Reinforcement mechanism of functionalized polypropylene containing hydroxyl group nanocomposites studied by rheo-optical near-infrared spectroscopy

European Polymer Journal 92 (2017) 86–96

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European Polymer Journal journal homepage: www.elsevier.com/locate/europolj

Macromolecular Nanotechnology

Reinforcement mechanism of functionalized polypropylene containing hydroxyl group nanocomposites studied by rheo-optical near-infrared spectroscopy

MARK

⁎

Ryota Watanabe, Hideyuki Shinzawa , 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 3058565, Japan

AR TI CLE I NF O

AB S T R A CT

Keywords: Near-infrared (NIR) Two-dimensional [2D] correlation spectroscopy Reinforcement of nanocomposites

Functionalized polypropylene with hydroxyl group (PPOH) nanocomposites with varying chemical structures in filler-matrix interfaces were examined by rheo-optical near-infrared measurements combined with two-dimensional correlation spectroscopy, to provide a better understanding of the reinforcement mechanism of nanocomposites. PPOH nanocomposites were fabricated by melt-mixing (i) silica nanospheres (SNS) and (ii) SNS modified with surfactant (mSNS). The decrease in the asynchronous correlation intensity from PPOH to SNS/PPOH indicated that the tensile deformation of amorphous and crystalline occurred simultaneously. These observations implied that the high interfacial adhesion between filler and matrix restricted the deformation of amorphous. This mechanism seems to contribute to the improvement in mechanical strength of the composite. The asynchronous correlation intensity of mSNS/PPOH was higher than that of SNS/PPOH, indicating that a certain level of deformation of the amorphous was allowed. It is therefore likely that the surface modification of SNS prevented filler-matrix interactions, leading to relatively low tensile strength of mSNS/PPOH. Thus, the techniques based on NIR spectroscopy clarify the effects of the filler-matrix interfacial adhesion of nanocomposites on the behavior of the crystalline and amorphous components during tensile testing, and may be generally applicable to evaluate filler-matrix adhesion behavior of a wide range of polymer nanocomposite systems.

1. Introduction Extensive work has been performed to develop polypropylene (PP) nanocomposites by dispersing nanofillers into the material due to the improved mechanical strength of such combinations over original PP [1–12]. To this end, several factors influencing the mechanical properties of the nanocomposites must be taken into account, including filler dispersion and filler-matrix interfacial adhesion; both of these are significantly influenced by chemical structures of filler-matrix interfaces [9–14]. We have recently developed PP functionalized with hydroxyl groups (PPOH) [15–18] and PPOH/monodisperse silica nanospheres (SNS) nanocomposites [19]. Here, the SNS were highly dispersed into single nanoparticles within the PPOH by simple melt-mixing due to the high affinity between SNS and PPOH. In addition, hydrogen bonding interactions between hydroxyl groups of the PPOH and silanol groups on the SNS surface enhanced the filler-matrix interfacial adhesion of the SNS/PPOH. While the incorporation of SNS noticeably ⁎

Corresponding authors. E-mail addresses: [email protected] (H. Shinzawa), [email protected] (H. Hagihara).

http://dx.doi.org/10.1016/j.eurpolymj.2017.04.032 Received 16 February 2017; Received in revised form 18 April 2017; Accepted 20 April 2017 Available online 22 April 2017 0014-3057/ © 2017 Elsevier Ltd. All rights reserved.

