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Title

Influences of polypropylene grafted to SiO̲2nanoparticles on the crystallization behavior andmechanical properties of polypropylene/SiO̲2
nanocomposites

Author(s)Umemori, Masaki; Taniike, Toshiaki; Terano,
MinoruCitation

Polymer Bulletin, 68(4): 1093‑1108

Issue Date

2011‑09‑09

Type

Journal Article

Text version

author

URL

http://hdl.handle.net/10119/10717

RightsThis is the author‑created version of Springer,
Masaki Umemori, Toshiaki Taniike, Minoru Terano,
Polymer Bulletin, 68(4), 2011, 1093‑1108. The
original publication is available atwww.springerlink.com,
http://dx.doi.org/10.1007/s00289‑011‑0612‑yDescription

Japan Advanced Institute of Science and Technology

Influences of polypropylene grafted to SiO2 nanoparticles on the
crystallization behavior and mechanical properties ofpolypropylene/SiO2 nanocomposites
Masaki Umemori, Toshiaki Taniike and Minoru Terano*
School of Materials Science, Japan Advanced Institute of Science and Technology, 1-1 Asahidai, Nomi,
Ishikawa, 923-1211, Japan

Fax: (+81)761 51 1625; E-mail: terano@jaist.ac.jp

Abstract
Influences of polypropylene (PP) grafted to SiO2 nanoparticles (7 nm) were studied on the
crystallization behavior and the mechanical properties of PP/SiO2 nanocomposites.

PP for the matrix

and grafting was synthesized in order to have an identical primary structure, aiming at theirco-crystallization and resulting reinforcement of filler-matrix interfaces. The grafted PP chains
improved the dispersion of SiO2, and notably accelerated nucleation in crystallization. It was plausiblethat the grafted chains whose one chain end was pinned to SiO2 became nuclei of the crystallization(co-crystallization between the matrix and grafted chains), thus directly bridging between the matrix and
SiO2 nanoparticles.

The Young’s modulus and tensile strength were most improved by the grafted PP

chains at low filler contents such as 2.3 wt%, whose origin was attributed to effective load transfer to
SiO2 through the co-crystallization-mediated bridging.
1. Introduction
Polypropylene (PP) materials have been widely used for a variety of fields such as
automobile parts, packaging and containers because of the advantages of low price, lightweight, high melting temperature, and excellent processability & mechanical properties.
They

are also environmentallyfriendly

with

low

production energy

and

high

reusability/recyclability. Due to these advantages, the production and application of PP
materials have continuously expanded, in which compounding with other materials has
played a key role.

Among a variety of PP compounds, nanocomposites containing

nano-sized inorganic fillers have attracted a great attention in the last two decades because
they are expected to offer properties that are not realizable with micron-sized fillers [1]Usuki et al. reported the first successful polymer nanocomposite, Nylon 6/clay, with

excellent properties: 87C higher heat distortion temperature and 4 times higher flexual
modulus than those of the pristine polymer were achieved by adding only 5 wt% of nanoclay[2]. With nano fillers, strong interaction is achieved between a matrix polymer and fillers
or between fillers at a small content (usually less than 5 wt%) of fillers. For example,
percolation thresholds for nanocomposites are considerably lower than those for micro-sizedcomposites, which facilitates much more effective reinforcement with a small amount offillers [3]. The glass transition behavior is also modulated by a small amount of nanofillers
as a result of chain confinement [4].

For semi-crystalline polymers with higher-order

structures, the size of nanofiller is smaller than or comparable to the size of spherulites orlamellar morphology, and therefore the addition of nanofiller can greatly change thesestructures. Nitta et al. have reported that SiO2 nanoparticles in PP depress the spherulite
growth rate as the particle size decreases and as the SiO2 content increases, while
micro-sized SiO2 (diameter of 51 m) did not affect the spherulite growth rate [5]. Moreover,spherulites were never formed when the theoretical interparticle distance between SiO2
nanoparticles became shorter than the end-to-end distance of PP chains.

