that the elongation at break (EB) increases with increased hydroxyl contents in the matrix polymer. Mechanical properties estimated from tensile testing of the nanocomposites melt-mixed with SNS49 are exhibited in Fig. 3. Elastic moduli of all nanocomposites were continuously increased with increased SNS49 content (Fig. 3a). Furthermore, this trend is dependent on the hydroxyl contents in the PP polymer chain. Fig. 3b exhibits relationships between rates of increase in elastic modulus of the nanocomposites melt-mixed with 5e30 wt% of SNS49 over original polymer and hydroxyl groups in PP polymer chain. The rates of increase in elastic modulus was calculated using Eq. (2)

Rate of increase in elastic modulus ¼

Ec  Em  100½% Em

(2)

where Em, Ec are elastic moduli of the matrix and the composite, respectively. The rate of increase in elastic modulus from loading SNS49 went up with the increased hydroxyl contents. When 30 wt% of SNS49 was added to PPOH6.4, the rate of increase in elastic modulus reached up to 120% from pure PPOH6.4. In contrast, the

Fig. 2. Representative stress-strain curves for nanocomposites melt-mixed with 10 wt % of SNS49; inset shows the expansion in the initial strain area.

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Fig. 3. Mechanical properties such as (a) elastic moduli, (b) rate of increase in elastic modulus, (c) yield strengths and (d) elongation at break of PP, PPOH1.3 and PPOH6.4 meltmixed with 0e30 wt% of SNS49.

addition of SNS49 to PP and PPOH1.3 resulted in just 42 and 49% improvements of elastic moduli, respectively. Thus, it was revealed that increasing the amount of hydroxyl groups in the matrix polymer enhances elastic modulus of the composite, while stiffening effect provided by SNS49 is major factor for improvement of elastic modulus. The yield strength of PP/SNS49 continuously decreased with increased amounts of SNS49. In contrast, the strengths of the PPOH1.3/SNS49 and PPOH6.4/SNS49 nanocomposites improved with increased SNS49 content (up to 20 wt%), after which the values were almost unchanged by further increases in SNS49 contents (Fig. 3c). The EB for PP nanocomposites with over 10 wt% of SNS49 decreased to less than 4%, much smaller than the EB of original PP (36%). The EB for both original PPOH1.3 (EB ¼ 423%) and PPOH6.4 (EB ¼ 556%) lie close together, and were much higher than that of original PP (EB ¼ 36%) (Fig. 3d). The EB for PPOH1.3 was significantly reduced to 84% by the addition of 10 wt% SNS49, while the EB for PPOH6.4 was 410% even after the addition of 10 wt% of SNS49. However, the EB for PPOH6.4 was gradually reduced by the addition of more than 20 wt% SNS49, much higher than that of PP and PPOH1.3. Fig. 4 shows the influence of hydroxyl contents in a matrix polymer on its EB and impact tensile strength, for PP, PPOH1.3 and PPOH6.4 melt-mixed with 10 wt% of SNS49 in this study. Note that

both EB and impact tensile strength of the nanocomposites increased with additional hydroxyl groups in the PP polymer chain. Table 2 and Fig. S2 show that the crystallinity and melting point of PP, PPOH1.3 and PPOH6.4 remained almost constant despite the changing amounts of SNS49, from 0 to 30 wt%. This indicates that the effects of matrix crystallinity on the mechanical properties can be neglected. Therefore, the changes in mechanical properties for the nanocomposites produced in this paper were possibly due to differences in matrix-filler interfacial adhesion force. PPOH6.4/SNS successfully maintained EB over 100% even after the addition of 30 wt% SNS49, whereas the ductility of PP and PPOH1.3 was decreased to 2 and 4%, respectively, by the addition of the same amount of SNS49. 3.3. Evaluation of matrix-filler interfacial adhesion The matrix-filler interfacial adhesion of the nanocomposites was estimated using previously reported theoretical models [5,20]. Based on a rigid particle assumption, Einstein’s equation for estimating the elastic modulus of particulate composites is

Ec ¼ 1 þ BVf Em

(3)

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Fig. 4. Effects of OH-comonomer contents in the matrix on (a) elongation at break and (b) impact tensile strength of nanocomposites melt-mixed with 10 wt% of SNS49.

