Composites: Part B xxx (2012) xxx–xxx

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Composites: Part B

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Thermal and mechanical characterization of poly(methyl methacrylate)nanocomposites filled with TiO2 nanorods

Niranjan Patra a,⇑, Marco Salerno a, P. Davide Cozzoli b, Alberto C. Barone a, Luca Ceseracciu a,Francesca Pignatelli a, Riccardo Carzino a, Lara Marini a, Athanassia Athanassiou a,b,c

a Istituto Italiano di Techologia, Via Morego 30, 16163 Genova, Italyb National Nanotechnology Laboratory, CNR – Istituto di Nanoscienze, Via Arnesano, 73100 Lecce, Italyc Center for Biomolecular Nanotechnologies@UniLe, Istituto Italiano di Techologia, Via Barsanti, 73010 Arnesano, Lecce, Italy

a r t i c l e i n f o a b s t r a c t

Article history:Received 10 October 2011Received in revised form 22 February 2012Accepted 18 April 2012Available online xxxx

Keywords:A. Nano-structuresA. Particle-reinforcementB. ThermomechanicalB. Thermal properties

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⇑ Corresponding author. Tel.: +39 010 71 781 756;E-mail address: patraji@gmail.com (N. Patra).

Please cite this article in press as: Patra N et al.nanorods. Composites: Part B (2012), http://dx.

Thick films of nanocomposites made of poly(methyl methacrylate) matrix and colloidal anatase TiO2

nanorods fillers were prepared by solvent mixing and solution drop casting. Different concentrationsof nanorods were tested in order to examine the influence of the nanoscale fillers on the compositesmaterial properties and structure. The thermal properties of the samples were investigated through ther-mogravimetric analysis, which showed an increase in thermal stability of the nanocomposites on increas-ing nanorods concentration, for the range of concentrations used. The viscoelastic properties wereinvestigated through dynamic mechanical analysis, which showed an increase in both the storage andloss modulus on increasing nanorods concentration. The in-depth distribution of the TiO2 nanorods inthe matrix was evaluated through cross-sectional transmission electron microscopy, which pointed outa uniform dispersion of mesoscale nanorods agglomerates with increasing diameter of 100–200 nm rangeon increasing nanorods concentration.

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1. Introduction

Over the last years hybrid nanocomposite materials became ofextraordinary interest for the scientific community because ofthe wide range of properties that can arise from the combinationof the peculiar characteristics of the employed nanoparticles(NPs) and polymeric matrices. Indeed, the size-dependent physicaland chemical properties of the inorganic NPs, along with the highprocessability, long-term durability and tunable chemical compo-sition and structure of the organic matrix, may result in materialswith unique characteristics that cannot be achieved by the individ-ual components alone [1]. Furthermore, when a high control of thelocal microstructural arrangement of the NPs in the polymer is ob-tained, the material properties can be tuned not only inside therange of those of the organic and inorganic constituents, but evennovel properties, not fully envisioned from the properties of thesingle components, may appear [2].

Compared to composites filled with microsized particles [3–5],nanocomposites have showed increased mechanical and rheologi-cal properties, reduced gas permeability, enhanced thermal stabil-ity, and self-extinguishing fire retardant characteristics [6,7]. Forexample, a twofold enhancement of the tensile modulus and of

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the heat distortion temperature, up to 100 �C, can be achieved fornanocomposites with as little as 2 wt.% of inorganic content [8].Recently, biodegradable polymer based nanocomposites have alsobeen developed [9,10], targeting applications such as cell growthand packaging.

Titanium dioxide (TiO2) based nanostructured materials haveemerged in the past decades as a platform on which a variety ofappealing physical–chemical properties coexist with biocompati-bility [4]. Presently investigated applications include photocata-lytic systems relying on controlled spatial organization of titaniapolymorphs [11], and light-responsive coatings with simultaneousantireflective, antibacterial, self-cleaning, and antifogging behavior[12–18]. Such nanocomposites, moreover, present in most casesimproved mechanical properties, mainly in terms of elastic modu-lus [19–22] or creep resistance [23].

The present work investigates the thermo-mechanical proper-ties of nanocomposites of poly(methyl methacrylate) (PMMA) ma-trix mixed with colloidal anatase TiO2 nanorods (NRs) prepared bysolution drop casting. The thermal and mechanical properties ofthe nanocomposites were characterized by means of thermogravi-metric analysis (TGA) and dynamic mechanical–thermal analysis(DMTA), respectively. The obtained results clearly demonstratethat the produced nanocomposites appear to be thermallymore stable and exhibit increased elastic modulus compared tothe pure polymeric matrix. The structural homogeneity of the

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Table 1Characteristic values of PMMA.

Characteristics of PMMA Values

Mol. wt (Mw) by GPC 120,000

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nanocomposites, which is crucial for the enhancement of thephysical properties of any composite material [24], is also investi-gated by transmission electron microscopy on cross sections of thesamples.