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improves the mechanical properties of the mixture, the detailed reinforcement mechanism of SNS/PPOH as it pertains to its fillermatrix interfacial adhesion is not fully understood, due to the absence of characterization techniques that can probe the dynamic deformation of the polymer system at the molecular level during mechanical processes. Rheo-optical characterization can provide information on the molecular-level deformation behavior of polymers undergoing transient structural changes during mechanical processes [20–23]. For example, in tensile tests, rheo-optical techniques can capture spectral intensity changes caused by the orientation of the polymer chain along the tensile direction. Comparisons between the spectral features and the corresponding stress or strain behavior provides key information on the evolution of the polymer structures that govern the macroscopic deformation of polymer samples. We previously introduced rheo-optical techniques based on nearinfrared (NIR) spectroscopy, which was used to reveal the molecular-level differences in the tensile deformations of crystalline and amorphous components in semicrystalline polyolefins such as polyethylene (PE) [21] and PP [22]. These studies revealed that the deformation of these polyolefins involved separate steps of elongation of amorphous chains prior to the onset of disintegration in the crystalline lamellae. Rheo-optical NIR measurements are also promising for analyzing polymer nanocomposites. For example, the measurement of a reinforced PP melt-mixed with clay revealed that nanoclay layers dispersed into the PP matrix restrict the elongation of amorphous chains [22]. Therefore, rheo-optical characterizations of nanocomposites with well-controlled chemical structures at the filler-matrix interface can clarify the effects of the filler-matrix interfacial adhesion on the behavior of the crystalline and amorphous components during tensile testing, and provide useful information about the reinforcement mechanism related to the filler-matrix interfacial adhesion. Herein, we aimed to provide a better understanding of the reinforcement mechanism of polymer nanocomposites by combining results obtained from rheo-optical NIR measurements, tensile tests and dynamic mechanical analyses (DMA) of model PPOH nanocomposites with varying chemical structure in the filler-matrix interface. Two types of model nanocomposites were fabricated; (i) SNS/PPOH and (ii) PPOH melt-mixed with SNS modified with a cationic surfactant (CTAB, cetyltrimethyl ammonium bromide) denoted as mSNS/PPOH. In addition, we investigated whether the techniques based on NIR spectroscopy could be possibly applied to examine filler-matrix interfacial adhesion of nanocomposites. 2. Method 2.1. 2D correlation spectroscopy The basic idea of generalized 2D correlation spectroscopy is the analysis of synchronicity and asynchronicity of dynamic spectra, which are induced by the applied external perturbation, such as concentration [14–16]. Assume a set of dynamic spectra at m evenly spaced points in time t between Tmin and Tmax represented as follows;

⎡ y∼(ν,t1) ⎤ ⎢∼ ⎥ ∼ y(ν ) = ⎢ y (ν,t2 ) ⎥ ⎢ ··· ⎥ ⎢⎣ y∼(ν,tm )⎥⎦

(1)

⎧ y(ν,t )−y (ν ) for Tmin ⩽ t ⩽ Tmax y∼(ν,t ) = ⎨ ⎩0 otherwise

(2)

where y (ν ) is the reference spectrum of the system. In most cases, an averaged spectrum is customary given as follows:

y (ν ) =

1 Tmax−Tmin

∫T

Tmax

y(ν,t )dt

(3)

min

Synchronous and asynchronous correlation spectra are obtained from the dynamic spectra by

Φ(ν1,ν2 ) =

1 ∼ t∼ y(ν1) y(ν2 ) m−1

(4)

Ψ(ν1,ν2 ) =

1 ∼ t ∼ y(ν1) N y(ν2 ) m−1

(5)

where N is the Hilbert-Noda transformation matrix given by

⎡ 0 1 1/2 1/3 ⋯⎤ ⎢ 0 1 1/2 ⋯⎥ 1 ⎢ −1 ⎥ N= −1/2 −1 0 1 ⋯⎥ π⎢ ⎢⎣ −1/3 −1/2 −1 0 ⋯⎥⎦ ⋯ ⋯ ⋯ ⋯ ⋯

(6)

The intensity of a synchronous correlation spectrum Φ(ν1,ν2 ) represents simultaneous changes of two spectral intensity variations measured at v1 and v2 during the interval between Tmin and Tmax. On the other hand, an asynchronous correlation spectrum Ψ(ν1,ν2 ) means sequential changes of spectral intensity measured at v1 and v2. Fig. 1 illustrates schematic contour maps of synchronous and asynchronous correlation spectra. Shaded area indicate negative correlation intensity. Plot of reference spectrum is placed at the top and side of contour map. Peaks appeared on a synchronous 87

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Fig. 1. Schematic illustration of contour maps of (a) synchronous and (b) asynchronous correlation spectra.