Table 2 Mechanical properties, crystallinity and melting point of nanocomposites melt-mixed with 0 and 30 wt% SNS49. Matrix

Content of SNS49 [wt%]

Elastic modulus [MPa]

Yield strength [MPa]

PP PP PPOH1.3 PPOH1.3 PPOH6.4 PPOH6.4

0 30 0 30 0 30

1858 ± 31 2213 ± 35 1314 ± 82 1794 ± 36 439 ± 11 852 ± 20

33.6 24.7 27.8 29.8 12.4 16.4

a

± ± ± ± ± ±

0.5 0.8 0.2 0.5 0.2 0.1

Elongation at break [%]

Crystallinitya [%]

Melting pointa [ C]

36 ± 27 2±0 424 ± 70 4±1 556 ± 24 151 ± 17

46 45 37 37 19 18

167 167 144 146 114 116

Determined by DSC for matrix resin.

where B is a constant parameter related to the matrix-filler interfacial adhesion, and Vf is the filler volume fraction which can be calculated using Eq (4)

wf

rf

Vf ¼ wf

!

(4)

wm

rf þ rm

where wf and wm are weights of the filler and matrix, and rf and rm are densities of the filler (2.2 g/cm3) and matrix (0.9 g/cm3), respectively. As can be seen from Fig. 5, increase in the content of hydroxyl group in matrix polymer led to higher values of Ec/Em for all volume fractions. Although Eq (3) gives no consideration to the filler size and surface area of filler, a qualitative comparison between samples that use the same fillers is possible by this method. In Fig. 5, the linear relationship between Ec/Em and Vf, which indicate a constant B parameter, were observed in the lower Vf range. The B for PPOH6.4/SNS, PPOH1.3/SNS and PP/SNS, in range that show the linear relationship, were 15.6, 4.0 and 1.0, as shown in Fig. 5, respectively. Thus, the value of B for the nanocomposites increased with as the amount of hydroxyl groups increased in the PP polymer chain, indicating that higher amounts of hydroxyl groups enhanced the matrix-filler adhesion force. On the other hand, in higher Vf range, the Ec/Em of PPOH1.3/SNS and PPOH6.4/ SNS were located under the line. It is plausible that the small Ec/Em in higher Vf range is attributed to partly aggregation of SNS, which inhibits contact of polymer matrix with SNS surface. These results suggest that the hydroxyl groups in the matrix play a key role in the high matrix-filler adhesion of PPOH6.4 nanocomposites, leading to enhanced yield strength, impact tensile strength and EB.

Fig. 5. Relative elastic modulus (Ec/Em) of the nanocomposites as a function of volume fraction of SNS49. The solid line indicates the constant parameter B as indicated in Eq (3).

3.4. Chemical structures of matrix-filler interface Interactions between the hydroxyl groups in the matrix and the surface silanol groups of SNS49 potentially play a key role in the high matrix-filler interfacial adhesion of PPOH6.4/SNS49, as compared with PP/SNS49. The cause of this higher matrix-filler