Thus,

nanocomposites exhibit several aspects that are not present in the conventional composites,
being expected to give a new class of high-performance materials.
A huge demand from industry has motivated enormous efforts for the development ofPP-based nanocomposites. However, the results obtained until now is far from satisfactory,compared with those achieved in other polymers such as nylon 6. One of the most serious
problems for PP-based nanocomposites is poor compatibility between PP and inorganic
fillers: polar inorganic fillers easily form large aggregates in apolar PP matrix.

Early

efforts were directed to achieve homogeneous dispersion of fillers, by adding a compatibilizer
and/or by an organic modification of the filler surface, for example, using silane coupling
agents or organic ammonium salts for clay. Usuki et al. reported that the combination of
maleic anhydride modified PP (MAPP) and organic ammonium salts resulted in siginificant
exfoliation of nanoclays in PP [6]. However, the degree of mechanical reinforcement was
not sufficient with respect to potentially expected reinforcement. Garcia et al. studied theeffects of the filler dispersion on the mechanical properties of PP/SiO2 nanocomposites [7].
They found that the improvement of the dispersion of SiO2 hardly affected the tensile
strength.

Bikiaris et al. also showed that the addition of MAPP as a compatibilizer

decreased the size of SiO2 aggregates in PP, but only moderately improved mechanicalproperties [8]. Accordingly, it is concluded that good dispersion of nanofiller is not enoughto realize dramatic reinforcement of PP-based nanocomposites.

In addition, the mostcommonly used MAPP compatibilizer is known to accelerate oxidative degradation of PP [9].
In-situ composite formation is an alternative method to achieve good dispersion offillers without any compatibilizer and organic modification [10]. In this method, matrix
polymer is synthesized in the presence of nanoparticles to automatically facilitate thedispersion of the nanoparticles. Although several authors reported good dispersion of clay,
SiO2 or carbon nanotube in PP [10-12], the mechanical properties of the resulting
nanocomposites were not sufficiently investigated. It is notable that the in-situ method
does not include a strategy to reinforce weak interface between PP and fillers.
One promising method to reinforce the polymer-filler interface is grafting of polymer
chains to fillers, which has advantages of improved compatibility between fillers and matrix
polymer and strengthed interfacial interaction through entanglements between grafted andmatrix polymer chains. In addition, grafted polymer chains may decrease the aggregation
tendency of fillers by their steric hinderace. Physical entanglement between grafted andmatrix chains is quite attractive, since it is almost impossible to get effective chemicalinteraction between PP and fillers. Rong and Friedrich et al. have synthesized PP-based
nanocomposites using SiO2 nanoparticles with a variety of grafted polymer chains such aspolystylene (PS), polybutylacrylate, polyvinylacrylate and so on [13-16]. However, thesegrafted chains did not induce large improvements of the mechanical properties of the formed
nanocomposites. It is supposed that the immiscibility between PP and the grafted chainsmight not allow sufficient entanglement.
Based on these backgrounds, it must be preferable to graft PP chains onto nanofiller inorder to acquire good miscibility with PP matrix. In employing PP itself as grafted chains,
it should be noted that even PP chains can be immicible with each other when they havedifferent primary structures.

Maier et al. and Phillips reported immisivility between

syndiotactic and isotactic PPs [17,18].

Silvestri et al. reported possible immisibility

between isotactic and atactic PPs, depending on Mw of each component [19]. In this sence,stereo- and regio-structures of grafted PP chains should be identical to those of matrix
chains. By fulfilling these requirements, resultant nanocomposites are expected to showmuch higer mechanical properties not only through entanglement but also throughincorporation of grafted PP chains into crystallites of the matrix based on co-crystallization).
Such nanocomposites with same grafted and matrix polymer chains were reported for PS
and polyethylene (PE) [20-22]. In particular, a PE/PE-g-carbon nanotube nanocomposite
showed exccellent improvements of mechanical properties by the grafted PE chains: 1.6
times higher Young’s modulus and 6.9 times higher elongation at breat than those of theungrafted nanocomposite [22]. Differently from polymers such as PS and PE, radical graftpolymerization is not avairable for PP, thus PP-grafted filler is not easy to synthesize.
In this study, terminally hydroxylated isotactic PP (PP-OH) having an identicalstructure with matrix PP was synthesized and grafted onto SiO2 nanoparticles, so as toobtain the first PP-based nanocomposite having the same primary structure for the matrix
and grafted PP chains.

