This content has been downloaded from IOPscience. Please scroll down to see the full text.

Download details:

IP Address: 130.203.222.114This content was downloaded on 18/09/2014 at 22:22

Please note that terms and conditions apply.

Influence of dispersion states on the performance of polymer-based nanocomposites

View the table of contents for this issue, or go to the journal homepage for more

Home Search Collections Journals About Contact us My IOPscience

Influence of dispersion states on theperformance of polymer-basednanocomposites

Payam Khodaparast1 and Zoubeida Ounaies2

1Department of Material Science and Engineering, Pennsylvania State University, University Park, PA16802, USA2Department of Mechanical Engineering, Pennsylvania State University, University Park, PA 16802, USA

E-mail: zxo100@psu.edu

Received 21 February 2014, revised 17 July 2014Accepted for publication 21 July 2014Published 12 September 2014

AbstractAlthough nanoparticle-modified polymers have tremendous promise in many applications,particularly dielectric energy storage, true nanoscale dispersion is extremely difficult to achieve.In this paper, we carefully engineer various dispersion states of titania nanoparticles inpolyvinylidene fluoride and analyze their impact on dielectric behavior and energy storageability. In particular, we compare nanocomposites prepared using commercially availablenanoparticles to those we prepared using in situ and ex situ synthesis of nanoparticles. SEM andTEM studies showed that the in situ case leads to the best dispersion. Interestingly, dielectricpermittivity was most influenced by dispersion state where the in situ case showed a higherincrease, however, dielectric breakdown and energy storage density were less affected bydispersion and more affected by procedure that minimized residues and impurities. The in situtechnique, in particular, showed nanoscale dispersion, low dielectric loss and higher energystorage density. In terms of mechanical behavior, all three cases showed a similar performance inthe rubbery region, whereas the impact of dispersion was more pronounced in the glassy region.In fact, the trend was opposite to the dielectric permittivity where nanoscale dispersion resultedin a lower storage modulus likely due to the lower effective mechanical load transfer going to thenanoscale. The results of our study shed some light on the role of dispersion quality andprocessing techniques in affecting the final dielectric, mechanical and breakdown behavior ofTiO2-based polymer nanocomposites.

Keywords: polymer nanocomposites, dielectric, dispersion

(Some figures may appear in colour only in the online journal)

1. Introduction

The demand for lighter and smaller capacitors with higherenergy storage capability has created a strong motivation forresearch and development in the field of dielectric capacitors.Polymer composites can potentially offer the combination oflightweight and high dielectric breakdown exhibited byorganics and high dielectric permittivity offered by inor-ganics. Motivated by this promising combination, researchershave focused on microparticle-modified polymers (herein

referred to as polymer microcomposites) where a polymermatrix was filled with microscopic inorganic particles [1–7].Results were mixed where the increase in dielectric permit-tivity owing to presence of microparticles was essentiallyovercome by the decrease in dielectric breakdown, resultingin overall decrease in energy storage density in most cases.The advent of nanotechnology and the promise of highinterfacial area in nanocomposites energized research in so-called nanodielectrics; researchers saw the opportunity tosurpass dielectric behavior predicted by effective media

Smart Materials and Structures

Smart Mater. Struct. 23 (2014) 104004 (13pp) doi:10.1088/0964-1726/23/10/104004

0964-1726/14/104004+13$33.00 © 2014 IOP Publishing Ltd Printed in the UK1

approaches and potentially achieve significant improvementin not only dielectric permittivity and breakdown strength butalso in mechanical properties [8–11]. To date however,research in polymer nanocomposite brought significant con-tradictory results to the dielectric literature where unforeseentrends in dielectric properties were observed such as decreasein dielectric permittivity by going to nanoscale [8, 12–16].Furthermore introduction of nanoparticles into polymers wasaccompanied with aggregation and dispersion issues, whichmade it particularly challenging to achieve true nanoscaledispersion [17–22]. Presence of aggregates will mask anypotential nanoscale advantage and may lead to pooreffective properties due to introduction of dispersion-relateddefects. Therefore, promising advantages of nanoscale overmicrocomposites can only be attained when aggregationchallenges are tackled effectively and true nano-dispersion isachieved. Although there are many studies on dielectricpolymer nanocomposites, little attention has been paid tothe effect of dispersion and aggregation on their dielectricproperties. The goal of this work is to address the challengeof aggregation in dielectric polymer nanocompositesin order to reach true nanoscale dispersion so that apossible tailored dielectric response through novel chargetrapping at interfaces can be used to lead to new generation ofcapacitors.

In this paper we investigate TiO2-polyvinylidene fluoridenanocomposites using three different techniques that result invarious degrees of dispersion. Polyvinylidene fluoride,PVDF, is chosen as the polymer matrix because it has arelatively high dielectric permittivity and dielectric break-down; it is also amenable to solvent-based processing. TiO2 isused as the nanofiller due to its high dielectric permittivityand its amenability to sol-gel synthesis as well as its wideavailability commercially in a variety of sizes and crystalstructures. We first use in situ sol-gel synthesis of TiO2

nanoparticles in the polymer solution due to its promise toachieve uniform dispersion at the nanoscale [23–26]. Inaddition, for comparison, we also study ex situ sol-gel syn-thesized nanoparticles and commercially obtained nano-particles. For the ex situ sample, TiO2 nanoparticles are firstsynthesized in an aqueous solution using sol-gel techniqueand then added to PVDF using solution mixing. The dielectricbehavior of in situ and ex situ TiO2/PVDF samples arecompared with TiO2/PVDF composites made with commer-cially available TiO2 nanoparticles. We will draw importantconclusions on the relationship between the three dispersionstates and the resulting dielectric permittivity, dielectricbreakdown and mechanical behavior.