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SNS49 with the matrices, the sharp band of (Ⅰ) disappeared completely, while the bands attributed to (Ⅲ) and (Ⅳ) increased (Fig. 6bed). These results indicate that the surface silanol groups of SNS49 interact with the matrices via hydrogen bonding. The spectrum of SNS49 with inert PP was similar to that of SNS49 added in PPOH1.3, presumably due to formation of hydrogen bonds between the surface silanol groups and antioxidants in the asreceived PP pellet. An absorption maximum shift from 3420 to 3340 cm1 was also observed in the spectrum of SNS49 mixed with PPOH6.4, suggesting the widespread formation of hydrogenbonded silanol groups assigned to types (Ⅲ) and (Ⅳ) (Fig. 6d). XPS measurements were carried out to investigate the possibility of forming interfacial covalent bonds. Si 2p XPS spectra of the nanocomposites were unchanged compared to the spectrum of neat SNS49, and did not contained any peaks attributed to interfacial covalent bonds such as SieOeC at 101.9 eV [37] (Fig. S4). This indicates that interfacial covalent bonds were absent in the nanocomposites. Therefore, the high matrix-filler adhesion of PPOH/ SNS49 nanocomposites was attributed to the formation of hydrogen bonds between hydroxyl groups in the matrix polymer and surface silanol groups in SNS49. Fig. 6. FTIR spectrum of SNS and difference spectra obtained by subtracting each spectra of the matrix polymer from those of the corresponding composites after normalized; (a) SNS49, (b) PP/SNS49, (c) PPOH1.3/SNS49, (d) PPOH6.4/SNS49.

interfacial adhesion was investigated by analyzing the chemical structures of the hydroxyl and silanol groups using FTIR. In order to analyze the silanol groups, FTIR data was collected from the neat SNS49, and difference FTIR spectra were obtained by subtracting the spectra of PP, PPOH1.3 and 6.4 (as reference) from those of their corresponding nanocomposites; these are shown in Fig. 6 (spectra of neat polymers and composites are in Fig. S3). The absorption band in the OeH stretching region is composed of (Ⅰ) an isolated silanol group at ca. 3740 cm1, (Ⅱ) a hydrogen-bonded silanol group at the oxygen atom (donor) at ca. 3640 cm1, (Ⅲ) a hydrogenbonded silanol group at the hydrogen atom (acceptor) at 3300e3500 cm1 and (Ⅳ) a hydrogen-bonded silanol group at both the oxygen and hydrogen atoms below 3300 cm1 [35,36]. The frequencies of the corresponding OeH bands decrease in order of (Ⅰ) > (Ⅱ) > (Ⅲ) > (Ⅳ). In the spectrum of neat SNS49, a sharp band at 3740 cm1 (assigned to (Ⅰ)), a shoulder band at 3683 cm1 (assigned to (Ⅱ)) and a broad band centered at 3420 cm1 (assigned to (Ⅲ) and (Ⅳ)) are all present (Fig. 6a). It is indicated that the hydrogenbonded silanol groups were then formed inside SNS49. By mixing

3.5. Effect of particle size of SNSs Effects of the SNS particle size on the mechanical properties of the nanocomposites were investigated by synthesis of nanocomposites melt-mixed with SNSs having controlled particle sizes from 49 to 282 nm. The stress-strain curves of tensile tests performed on nanocomposites containing 20 wt% SNS49, -103, -160 and -282 indicated that the mechanical properties (such as elastic modulus, yield strength and EB) of the nanocomposites depend on the SNS particle size (Fig. S5). The crystallinity and melting points were almost unchanged by varying the SNS particle size. (Fig. S6), while elastic modulus of all nanocomposites decrease with the increase in the particle size of SNS (Fig. 7a). The rates of increase in elastic modulus by SNS loading calculated by Eq (2) were increased with increased hydroxyl contents and decreased SNS particle size (Fig. S7a). Fig. 7b shows yield strengths of the nanocomposites; the yield strengths of the PP/SNSs were hardly affected by the SNS particle size. On the other hand, the yield strengths of PPOH1.3/ SNSs and PPOH6.4/SNSs decreased as the SNS particle size increased. Thus, elastic modulus and yield strength strongly depend on the SNS particle size and hydroxyl contents in the matrix. The polymer reinforcement by compositing SNSs is enhanced with decreasing SNS particle size. Fig. 7c shows the relationship between EB of the nanocomposites and the SNS particle size. EB of

Fig. 7. Effects of particle size on (a) elastic modulus, (b) yield strength, and (c) elongation at break of the nanocomposites. All nanocomposites were melt-mixed with 20 wt% of SNSs.

