xrd fluor.pdf

Nanoindentation, Nanoscratch, and Nanotensile Testing ofPoly(vinylidene fluoride)-Polyhedral Oligomeric SilsesquioxaneNanocomposites

Fanlin Zeng,1 Yizhi Liu,1 Yi Sun,1 Enlai Hu,1 Yu Zhou2

1Department of Astronautic Science and Mechanics, Harbin Institute of Technology, Harbin, Peoples Republic of China

2Material Science and Engineering, Harbin Institute of Technology, Harbin, Peoples Republic of China

Correspondence to: F. Zeng (E-mail: [email protected])

Received 19 June 2012; revised 9 August 2012; accepted 13 August 2012; published online 14 September 2012

DOI: 10.1002/polb.23159

ABSTRACT: Nanocomposites composed of a poly(vinylidene flu-

oride) (PVDF) matrix and 0, 3, 5, and 8 wt % fluoropropyl polyhe-

dral oligomeric silsesquioxane (FP-POSS) were prepared by

using the solvent evaporation method. The morphology and the

crystalline phase of the nanocomposites were investigated by

digital microscopy, scanning probe microscopy, X-ray diffrac-

tometer, and Fourier transform infrared spectroscopy. FP-POSS

acted as nucleating agent in PVDF matrix. A small content of FP-

POSS resulted in an incomplete nucleation of PVDF and gener-

ated bigger spherical particles, whereas higher contents led to a

complete nucleation and formed more separate and less-cross-

linked particles. Nanoindentation, nanoscratch, and nanotensile

tests were carried out to study the influence of different contents

of FP-POSS on the key static and dynamic mechanical proper-

ties of different systems. The nanocomposite with 3 wt % FP-

POSS was found to possess enhanced elastic properties and

hardness. However, with the increase of the FP-POSS content,

the elastic modulus and hardness were found to decrease, and

the improvement on stiffness was negative at contents of 5 and

8 wt %. Compared with neat PVDF, the scratch resistance of the

PVDF/FP-POSS nanocomposites was decreased due to a

rougher surface derived from the bigger spherulites. Nanoten-

sile testing results showed both the stiffness and toughness of

PVDF-FP3% were enhanced and further additions of FP-POSS

brought dramatic enhancements in toughness while associated

with a decline in stiffness. Dynamical mechanical properties

indicated the viscosity of the nanocomposites increased with

the increasing FP-POSS contents. VC 2012 Wiley Periodicals, Inc.

J Polym Sci Part B: Polym Phys 50: 15971611, 2012

KEYWORDS: mechanical properties; morphology; nano-compo-

sites; polyhedral oligomeric silsesquioxanes; poly(vinylidene

difluoride)

INTRODUCTION As a semicrystalline polymer, poly(vinyli-dene difluoride) (PVDF) is widely used in smart structuresensors, actuators, and transducers because it exhibits excel-lent piezoelectric and pyroelectric properties which werederived from the two polar phases b and c of PVDF.13

Because PVDF is not as hard as some other piezoelectricmaterials, such as lead zirconate titanate (PZT), PbTiO3, andso on, it can easily be prepared to the soft and thin trans-ducers which are dictated by the shape of the structuresthey coat. This property is very useful in fields of energyharvesting4,5 and adaptive inflatable structures in aerospaceengineering.6,7 But unfortunately, neat PVDF cannot com-pletely meet the mechanical, thermal, and oxidation resist-ance property requirements of some harsh environments.79

Many efforts have been taken to improve the properties ofPVDF. For example, incorporation of organic polymer or inor-ganic fillers into the PVDF matrix to produce composites hasbeen extensively studied with the objective of furtherimproving its properties, such as mechanical performance,

thermal stability, and dielectric properties.1012 Specifically,polyaniline (PANI),12 PZT,13 polyamide,14 LiClO4, and TiO2,

15

thermoplastic polyurethane (TPU),16 poly(methylmethacry-late) (PMMA),17 and so on have been used to improve themechanical properties of PVDF. It appears from current ex-perimental investigations that improvements in mechanicalproperties of PVDF composites depend on a set of factors interms of filler size, shape, mass ratio, degree of dispersion,and so on. However, it is difficult to find plausible correla-tions between specific factors, for example, the filler mass ra-tio, and mechanical properties from the reported results. Itseems that the improvement effects depend not only on thefiller mass ratio but also on the filler type.

Polyhedral oligomeric silsesquioxanes (POSS) are a uniqueclass of organicinorganic hybrid materials that can bedepicted by the formula (RSiO1.5)n (where n is an even num-ber and R H, Cl, or a variety of organic groups). Becauseof the robust inorganic cage-like core structure with theSiAO atoms, POSS exhibits many superior thermomechanical

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properties in terms of wearability, thermal stability, oxidationresistance, and high strength1820 among others. FluorinatedPOSS are a new kind of POSS compounds with high molecu-lar weight and density.21 Theoretically, PVDF/POSS nanocom-posites with highly improved mechanical properties areexpected to be produced if fluorinated POSS is incorporatedinto PVDF, assuming they are miscible. For a polymer matrix,it is widely accepted that inorganic fillers tend to improve instiffness while organic ones in toughness. But as a kind oforganicinorganic hybrid compounds, how will different con-tents of fluorinated POSS influence the key mechanical prop-erties of PVDF is hard to evaluate and needs to beresearched by experiments.

As a standard mechanical test method, nanoindentation hasbeen developed and widely used to characterize some im-portant mechanical properties (hardness, elastic modulus,and so on) of nanostructural materials and thin films orcoatings.2225 Nanoindentation method measures the hard-ness and elastic modulus (Youngs modulus) of a materialfrom indentation load-displacement data obtained duringone cycle of loading and unloading.23 Besides the standardmethod for hardness and elastic modulus, other nanoinden-tation methods have also been devised for evaluating someother properties. For example, although not well developed,the yield stress and strain-hardening characteristic of met-als;26,27 parameters characteristic of damping and internalfriction in polymers, such as the storage and loss modu-lus;28,29 and the activation energy and stress exponent forcreep30,31 have been determined by these methods. A nano-scratch test with the nanoindentation instrument is widelyused to characterize the films and coatings for scratch resist-ance.25,32 The scratch resistance is measured from the depthat a given load or the load at which material fails (the criti-cal load for a thin film adhering on a substrate). There arethree scan steps in a typical scratch procedure: (1) prescan,the indenter tip only approaches and scans from the scratcharea to achieve the original topography of the surface; (2)scratch scan, the tip scratches on the sample with rampingloads; and (3) postscan, the tip unloads from the scratch endto obtain the residual deformation. From the original topog-raphy of the surface, morphology compensation can be ful-filled at the latter two steps and the effective penetrationdepth can be obtained. A tensile test has been widely usedto determine some important mechanical properties such asYoungs modulus, fracture toughness, and yield stress ofmaterials. However, a traditional tensile test is not capable ofa sample with nanoscale size, where a very small tensileload and very precise measure equipments are needed. Inrecent years, the nanotensile test has been developed andused to characterize the tensile behaviors of very small sam-ples such as the nanotubes, silk, fibers, and so on.3335 Thenanotensile test is also used to determine some dynamic me-chanical properties by adding a very small harmonic forcewith specific frequency on the sample.36