The crystallization behavior and mechanical properties of thenanocomposite were investigated by various techniques such as differential scanningcalorimetry (DSC), polarized optical microscopy (POM), a tensile test and dynamic
mechanical analyses (DMA).
2. Experimental

2.1. Materials
rac-[dimethylsilylenebis(4,5,6,7)-tetrahydro-1-indenyl)]zirconium

dichloride

waspurchased from Wako pure chemical Co. Toluene and tetradecane were dried over themolecular sieve 13X before usage.

Propylene gas of research grade and modified

methylalumoxane (MMAO) were donated by Mitsubishi Chemical Co. and Tosoh FinechemCo., respectively. Aerosil 300 (the diameter of 7 nm and the surface area of 300 m2g1) was
employed as filler.

2.2. Synthesis of PP and PP-OH
PP used as a matrix was synthesized with rac-[dimethylsilylenebis(4,5,6,7)-tetrahydro1-indenyl)]zirconium dichloride and MMAO in toluene (Al/Zr = 5000 and [Zr] = 4 mol/L) at
0C for 1 h under a continuous propylene flow at 1 atm.

PP-OH for grafting wassynthesized by hydroxylation of the chain end of PP [23].

In detail, propylene waspolymerized for 1 h under the above-mentioned conditions, and then oxygen gas was
bubbled in the reactor at 0C for 1 h, followed by the addition of 1 M NaOH and 35% H2O2
aqueous solution. The slurry was stirred for 30 min, to accomplish terminal hydroxylationof PP. The stereoregularity, weight average molecular weight (Mw) and molecular weight
distribution (MWD), and melting point of the obtained PP and PP-OH were measured by
13C-NMR,

high-temperature size exclusion chromatography, and differential scanningcalorimetry (DSC). The synthesized PP had a high isotacticity (91 mol% in mmmm), Mw of
307,000 with relatively narrow MWD about 2.6, and a high melting point (Tm = 160C).
PP-OH was found to have the primary structure and thermal properties completely identicalto those of the matrix PP.

2.3. Synthesis of PP-grafted SiO2
The terminal OH group of PP-OH was reacted with the silanol groups of SiO2 at 200C
in tetradecane for 7 h with stirring to obtain PP-grafted SiO2 (PP-g-SiO2). An adequate
amount of AO-50 (donated by ADEKA Co.) was added in tetradecane to avoid thermaldegradation of PP during the reaction. The made PP-g-SiO2 was washed repeatedly by hot
filtration with o-dichlorobenzene at 145C to completely remove ungrafted PP chains, thenrinsed by methanol, and finally dried in vacuo at 60C for 6 h. Successful grafting was
confirmed by Fourier-transform infrared spectroscopy (FT-IR, FT/IR-6100, JASCO). The
amount of grafted PP was evaluated by thermogravimetric analysis (TG, TG-50, METTLER),
where the sample temperature was kept at 200C for 30 min to remove physisorbed water,
and then increased up to 650C at 20C/min. The grafted amount was estimated from thedifference of the weight loss in 200-650C between PP-g-SiO2 and unmodified SiO2. SiO2
was also modified by octadecyltricholorosilane for the comparison (C18-SiO2).

The C18

modification lowers the surface energy of SiO2 as the PP grafting does, while the C18 group isunable to entangle and co-crystallize with the matrix PP. Thus, C18-SiO2 was utilized to
exploit the relative importance between surface bulk properties (surface energy) and themolecular structure of grafted chains in the design of nanocomposites.