correlation spectrum represent simultaneous changes of spectral intensities measured at v1 and v2. The sign of a synchronous cross peak becomes positive if the spectral intensities at v1 and v2 are either increasing or decreasing during the observation period. On the other hand, a negative sign of the peak indicates that one of the spectral intensities is increasing while the other is decreasing. In Fig. 1a, the negative cross peak between spectral coordinate A and B indicate that spectral intensity at one coordinate is increasing while that of the other is decreasing. An asynchronous cross peak develops only if the intensities at v1 and v2 change out of phase, e.g., delayed or accelerated. Cross peak observed at the coordinate (v1, v2) reveals that spectral intensity at v1 predominantly occurs before that at v2. However this sign rule is reversed if the synchronous correlation intensity at the same coordinate becomes negative. For example, the positive asynchronous correlation peak at the coordinate (A, B) indicates predominant intensity change at B before that at A since the corresponding synchronous correlation peak is negative. 3. Experimental 3.1. Polymer matrices PPOH was synthesized using previously reported methods [16]. Poly(5-hexen-1-ol-co-propylene) containing 6.4 mol% of comonomer (PPOH6.4; Mw = 146,000, Mw/Mn = 1.95, mm = 98.4%) was used. 3.2. Syntheses of SNS and mSNS Silica nanospheres (SNS): SNS 49 nm in size were synthesized using previously reported methods [19,24]. Surface-modified SNS with CTAB (mSNS): 0.5 g of CTAB (Wako, > 98.0%) was added to 100 g of a SNS synthesis solution with stirring at room temperature for 5 min, and white precipitates were readily formed. This precipitate is an aggregation of mSNS formed by “hydrophobic effects” [25], and was collected using centrifugation for 5 min at 3000 rpm. Finally, the product was dried in an oven at 70 °C for 10 h. 3.3. Fabrication of nanocomposites The SNS/PPOH and mSNS/PPOH nanocomposites were prepared by melt-mixing using a Labo-Plastmill kneading 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. An original PPOH sample without additional nanofillers was also melt-mixed to determine a coherent thermal history. The total amount of polymer material, including 10 wt% SNS or mSNS, was 3 g during each process. These melt-mixing processes were repeated to obtain sufficient amounts of each nanocomposite. Sample sheets of 1 mm thickness (for rheo-optical NIR measurements) and of 0.5 mm thickness (for structural analyses and mechanical tests) were prepared 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. 3.4. Analyses Filler dispersion in the nanocomposites was evaluated using field-emission scanning electron microscopy (FE-SEM, S-4800, Hitachi High-Tech Science Corporation, Japan) operated at 1 kV. Freeze-fracture surfaces of the nanocomposites were observed in 88

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Scheme 1. Schematic illustration of the rheo-optical NIR spectroscopy setup. (Reproduced with permission from H. Shinzawa, W. Kanematsu, I. Noda, Rheo-optical near-infrared (NIR) spectroscopy study of low-density polyethylene (LDPE) in conjunction with projection two-dimensional (2D) correlation analysis, Vib. Spectrosc. 70. (2014) 53–57.)