Specifically, the crystallization behaviors, morphology, andphysical properties of PVDF or its nanocomposites have beeninvestigated by many researchers. Chang et al.,37 for exam-

ple, conducted an experimental study to address the thermo-mechanical and optical characteristics of PVDF for flexibleelectronic applications. They found that thermomechanicalcharacteristics, thermal elongation, and expansion weregreatly influenced at stretching ratios of over four in thestretching direction, and the optical properties were greatlyinfluenced in stretched films. Chae et al.38 mixed differentmass ratios of silver nanoparticles into PVDF and testedtheir physical properties, the crystallization behavior undershear, and the consequential crystalline morphology. Theyclaimed that the overall crystallization process of PVDFunder shear was accelerated by reducing both inductiontime and crystallization time, when Ag loading level wasincreased. Guney,39 who researched the influence of temper-ature on the elastic properties of PVDF, found that the relax-ation behavior of PVDF was affected from the form of me-chanical disturbance. Ma et al.40 mixed different smallamount of nanoparticles, such as montmorillonite (MMT),SiO2, CaCO3, or polytetrafluoroethylene (PTFE), into PVDFand researched their crystallization and melting behaviors.They found that addition of small amount of these four typesof nanoparticles would not affect the original crystallinephase obtained in the neat PVDF sample, but accelerated thecrystallization rate because of the nucleation effect. In thesefour blend systems, MMT or PTFE nanoparticles could beapplied well for PVDF nanocomposite preparation because ofstronger interactions between particle surface and PVDFmolecules. The nucleation enhancement and the growth rateof the spherulites were decreased in the order SiO2 >CaCO3 > PTFE > MMT. Fang et al.41 conducted an uniaxialtension experiment to study the deformation and fracturebehavior of poly(vinylidene fluoride-trifluorethylene) ferro-electric copolymer films. They found that the polymer filmsamples prepared by stretching the solution-cast films andthen annealing fracture at a much larger maximum strainand a higher tensile strength than those prepared by solu-tion casting and then annealing. Salimi et al.42 researchedthe conformational changes and phase transformation mech-anisms in PVDF solution-cast films. They found that the low-temperature crystallization of PVDF in dimethylacetamide so-lution, mainly resulted in the formation of trans states (band c phases), whereas at higher temperatures, gauche statesbecome more populated (a phase). Moreover, the uniaxialstretching greatly enhanced piezoelectric properties of thefilms, due to formation of oriented b phase crystals, whichwere of more uniform distribution of dipole moments. Veryrecently, Martins et al.43 conducted a study about the mor-phological, viscoelastic, and thermal properties of PVDF/POSS nanocomposites. They found that the crystallinity ofPVDF was little influenced at low POSS contents, and themethacryl POSS were acting as lubricant in the nanocompo-site system. It seems from current investigations that thecrystallization behaviors, morphological, and physical proper-ties of PVDF and its blends will be influenced by many dif-ferent factors. In addition, these properties of PVDF andPOSS nanocomposites, however, were seldom researched.

In previous work, we have researched the miscibility in mix-tures of PVDF and several kinds of POSS compounds

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including four kinds of fluorinated POSS compounds usingmolecular simulations.44 The simulation results show fourkinds of fluorinated POSS compounds and PVDF are fullymiscible at any temperature and the miscibility is derivedfrom the polar CAF bonds and the electrostatic interactionsin the POSS and PVDF molecules. We have also investigatedthe elastic properties of mixtures of PVDF and the (3,3,3-tri-fluoropropyl)8Si8O12 (FP-POSS)

21 at different temperaturesusing molecular dynamics (MD) simulations.45 We found thatthe glass transition temperature of PVDF was significantlyimproved with FP-POSS, and the moduli of PVDF wereimproved and the improvement effect, in general, nearlydecreased with the increase of the mass ratio of FP-POSS.However, experiments are indispensable to this work. On theone hand, simulation results need to be verified from experi-ments, and on the other hand, some other key mechanicalproperties, such as the scratch resistance, fracture toughness,and some dynamic mechanical properties can also beobtained. In this work, the PVDF/FP-POSS nanocompositeswere first prepared by the solvent evaporation method. Andthen, the microscopy characterization, nanoindentation,nanoscratch, and nanotensile testing were carried out. Thefocus of this work is on the effects of the FP-POSS contentson the morphology, the key mechanical properties of PVDF/FP-POSS nanocomposites, such as the hardness, elastic mod-ulus, scratch resistance, tensile strength, fracture toughness,loss tangent, and so on, and the aim is to determine howthese properties would be influenced by varying the amountof FP-POSS. Furthermore, the deformation mechanism of thenanocomposites was analyzed, and it would be helpful tounderstand why different materials exhibited different me-chanical properties from microscale.