2.4. Preparation of nanocomposites
Nanocomposites were prepared by melt mixing using a two-roll mixer at 20 rpm. PP
pellets were kneaded at 185C for 5 min and then a specified amount of unmodified SiO2,
PP-g-SiO2 or C18-SiO2 was added. The mixture was further kneaded at 185C for additional
10 min. Thus produced nanocomposites with a variety of filler contents (0-10 wt%) were
hot-pressed into sample films with the thickness of 200 m at 230C and 20 MPa, and thenquenched at 100C.

2.4. Instrumentation
The dispersion of SiO2 nanoparticles in the matrix was monitored by a transmission
electron microscope (TEM, Hitachi, H-7100). TEM specimens with the thickness of 100 nm
were prepared by an ultramicrotome (Leica, ULTRACUTS FCS) equipped with a diamondknife (Diatome). The obtained TEM images were analyzed with a image analysis software(Image J software, NIH) to quantify the dispersion.
The crystalline structure and crystallinity of the nanocomposites were obtained by thewide-angle X-ray diffraction (Rigaku, Rint2500).

Wide-angle X-ray diffraction (WAXD)

measurements were performed in a reflection mode at room temperature with Cu K
radiation operating at 40 kV and 30 mA. The scanning rate was 1/min over 2 of 10-30.
Sample films with a thickness of ca. 200 m were used.
Isothermal crystallization was conducted at 128C with DSC (METTELER, DSC-822).
Samples were kept at 200C for 5 min to erase a thermal history, and then cooled down to128C at a rate of 20 C/min. A crystallization rate was determined as an inverse of thehalf time of the crystallization (denoted as t1/2).

A nuclear density during isothermal

crystallization at 130C was evaluated by POM (Leica, DMLB HC) equipped with anautomated hot stage (METTELER, FP82HT).
Tensile tests were carried out at room temperature using a dumbbell-shaped specimen
at a crosshead speed of 1.0 mm/min by a tensile tester (DAT-100, Abecks Inc.).

Fivemeasurements were performed for each sample. DMA (UBM, E4000) was conducted for thetemperature dependence of oscillatory tensile moduli in the solid state (E’, tan).
Measurements were performed in the temperature range from 80 to 170C with a heating
rate of 3C/min and the frequency of 10 Hz. The size of a sample specimen was 2 mm x 8mm x 0.2 mm. The frequency dependence of oscillatory shear moduli in the molten state(G’, G”) was measured by a parallel-and-plate rheometer (TA, AR2000ex) at 180C with afrequency range from 100 to 0.01 rad/sec under N2 atmosphere.

The diameter of theparallel plates was 25 mm. Each measurement was performed within a linear viscoelastic
region.
3. Results and discussion

3.1 Preparation of PP-g-SiO2
The FT-IR spectrum of synthesized PP-g-SiO2 was measured after careful removal ofungrafted chains with hot o-dichlorobenzene (Figure 1). The peaks at 2920 and 2950 cm1
(corresponding to the C-H stretching vibrations for the -CH2- and -CH3 groups) indicate thesuccessful grafting. TG analyses were performed to evaluate the amount of PP chainsgrafted on SiO2 (Figure 2). PP-g-SiO2 exhibited 12.4 wt% of the weight loss between 200
and 650C.

Referencing the weight loss of 11.1 wt% for unmodified SiO2 in the same

temperature range, the amount of the grafted PP chains was estimated as 1.3 wt%.
Assuming a poreless and perfect sphere for the SiO2 particles, the number of the graftedchains was calculated as 1 chain/particle. The amount of the octadecyl group grafted on
SiO2 was similarly obtained to be 7.5 wt%.