back-scattered electron modes respectively, without any pretreatment such as polishing and metal coating. The densities of the original polymers were measured using an electronic densimeter (MDS-300, Alfa Mirage, Japan). Tensile tests were carried out by a multi-purpose stretching tester (EZ-L, Shimadzu, Japan) according to ISO527-1, using dumbbell-shaped specimens (0.5 mm in thickness, 4 mm in width and 15 mm of length in the parallel region) at a crosshead speed of 10 mm/min at room temperature. Three specimens were tested for each sample. Elastic moduli were determined from data in a strain range of 0.1–0.2 mm. Dynamic mechanical analyses (DMA) was performed by Q800DMA (TA Instrument, US) using rectangular specimens (40 × 5 × 0.5 mm) in a temperature range from −20 to 80 °C at 3 °C/min, with a 1 Hz frequency and 15 μm vibration width in the tensile mode. Rheooptical NIR measurements were performed using an acousto-optic tunable filter (AOTF) NIR spectrometer (Systems Engineering, Japan) equipped with a tensile testing machine (MX2-2500N, Imada, Japan). The specimens were gradually deformed by the mechanical stretcher while being probed with a NIR beam polarized perpendicular to the deformation direction (Scheme 1). The 50 × 20 × 1 mm samples were stretched at a speed of 0.5 mm/min. Sets of NIR spectra were collected every 8 s by co-adding 512 scans over the 1300–2000 nm region, and the corresponding load and strain were also recorded simultaneously. 4. Results and discussion 4.1. Fabrication of PPOH nanocomposites In order to control the chemical structure of the filler-matrix interface of the PPOH nanocomposites, we prepared two types of nanofillers; (i) SNS with an unmodified silica surface and (ii) SNS modified with CTAB (mSNS). SEM images of the SNS and mSNS confirmed the formation of highly uniform-sized nanospheres (Fig. S1). The mean particle sizes of SNS and mSNS were 49 and 63 nm, respectively. The SNS particle size was slightly smaller than that of mSNS, due to the calcination that occurs when synthesizing SNS for further condensation of the silanol. According to prior results, such a slight difference in the particle size does not influence the mechanical properties of the different nanocomposites [19]. Thermogravimetric (TG) and FTIR studies revealed that mSNS contained 4 wt% CTAB, indicating that about 10% of the silanol interacted with CTAB (Figs. S2 and S3). There are no differences in the size of initial agglomerates between SNS and mSNS. SEM images of the fracture surfaces of SNS/PPOH and mSNS/PPOH are shown in Fig. 2a and b. Both SNS and mSNS were uniformly dispersed into PPOH, and the dispersions of SNS and mSNS into PPOH were nearly identical. The initial agglomerates were not found in both composites. In addition, differential scanning calorimetry (DSC) measurements confirmed that the crystallinities of PPOH, SNS/PPOH and mSNS/PPOH estimated by heat of fusion of the melting peaks were nearly identical, at 20, 20 and 21%, respectively. The wide-angle X-ray diffraction (WAXD) patterns of PPOH, SNS/PPOH and mSNS/PPOH are shown in Fig. S4. There were no significant differences between the patterns. These results suggest that the differences between the filler dispersion in the matrix, the crystallinity and the crystal structure of the polymer matrix may be negligible for the following discussion. Plausible schematic illustrations of the polymer-filler interfaces of SNS/PPOH and mSNS/PPOH are shown in Fig. 2c and d. The SNS/PPOH exhibits high filler-matrix interfacial adhesion due to hydrogen bonding interactions between SNS and PPOH [19]. However, the bulky CTAB molecules electrostatically adsorbed on the mSNS surface likely work as spacers between mSNS and PPOH, reducing the filler-matrix interfacial adhesion. 4.2. Mechanical properties of PPOH nanocomposites Fig. 3 shows stress-strain curves of PPOH, SNS/PPOH and mSNS/PPOH based on tensile testing in the rheo-optical system. The curves show an increase in load at the initial strain area of the tensile test, indicating elastic deformation. Further tensile loading 89

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Fig. 2. SEM images of (a) SNS/PPOH and (b) mSNS/PPOH containing 10 wt% fillers. Schematic illustrations of polymer-filler interfaces of (c) SNS/PPOH and (d) mSNS/PPOH.

resulted in irreversible plastic deformation. Yield strengths were enhanced by melt-mixing the fillers into PPOH. The strength of SNS/ PPOH was higher than that of mSNS/PPOH, implying more favorable interactions between SNS and the PPOH matrix [9]. Such effects, on the other hand, are less acute for the mSNS. It is therefore likely that surface modification of SNS substantially prevents the development of interactions between SNS and PPOH. The mechanical properties estimated from the tensile testing according to ISO527-1 are summarized in Table 1. The yield

Fig. 3. Stress-strain curves of original PPOH, SNS/PPOH and mSNS/PPOH using the tensile strength machine of the rheo-optical system.

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Table 1 Mechanical properties of PPOH, SNS/PPOH and mSNS/PPOH. Sample

PPOH SNS/PPOH mSNS/PPOH

Rheo-optical NIR measurementa

ISO527-1b

Filler content [wt%]

Yield strength [MPa]

Yield strength [MPa]

Elastic modulus [MPa]

Constant parameter B

0 10 10

10.2 13.7 11.1

12.4 ± 0.2 16.7 ± 0.1 15.1 ± 0.3

439 ± 11 686 ± 6 528 ± 16

– 12.9 4.7

a

Rectangular specimens (20 × 50 × 1 mm) were tested at a speed of 0.5 mm/min. Dumbbell-shaped specimens (0.5 mm in thickness, 4 mm in width and 15 mm in length of the parallel region) were tested at a speed of 10 mm/min according to ISO525-1. b

strengths were different depending on the measuring methods because of the differences in the cross head speed and the specimen shape. However, the yield strengths measured in this manner showed a similar trend with the yield strengths measured by tensile testing with the rheo-optical system. Elastic moduli were also enhanced by melt-mixing the fillers, and the elastic modulus of SNS/ PPOH was higher than that of mSNS/PPOH. The filler-matrix interfacial adhesion force of the nanocomposites was evaluated using Einstein’s equation as previously reported [12,19]. A constant parameter B related to the filler-matrix interfacial adhesion can be calculated using Eq. (7)