EXPERIMENTAL

Materials and Sample PreparationCommercial PVDF (Mw 534,000; 99.99% purity; Aldrich),(3,3,3-trifluoropropyl)8Si8O12 (FP-POSS; Hybrid Plastics cod.FL0578), and N,N-dimethylformamide (DMF; 99%; Aldrich)were used to prepare the PVDF/FP-POSS mixtures. PVDF

and FP-POSS powder were first dissolved in DMF in abeaker; next, the solution was kept at room temperatureunder mechanical stirring until PVDF and FP-POSS were dis-solved, and air was removed fully. The solution containingPVDF and FP-POSS was homogeneously transparent. Then,the total evaporation was carried out at 70 C to remove thesolvent, and composite films with thicknesses of around 70lm were obtained. The whole process of the preparationwas carried out in a vacuum. Samples for nanoindentation,nanoscratch, and nanotensile testing were cut from thesefilms. The size of the samples is about 5 5 mm2 for nano-indentation and nanoscratch, and about 100 lm 10mm for nanotensile testing (shown as Fig. 1, but not thesame size in each test). In order to be comparable to theprevious simulation work,45 the mass ratios of FP-POSSadded to the PVDF matrix were 0, 3, 5, and 8%, and nano-composites with different FP-POSS contents are denoted asPVDF-FPi% (i 0, 3, 5, and 8).Morphology and Structure CharacterizationOptical microscopy was performed on the nanocompositefilms using a VHX-600E digital microscope (Keyence Com-pany) with a real-time depth composition, two/three-dimen-sional functions, and 20 to 5000 zoom. Scanning probemicroscopy (SPM) was performed using a SPM9500J2 micro-scope (Shimadzu Company) equipped with software (version2.30) for images processing and profile analysis. During thescanning, the contact scanning mode was used and both thedeflection and height traces were obtained.

X-Ray diffractometer (XRD) analysis was applied using aRigaku D/max-rB rotating anode XRD with CuKa radiation, k 0.15418 nm, at a generator voltage of 60 kV and a currentof 200 mA. The data were collected from 10 to 80 inter-vals. Fourier transform infrared (FTIR) analysis was carriedout using a FTIR spectrometer (Avatar360, Nicolet), with ascan range 4000500 cm1 and a resolution of 4 cm1.

Nanoindentation and Nanoscratch TestingThe hardness and elastic modulus of the nanocompositefilms were determined by using the OliverPharr

FIGURE 1 The nanotensile testing system UTM T150 (a) and the sample image measured by optical microscopy (b).



method23,4648 in the nanoindentation testing with a G200Nano Indenter (Agilent Technologies). A diamond triangularpyramid Berkovich indenter (TB20114 ISO) with a tip radiusof about 20 nm, face angle (the centerline-to-face angle) of65.3 , Youngs modulus of 1141 GPa, and Poissons ratio of0.07 was used. From the simulation results,45 the Poissonsratio of the nanocomposite films was estimated as 0.3. Itmay seem counterintuitive that one must know the veryaccurate Poissons ratio of the material to compute its modu-lus, even a rough estimate. However, in fact, a very precisevalue of this parameter is not necessary, because when it isgiven as 0.25 6 0.1, produces only about a 5% uncertaintyin the calculated value of the elastic modulus for most mate-rials.46 Two kinds of specific test methods: (1) G-Series Ba-sic Hardness, Modulus at a Depth; and (2) G-Series Hard-ness and Modulus via Cycles Load Control were used todetermine the hardness and elastic modulus at differentmaximum depths and loads. In the former method, testswith different maximum depths of 500, 1000, 1500, and2000 nm were carried out on each sample. During the test,the loading speed was controlled as the same strain rate of0.05 s1, which means the indenter was driven into the sam-ple at a same speed; thus, the loads would increase veryslowly at the initial stage and very fast at the end. In the lat-ter method, the maximum load of 10 mN and eight loadingcycles were carried out at each test. Thus, eight relevanttests were carried out at the same position, and the maxi-mum loads at each test increased as an exponential function(0.075, 0.15, 0.31, 0.62, 1.25, 2.5, 5.0, and 10.0 mN for eachtest, respectively). In this method, the size effect can beclearly displayed; from which one can judge what maximumloads are appropriate to obtain the reliable results. Duringthe test in this method, the loading speed was kept as a con-stant force increasing per second (loading time of 20 sadopted here), thus, the indenter was driven into the sampleat different speeds, fast at the initial stage and slowly at theend. Both in the two methods, each kind of test wasrepeated at six different positions with intervals of 50 lm(defined as a 3 2 array). The peak-hold time, that is, thetime the indenter held at the maximum load (load control)or depth (depth control) was set as 10 s. The thermal driftwas considered after the tip detected the surface with a cus-tom stiffness criteria of 125N/m at each test. The allowabledrift rate was set as 0.5 nm/s, which means the test will notbegin until the drift of the indenter influenced by the ther-mal fluctuation is less than 0.5 nm/s. At the end stage of theunloading (90%), the indenter would be held for about 1minute again to compute and correct the final drift. All testswere carried out in a clean-air environment with a relativehumidity of 30%, while the temperature was kept constantat 20 6 0.5 C. Hardness and elastic modulus were meas-ured from indentation load-displacement data obtained dur-ing one cycle of loading and unloading. From the maximumload and the corresponding projected contact area hardnesscan be determined and by measuring the elastic contact stiff-ness (measured from the upper portion of the unloadingdata) and the projected contact area the elastic modulus canthus be derived. In the real testing work, the projected con-

tact area A has been calibrated as A 24.0808h2 382h (his the contact depth). In addition, the two main factors, theadhesion between the tip and the specimen surface and thepile-up (or sink-in) of the specimen surface under the tip,will bring spurious results in the nanoindentation testing.For the adhesion, its effect on the nanoindentation results isonly obvious at low loads or low indentation depths. Thus,the overestimation in the measured indentation moduli/hardness can be corrected by performing indentations withloads high enough so that the modulus/hardness is inde-pendent of the applied load. When pile-up (or sink-in)occurs, the contact area is greater (or less) than that pre-dicted by the OliverPharr method, and both the hardnessand the modulus are overestimated.23 Fortunately, Pharret al.23 found that when hf/hmax < 0.7 (hf, the final depth,the permanent depth of penetration after the indenter isunloaded fully; hmax, the maximum displacement), very littlepile-up (or sink-in) is observed no matter what the work-hardening behavior of the material. Thus, the contact areasare independent of the work-hardening characteristics in thiscase. Therefore, in an indentation experiment, care must betaken when hf/hmax > 0.7, as using the OliverPharr methodcan lead to large errors in the contact area.