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PPOH6.4 nanocomposites (ca. 300%) was much higher than those of PP and PPOH1.3 nanocomposites (below 8%) and did not drastically differ at varying particle sizes. In addition, the SNS particle size did not affect the impact strength of the PPOH6.4 nanocomposites (Fig. S7b). Thus, the mechanical properties of the nanocomposite depend on the SNS particle size, and were improved by reducing the particle size. 4. Conclusions PP-based nanocomposites were fabricated by highly dispersing uniform-sized SNSs with controllable particle sizes (ranging from 49 to 282 nm) into PP, PPOH1.3 and PPOH6.4 by melt-mixing. The effects of hydroxyl groups in the matrix polymer on the interfacial adhesion force between the matrix and the filler was evaluated by examining the mechanical properties and chemical structure of the resulting nanocomposite. Tensile tests confirmed that the improvement of elastic modulus for PPOH6.4 nanocomposite was higher than that of PP and PPOH1.3. Furthermore, the ductility of PPOH6.4 remained high after addition of SNSs, whereas the ductility of PP and PPOH1.3 was drastically decreased. The matrixfiller interfacial adhesion force of the nanocomposites was investigated using Einstein’s equation to represent a relationship between elastic moduli of the matrix and nanocomposite. The high ductility of PPOH6.4/SNS49 was likely achieved through improved matrix-filler interfacial adhesion forces. Interfacial characterization using FTIR and XPS implied that the high interfacial adhesion of PPOH6.4/SNS49 was due to hydrogen bonding between the hydroxyl groups in the matrix polymer and surface silanol groups on the SNSs. The relationship between ductility and matrix-filler interfacial adhesion of PP and PPOH6.4 nanocomposites, discussed above, is illustrated in Scheme 1. Moreover, the effects of SNS particle size on the mechanical properties were evaluated by fabrication and examination of mechanical properties from nanocomposites melt-mixed with different SNS particle sizes. Elastic modulus and yield strength of PPOH6.4 nanocomposites were improved at decreased particle sizes, without any significant changes in ductility. We believe that appropriate improvements in the interfacial adhesion and decreased particle size will bring further reinforcement to PP-based nanocomposites without any losses of ductility. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.polymer.2016.06.054. References [1] S.S. Ray, M. Okamoto, Polymer/layered silicate nanocomposites: a reviews from preparation to processing, Prog. Polym. Sci. 28 (2003) 1539e1641. [2] A. Okada, A. Usuki, Twenty years of polymer-clay nanocomposites, Macromol. Mater. Eng. 291 (2006) 1449e1476. [3] C.M. Chan, J.S. Wu, J.X. Li, Y.K. Cheung, Polypropylene/calcium carbonate nanocomposites, Polymer 43 (2002) 2981e2992. [4] N. Hasegawa, H. Okamoto, M. Kato, A. Usuki, Preparation and mechanical properties of polypropylene-clay hybrids based on modified polypropylene and organophilic clay, J. Appl. Polym. Sci. 78 (2000) 1918e1922. [5] S.Y. Fu, X.Q. Feng, B. Lauke, Y.W. Mai, Effects of particle size, particle/matrix interface adhesion and particle loading on mechanical properties of particulate-polymer composites, Compos. Part. B Eng. 39 (2008) 933e961. [6] S. Iwamoto, S. Yamamoto, S.H. Lee, T. Endo, Mechanical properties of polypropylene composites reinforced by surface-coated microfibrillated cellulose, Compos. Part A Appl. Sci. Manuf. 59 (2014) 26e29. [7] V.N. Dougnac, R. Alamillo, B.C. Peoples, R. Quijada, Effect of particle diameter on the permeability of polypropylene/SiO2 nanocomposites, Polymer 51 (2010) 2918e2926. [8] M.Z. Rong, M.Q. Zhang, Y.X. Zheng, H.M. Zeng, K. Friedrich, Improvement of tensile properties of nano-SiO2/PP composites in relation to percolation, Polymer 42 (2001) 3301e3304.