Ec = 1 + BVf Em

(7)

where Em and Ec are the elastic moduli of the matrix and composite, respectively. Vf is a filler volume fraction, which can be calculated using Eq. (8)

Vf =

wf ρf

(

wf ρf

+

wm ρm

)

(8)

where wf and wm are the weights of the filler and matrix, and ρf and ρm are the densities of the filler (2.2 g/cm ) and matrix (0.9 g/ cm3), respectively. The parameter B of SNS/PPOH was higher that of mSNS/PPOH, suggesting that the filler-matrix interfacial adhesion of SNS/PPOH was higher than that of mSNS/PPOH. Fig. 4 shows storage moduli (E′) of PPOH, SNS/PPOH and mSNS/PPOH measured using DMA. The E′ of SNS/PPOH was higher than those of PPOH and mSNS/PPOH at 0–50 °C, consistent with the elastic moduli estimated from tensile testing. However, the modulus of SNS/PPOH approaches that of mSNS/PPOH by increasing the temperature above 50 °C, presumably due to the weakened hydrogen bonding interactions between SNS and PPOH. Again, the surface modification of the SNS as well as the presence of nanospheres is closely related to the physical behavior of the polymer system. Fig. S5 shows the loss factor tan δ of PPOH, SNS/PPOH and mSNS/PPOH as a function of temperature. The major tan δ peak is caused by the segmental motion of the main chain backbone in amorphous part. Therefore, the peak temperature on tan δ is often used as to denote the glass transition temperature Tg, although that is slightly higher than Tg measured by DSC in general [26]. The shape of tan δ curve was slightly changed by addition of SNS or mSNS, indicating that molecular mobility of some part of amorphous was influenced by the fillers. Nevertheless, the peak temperature on tan δ was almost unchanged (that of PPOH, SNS/PPOH and mSNS/PPOH are 26.9, 27.3 and 27.0 °C, respectively), suggested that thermodynamics of most part of amorphous was not affected by fillers. 3

Fig. 4. Dynamic storage moduli of PPOH, SNS/PPOH and mSNS/PPOH.

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Fig. 5. Strain-dependent NIR spectra of (a) original PPOH, (b) SNS/PPOH and (c) mSNS/PPOH.

Consequently, the two types of nanocomposites with different chemical structures at their filler-matrix interfaces show different mechanical properties and filler-matrix interfacial adhesion. As such, molecular-level observations of the deformation behaviors based on rheo-optical spectroscopy becomes an important task to provide greater insight into the reinforcement afforded by enhancing the filler-matrix interfacial adhesion. The strong adhesion between filler and matrix might restrict the displacement or deformation of the amorphous and/or crystalline regions of the polymer. 4.3. Strain-dependent NIR spectra Fig. 5a–c shows strain-dependent NIR spectra of PPOH, SNS/PPOH and mSNS/PPOH. The spectra were measured with a NIR beam polarized perpendicular to the deformation direction. Entire features of these spectra are complicated by the presence of overlapped contributions from overtone modes associated with CH2 or CH3 groups in the polymer chains [27]. It is important to point out comparison of the spectra of the samples reveals that the inclusion of the SNS and mSNS does not induce apparent change in the 92

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Fig. 6. (a) Synchronous and (b) asynchronous correlation spectra derived from strain-dependent spectra of PPOH. Negative correlation intensity areas of the contour maps are represented by shading.

spectral feature of the PPOH probably due to the inclusion of small amount of the SNS and CTAB. Semicrystalline PP samples prepared from the melt show a complex polymer structure consisting of crystal lamellae embedded in a liquid-like amorphous matrix. The elongation of the sample results in the displacement of the crystalline and amorphous structures. This induces the polymer chain to orient parallel to the tensile direction, eventually leading to a decrease in the NIR absorbance for both the crystalline and amorphous components. In other words, the intensity variations of these NIR bands can be used to empirically detect changes in the amorphous and crystalline components. Rheo-optical NIR measurements enable the determination of spectral intensity changes derived from the orientation of polymer chains [22]. However, changes in the sample thickness and light scattering can also cause substantial variation of the spectral intensity, overwhelming the main features of the NIR spectra [22]. For example, the spectral baseline changes from one spectrum to another (Fig. 5); it is obvious that these changes in spectral intensity caused by the orientation of polymer structures is overwhelmed by exceptionally pronounced baseline changes. In addition, peaks arising from the crystalline and amorphous structures of the PPOH are located close to each other. Neighboring peaks tend to merge into continuous peaks with less demarcation, making the identification of individual peaks and determination of the sequential order of their intensities variations difficult.