Scratch testing was carried out using a Nano Indenter G200instrument with the method of G-Series Ramp Load Scratchwith Topography Compensation. The test procedure wassimilar to that presented elsewhere,25,49 but the topographycompensation was considered here because the samples sur-face was not very smooth. Before the real scratching, theoriginal topography and surface roughness of the scratchingarea were detected by prescanning the surface with the in-denter under a profiling load 20 lN. In the real scratching,the indenter tip scratched on the sample with linear rampingloads from 0 to maximum value of 100 mN. The scratchlength and speed were set as 250 lm and 10 lm/s, respec-tively. Depths of scratches with increasing scratch load weremeasured by profiling the surface during and after thescratch test, resulting in a total length of about 350 lm foreach test (including 50-lm prescanning and postscanning atthe two ends of the real scratch path), whereas the realeffective scratch length was 250 lm as applied to all speci-mens. The effective penetration depth during (scratch depth)and after (residual depth) the scratch test at differentscratch distance then could be obtained. Three independentscratch tests were carried out for each sample, and the dis-tance interval of different scratches was 100 lm. The testenvironment and the allowable thermal drift rate were simi-lar to those in the nanoindentation test.

Nanotensile TestingA commercial nanotensile testing system (Nano UTMTM Uni-versal Testing System T150, Agilent Technologies) with themethod of UTM-Bionix Standard Toecomp CDA was used toconduct the tensile test. The instrument consists of a loadingframe, moving crosshead, two grips, and a nanomechanicalactuating transducer (NMAT) as shown in Figure 1(a). Thenominal maximum load of this system is 500 mN (can reachto 750 mN in real test), and the maximum crosshead


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extension is 150 mm (nominal, stages with larger extensionscan be obtained on a custom basis). There is a sensitive springinstalled in NMAT that can provide a very small harmonicforce with specific frequency on the sample; thus, the dynamicmechanical properties, meanwhile, can be obtained during thetensile procedure. The samples were prepared as slender beltscut from PVDF-FPi% films and fixed on a perforated cardboardas shown in Figure 1(b). The length and height of belt-likesamples were measured by a caliper and screw micrometer,and the width was measured by the optical microscopy men-tioned in Morphology and Structure Characterization section.Three individual samples have been prepared and tested ineach category (12 in all). The tensile strain rate was set as 1.0 103 s1, and the harmonic force and the frequency weretypically 4.5 mN and 20 Hz, respectively.

RESULTS AND DISCUSSION

Microscopic MorphologyThe morphologies of the free surface of PVDF-FPi% films,that is, the surface opposite to the one in contact with theglass substrate, were analyzed. Figure 2 shows the micro-structural conformation of PVDF-FPi% measured by opticalmicroscopy. Samples of neat PVDF [Fig. 2(a)] show a micro-porous structure composed by micron-sized spherulites. Thisstructure is very similar to that in the work reported byNunes et al.50 and Serrado et al.,51 where it was identified asthe porous b-PVDF. Compared with the image of neat PVDF,spherulites in samples of PVDF-FPi% (i 3, 5, and 8) arevisibly bigger. Maybe it would be considered to the adhesionof FP-POSS on PVDF spherulite, but the content of FP-POSSobviously determined that there is no so much FP-POSS can

be provided. Further microscopy details would be needed toexplain this satisfactorily. Furthermore, the nonuniformity ofspherulites in Figure 2(bd), that is, some spherulites lookmuch smaller than others, may be noticed. POSS aggregateshave taken place in these samples. However, we want toexplain that it is not this case since these images wereobtained at a very high magnification (2000 zoom) and avery shallow depth of field. The smaller spherulites arejust those with only a small top part located on the focalplane. These results also show that the samples of PVDF-FPi% (i 3, 5, and 8) have bigger spherulites and roughersurfaces compared with neat PVDF.

The surface SPM height trace images taken within one spher-ulite for PVDF-FPi% at scan ranges of 2.0 2.0 lm2 areshown in Figure 3(ad), respectively. Height data correspondto the change in piezo height needed to keep the cantileverdeflection constant, which is always used to reflect the to-pography of the surface. It can be observed clearly that allspherulites consist of small particles with different sizes fordifferent samples. Compared with the image of neat PVDF[Fig. 3(a)], particles in PVDF-FPi% (i 3, 5, and 8) are simi-lar but possess larger size and more uniform dispersion. NoPOSS aggregate could be observed in all PVDF-FPi% (i 3, 5,and 8). It seems that the effect of FP-POSS is just to formbigger PVDF particles and very similar to the nucleatingagents in some PVDF mixtures.40,52,53 This may be under-stood from the simulation results:44 strong electrostaticforces arise owing to the same polar CAF bonds on PVDF,and the fluorinated POSS molecules tend to make morePVDF molecules be nucleated around FP-POSS. Further addi-tion of nucleating agents probably has little effect on forming

FIGURE 2 Optical microscopy surface images of PVDF-FPi% obtained by solution, i (a) 0, (b) 3, (c) 5, and (d) 8.



even bigger particles since they possess similar size forPVDF-FPi% (i 3, 5, and 8). However, some changes on themorphology can still be observed, particles are exhibitedmore clearly with the increasing of the FP-POSS content.This may be derived from a more complete nucleation dueto a relative high content of FP-POSS forming stronger nucle-ating agents. The surface SPM deflection trace images takenwithin one spherulite for PVDF-FPi% at scan ranges of 5.0 5.0 lm2 are shown in Figure 4(ad), respectively. Deflectiondata come from the differential signal off of the top and bot-tom photodiode segments. They are collected with low gainfeedbacks so the piezo remains at a constant position rela-tive to the sample. In this case, the tip and cantilever aredeflected by the features on the sample surface. The outputfluctuations in the cantilever deflection voltage from the topand bottom photodiode segments are recorded as a measureof the variation in the sample surface. Thus, the surfacedetails can be shown more clearly from the deflection trace.It can be found clearly from these figures that the particlesare more separate and the domain size of the particlesbecomes larger with the increasing of the FP-POSS content.As mentioned earlier, it should be derived from the more

complete nucleation of PVDF due to the higher content ofFP-POSS. These morphology changes would greatly influencesome mechanical properties of PVDF/FP-POSS nanocompo-sites, which will be discussed later.