2. Experimental details

2.1. Materials

Commercial titania nanoparticles were purchased fromNanostructured & Amorphous Materials (Houston, TX).Average particle diameter per manufacturer data sheet isgiven as 15 nm. PVDF powder (trademark Kynar 301) was

provided by Arkema Dimethylacetamide, DMAc, (99% pure)was purchased from Sigma Aldrich (St. Louis, MO). Tita-nium butoxide (TiBuO4), 99.9% purity, HNO3, 70% con-centrated and acetyleacetone were all purchased from SigmaAldrich (St. Louis, MO).

2.2. Synthesis and processing

2.2.1. In situ TiO2/PVDF nanocomposites. The amounts ofchemicals used in the sol-gel reaction are described in table 1.Titanium butoxide (TiBuO4), was chosen as the precursor andwas added along with deionized water to the PVDF solution.Deionized water hydrolyzes TiBuO4 alkoxide end groups andresults in titanium hydroxide groups. Subsequently,hydroxide groups undergo polycondensation reaction andfinally form titanium oxide (Ti–O–Ti) bonds [23, 27].Figure 1 contains the scheme of TiO2 polycondensationreaction in the presence of water.

Since the water hydrolysis rate of most titaniumorganometallic compounds is extremely fast, making itdifficult to control the reaction, a chelating agent, namelyacetylacetone was used to slow down the reaction rate. Inaddition to chelating agent, nitric acid (HNO3) was also usedin order to control the pH of the solution; high concentrationof ions on the nanoparticles’ surface increases the nanopar-ticle surface charge and reduces the chance of particlecoagulation [23, 27]. Thus specific amounts of chelatingagent and nitric acid were used to control the final particlesize in the system (see table 1). Figure 2 shows the schematicof the in situ processing steps.

2.2.2. Ex situ TiO2/PVDF nanocomposites. The synthesis ofex situ TiO2 particles starts with addition of acetylacetone toDI water, followed by nitric acid (figure 3). In the next step,titanium butoxide (TiBuO4) is added to the solution andstirred at room temperature for 24 h. Then sol-gel TiO2

particles were centrifuged, removed from the rest of thesolution, then washed and rinsed multiple times with DI waterand DMAc. In the next step, dispersed TiO2 in DMAcunderwent high power probe sonication for 1.5 h. The highpower probe sonication procedure promoted particle

Figure 1. Schematic representation of sol-gel reaction of titaniumdioxide.

2

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

crystallization in the colloidal phase through localizedheating. In the nanomaterial literature, there is strongevidence that ultrasonication is an effective method toinduce crystallinity in particle suspension systems [28–31].Probe-sonicated TiO2 in DMAc was then centrifuged for10 min at 7000 rpm to remove micron-sized aggregates. Thesupernatant containing nanoscale ex situ TiO2 particles wasthen collected and added to PVDF solution to make ex situTiO2/PVDF nanocomposites.

In order to compare in situ and ex situ nanocompositeswith conventional TiO2/PVDF composite, commerciallyavailable anatase TiO2 nanoparticles in powder form werepurchased from Nanoamor and embedded in the PVDF resin.A comprehensive description of fabrication method ofcommercial nanoTiO2/PVDF is explained elsewhere [13].

2.3. Characterization

2.3.1. Particle size analysis. Average nanoparticle size wascharacterized in the solvent by particle size analysis(Nanosizer-S by Malvern) using dynamic light scatteringtechnique.

2.3.2. X-ray photoelectron spectroscopy. The chemistry ofsol-gel TiO2 particles was characterized using x-rayphotoelectron spectroscopy (XPS). The XPS is a KratosAxis Ultra XPS with a monochromatic Al Kα x-ray sourceoperated at 14 kV, 20 mA in hybrid slot mode and with0.01–0.5% atomic sensitivity and 0–10 nm sampling depth.Samples were sputtered with 4 keV Ar+ ions rastered over a

Table 1. List and amount of chemicals used in sol-gel reaction of titanium oxide.

Processing method/materials TiBuO4 DI Water HNO3 (70% concentrated) Acetylacetone (ACAC) PVDF DMAc

In situ Weight (gr) 0.8 0.2 0.2 0.2 3.8 47Mole 0.002 0.010 0.004 0.002 — 0.54

Ex situ Weight (gr) 0.8 100 0.2 0.2 3.8 47Mole 0.002 5 0.004 0.002 — 0.54

Figure 2. Flowchart of in situ TiO2/PVDF processing steps.Figure 3. Flowchart of ex situ TiO2/PVDF processing.

3

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

3 mm2 area for 30 min The total depth is estimated to beapproximately 675 Å.

2.3.3. Dielectric permittivity measurements. A custom-madedielectric measurement setup using a LCR meter coupled witha lock-in amplifier (Stanford Research Systems) was used tomeasure the dielectric permittivity from 10!2 to 106 Hz.Sample geometry was such that a parallel plate capacitorconfiguration is used; the electrodes were evaporated silver,1 cm in diameter and 100 nm in thickness.

2.3.4. Dynamic mechanical analysis. A state of the artdynamic mechanical analyzer, DMA RSA G2 by TAInstruments, was used to measure mechanical behavior ofthe nanocomposite samples under dynamic loading.Temperature sweep tests from !90 °C to +90 °C were usedwith temperature ramp of 3 °Cmin!1. All experiments wererun at a frequency of 1 Hz.

2.3.5. Scanning electron microscopy. Scanning electronmicroscopy was performed using an FE-SEM (NanoSEM630 by FEI). Nanocomposite films were first fractured underliquid nitrogen to achieve a brittle fracture and then sputtercoated by an Eminec Sputter coater using Iridium as thecoating metal. Cross-section of the fractured samples wascharacterized under FE-SEM to study the morphology as wellas particle dispersion in nanocomposite films.