[9] P. Song, Z. Cao, Y. Cai, L. Zhao, Z. Fang, S. Fu, Fabrication of exfoliated graphene-based polypropylene nanocomposites with enhanced mechanical and thermal properties, Polymer 52 (2011) 4001e4010. [10] T. Kashiwagi, E. Grulke, J. Hilding, K. Groth, R. Harris, K. Butler, J. Shilds, S. Kharchenko, J. Douglas, Thermal and flammability properties of polypropylene/carbon nanotube nanocomposites, Polymer 45 (2004) 4227e4239. [11] J.W. Gilman, C.L. Jackson, A.B. Morgan, R. Harris Jr., Flammability properties of polymer-layered-silicate nanocomposites. Polypropylene and polystyrene nanocomposites, Chem. Mater. 12 (2000) 1866e1873. [12] V. Vladimirov, C. Betchev, A. Vassiliou, G. Papageorgiou, D. Bikiaris, Dynamic mechanical and morphological studies of isotactic polypropylene/fumed silica nanocomposites with enhanced gas barrier, Compos. Sci. Technol. 66 (2006) 2935e2944. [13] K. Kalaitzidou, H. Fukushima, L.T. Drzal, Multifunctional polypropylene composites produced by incorporation of exfoliated graphite nanoplatelets, Carbon 45 (2007) 1446e1452. [14] D.N. Bikiaris, G.Z. Papageorgiou, E. Pavlidou, N. Vouroutzis, P. Palatzoglou, G.P. Karayannidis, Preparation by melt mixing and characterization of isotactic polypropylene/SiO2 nanocomposites containing untreated and surface-treated nanoparticles, J. Appl. Polym. Sci. 100 (2006) 2684e2696. [15] K.A. Iyer, J.M. Torkelson, Importance of superior dispersion versus filler surface modification in producing robust polymer nanocomposites: the example of polypropylene/nanosilica hybrids, Polymer 68 (2015) 147e157. [16] D. Bikiaris, A. Vassiliou, K. Chrissafis, K.M. Paraskevopoulos, A. Jannakoudakis, A. Docoslis, Effect of acid treated multi-walled carbon nanotubes on the mechanical, permeability, thermal properties and thermo-oxidative stability of isotactic polypropylene, Polym. Degrad. Stab. 93 (2008) 952e967.
tink, R.J. Gaymans, Toughening of [17] W.C.J. Zuiderduin, C. Westzaan, J. Hue polypropylene with calcium carbonate particles, Polymer 44 (2003) 261e275. [18] S.C. Tjong, Structural and mechanical properties of polymer nanocomposite, Mater. Sci. Eng. R. 53 (2006) 73e197. [19] S.L. Phua, L. Yang, C.L. Toh, S. Huang, Z. Tsakadze, S.K. Lau, Y.W. Mai, X. Lu, Reinforcement of polyether polyurethane with dopamine-modified clay: the role of interfacial hydrogen bonding, ACS Appl. Mater. Interfaces 4 (2012) 4571e4578. [20] D.N. Bikiaris, A. Vassiliou, E. Pavlidou, G.P. Karayannidis, Compatibilisation effect of PP-g-MA copolymer on iPP/SiO2 nanocomposites prepared by melt mixing, Eur. Polym. J. 41 (2005) 1965e1978. [21] W. Li, J. Kagar-Kocsis, R. Thomann, Compatibilization effect of TiO2 nanoparticles on the phase structure of PET/PP/TiO2 nanocomposites, J. Polym. Sci. Part B Polym. Phys. 47 (2009) 1616e1624. [22] H. Hagihara, T. Ishihara, T.B. Hoang, T. Shiono, Precise control of microstructure of functionalized polypropylene synthesized by the ansa-zirconocene/ MAO catalysts, J. Polym. Sci. Part A Polym. Chem. 46 (2008) 1738e1748. [23] Y. Iizuka, J. Sugiyama, H. Hagihara, Unexpected mechanical properties of functionalized polypropylene: tensile test, Charpy impact tensile test, DSC, and WAXD analysis of poly(5-hexen-1-ol-co-propylene), Macromolecules 42 (2009) 