Both optical and SPM images for all samples show that allspherulites and particles are uniformly dispersed and noPOSS aggregates could be observed. All these proved thatFP-POSS compounds and PVDF are fully miscible, which is ingood agreement with the simulation results.44

Phase AnalysisXRD patterns for PVDF-FPi% (i 0, 3, 5, and 8) and neat FP-POSS are shown in Figure 5. It can be found that the b and aphases are obviously exhibited in both neat PVDF and PVDF-FPi% (i 3, 5, and 8) via reflections at (110) and (020) crys-tal planes, occurring at 2y 20.6 and 18.5, respectively.No obvious difference can be found in the diffraction pat-terns between neat PVDF and PVDF-FPi% (i 3, 5, and 8),which means there may be similar crystallization process infour different systems, and the incorporation of FP-POSS haslittle influence on the crystalline phase of PVDF. Although avery intense main diffraction peak at 2y 19.4 reveals that

FIGURE 3 SPM height trace images taken within a spherulite for PVDF-FPi%, i (a) 0, (b) 3, (c) 5, and (d) 8, at scan rangesof 2.0 2.0 lm2.


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the neat FP-POSS is highly crystalline, and the XRD spectraof PVDF-FPi% (i 3, 5, and 8) show their crystalline featuresthat are only referable to neat PVDF matrix. This result pro-vides evidence that FP-POSS with mass ratio of 3 and 5%,

even up to 8%, are almost dispersible in PVDF matrix, justas already observed from AFM images.

The FTIR spectra for PVDF-FPi% (i 0, 3, 5, and 8) and neatFP-POSS are shown in Figure 6. It can be found that all FTIRcurves for PVDF-FPi% (i 3, 5, and 8) are very similar to

FIGURE 4 SPM deflection trace images taken within a spherulite for PVDF-FPi%, i (a) 0, (b) 3, (c) 5, and (d) 8, at scan ranges of5.0 5.0 lm2.

FIGURE 5 XRD patterns for PVDF-FPi% (i 0, 3, 5, and 8) andneat FP-POSS.

FIGURE 6 FTIR spectra for PVDF-FPi% (i 0, 3, 5, and 8) andneat FP-POSS. The peaks for a phase, b phase, and SiAOASi inFP-POSS are located.



that of neat PVDF, indicating that FP-POSS has not changedthe crystalline phase of PVDF. The band at 1120 cm1 isattributed to the stretching vibration of SiAOASi in POSS.54

However, it is not obvious in PVDF-FPi%, revealing that noPOSS agglomerates in the nanocomposites. The characteristictransmission peaks of a phase are 611, 766, and 1404 cm1,and the b phase can be proven through the presence of 870,1064, and 1182 cm1 transmission peaks.43 Compared withother nanocomposites, PVDF-FP5% presents more intensepeaks for the same a and b phases, indicating FP-POSS maybe leading to the formation of more a and b phase crystalswith 5 wt % content, which is very similar to the result inliterature.43

Hardness and Elastic ModulusThe local values of hardness and elastic modulus of PVDF-FPi% as a function of the maximum load and indentationdepth are shown in Figures 7 and 8, respectively. The size

effect revealed in Figure 7 is so evident that it is hard tojudge what values are reliable enough as determined fromthe G-Series Hardness and Modulus via Cycles Load Controlmethod. It may be related to the rough surface of PVDF-FPi%films and constant loading rate derived from this method. Arough surface determines that the real projected contactarea is always different with that computed from the pene-tration depth. And a constant loading rate means a changingstrain rate during the test, that is, the indenter is driven intothe sample with different speed as explained previously, willbring strain rate effect to the test results,55 especially tosome strain rate sensitive polymers and viscoelastic materi-als.56 Even so, approximate orders of PVDF-FP3% > PVDF-FP0% > PVDF-FP5%/PVDF-FP8% for the values of hardnessand elastic modulus are shown definitely in Figure 7. Theorder for PVDF-FP5% and PVDF-FP8% is hard to give sincethese two pairs of curves are so intertwined. It appears thatthe hardness and elastic modulus obtained by the G-Series

FIGURE 7 Influence of different FP-POSS contents on (a) hardness and (b) elastic modulus of PVDF-FPi% (i 0, 3, 5, and 8) asdetermined from the G-Series Hardness and Modulus via Cycles Load Control method.

FIGURE 8 Influence of different FP-POSS contents on (a) hardness and (b) elastic modulus of PVDF-FPi% (i 0, 3, 5, and 8) asdetermined from the G-Series Basic Hardness, Modulus at a Depth method.


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Basic Hardness, Modulus at a Depth method (Fig. 8) aremuch more uniform than the values in Figure 7. As the samevery low strain rate of 0.05 s1 was carried out at each test,the strain rate effect could be ignored. In addition, the sizeeffect is likely to be eliminated by indentation depths with atleast 500 nm. Compared with Figure 7, we think the resultsin Figure 8 are more reliable, which were used to computethe intrinsic average values of hardness and elastic modulusof PVDF-FPi%. Similar orders for the values of hardness andelastic modulus are shown in Figure 8 and proved this resultis not accidental. This needs to be identified by specificvalues.

Table 1 lists the values of hardness (H), elastic modulus (E),plasticity index (E/H),25,57 and H/E ratio of PVDF-FPi% aver-aged from the results as determined from the G-Series BasicHardness, Modulus at a Depth method (Fig. 8). It is clearthat the hardness and elastic modulus values follow theorders PVDF-FP3% > PVDF-FP0% > PVDF-FP5% > PVDF-FP8%. PVDF-FP3% possesses enhanced hardness and elasticproperties, which is in good agreement with the simulationresults.45 This may be understood from the role of FP-POSS.A small content of FP-POSS probably made the particles big-ger and more condensed, and thus, the mechanical proper-ties were improved in this case. While a large content of FP-POSS addition perhaps resulted in fully PVDF nucleation andformed more separate spherical particles, which wouldlargely decrease the crosslinkage among PVDF particles.Thus, although stronger particles were probably brought inthis case, the whole mechanical properties dropped. In addi-tion, all the experimental values are substantially less thanthe simulated ones, probably caused by the amorphous mod-els used in the simulation work, in which the nucleation wasnot considered thus the models are different from the struc-tures in this work. Plasticity index values do not showobvious differences or changing rules, and indicate that theplasticity of PVDF was not observably influenced by FP-POSS.