2.3.6. Transmission electron microscopy. TEMinvestigations were carried out on a JEOL JEM-2010(200 kV) microscope. Polymer nanocomposite films weremicrotomed to 100 nm-thick slices under cryogenicconditions.

2.3.7. Electrical breakdown measurements. Electricalbreakdown experiments were done using a Trek highvoltage amplifier in galden oil. Around 30 different pointswere selected on the surface of each film and high voltagewas applied to each point with a ramp rate of 500 V s!1 fromzero up to the failure point of the sample.

2.3.8. FT-IR spectroscopy. A Bruker V70 Fouriertransformed infrared spectroscopy (FT-IR) spectrometer inattenuated total reflection (ATR) mode was used to gatherspectra data of TiO2 powders and also TiO2/PVDFnanocomposite films. FT-IR spectrum data can beparticularly useful in studying chemical species present insol-gel-processed TiO2 particles as well as identifying thedifferent crystalline phases of PVDF-basednanocomposite films.

3. Results and discussion

3.1. Characterization of sol-gel TiO2 particles

Figure 4 shows an FE-SEM image (right) as well as particlesize distribution of TiO2 particles (left) made by sol-geltechnique, as described in section 2.2. Particles are all sphe-rical and the size distribution is bimodal with two peaks at200 nm and 30 nm. Particle size analysis of diluted colloids ofTiO2 particles confirmed the bimodal size distribution of thenanoparticles.

Figure 5 shows XPS results of binding energy of titaniumand oxygen elements on the surface of sol-gel TiO2 particlesand also for argon-etched samples. The main Ti peak at458.3 eV in XPS high resolution spectrum corresponded tobinding energy of Ti2p in TiO2 compound [32]. The oxygenpeak at 530 eV was also assigned to O1s in TiO2 compound.There is a shoulder at 532 eV for oxygen corresponding tosurface hydroxyl group and molecular water that almostcompletely vanishes after etching with argon ionbeam [32–34].

Figure 6 shows the FTIR spectrum of both in situ andex situ sol-gel TiO2 samples, along with commercial TiO2

sample. The FTIR spectrum of commercial TiO2 powder wasas expected with a main broad peak of Ti–O–Ti from 500 to900 cm!1 wavenumber. In contrast, the sol-gel TiO2 powdershowed many extraneous peaks that could be related toresidues of the sol-gel process. In particular, peaks at 482,584, 745 and 972 cm!1 possibly correspond to the residues ofTiO2 precursor; peaks at 780, 1167 and 1600 cm!1 areassigned to the residues of acetylacetone and finally peaks at

Figure 4. Particle size analysis and FE-SEM images of sol-gel nano TiO2 particles.

4

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

1126, 1282, 1367, 1600 and 1530–1650 cm!1 can be relatedto the TiBuO4-acetylacetone complex structure [35–37]. Inaddition, by comparing the FTIR spectrum of ex situ andin situ, it is found that most of the peaks related to acet-yleacetone and titanum butoxide complex are not presentanymore in the ex situ case. The only peaks left in the ex situ

sample are at 1330, 1400 and 1628, which can be assigned toTi–O–Ti, residue of solvent and bending vibration of absor-bed molecular water, respectively [38]. Both presence ofsolvent and water can be further controlled through highertemperature processing.

The XRD results of sol-gel TiO2 particles (figure 7)shows an amorphous halo with no significant crystallinesignature; whereas XRD of the commercial TiO2 powderindicated TiO2 was in the anatase crystalline form. It is worthmentioning that TiO2 particles in crystalline form have ahigher dielectric constant than TiO2 in amorphous form. Theaverage dielectric constant for amorphous TiO2 reported inthe literature is less than 20 [39, 40]; whereas for anatase andrutile crystalline forms, values of around 50–70 and 100–120are reported, respectively [41–43]. Thus it seems advanta-geous to have crystalline TiO2 particles when higher effectivedielectric constant is desired. Figure 7 also shows the XRDspectrum of ex situ TiO2 sample; the peaks confirm anatasecrystalline structure.

Figure 5. XPS results for sol-gel TiO2: (a) survey scans over a widerange of binding energy, (b) high-resolution scan of Ti and (c) high-resolution scan of O.

Figure 6. Ft-IR spectra of sol-gel TiO2 as well as commercial TiO2purchased from Nanoamor.

Figure 7. XRD of sol-gel TiO2 particles.

5

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

3.2. Characterization of TiO2/PVDF nanocomposite films

3.2.1. SEM. Figure 8 shows the FE-SEM image of 5 wt%commercial (figure 8(a)), in situ (figure 8(b)) and ex situ

(figure 8(c)) processed TiO2/PVDF composite films. In thecase of commercial nanoTiO2 in PVDF, the nominal size ofindividual nanoparticles reported by Nanoamor is 15 nm.However as it is shown in figure 8(a), micron size aggregates

Figure 8. (a) FE-SEM image of 5 wt% commercial nanoTiO2 in PVDF (b) in situ sol-gel TiO2/PVDF nanocomposites and (c) ex situ sol-gelTiO2/PVDF.

Figure 9. (Left column) TEM images of in situ sol-gel TiO2/PVDF and (right column) ex situ sol-gel TiO2/PVDF nanocomposite.

6

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

are seen throughout the sample, demonstrating a poordispersion in PVDF. On the other hand, uniformdistribution of TiO2 particles was achieved using in situ sol-gel technique, as can be seen in figure 8(b). Ex situ sample(figure 8(c)) also showed uniform dispersion and distributionof nanoparticles throughout PVDF matrix and therefore, forboth in situ and ex situ cases, there is a much better dispersionthan in the commercial TiO2 sample.