3.2. Dispersion State of fillers
The dispersion of unmodified SiO2, PP-g-SiO2 and C18-SiO2 in the matrix was studied
by TEM (Figure 3). All the fillers formed small aggregates of 20-100 nm in the matrix. Up
to the filler content of 4.0 wt%, the dispersion of PP-g-SiO2 and C18-SiO2 was better thanthat of unmodified SiO2. At the filler content of 10 wt%, the three kinds of fillers similarly
formed large and extended aggregates. The better dispersion of PP-g-SiO2 and C18-SiO2
was quantitatively confirmed by an image analysis. Table 1 shows the average areas of theaggregates, which are 50% smaller for PP-g-SiO2 and 30% for C18-SiO2.

The average

diameters are about 30% smaller for both PP-g-SiO2 and C18-SiO2. This result proved theeffectiveness of the grafting of PP chains in improving the dispersion of fillers in the PP
matrix, even though the grafted amount was as small as 1 chain/particle. The observed
dispersion for the three kinds of fillers might be correlated with their cohesion abilities.
The modification by short alkyl groups is known to largely reduce the surface energy of SiO2[24], leading to better dispersion. On the other hand, the largely modified dispersion for
PP-g-SiO2 can be explained not only by reduction of the surface energy through the grafting,but also by steric hindrance from the grafted PP chains [13]. The way of the aggregation
clearly depended on the kinds of the fillers. The aggregates of unmodified SiO2 were larger
and had an open or extended shape, while those of PP-g-SiO2 and C18-SiO2 were muchsmaller and more compact. Bartholome et al. reported similar effects of grafted PS chainson the dispersion of SiO2 in PS [21].

3.3. Crystalline structure and crystallization behavior
The WAXD patterns of pristine PP, PP/SiO2, and PP/PP-g-SiO2 (Figure 4) exhibited typicalcharacteristics of the -form crystal with the peaks of 2 = 14.07, 16.85, 18.51, 21.08 and21.77 corresponding to (110), (040), (130), (111) and (131) reflections, respectively [25].
The other crystalline forms were never observed.

Thus, SiO2 nanoparticles and theirmodification had few influences on the crystalline structure of the PP matrix. The WAXD
crystallinities (c) summarized in Table 2 were quite similar for all the samples, whichallowed us to fairly compare the mechanical properties of the different samples.
Isothermal crystallization of the nanocomposites was tracked by DSC to examine
influences of the fillers on the crystallization behavior of the matrix (Figure 5). Surfaceorganic modification of fillers sometimes enhances a nucleating ability of fillers. It was
found that the crystallization of PP was considerably accelerated by the addition of
PP-g-SiO2: the crystallization rate was 0.14 min1 for PP/SiO2 at the SiO2 content of 2.3 wt%,
while that of PP/PP-g-SiO2 was 0.30 min1 at the same content. The C18 modification of
SiO2 did not give such that significant enhancement of the nucleating ability. Thus, thehigh nucleating ability was originated by the molecular structure of grafted chains rather
than the surface bulk nature such as surface energy.
The origin of the high nucleating ability of PP-g-SiO2 was examined by POM, where
the time-course increase of the nuclear densities was plotted during isothermal
crystallization at 130ºC (Figure 6). The nuclear density became more than twice by theaddition of PP-g-SiO2, indicating that the increase of the crystallization rate by PP-g-SiO2
was owing to the acceleration of the nucleation rather than the acceleration of the spherulite

growth. On the other hand, unmodified SiO2 hardly altered the nuclear density of pristine
PP. Based on the result that the grafted PP chains having the same primary structure withthe matrix chains significantly enhanced the nuclear density, it was plausible that the
crystallization of the matrix was initiated by the grafted chains with one chain end pinned
to SiO2.

3.4. Mechanical properties
Results of Tensile tests are summarized in Figure 7 and Table 2. Unmodified SiO2
hardly affected the Young’s modulus and tensile strength at the lowest filler content, while
increased these values above 4.0 wt%. The onset of the reinforcement at 4.0 wt% coincidedwith the formation of the percolation (described later). PP-g-SiO2 showed the most notablereinforcement at the lowest content compared with unmodified SiO2. The degree of thereinforcement became closer to that for unmodified SiO2 at higher contents, but always
larger.