Typical loadunload cycles for PVDF-FPi% as obtained fromtwo different methods are shown in Figures 9 and 10,respectively. First, hf/hmax in all the curves is distinctly lessthan 0.7, indicating little pile-up occurred; thus, its influenceon the computing of contact areas could be ignored. In themethod G-Series Hardness and Modulus via Cycles LoadControl, as the same maximum loading was carried out ineach test and the loadunload cycles are similar thus thehysteresis (area between load and unload curves) can be

used to measure the plastic deformation produced duringthe loading part of the cycle. First, the area of PVDF-FP3% issmallest, indicating a small plastic deformation and betterhardness for this film. Then the areas of PVDF-FPi% followthe order PVDF-FP3% > PVDF-FP0% > PVDF-FP8% > PVDF-FP5%, very similar to the variation trend of the hardness andelastic modulus as listed in Table 1. In the method G-SeriesBasic Hardness, Modulus at a Depth, as the same maximumindentation depth was kept in each test and the projectedcontact areas are similar, then the loading at this point is ameasure of hardness, and the slope of the upper portion ofthe unload curve corresponds to the elastic modulus.48 Itcan be found from Figure 10 that the curves of PVDF-FP3%always exhibit the maximum loading among others, indicat-ing a better hardness for it. It is hard to compare the slopeof the upper portion of the unload curve directly; however,PVDF-FP3% and PVDF-FP0% always reveal the bigger slope,indicating the elastic moduli for these two films are greaterthan those of the other two. This is in accordance with theresults in Table 1.

Nanoscratch ProfilesThe influence of different FP-POSS contents on the effectivescratch profiles of PVDF-FPi% is shown Figure 11(ad),where the original morphology profile in prescan is notshown since it has been compensated for computing theeffective scratch and residual depth in the scratch-scan andpostscan, respectively. The scratch-scan and postscan profilescan be divided into two regions in the horizontal displace-ment: (1) the superficial region (050 lm), in which the tipdoes not penetrate the sample and the case is very similarto the prescan; (2) the scratch (for scratch-scan) or residual(for postscan) depth region, in which the profile is essen-tially determined by the scratch force, hardness, and plastic-ity, showing the scratch resistance, recovery behavior, andthe elastoplastic deformation of a sample.58 In case of a

TABLE 1 The Hardness (H), Elastic Modulus (E), Plasticity Index

E/H, and H/E Ratio of PVDF-FPi% (i 5 0, 3, 5, and 8)

FP-POSS (wt %) H (GPa) E (GPa) E/H H/E

0 0.174 (0.012) 2.61 (0.12) 14.99 0.0667

3 0.177 (0.018) 2.77 (0.16) 15.63 0.0640

5 0.158 (0.018) 2.33 (0.26) 14.68 0.0681

8 0.130 (0.036) 2.21 (0.26) 16.94 0.0590

Values in parentheses indicate (6) standard deviations.

FIGURE 9 Loadunload cycles for PVDF-FPi% (i 0, 3, 5, and 8)as obtained from the G-Series Hardness and Modulus via

Cycles Load Control method.



coating or thin film on a substrate, there should be anothermore region in which the fluctuant profile is related to theadhesive strength between the film and the substrate.59 Thisis always used to compute the critical load (Lc), which isvery valuable to characterize some failure behaviors causedby scratch testing, such as the adhesion strength, coatingcracking, delamination and brittle fracture, and so on.60 Inthis work, only the former two regions are involved and noany critical load values were obtained.

The effective scratch depth reveals the deformation resist-ance and hardness of the samples at a specific load, theharder the sample is and the shallower the depth should be.However, nearly all the samples possess similar effectivescratch depth [Fig. 11(ad)] (within the range of 50300lm), which indicates FP-POSS is probably not helpful toenhance the scratch resistance of PVDF. The effective resid-ual depth show that the additions of FP-POSS even reducethe elasticity recovery of PVDF because larger plastic defor-mations (residual depth) can be found in Figure 11(c, d). Itseems that the content of FP-POSS has brought little differ-

ence to this influence because the areas between the residualdepth curve and the upper horizontal axis in these three fig-ures almost are the same size. Specific values of residualscratch depth and pile up height at cross profile (located atthe scratch load of 0.5 mN) obtained after the test (listed inTable 2) have provided further evidence for this.

The additions of FP-POSS will not increase but probablyweaken the scratch resistance of PVDF to some extent,seems to be against the nanoindentation results in hardnessbecause PVDF-FP3% exhibits the best hardness and elasticmodulus properties among all samples. This should bederived from the surface roughness, which has been provedto affect the abrasion or scratching profile of a film.25,61,62

The rougher the surface is, the weaker the scratch resistancewould be exhibited.63 All the surface roughness values forPVDF-FPi% are listed in Table 2. Not surprisingly, PVDF-FP3%possesses the roughest surface, which has been confirmedby several more tests at different positions. Although thebest hardness is shown for PVDF-FP3%, the roughest surfacedetermines a weaker scratch resistance for it. To PVDF-FP5%

FIGURE 10 Loadunload cycles for PVDF-FPi% (i 0, 3, 5, and 8) at maximum indentation depth of (a) 500 nm, (b) 1000 nm, (c)1500 nm, and (d) 2000 nm as obtained from the G-Series Basic Hardness, Modulus at a Depth method.


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and PVDF-FP8%, their relative softer hardness makes it beeasily understood that they possess weaker scratch resist-ance. Moreover, neat PVDF (PVDF-FP0%) possesses the bestscratch resistance may be caused by two factors: a smoothersurface and a higher crosslinkage among particles, which isvery likely to exist in PVDF.

Engineering StressStrain ResponseThe engineering stressstrain responses for PVDF-FPi% asobtained from the nanotensile testing are shown in Figure12. Some key points on these curves, such as the offset yieldstrain (0.2% offset), peak stress, and break point, have beenlabeled as letters Y, P, and B, respectively. The subscriptnumbers point out which sample these values belong to. Thefull stressstrain curves can be divided into three regions:(1) the elastic region (from the start point to Y), in whichrecoverable (pure elastic) deformation happens and theHookes law is satisfied between the stress and strain, andthus, the Youngs modulus can be computed. The stress atthe offset yield strain point (ry) is called the tensile yieldstrength,63 at which a material begins to deform plastically;(2) the strain hardening region (Y!P), in which the stress isno longer proportional to strain and permanent nonrecover-able (plastic) deformation occurs. The stress at the maximumon the engineering stressstrain curve (point P) is defined asthe tensile strength or ultimate strength.64 The stressstrainbehavior in this region is always described by the formula r Ken,63 where K (the strength coefficient) and n (the strain

hardening exponent) are constants. The value of the strainhardening exponent lies between 0 and 1. A value of 0means that a material is a perfectly plastic solid, while avalue of 1 represents a 100% elastic solid; (3) the unstablefailure region (P!B), in which a small constriction or neckbegins to form at some point, and fracture ultimately occursat the neck. The fracture strength63 corresponds to the stressat fracture and the area under the stressstrain curve up tothe point of fracture is always used to ascertain the fracturetoughness [for the quasi static (low strain rate) situa-tion],52,63 which is a measure of the ability of a material toabsorb energy up to fracture and always used to describethe brittleness or ductility of a material. All these key me-chanical properties, for PVDF-FPi% in the three regions, aswell as the brittleness index (H/Kc)

65 are listed in Table 3.