3.2.2. TEM. The transmission electron microscope (TEM)images of in situ and ex situ nanocomposites samples areshown in figure 9. In situ sample shows better dispersion andmore uniform particle distribution throughout the PVDFmatrix with individual nano-TiO2 particles observed. On theother hand, ex situ sample shows clusters of nanoparticles inthe range of 100–200 nm but well dispersed in the matrix.Films prepared using commercial TiO2 particles were notcharacterized using TEM due to the large agglomerates seenunder SEM.

3.2.3. Differential scanning calorimetry (DSC) ofnanocomposites. The percent crystallinity is calculated bytaking the ratio of the area under the melting peak of thesample Δ( )Hm to the enthalpy of fusion of 100% crystalline

sample Δ( )Hm0 , as shown by equation (1):

ΔΔ

= !XH

H% 100. (1)c

m

m0

The value of the enthalpy of fusion for 100% crystallineα-phase PVDF Δ( )Hm

0 is 104.4 J g!1 [61]. In the case of thenanocomposites and due to the presence of titania particles, acorrection factor is applied using equation (2) taking intoconsideration the true weight fraction (wf) of the polymer:

=X correctedX

wf_ %

%. (2)c

c

polymer

Differential Scanning calorimetry (DSC) results of TiO2/PVDF nanocomposites are shown in figure 10 and table 2. Allsamples show the same melting behavior at Tm (!160 °C).The percent crystallinity of nanocomposites was calculatedfrom the area under the melting peak and the results areshown in table 1. Except the in situ sample, which showedalmost 10% drop in percent crystallinity, all the other samplesshowed the same range of crystallinity as pure PVDF sample.At this point, it is speculated that the drop in percentcrystallinity of in situ sample could be related to theconfinement of polymer chains; the nanoscale dispersion inthis case results in a very low inter-particle distance, andtherefore smaller effective particle size and highereffective number of particles as compared to the other twocases. As a result, the arrangement of polymer chains intocrystalline structure is hindered, and the percent crystallinityis reduced.

3.2.4. DMA of nanocomposites. Dynamic mechanicalanalysis (DMA) of nanocomposite samples is presented in

terms of storage modulus and tan(δ) as functions oftemperature in figure 11 (left) and (right), respectively.Glassy elastic modulus was higher in commercial samplescompared to both in situ and ex situ samples. The bigger sizeof aggregates in commercial samples facilitates mechanicalload transfer from softer polymer matrix to stiffer ceramicparticles. In the rubbery state (above Tg) the difference inelastic moduli for all filled cases fades away and they all showsimilar mechanical behavior. As to the loss behavior, tan(δ) ofnanocomposites showed similar behavior at the glasstransition temperature (Tg or αa) measured at !40 °Cfor all cases. The second relaxation peak is αc (around80 °C), corresponding to thermo-mechanical relaxation ofcrystalline moieties. Even though the location of both αaand αc in the temperature domain stays the same for allnanocomposite samples as well as pure PVDF, the magnitudeof αc peak differs when comparing in situ samples withthe rest. This difference can be explained by the lowerpercent crystallinity of in situ samples, as shown by DSCexperiments.

3.2.5. Dielectric permittivity of nanocomposites. Dielectricpermittivity behavior of the TiO2/PVDF nanocompositefilms is plotted in figure 12(a). Significant increase indielectric permittivity was observed in the in situ TiO2/PVDF composite samples compared to commercialTiO2/PVDF samples. Particularly, dielectric permittivity ofin situ 5 wt% sample was 65% higher than the commercialTiO2/PVDF sample at 1 kHz. However the dielectric loss

Figure 10. DSC profile of PVDF based nanocomposites.

Table 2. Percent crystallinity (X%), melting temperature (Tm) andcrystallization temperature (Tc) in TiO2/PVDF nanocomposites.

Xc% Tm (°C) Tc (°C)

Pure PVDF 46 159 132Commercial TiO2/PVDF 46 158 134In situ TiO2/PVDF 37 157 132Ex situ TiO2/PVDF 47 158 135

7

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

factor of the in situ sample (figure 12(b)) is also muchhigher. Dielectric permittivity of the ex situ sample liesabove the commercial sample but below the in situ sample.As figure 12(b) shows, the loss factor of the ex situ samplewas reduced significantly compared to the in situamorphous sample.

To better understand the dielectric permittivity behaviorof in situ samples, the AC conductivity of all samples weremeasured as a function of frequency (figure 13). The ACconductivity of in situ TiO2/PVDF sample showed up to threeorders of magnitude increase in the low frequency regimecompared to the other cases, with a behavior that reaches aplateau as it approaches lower frequencies. This behavior isindicative of ionic conductivity and is the reason for thehigher dielectric loss of in situ samples at low frequencies[44–47]. It is believed that considerable number of ionicspecies are trapped in the in situ nanocomposite samplesfollowing the sol-gel reaction, such as byproducts and residueof water and precursors. Figure 6 also confirms presence ofthese ionic species in the FTIR spectrum. Upon applying anelectric field, these ionic species respond in form of chargecarriers, as they ionize and drift toward the opposite electrode.At low frequencies, higher numbers as well as heavier ionsare able to drift and carry current across the sample so that ACconductivity and dielectric loss increase significantly. In case

of commercial TiO2/PVDF, the concentration of ionic speciesis much lower and AC conductivity does not increase asmuch. Finally ex situ TiO2/PVDF sample showed the lowestAC conductivity at low frequencies and almost no changecompared to pure PVDF. First, it is believed that washing andrinsing of ex situ TiO2 particles as part of the process

Figure 11. DMA results of TiO2/PVDF nanocomposites in terms of storage modulus (left) and tan(δ)(right).

Figure 12. (a) Dielectric permittivity and (b) dielectric loss factor as a function of frequency for in situ TiO2/PVDF as well as ex situ andcommercial nanoTiO2/PVDF.