4.4. 2D correlation analysis The examination of rheo-optical NIR spectra is often combined with 2D correlation spectroscopy by taking advantage of resolution enhancement in 2D correlation spectra [21,22,28]. Analysis of the signs of 2D correlation peak intensities leads to simple and practical assignments of the order of sequential events represented by the spectral signal variations. When the spectral intensity variations caused by perturbations (e.g., tensile deformation) occur asynchronously rather than simultaneously, one can deduce the order of this sequence from the signs of cross peaks appearing in the 2D correlation spectra. Such advantageous features often provide an additional opportunity in the analysis of rheo-optical NIR spectra. Fig. 6 exhibits synchronous and asynchronous correlation spectra generated from the NIR spectrum of PPOH in Fig. 5a. Note that the synchronous and asynchronous correlation spectra were calculated from the original NIR spectra using well-established 2D correlation equations [29,30]. Prior to calculating the 2D correlation spectra, the spectra were subjected to baseline correction with mathematical projection based on offset collection. In this study, this baseline treatment consists of the removal of the offset deviation by subtracting the intensities at 1650 nm and subsequent 93

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elimination of multiplicative scatter factor by orthogonal projection with the intensities at 1800 nm [21,22,28]. All the calculations for 2D correlation spectra were performed by in-house programs in MATLAB (The Math Works, USA). In Fig. 6, the plot of the average spectrum is placed at the top and side of the contour map for reference. Note that negative correlation intensity areas of the contour maps are represented by shading. The entire plane of a synchronous correlation spectrum is coved with positive value. This indicates that the spectral intensities in this spectral region change to the same direction during the tensile deformation of the sample. A correlation peak appearing at a diagonal position is referred to as an autopeak, revealing that spectral intensities in this wavelength region is especially acute. For example, the generation of a distinct auto correlation peak around 1720 nm indicates a gradual decrease in the spectral intensities in this NIR region. This can be interpreted as a dominant change in the spectral intensities of the PPOH bands caused by the elongation of the sample, reflecting mechanical deformation of the polymer structures. The asynchronous spectrum consists of cross peaks located at off-diagonal positions. An asynchronous cross peak develops only if the intensities of two spectral features change out of phase (i.e., delayed or accelerated) with each other. The features of an asynchronous spectrum are especially useful in differentiating overlapped bands arising from spectral signals of different origins. Even if bands are located close to each other, as long as the signatures or the pattern of sequential variations of spectral intensities are substantially different, asynchronous cross peaks will develop between their spectral coordinates. For example, the cross peaks at the coordinates (1720, 1733) and (1720, 1700) indicate that the spectral intensity change at 1720 nm occurs before that at 1700 or 1733 nm. In other words, the generation of the asynchronous correlation peaks reveal the presence of different deformation mechanism undergoing different defamation during the elongation, namely elastic and plastic deformations. It is hence reasonable to conclude that the first deformation represented with the intensity at 1720 nm is associated with the elastic amorphous component of the PPOH. The later deformation process indicated by the intensity at 1700 or 1733 nm is mostly due to the plastic deformation caused by the displacement of the hard crystalline structure. Consequently the development of positive correlation peaks can be interpreted to mean that the decreased spectral intensity of the amorphous band predominantly occurs before that of the crystalline bands. Namely, the crystalline lamellae are aligning in the direction of strain and the CH2 groups and their corresponding overtone bands now have electric dipole-transition moments oriented perpendicular to the helical chains in the lamellae. It is likely the elastic deformation by elongation of amorphous chains of PPOH is induced at the initial stage of the tensile test, and further stretching causes plastic deformation by displacement and disintegration of the crystalline lamellae. In fact, such a sequence agrees with the actual physical model describing tensile deformation of semicrystalline polymers [21,22,27]. Fig. 7 summarizes the asynchronous correlation intensities of the cross peak at the coordinate (1720, 1700) derived from the PPOH and PPOH nanocomposites. Note that the correlation intensities of the SNS/PPOH and mSNS/PPOH were also calculated in a similar manner to the case of PPOH. The asynchronous correlation intensity can be interpreted as an index representing the delay in the deformation. For example, a sufficient asynchronous correlation intensity for PPOH indicates an obvious delay in the deformation of its crystalline structure. The deformation of rigid crystals requires even greater mechanical force compared to the elastic amorphous component. The apparent decrease in the correlation intensity of the SNS/PPOH suggests a less significant delay in the deformation of the amorphous and crystalline components. In other words, the tensile deformation of these components occurs more or less simultaneously. This implies the generation of interfacial adhesion between the SNS and polymer matrix, which in turn restricts the deformation of crystalline as well as amorphous structures. The correlation intensity of the mSNS/PPOH shows a slight increase while it is smaller than that of PPOH. It is therefore likely that surface modification of the SNS prevents interactions between the mSNS and PPOH, allowing a certain level of deformation of the amorphous component. By combining the results from the tensile tests, DMA and rheo-optical NIR measurements, we propose a possible mechanism for the tensile deformation of PPOH and PPOH nanocomposites, shown in Scheme 2. The hydrogen bonding interactions between silanol groups on the surface of SNS and hydroxyl groups (present mainly in the amorphous chains of PPOH) induce strong mobility restriction in the amorphous chains. As such, the amorphous chains tend to deform along with the crystal lamellae attached to the amorphous chain during tensile testing of the SNS/PPOH. Such synchronous deformation of the amorphous chains and crystalline lamellae would, in turn, result in improved mechanical properties. Furthermore, the modification of SNS with CTAB molecules reduces the filler-matrix interactions and causes low mobility restriction of the amorphous chains, resulting in the lower