FIGURE 11 Influence of different FP-POSS contents on the scratch profiles of PVDF-FPi%, i (a) 0, (b) 3, (c) 5, and (d) 8.

TABLE 2 The Surface Roughness (Ra), Residual Scratch Depth

(Dr), Pile-Up Height (Hp), and Scratch Width (Wr) at cross

profile of PVDF-FPi% (i 5 0, 3, 5, and 8)

FP-POSS

(wt %) Ra (nm) Dr (nm) Hp (nm) Wr (um)

0 655 (138) 1,130 (350) 469 (322) 41.7 (9.7)

3 1,031 (163) 1,552 (1,191) 744 (806) 32.6 (17.2)

5 587 (143) 1,065 (1,032) 789 (682) 53.9 (21.1)

8 689 (103) 1,366 (182) 635 (232) 45.7 (10.1)




The Youngs modulus of PVDF-FPi% shows the same varia-tion trend with the indentation results, which follows theorders PVDF-FP3% > PVDF-FP0% > PVDF-FP5% > PVDF-FP8%. However, the specific values are all much lower here.This is probably resulted from two factors: (1) the uncer-tainty in measuring the sizes of samples, which alwaysbrings positive errors (greater than the real value) in com-puting the equivalent diameters of cross sections becausethe deformations and fracture always occur at the positionwith a minimum size of the cross section but hard to befound when the sizes of the samples were measured; (2)

some micro defects in the samples, which will influence theresults in a tensile test while perhaps have no effect on ananoindentation one because only some small local areas areinvolved in such case. All these two factors are likely tomake the test values of Youngs modulus are lower than realones. Anyway, suppose these influences are similar to allsamples, it is obvious that the elastic modulus of PVDF wasreinforced by a small content of FP-POSS (3 wt %) but weak-ened by further additions of FP-POSS (5 and 8 wt %). Thevalues of yield strength (YS), tensile strength (TS), and frac-ture strength (FS) almost show the same trend and indicate

FIGURE 12 Engineering stressstrain responses for PVDF-FPi%, i (a) 0, (b) 3, (c) 5, and (d) 8 at strain rate of 103 s1. The lettersY, P, and B locate the positions of offset yield strain, peak stress, and break point, respectively.

TABLE 3 The Youngs Modulus (E), Offset Yield Strain (ey), Yield Strength (YS), Tensile Strength (TS), Fracture Strength (FS),

Fracture Strain (ef), Fracture Toughness (Kc), and Brittleness Index (H/Kc) of PVDF-FPi% (i 5 0, 3, 5, and 8) as obtained from

stressstrain responses and hardness values

FP-POSS (wt %) E (GPa) ey (%) YS (MPa) TS (MPa) FS (MPa) ef (%) Kc (MPa) H/Kc

0 1.71 (0.13) 1.77 (0.15) 26.67 (1.13) 36.51 (2.28) 35.76 (2.07) 5.90 (1.00) 1.70 (0.40) 102.35

3 2.00 (0.11) 1.77 (0.21) 30.44 (1.36) 40.75 (2.28) 39.61 (2.37) 7.23 (1.03) 2.42 (0.51) 71.43

5 1.59 (0.12) 1.97 (0.31) 27.94 (4.11) 36.96 (4.86) 32.98 (5.32) 13.87 (3.02) 4.52 (1.38) 34.96

8 1.17 (0.26) 1.90 (0.26) 19.96 (2.74) 30.64 (2.22) 26.63 (2.70) 24.20 (1.61) 6.50 (0.93) 20



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the FP-POSS contents have the same effect based on PVDF.The values of offset yield strain (ey) are almost the same forPVDF-FP3% and PVDF-FP0% and a little higher for PVDF-FP5% and PVDF-FP8%, shows FP-POSS will not decrease butslightly increase the elastic deformation range of PVDF.

The elongation at break (fracture strain) and fracture tough-ness values reveal an undoubtedly influence of FP-POSS onPVDF. The fracture strain and fracture toughness increasewith the increasing content of FP-POSS. The value of efincreases from 5.9 to 24.2%, giving PVDF-FP8% the tough-ness 300% higher than that of neat PVDF. It can also beseen that the brittleness index of PVDF-FP3%, PVDF-FP5%,and PVDF-FP8% are about 43, 193, and 412% lower thanthat of neat PVDF, respectively. Dramatic enhancements infracture toughness have been observed when the FP-POSScontent reaches to some extent. This result is very valuablein some applications of PVDF, such as energy harvesting,adaptive inflatable structures in aerospace engineering, largedeformational piezoelectric actuators, and so on, in which abetter toughness is vital and needs to be strictly satisfied.The mechanism to the toughness enhancements can prob-ably be understood from the morphology change in differentPVDF-FPi% images. In contrast to the neat PVDF, more sepa-rate and less crosslinked particles in PVDF nanocomposites(especially in PVDF-FP5% and PVDF-FP8%) lead to a structuremuch more conductive to plastic flow under applied stressand probably result in a more efficient energy dissipation,thereby reducing crack formation. This is similar to the workby Shah et al.52 and may be useful to explain the research byMartins et al.,43 where they found that POSS were acting aslubricant in PVDF matrix. More interpretations can bequoted from MD studies.66 PVDF nanocomposites might pos-sess remarkable toughness derived from the FP-POSS nano-particles to dissipate energy due to their mobility underapplied stress. Comparable time scales for motion of FP-POSS molecules and PVDF chains results in the enhancedtoughness. Our results can also provide more evidence of agood nanoscale dispersion of FP-POSS in PVDF matrixbecause it has been proved that good nanoscale dispersionis critical for improved toughness.52

It should be noticed that the properties of PVDF-FP3% areunique, as it shows increases in both stiffness and toughness.Most rigid fillers produce increase in stiffness versus that ofthe unfilled polymer. However, it is generally associated witha significant decline in toughness.67 On the contrary, animprovement in toughness is always accompanied by adecline in stiffness for soft organic fillers.68 It seems that FP-POSS has acted as both rigid and soft fillers in PVDF matrixsimultaneously. However, this has been influenced greatly bythe content of POSS. When the content is low, the role of FP-POSS is thus the rigid fillers to stiffen the material. Whilewhen the content is high, the toughness enhancements aredominated, and the stiffness declines; thus, the role of FP-POSS is the soft fillers. Similar results were obtained byKopesky et al.69 when cyclohexyl-POSS was incorporatedinto the poly(methyl methacrylate) matrix; however, themechanism was not analyzed. It seems that this effect is, to

some extend, a common pattern of POSS when it is used asfillers to a polymer matrix.