Figure 13. AC conductivity as a function of frequency for TiO2/PVDF nanocomposites.

8

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

eliminated most of the ionic species, as confirmed by FTIR(figure 6). Second, better dispersion and no aggregation in theex situ samples significantly increased effective interfacialarea between nanoparticles and polymer matrix; there isevidence that high volume content of interface can trap chargecarriers in the systems containing large dipole moments[48, 49]. This also explains the lower loss behavior of ex situTiO2/PVDF sample compared to commercial TiO2/PVDF.

Figure 14 shows dielectric loss as a function of frequencyat different temperatures for all nanocomposite samples. Theloss behavior of pure PVDF sample displays two majorrelaxation processes, as also reported in the literature [50–54].The first one, αa relaxation, happens in the higher frequencyregime and is associated with the dielectric response of PVDFat the glass transition temperature. The second relaxationprocess is the !c relaxation and is associated with dielectricrelaxation of crystalline moieties in the polymer.

Both relaxation processes move toward lower frequen-cies with increasing temperature and follow an Arrheniustemperature dependence [54, 55]:

φ=⎛⎝⎜

⎞⎠⎟f f

k Texp (3)max o

B

where fmax and T are the frequency (Hz) at max loss value andtemperature (K), respectively; kB is Boltzmann factor and φ isthe activation energy for the specific relaxation process.

Therefore, by plotting fmax in natural logarithm format as afunction of (1/T), the activation energy can be calculated asthe slope of the curve obtained by linear regression of data.These results are plotted in figure 15 for αa relaxation process.The calculated activation energies for all samples wereapproximately φ= 1 eV, except for the ex situ TiO2/PVDFsample where the activation energy rose to φ= 1.2 eV. Theincrease in the activation energy of αa relaxation process

Figure 14. Dielectric loss as a function of frequency for TiO2/PVDF nanocomposites at different temperatures: (a) pure PVDF (b) in situTiO2/PVDF (c) ex situ TiO2/PVDF (d) commercial TiO2/PVDF.

Figure 15. Arrhenius behavior of dielectric relaxation process inPVDF and its nanocomposites obtained from the dielectric loss data.

9

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

could be explained by the significant increase in specificsurface area between nanoparticles and polymer in the case ofex situ sample, where more energy is needed for the polymerdipoles to respond and orient themselves to the appliedelectric field. It is believed that presence of ionic impurities inthe case of in situ nanocomposite samples counter-balancesthe effect of high surface area so that the activation energy ofin situ sample stays the same as that of pure PVDF.

Dielectric loss of nanocomposites can also be presentedas a function of temperature in order to study the effect ofnanoparticles on Tg. Figure 16 shows dielectric loss withrespect to temperature for nanocomposites as well as purePVDF sample measured at 20 Hz and 1 kHz. The Tg of PVDFwas determined to be !38.5 °C at 20 Hz. On the other hand Tgdecreased in the case of in situ TiO2/PVDF (to !39.5 °C at20 Hz) and commercial TiO2/PVDF (to !42.5 °C at 20 Hz).However, in the case of ex situ TiO2/PVDF, an increase in Tgup to !30.4 °C was observed. The same trend also occurred inother frequencies including 1 kHz. The reason for thesignificant increase in Tg in the ex situ sample can beexplained by the better dispersion and smaller size ofaggregates as opposed to the commercial TiO2/PVDF sample.In addition to that, by comparing ex situ sample with in situsample, the Tg is significantly lower in the in situ case, due topresence of impurities at the interphase which facilitatesmobility of the interfacial dipoles upon application ofelectric field.

3.2.6. Dielectric breakdown of nanocomposites. Dielectricbreakdown results of nanocomposites are shown in figure 17and table 3. Weibull cumulative failure probability is used tofit the breakdown data and extract electrical breakdownvalues. Weibull model is the most common statistical analysistool used for dielectric breakdown studies and it isrepresented as [56–58]:

= " "β⎪

⎪⎪⎪

⎧⎨⎩

⎛⎝⎜

⎞⎠⎟

⎫⎬⎭

P EEE

( ) 1 exp . (4)b

Where P is the cumulative probability as a function of E.The parameter β is the shape factor; it is analogous to the

inverse of the standard deviation of the normal distribution. Inother words, higher shape factor values guarantee higherreliability of measurements. Finally Eb is the scale parameterand it acounts for electric field at failure or statistical electricalbreakdown of the sample [56, 58]. In figure 17, Weibullcumulative failure probability (P) is plotted as a function ofelectrical breakdown E (MV/m) for pure PVDF as well asnanocomposites. The values of statistical electrical break-down were obtained at failure probability of P = 0.6321 fordifferent samples and are presented in table 3.

Dielectric breakdown of in situ TiO2/PVDF nanocompo-sites decreased compared to pure PVDF by 7%. This decreasestems from presence of impurities in the in situ samples. Themechanism of dielectric breakdown in polymers is explainedby electrical treeing propagation phenomenon [9, 59].Electrical tree mechanism results in multiple electricallyconductive paths and branches in the polymer matrix that areconnected by defect centers such as microvoids, impurities,trapped ions and free radicals. Presence of defects of thisnature facilitates electrical treeing propagation in the polymermatrix by enhancing electron-hopping mechanism from onedefect to another and hence results in lower dielectricbreakdown strength. In the case of the in situ samples,

Figure 16. Dielectric loss as a function of temperature for different PVDF nanocomposites (left) 20 Hz and (right) 1 kHz.

Figure 17.Dielectric breakdown data in terms of Weibull cumulativefailure probability (P) for TiO2/PVDF nanocomposites.