Fig. 7. The asynchronous correlation intensities of the cross peak at the (1720, 1700) of PPOH, SNS/PPOH and mSNS/PPOH.

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Scheme 2. A schematic illustration of the reinforcement mechanism of PPOH/SNS by hydrogen bonding interactions between SNS and amorphous PPOH.

reinforcement of the mSNS/PPOH compared to SNS/PPOH. While this study is a characterization of a specific nanocomposite system to evaluate the effect of matrix-filler adhesion on the mechanical properties of nanocomposites, the discovered reinforcement mechanism related to the filler-matrix interfacial adhesion should be generally accepted for the other nanocomposites.

5. Conclusions PP-based nanocomposites with various chemical structures at the filler-matrix interface were fabricated by melt-mixing of uniform-sized SNS and mSNS (SNS modified with CTAB) into PPOH. Tensile tests and DMA revealed that the yield strength and elastic modulus of SNS/PPOH were higher than those of mSNS/PPOH, owing to greater filler-matrix interfacial adhesion of SNS/ PPOH compared to mSNS/PPOH. The deformation behaviors of amorphous and crystalline components in the PPOH and the different PPOH nanocomposites were evaluated by rheo-optical NIR measurements combined with 2D correlation spectra. The amorphous chains and crystalline lamellae of pure PPOH were deformed in multiple stages, in that the amorphous chains were deformed prior to the crystalline lamellae being deformed. On the other hand, the amorphous and crystalline components of the nanocomposites tended to deform simultaneously, more noticeably for SNS/PPOH. The interaction between the silanol groups on the SNS surface and the hydroxyl groups present in the amorphous chains of PPOH strongly restricts the mobility of the amorphous chains in PPOH. The amorphous chains then tend to deform along with the crystal lamellae attached to the amorphous chains during tensile testing. The synchronous deformation of amorphous chains and crystalline lamellae during tensile testing would result in highly improved yield 95

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strength and elastic modulus of the SNS/PPOH. Moreover, the mobility restriction of the amorphous chains was decreased by modification of the SNS with CTAB, resulting in the relatively low reinforcement of mSNS/PPOH. Thus, the reinforcement mechanism of these different polymer nanocomposites was demonstrated by evaluating newly developed SNS/PPOH and mSNS/ PPOH with rheo-optical NIR spectroscopy. It is important to point out that, the techniques based on NIR spectroscopy may be generally applicable to examine filler-matrix adhesion behavior of a wide range of polymer nanocomposite systems whose mechanisms are not fully understood yet. Therefore, we believe that our findings will contribute to greater reinforcement of new polymer nanocomposites. Appendix A. Supplementary material Supplementary material is available on the publisher’s web site along with the published article. FE-SEM images and FTIR data of SNS and mSNS, and WAXD and DMA measurements of nanocomposites are included. Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.eurpolymj.2017.04.032. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30]

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