Dynamic Mechanical PropertiesDynamic mechanical properties of PVDF-FPi%, such as thestorage modulus, loss modulus, and loss tangent can also beobtained in the nanotensile testing. Dynamic mechanicalproperties are helpful in understanding and confirming theenergy dissipation mechanism for PVDF-FPi% as analyzedpreviously. Table 4 lists the values of storage modulus (G0),loss modulus (G00), and loss tangent (tan d) for PVDF-FPi% (i 0, 3, 5, and 8) obtained from the nanotensile testing atthe harmonic force of 4.5 mN and the frequency of 20 Hz. Itcan be observed that all the values of storage modulus arehigher than the corresponding values of Youngs modulus asobtained from the elastic portion of stressstrain responses(listed in Table 3). This is not surprising as the storage mod-ulus is averaged within the whole range of stressstrainresponses, in which modulus with higher values is kept afterthe yield strain due to a full elastic deformation. Loss modu-lus values for PVDF-FPi% (i 3, 5, and 8) show that theaddition of FP-POSS obviously increases the viscosity ofPVDF because they are much higher than that of the neatPVDF. Compared with the storage and loss modulus, loss tan-gent is more reliable to measure viscoelastic propertiesbecause it is the ratio of loss to storage modulus thus lessinfluenced by the uncertain factors in measuring or someunpredictable defects in materials. The values of loss tangentin Table 4 show an evidently viscosity increasing with theincreasing content of FP-POSS. The variation of loss tangentis in good agreement with the previous results in toughness.A higher value of the loss tangent, reflecting greater viscousflow, is corresponded with a more efficient energy dissipa-tion. This is expected, since particles with lower crosslinkagelevel resulted from further addition of FP-POSS lead to astructure much more favorable for plastic flow under appliedstress.

CONCLUSIONS

In this article, the morphology and some key mechanicalproperties of PVDF and FP-POSS nanocomposites wereexplored using experiments. PVDF/FP-POSS nanocompositescan be prepared by the solvent evaporation method. Opticaland SPM microscopy results showed FP-POSS and PVDFwere fully miscible and FP-POSS acted as the nucleating

TABLE 4 The Storage Modulus (G0), Loss Modulus (G00), andLoss Tangent (tan d) of PVDF-FPi% (i 5 0, 3, 5, and 8) as

obtained from the nanotensile testing at a harmonic force of

4.5 mN and a frequency of 20 Hz

FP-POSS (wt %) G0 (GPa) G00 (MPa) tan d

0 1.80 (0.11) 159.6 (31.5) 0.089 (0.022)

3 2.28 (0.18) 221.3 (48.0) 0.097 (0.021)

5 2.13 (0.20) 232.3 (56.7) 0.109 (0.025)

8 1.86 (0.26) 207.9 (28.9) 0.113 (0.015)




agent in PVDF matrix. A small content of FP-POSS couldmake PVDF chains be nucleated to form bigger spherical par-ticles and a further addition of FP-POSS formed more sepa-rate and less crosslinked particles due to a more completenucleation. XRD patterns and FTIR spectra showed that thereare two different crystalline phases in the nanocompositesand FP-POSS has not changed the crystalline phase of PVDF.As in XRD patterns, up to 5-wt % FP-POSS content, therewas an increase in both a and b phase. Hardness and elasticmodulus of PVDF/FP-POSS nanocomposites can be measuredby nanoindentation test with a very low strain rate and theadequate indentation depths. PVDF-FP3% possessed the bestelastic properties and hardness and further additions of FP-POSS made the stiffness of PVDF/FP-POSS nanocompositesdecrease. Nanoscratch testing can be used to measure thescratch resistance of PVDF/FP-POSS nanocomposites. Thescratch-testing profiles were clearly influenced by differentFP-POSS contents in PVDF-FPi%. Neat PVDF exhibited the bestscratch resistance while the additions of FP-POSS weakenedthe scratch resistance due to a rough surface or a less cross-linked structure. Nanotensile testing results showed both thestiffness and toughness of PVDF-FP3% were enhanced and fur-ther additions of FP-POSS brought surprisingly enhancementsin toughness of the nanocomposites while associated with adecline in stiffness. Dynamical mechanical properties indi-cated the viscosity of the nanocomposites increased with theincreasing FP-POSS contents in PVDF-FPi%. Because of theunique organicinorganic structure, FP-POSS probably actedas rigid inorganic fillers in PVDF matrix at a low content whilesoft organic fillers at relative high contents. The greatimprovement effects on toughness and the increasing viscos-ity resulted in further additions of FP-POSS should be derivedfrom the more separate and less crosslinked particles in thenanocomposites, and the nanoscale size of POSS compounds,which made the materials, are easier to form plastic flow andmore efficient on energy dissipation. The results in this workare helpful in realizing the improvement effects of differentFP-POSS contents on PVDF and understanding the mechanismfor these different effects. Another important advantage of FP-POSS on PVDF, we found recently, is that the piezoelectricproperties of PVDF are only very slightly influenced probablydue to the similar CAF bonds in their molecules. This is veryimportant to the application of PVDF and will be discussed inour future work.

ACKNOWLEDGMENTS

The authors thank the National Natural Science Foundation ofChina (11102053), the Science and Technology Innovation Tal-ents Special Fund of Harbin (Grant No. 2012RFQXG001), andthe Fundamental Research Funds for the Central Universities(Grant No. HIT. NSRIF. 2010070) for the financial support ofthis research.

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