10

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

because of presence of large amount of ionic impurities leftbehind after the sol-gel reaction, dielectric breakdownstrength has decreased. On the other hand, in case of ex situand commercial TiO2/PVDF composites, improvement indielectric breakdown strength was observed. This improve-ment can be explained by nanoparticles acting as obstacles tothe treeing mechanism. Since defects, acting as centers ofinterconnecting electrical trees, are inherently located in thepolymer matrix, presence of particles causes a longer pathbetween these defects and makes its more difficult for theelectric current to flow. There was not any relation foundbetween the size of aggregates and dielectric breakdownmechanism by comparing ex situ to commercial sample in thisstudy.

The energy density can be used as figure of merit in highpower pulse capacitors, calculated as the stored energy perunit volume. Assuming a linear dielectric behavior, theenergy storage density can be obtained from the followingequation [60, 61]:

ε ε=U ! E . (5)r02

Where U is the energy density, ε0 is permittivity ofvacuum, εr is the relative permittivity of the sample and E isthe maximum allowable electrical field in the sample. Thevalues of maximum allowable energy density of thenanocomposites in this study are calculated and presentedin table 3. Energy density of nanocomposites showed morethan 50% improvement over pure PVDF sample. Eventhough the highest improvement was observed in the in situnanocomposites, lower number of shape factor, β= 3,significantly limits fabrication of reliable capacitors by thein situ technique. On the other hand, ex situ nanocompositesamples showed over 60% improvement in energy densityand highest shape factor value, pointing to a promisingtechnique in fabricating polymer based nanocompositecapacitors.

4. Conclusion

In this paper, in situ sol-gel synthesis of TiO2 particles isutilized to achieve a uniform dispersion of nano-sized TiO2

particles in the PVDF matrix. However, we also show that thein situ TiO2/PVDF sample suffered from undesirably high

dielectric loss; FTIR and AC conductivity studies indicatethat the reason for the high dielectric loss is most likelypresence of impurities and byproducts left in the sample fromthe in situ sol-gel process, resulting in high ionic conduction.In order to eliminate this unwanted effect, sol-gel TiO2 par-ticles were synthesized using an ex situ technique whererinsing and cleaning steps are introduced to remove the left-over chemical species. The ex situ particles were then sub-jected to high-power probe sonication in order to obtaincrystalline TiO2 particles with higher dielectric constant.Finally, resulting crystalline ex situ TiO2 particles wereembedded in PVDF. The dielectric behavior of ex situ sam-ples showed considerably lower dielectric loss compared toin situ samples, and in general, the ex situ method is morepromising to fabricate high dielectric permittivity and lowdielectric loss polymer-based composites. The result of thisstudy specifically revealed that nanoscale dispersion couldinfluence dielectric permittivity by means of higher Maxwell-Wagner interfacial polarization at the boundaries of hetero-geneous phases. Interestingly, dielectric breakdown anddielectric loss did not show any appreciable trend with dis-persion quality and size of inclusions, whereas they weremore impacted by amount of ionic impurities in the system. Interms of mechanical behavior, all three cases showed a similarbehavior in the rubbery region, whereas impact of dispersionwas more pronounced in the glassy region. The results of ourstudy shed some light on the role of dispersion quality andprocessing techniques in affecting the final dielectric per-mittivity and breakdown as well as mechanical behavior ofTiO2-based polymer nanocomposites.

References

[1] Imai T, Sawa F, Nakano T, Ozaki T, Shimizu T,Kozako M and Tanaka T 2006 Effects of nano-and micro-filler mixture on electrical insulation properties of epoxybased composites IEEE Trans. Dielectr. Electr. Insul. 13319–26

[2] Chiang C, Popielarz R and Sung L 2001 Dielectric propertiesand morphology of ferroelectric ceramic-polymer compositefilms MRS Proc. (Cambridge: Cambridge University Press)

[3] Akram M, Javed A and Rizvi T Z 2006 Dielectric properties ofindustrial polymer composite materials Turk. J. Phys. 29355–64

Table 3. Dielectric permittivity and breakdown strength for TiO2/PVDF nanocomposites.

SampleSample thick-ness (μm)

Electrical break-down (MV m!1)

Shapefactor (β)

Activationenergy (αa)

"’(1 kHz)

Tandelta

(1 kHz)

Maximum allowableenergy den-sity (J cc!1)

Pure PVDF 38 ± 2.5 446 6 1.0 eV 6.8 0.014 5.98In situTiO2/PVDF

34 ± 3.4 415 3 1.0 eV 13.0 0.036 9.90

Ex situTiO2/PVDF

34 ± 2.7 460 7 1.2 eV 10.2 0.018 9.53

CommercialTiO2/PVDF

35 ± 5 468 6 1.0 eV 9.2 0.019 9.00

11

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

[4] Cao Y, Irwin P C and Younsi K 2004 The future ofnanodielectrics in the electrical power industry IEEE Trans.Dielectr. Electr. Insul. 11 797–807

[5] Youngs I 2000 Dielectric measurements and analysis for thedesign of conductor/insulator artificial dielectrics IEE Proc.,Sci. Meas. Technol. 147 202–8

[6] Chan C M 1996 A novel method to prepare high-resistivitypolymer composites Polym. Eng. Sci. 36 495–500

[7] Rao Y, Yue J and Wong C 2001 Material characterization ofhigh dielectric constant polymer-ceramic composite forembedded capacitor to RF application Advanced PackagingMaterials: Processes, Properties and Interfaces, 2001. Proc.Int. Symp. on: IEEE) pp 280–5

[8] Nelson J K and Fothergill J C 2004 Internal charge behaviourof nanocomposites Nanotechnology 15 586–95

[9] Tanaka T 2005 Dielectric nanocomposites with insulatingproperties IEEE Trans. Dielectr. Electr. Insul. 12 914–28