Only the exception was the tensile strength at 10 wt%, where the break inelongation occurred in a very early stage for PP-g-SiO2.

The C18-modification of SiO2

affected negatively, even though C18-SiO2 showed the better dispersion than unmodified
SiO2.
Rheological properties were evaluated to investigate interaction between the matrix
and fillers. The temperature dependences of E’ and tan are shown in Figures 8 and 9,
respectively. The glass transition temperatures (Tg) determined from E” increased about5C by modifications irrespective of the content and kind of the fillers (not shown). The E’values below Tg were almost the same among all the samples. The E’ values above Tg wereincreased by the addition of both unmodified SiO2 and PP-g-SiO2 (Figure 8). However, theways of the increases were different between the two kinds of fillers: increase of the E’ value
from pristine PP was about 2 times higher for PP-g-SiO2 compared with unmodified SiO2,
while the advantage of PP-g-SiO2 over PP/SiO2 became much smaller at higher contents.
These trends for E’ are in line with the results of tensile tests.

The peak intensity and areaof tan at the glass transition generally correlate with the portion of the amorphous phase
and/or the mobility of polymer chains therein [4]. Unmodified SiO2 and PP-g-SiO2 similarly
reduced the peak intensity of tan (compared with pristine PP) (Figure 9). As shown in thec values of Table 2, the amorphous portion was similar among all the samples. Hence, thepeak reduction by the addition of the fillers must be ascribed by reduced mobility of theamorphous chains, which is known as a confinement effect by nanoparticles [4, 26, 27]. The
frequency dependences of G’ and complex viscosity * were shown in Figure 10. A large
difference was observed between pristine PP and the nanocomposites in the low frequency
region (< ~0.1 rad/sec). Both G’ and * significantly increased by the addition of the fillers.

At the filler content of 2.3 wt%, the two nanocomposites had similar G’ values with terminal
flow.

However, the increase of G’ was much more for unmodified SiO2 at the higher

contents. A plateau appeared instead of the terminal flow at the content over 4 wt% forunmodified SiO2, and at the content of 10 wt% for PP-g-SiO2.

The appearance of theplateau in G’ accompanied drastic increase of * by an order of magnitude.

Thisphenomenon at higher filler contents is general for particulate nanocomposites, andindicates the formation of percolation network of fillers. The percolation has been mostprecisely studied for polymer/clay nanocomposites, and percolation networks through directconnections among platelets such as “edge-to-edge” or “edge-to-face” have been reported[28-30]. On the other hand, it has been considered for polymer/SiO2 nanocmomposites thatfractal structures formed by open aggregates of SiO2 nanoparticles are responsible for the
percolation [31]. As shown in Figure 3, unmodified SiO2 formed rather open aggregates in
the matrix, while PP-g-SiO2 and C18-SiO2 existed as smaller and compact aggregates. The
lower percolation threshold for PP/SiO2 is clearly explained by this fact, where thepercolation thresholds were between 2.3-4.0 wt% for unmodified SiO2 and 4.0-10 wt% forPP-g-SiO2.

3.5. Discussion
The obtained effects of the fillers are briefly summarized as follows. Modified SiO2
(PP-g-SiO2 and C18-SiO2) apparently improved the filler dispersion except at the highest
filler content (Figure 3). PP-g-SiO2 and C18-SiO2 made smaller but compact aggregates,
while unmodified SiO2 made network-like aggregates connected with each other.
PP-g-SiO2 showed the nucleation ability much higher than C18-SiO2 and SiO2 (Figure 5). It
was believed that the grafted PP chains initiate the crystallization of the matrix due to theirlower mobility and to the same primary structure with the matrix. PP-g-SiO2 showed thehighest reinforcement in the tensile test at the lowest filler content (Figure 7). On theother hand, the degree of the reinforcement was similar to that of unmodified SiO2 over 4.0
wt%. The reinforcement by C18-SiO2 was much poorer, in spite of the better dispersion of
C18-SiO2 than unmodified SiO2. SiO2 and PP-g-SiO2 similarly reduced the intensity of tanat the glass transition (Figure 9), indicating that the confinement of the amorphous chainsby SiO2 nanoparticles was not affected by the grafted chains.