2321e2323. [24] H. Hagihara, K. Ito, S. Kimata, Comprehensive study of altered amorphous structure in functionalized polypropylenes exhibiting high tensile strength, Macromolecules 46 (2013) 4432e4437. [25] H. Hagihara, K. Tsuchihara, K. Takeuchi, M. Murata, H. Ozaki, T. Shiono, Copolymerization of ethylene or propylene with a-olefins containing hydroxyl groups with zirconocene/methylaluminoxane catalyst, J. Polym. Sci. Part A Polym. Chem. 42 (2004) 52e58. [26] T. Taniike, M. Toyonaga, M. Terano, Polypropylene-grafted nanoparticles as a promising strategy for boosting physical properties of polypropylene-based nanocomposites, Polymer 55 (2014) 1012e1019. [27] T. Yokoi, Y. Sakamoto, O. Terasaki, Y. Kubota, T. Okubo, T. Tatsumi, Periodic arrangement of silica nanospheres assisted by amino acids, J. Am. Chem. Soc. 128 (2006) 13664e13665. [28] T. Yokoi, M. Iwama, R. Watanabe, Y. Sakamoto, O. Terasaki, Y. Kubota, J.N. Kondo, T. Okubo, T. Tatsumi, Synthesis of well-ordered nanospheres with uniform mesopores assisted by basic amino acids, Stud. Surf. Sci. Catal. 170B (2007) 1774e1780. [29] T. Yokoi, J. Wakabayashi, Y. Otsuka, W. Fan, M. Iwama, R. Watanabe, K. Aramaki, A. Shimojima, T. Tatsumi, T. Okubo, Mechanism of formation of uniform-sized silica nanospheres catalyzed by basic amino acids, Chem. Mater 21 (2009) 3719e3729. [30] R. Watanabe, T. Yokoi, E. Kobayashi, Y. Otsuka, A. Shimojima, T. Okubo, T. Tatsumi, Extension of size of monodisperse silica nanospheres and their well-ordered assembly, J. Colloid Interface Sci. 360 (2011) 1e7. [31] M.P. McDonald, I.M. Ward, The assignment of the infra-red absorption bands and the measurement of tacticity in polypropylene, Polymer 2 (1961) 341e355. [32] S. Brunauer, P.H. Emmett, E. Teller, Adsorption of gases in multimolecular layers, J. Am. Chem. Soc. 60 (1938) 309e319. [33] W. Hani, L. Vonna, H. Haidara, Experimental characterization of the nanoparticle size effect on the mechanical stability of nanoparticle-based coating, Nano. Lett. 15 (2015) 442e449. [34] The some mechanical properties of PP and PPOH differ from those of previous publication [23]. It is plausible that the differences is attributed to the differences of thermal history in preparing test specimens, which can affect morphology of the polymer and thereby change the mechanical properties. In this study, as with the composite materials, specimens of PP and PPOH were

R. Watanabe et al. / Polymer 99 (2016) 63e71 also made by melt-mixing at 180  C for 30 min, and then hot pressing at 180  C under 15 MPa for 10 min. In previous paper, on the other hand, all of the specimens were made by hot pressing at 190  C under 5 MPa for 3 min without melt-mixing process. €zinger, Novel aspects of mid and far IR Fourier spectros[35] P. Hoffmann, E. Kno copy applied to surface and adsorption studies on SiO2, Surf. Sci. 188 (1987)

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181e198. [36] A. Fidalgo, L.M. Ilharco, The defect structure of sol-gel-derived silica/polytetrahydrofuran hybrid films by FTIR, J. Non Cryst. Sol. 283 (2001) 144e154. [37] L. Yang, L. Wang, M. Xing, J. Lei, J. Zhang, Silica nanocrystal/graphene composite with improved photoelectric and photocatalytic performance, Appl. Catal. B Environ. 180 (2016) 106e112.