[10] Tanaka T, Kozako M, Fuse N and Ohki Y 2005 Proposal of amulti-core model for polymer nanocomposite dielectricsIEEE Trans. Dielectr. Electr. Insul. 12 669–81

[11] Kim P, Jones S C, Hotchkiss P J, Haddock J N, Kippelen B,Marder S R and Perry J W 2007 Phosphonic acid-modifiedbarium titanate polymer nanocomposites with highpermittivity and dielectric strength Adv. Mater. 19 1001–5

[12] Roy M, Nelson J K, MacCrone R K, Schadler L S,Reed C W and Keefe R 2005 Polymer nanocompositedielectrics-the role of the interface IEEE Trans. Dielectr.Electr. Insul. 12 629–43

[13] Khodaparast P and Ounaies Z 2013 On the impact offunctionalization and thermal treatment on dielectricbehavior of low content TiO2 PVDF nanocomposites IEEETrans. Dielectr. Electr. Insul. 20 166–76

[14] Nelson J and Hu Y 2005 Nanocomposite dielectrics—properties and implications J. Phys. D: Appl. Phys. 38 213

[15] Hallouet B, Desclaux P, Wetzel B, Schlarb A and Pelster R2009 Analysing dielectric interphases in compositescontaining nano-and micro-particles J. Phys. D: Appl. Phys.42 064004

[16] Singha S and Thomas M J 2008 Dielectric properties of epoxynanocomposites IEEE Trans. Dielectr. Electr. Insul. 1512–23

[17] Vrejoiu I, Pedarnig J, Dinescu M, Bauer-Gogonea S andBäuerle D 2002 Flexible ceramic–polymer composite filmswith temperature-insensitive and tunable dielectricpermittivity Appl. Phys. A 74 407–9

[18] Dillon D R, Tenneti K K, Li C Y, Ko F K, Sics I and Hsiao B S2006 On the structure and morphology of polyvinylidenefluoride–nanoclay nanocomposites Polymer 47 1678–88

[19] Manias E, Polizos G, Nakajima H and Heidecker M 2007Fundamentals of polymer nanocomposite technology FlameRetardant Polymer Nanocomposites ed A Morgan andC Wilkie 31–66

[20] Perrin-Sarazin F, Ton-That M-T, Bureau M and Denault J 2005Micro-and nano-structure in polypropylene/claynanocomposites Polymer 46 11624–34

[21] Dang Z-M, Wang H-Y and Xu H-P 2006 Influence of silanecoupling agent on morphology and dielectric property inBaTiO3/polyvinylidene fluoride composites Appl. Phys.Lett. 89 112902

[22] Dang Z-M, Wang H-Y, Zhang Y-H and Qi J-Q 2005Morphology and dielectric property of homogenous BaTiO3/PVDF nanocomposites prepared via the natural adsorptionaction of nanosized BaTiO3 Macromol. Rapid Commun. 261185–9

[23] Wang S, Wang M, Lei Y and Zhang L 1999 ‘Anchor effect’ inpoly (styrene maleic anhydride)/TiO2 nanocompositesJ. Mater. Sci. Lett. 18 2009–12

[24] Tuncer E, Sauers I, James D R, Ellis A R, Paranthaman M P,Aytu! T, Sathyamurthy S, More K L, Li J and Goyal A 2007

Electrical properties of epoxy resin based nano-compositesNanotechnology 18 025703

[25] Yogo T, Yamamoto T, Sakamoto W and Hirano S-I 2004In situ synthesis of nanocrystalline BaTiO3 particle–polymerhybrid J. Mater. Res. 19 3290–7

[26] Lu J and Wong C 2008 Recent advances in high-knanocomposite materials for embedded capacitorapplications IEEE Trans. Dielectr. Electr. Insul. 15 1322–8

[27] Su C, Hong B-Y and Tseng C-M 2004 Sol–gel preparation andphotocatalysis of titanium dioxide Catalysis Today 96119–26

[28] Jimmy C Y, Yu J, Ho W and Zhang L 2001 Preparation ofhighly photocatalytic active nano-sized TiO2 particles viaultrasonic irradiation Chem. Commun. 1942–3

[29] Huang W, Tang X, Wang Y, Koltypin Y and Gedanken A2000 Selective synthesis of anatase and rutile via ultrasoundirradiation Chem. Commun. 1415–6

[30] Zhu Y, Li H, Koltypin Y, Hacohen Y R and Gedanken A 2001Sonochemical synthesis of titania whiskers and nanotubesChem. Commun. 2616–7

[31] Guo W, Lin Z, Wang X and Song G 2003 Sonochemicalsynthesis of nanocrystalline TiO2 by hydrolysis of titaniumalkoxides Microelectron. Eng. 66 95–101

[32] Gao Y, Masuda Y, Peng Z, Yonezawa T and Koumoto K 2003Room temperature deposition of a TiO2 thin film fromaqueous peroxotitanate solution J. Mater. Chem. 13 608–13

[33] Masuda Y, Jinbo Y, Yonezawa T and Koumoto K 2002Templated site-selective deposition of titaniumdioxide on self-assembled monolayers Chem. Mater. 141236–41

[34] Erdem B, Hunsicker R A, Simmons G W, Sudol E D,Dimonie V L and El-Aasser M S 2001 XPS and FTIRsurface characterization of TiO2 particles used in polymerencapsulation Langmuir 17 2664–9

[35] Siuzdak D A, Start P R and Mauritz K A 2003 Surlyn®/titanatehybrid materials via polymer in situ sol–gel chemistryJ. Polym. Sci. B 41 11–22

[36] Cueto L F, Sánchez E, Torres-Martínez L M and Hirata G A2005 On the optical, structural, and morphologicalproperties of ZrO2 and TiO2 dip-coated thin films supportedon glass substrates Mater. Charact. 55 263–71