G’ of PP/SiO2 at low

frequencies was larger than that of PP/PP-g-SiO2 in the molten state (Figure 10), on thecontrary to the higher Young’s modulus and E’ of PP-g-SiO2 in the solid state.

The

disappearance of the terminal flow (corresponding to the formation of percolation network)
occurred at a lower filler content for unmodified SiO2 that for PP-g-SiO2. Hereafter, themechanistic origin of the higher reinforcement by PP-g-SiO2 in the tensile tests (Figure 7) isdiscussed, for which C18-SiO2 offers a useful reference.
Although the reinforcement mechanism in polymer-based nanocomposites has notbeen fully understood yet, the numerous efforts have identified plausible mechanisms into i)strong interfacial interaction between matrix chains and filler particles, and ii) percolationor network structures of filler particles.

Load transfer to fillers based on these two

mechanisms are also affected by the content, stiffness, aspect ratio, surface chemical natureand dispersion state of fillers.
The percolation threshold for PP/SiO2 lay within the filler content 2.3-4.0 wt%. The
results of tensile tests for unmodified SiO2 were well correlated with the dependence of G’ on
the filler content: little reinforcement at 2.3 wt% before the percolation and sharp increasesof the Young’s modulus and tensile strength from 4.0 to 10 wt% after the percolation. Thus,
the reinforcement for unmodified SiO2 is believed to be mainly based on the percolation.
On the other hand, PP-g-SiO2 showed the most prominent reinforcement in the tensile test(compared with unmodified SiO2) at the lowest filler content, i.e. before the percolation.
This suggests that some strong interaction between the matrix and filler particles was
responsible for the observed reinforcement. Better dispersion of fillers has been believed toresult in higher reinforcement because of more homogeneously distributed load transfer toparticles. In this viewpoint, PP-g-SiO2 was considered to be advantageous at lower filler
contents. However, as was described in introduction, good dispersion of fillers is not a
sufficient condition for the reinforcement. In fact, C18-SiO2 with better dispersion thanunmodified SiO2 never showed notable improvements. These results claim the importanceof the quality of the interaction between the matrix and filler particles. As in the results of
DSC and POM (Figures 5 and 6), the largest difference between PP-g-SiO2 and C18-SiO2
appeared in the nucleating ability of the fillers: the key factor for the nucleation was themolecular structure of grafted chains rather than the surface bulk natures.

The

considerably high nucleation ability of PP-g-SiO2 suggested the crystallization of the matrix
initiated from the grafted chains having the same primary structure with the matrix chains.
In this way, we propose that the grafted chains are preferentially co-crystallized together
with the matrix PP, and play a role of direct bridges between the matrix and filler particles
for effective load transfer. This mechanism is supported by the poorer reinforcement in G’
for PP/PP-g-SiO2 than that of unmodified SiO2 at the lower contents, where all crystallites
were molten.

The C18 modification was not effective for both the percolation andco-crystallization, leading to the poorest reinforcement.
A schematic image of PP/PP-g-SiO2 for both the solid and molten states is drawn in
Figure 11.

Since the PP chains covalently grafted to SiO2 particles preferentiallyco-crystallized with the matrix PP, the effective load transfer from the matrix to the fillers

occurred in the solid state, resulting in the better mechanical properties even at the lower
contents much before the percolation. On the other hand, the grafted PP chains gave better
filler dispersion by steric hindrance without a large increase of complex viscosity in themolten state.
4. Conclusion
In this study, a PP/SiO2 nanocomposite having the same primary structure for thematrix and grafted PP chains was synthesized.