[37] Krunks M, Oja I, T"nsuaadu K, Es-Souni M, Gruselle M andNiinistö L 2005 Thermoanalytical study of acetylacetonate-modified titanium (IV) isopropoxide as a precursor for TiO2films J. Therm. Anal. Calorimetry 80 483–8

[38] Singha S and Thomas M J 2009 Influence of filler loading ondielectric properties of epoxy-ZnO nanocomposites IEEETrans. Dielectr. Electr. Insul. 16 531–42

[39] Lee W G, Woo S I, Kim J C, Choi S H and Oh K H 1994Preparation and properties of amorphous TiO2 thin films byplasma enhanced chemical vapor deposition Thin SolidFilms 237 105–11

[40] Zhou L, Hoffmann R C, Zhao Z, Bill J and Aldinger F 2008Chemical bath deposition of thin TiO2-anatase films fordielectric applications Thin Solid Films 516 7661–6

[41] Fleming R J, Pawlowski T, Ammala A, Casey P S andLawrence K A 2005 Electrical conductivity and spacecharge in LDPE containing TiO2 nanoparticles IEEE Trans.Dielectr. Electr. Insul. 12 745–53

[42] Mikhelashvili V and Eisenstein G 2001 Effects of annealingconditions on optical and electrical characteristics oftitanium dioxide films deposited by electron beamevaporation J. Appl. Phys. 89 3256–69

[43] Campbell S A, Gilmer D C, Xiao-Chuan W, Ming-Ta H,Hyeon-Seag K, Gladfelter W L and Jinhua Y 1997MOSFET transistors fabricated with high permitivity TiO2dielectrics IEEE Trans. Electron Devices 44 104–9

[44] Mirsaneh M, Furman E, Ryan J V, Lanagan M and Pantano C2010 Frequency dependent electrical measurements of

12

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies

amorphous GeSbSe chalcogenide thin films Appl. Phys. Lett.96 112907–3

[45] Almasri A, Ounaies Z, Kim Y S and Grunlan J 2008Characterization of solution-processed double-walled carbonnanotube/poly (vinylidene fluoride) nanocompositesMacromol. Mater. Eng. 293 123–31

[46] Jonscher A K 1999 Dielectric relaxation in solids J. Phys. D:Appl. Phys. 32 R57

[47] Park C, Wilkinson J, Banda S, Ounaies Z, Wise K E, Sauti G,Lillehei P T and Harrison J S 2006 Aligned single-wallcarbon nanotube polymer composites using an electric fieldJ. Polym. Sci. B 44 1751–62

[48] Takada T, Hayase Y and Tanaka Y 2008 Space charge trappingin electrical potential well caused by permanent and induceddipoles for LDPE/MgO nanocomposite IEEE Trans.Dielectr. Electr. Insul. 15 152–60

[49] Murakami Y, Nemoto M, Okuzumi S, Masuda S, Nagao M,Hozumi N and Sekiguchi Y 2008 DC conduction andelectrical breakdown of MgO/LDPE nanocomposite IEEETrans. Dielectr. Electr. Insul. 15 33–9

[50] Sencadas V, Lanceros-Méndez S, i Serra R S, Balado A A andRibelles J G 2012 Relaxation dynamics of poly(vinylidene fluoride) studied by dynamical mechanicalmeasurements and dielectric spectroscopy Eur. Phys. J. E 351–11

[51] Menczel J D and Prime R B 2009 Thermal Analysis ofPolymers, Fundamentals and Applications (New York:Wiley)

[52] Sy J W and Mijovic J 2000 Reorientational dynamics of poly(vinylidene fluoride)/poly (methyl methacrylate) blends bybroad-band dielectric relaxation spectroscopyMacromolecules 33 933–46

[53] Bello A, Laredo E and Grimau M 1999 Distribution ofrelaxation times from dielectric spectroscopy usingMonte–Carlo simulated annealing: application to α-PVDFPhys. Rev. B 60 12764

[54] Chanmal J P J C V 2008 Dielectric relaxations in PVDF/BaTiO3 nanocomposites Express Polym. Lett. 2 294–301

[55] Kochervinskii V, Malyshkina I, Gavrilova N, Sulyanov S andBessonova N 2007 Peculiarities of dielectric relaxation inpoly (vinylidene fluoride) with different thermal historyJ. Non-Cryst. Solids 353 4443–7

[56] 2005 IEEE Guide for the statistical analysis of electricalinsulation breakdown data IEEE Std 930-2004 (Revision ofIEEE Std 930-1987) 0_1-41

[57] Zhou T, Zha J-W, Cui R-Y, Fan B-H, Yuan J-K and Dang Z-M2011 Improving dielectric properties of BaTiO3/ferroelectricpolymer composites by employing surface hydroxylatedBaTiO3 nanoparticles ACS Appl. Mater. Interfaces 3 2184–8

[58] Nelson J K 2010 Dielectric Polymer Nanocomposites (Berlin:Springer)

[59] Danikas M and Tanaka T 2009 Nanocomposites-a review ofelectrical treeing and breakdown IEEE Electr. Insul. Mag.25 19–25

[60] Guo N, DiBenedetto S A, Tewari P, Lanagan M T,Ratner M A and Marks T J 2010 Nanoparticle, size, shape,and interfacial effects on leakage current density,permittivity, and breakdown strength of metal oxide!polyolefin nanocomposites: experiment and theory Chem.Mater. 22 1567–78

[61] Dou X, Liu X, Zhang Y, Feng H, Chen J-F and Du S 2009Improved dielectric strength of barium titanate-polyvinylidene fluoride nanocomposite Appl. Phys. Lett. 95132904–3

13

Smart Mater. Struct. 23 (2014) 104004 P Khodaparast and Z Ounaies