PP-g-SiO2 achieved much higher

reinforcement than unmodified SiO2, at a low filler content (2.3 wt%). We have proposed areinforcement mechanism that the grafted PP chains trigger the crystallization of thematrix and are imparted in the lamellae of the matrix, thus directly bridging between thematrix and filler particles for effective load transfer. On the contrary, the increase of theviscosity in melt by the addition of SiO2 was partially suppressed with the grafted PP chains.
In this way, to graft PP chains having an identical structure to matrix chains can be a
successful strategy to achieve significant reinforcement of PP-based nanocomposites,
without sacrificing processability and light weight.

Further improvements in themechanical properties will be achieved by the optimization of the amount and structure of
the grafted chains.
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Fig. 1

Fig. 2

IR spectrum of PP-g-SiO2

TG analysis of PP-g-SiO2 and unmodified SiO2. The amount of grafted PP chains on SiO2 wasevaluated based on the weight loss between 200-650C.

Fig. 3wt%.

TEM images of prepared PP-based nanocomposites. Filler contents were varied from 2.3 to 10

Fig. 4

WAXD patterns of pristine PP and nanocomposites. Regardless of the filler contents, all thesamples showed characteristic peaks of -form and other forms were never observed.

Fig. 5

Fig. 6

DSC curves for isothermal crystallization at 128C

POM images in isothermal crystallization at 130C (t = 7 min) and the time dependence of thenuclear density calculated by the number of nuclei in the pictures at each timeFig. 7

Stress-strain curves (cross head speed of 1 mm/min) at the filler content of a) 2.3 wt% and b) 10

wt%.

c) The Yang’s modulus and d) the tensile strength were plotted against the filler content.

Fig. 8Temperature dependence of tensile storage modulus E’ for pristine PP and nanocomposites.

Measurements were performed in the temperature range from 80 to 170C with a heating rate of3C/min and the frequency of 10 Hz.

Fig. 9Temperature dependence of loss tangent for pristine PP and nanocomposites. Measurements

were performed in the temperature range from 80 to 170C with a heating rate of 3C/min and thefrequency of 10 Hz.

Fig. 10Frequency dependence of shear storage modulus G’ (upper) and (bottom) complex viscosity *for pristine PP and nanocomposites. Measurements were performed at 180C with a frequency range
from 100 to 0.01 rad/sec under N2 atmosphere.

Fig. 11

Schematic image of PP/PP-g-SiO2 in solid and molten states

Table 1

Size of SiO2 aggregatesa

Area / 104 nm2

Feret diameter / nm

PP/SiO2 (2.3%)

4.60

338

PP/SiO2 (4.0%)

5.03

340

PP/PP-g-SiO2 (2.3%)

2.15

218

PP/PP-g-SiO2 (4.0%)

2.62

230

PP/C18-SiO2 (2.3%)

2.82

250

PP/C18-SiO2 (4.0%)

3.52

250

Samplea Approximately 400 aggregates of SiO2 were analyzed and averaged by a software.

Table 2Mechanical properties and crystallinities

Young’s modulusa

Tensile strengtha

 cb

/ MPa

/ MPa

/ wt%

PP

347

23.4

55.1

PP/SiO2 (2.3%)

347

23.7

55.6

PP/SiO2 (4.0%)

364

25.3

52.9

PP/SiO2 (10%)

416

27.0

51.4

PP/C18-SiO2 (2.3%)

325

24.1

52.5

PP/C18-SiO2 (4.0%)

332

24.0

51.8

PP/C18-SiO2 (10%)

362

25.3

51.5

PP/PP-g-SiO2 (2.3%)

369

25.0

54.3

PP/PP-g-SiO2 (4.0%)

374

25.6

52.2

PP/PP-g-SiO2 (10%)

439

26.7

52.2

Samplea measured by tensile test, b determined by WAXD