Total impurities 62.0%Refractive index N20/D1.49Transition temperature Tg 105 �CDensity 1.188 g/mL at 25 �C

2. Experimental

2.1. Materials

All chemicals were of the highest purity available and wereused as received. Titanium tetraisopropoxide (TTIP, 97%), titaniumtetrachloride (TiCl4, 99.999%), trimethylamine N-oxide dihydrate(TMAO, 98%), oleic acid (OLAC, 90%), 1-octadecene (ODE, 90%),oleyl amine (OLAM, 70%), and PMMA of Mw 120,000 were pur-chased from Sigma–Aldrich (Milan, Italy). The detailed characteris-tics properties of PMMA are given in Table 1. All solvents usedwere of analytical grade and were also purchased from Sigma–Al-drich. Used water was bi-distilled (Millipore Q).

2.2. Synthesis of colloidal TiO2 nanorods

OLAC capped anatase TiO2 NRs were synthesized by a slightlymodified literature protocol using a standard Schlenk line setup[25,26]. Prolate NRs with mean equatorial diameter of 3–4 nmand mean length of 25–30 nm were obtained by low temperatureTMAO-catalyzed hydrolysis of TTIP [25]. In a typical synthesis,15 mmol of TTIP dissolved in 70 g of degassed OLAC was reactedwith 5 mL of an aqueous 2 M TMAO solution at 100 �C for 72 h.The TiO2 NRs were separated from their growing mixture upon2-propanol addition, and were subsequently subjected to repeatedcycles of redissolution in chloroform and precipitation with ace-tone to wash out surfactant residuals. Finally, optically clear NRstock solutions in chloroform were prepared to be used as the fillerfor the nanocomposites. As a result of the process, the NRs OLACcapping was specifically constituted by oleate anions, bonded toTi atoms on the surface of TiO2 (R-COOA) in a bidentedconfiguration.

2.3. Preparation of the nanocomposites

Composite solutions of PMMA and TiO2 NRs were prepared byblending the chloroform solution of the polymer (200 mg/mL) withthe chloroform solution containing the NRs (16 mg/mL). Variousvolume amounts of TiO2 NRs solution were added to a fixed vol-ume of PMMA solution. The viscosity of the final solution was keptconstant for all concentrations by adding the necessary chloroformvolume. The nanocomposite samples are here mentioned asPMMAXNRs, with ‘X’ values of 2, 4, and 8, respectively, represent-ing the weight percent concentration of the NRs in the polymer.

For the DMA analysis, thick PMMA and nanocomposite filmswere prepared by drop casting the 2, 4, and 8 wt.% PMMA–NRssolutions onto clean Teflon sheets to prevent the sticking of thefilm onto the substrate. All the samples were dried in a vacuumoven at 90 �C for 15 h, in order to remove the residual chloroform.The thickness of the films was �0.10 mm, as measured by a digitalmicrometer (Mitutoyo, USA).

2.4. Thermal analysis

TGA measurements were carried out on a Q500 TA apparatus(TA Instruments, New Castle, USA), working in N2 atmosphere be-tween 30 �C and 500 �C, with a heating rate of 10 �C/min and a flowrate of 60 mL/min. The same samples prepared for the DMA mea-surements were also used for the TGA measurements. The amountof solid sample was approximately 8 mg in weight. From the TGA

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traces, differential thermogravimetric (DTG) plots were calculatedas the first derivative of the TGA curve.

2.5. Mechanical analysis

All the measurements were performed on a Q800 dynamicmechanical–thermal analyser (TA Instruments, New Castle, USA).The thickness and width of the films was �0.10 mm and�6.40 mm, respectively. The span length of the films was alwayskept constant (�12 mm). The tests were carried out in uniaxialtensile mode, applying a sinusoidal deformation with a frequencyof 1 Hz and an amplitude of 5 lm. The temperature was rampedfrom 30 �C to �150 �C at a rate of 3 �C/min. This rate was main-tained throughout all the test runs, so that there was a minimumtemperature lag between the sample and the furnace environment.Isothermal creep tests were also carried out in tensile mode at con-stant temperature, with an initial static load of 0.001 N at 80 �C andthen at constant stress of 0.2 MPa for 10 min.

2.6. TEM analysis

The samples for TEM imaging were prepared by transversallycutting the nanocomposites thick films by means of a EM UC6ultramicrotome (Leica Mikrosystems, Wetzlar, Germany). A cleandiamond knife with cutting edge of 45 reaction generating a radicalwas used to obtain sample cross-sections of �70 nm thickness atambient temperature of �25 �C, which subsequently were placedon a 150-mesh carbon coated copper grid. The TEM imaging wasperformed with a JEM 1011 instrument (JEOL, Japan), operated atan accelerating voltage of 100 kV.

3. Results and discussion

3.1. Thermal behavior

In Fig. 1 the results of the thermal analysis under inert N2 atmo-sphere for both bare PMMA and PMMA–NRs nanocomposites arereported. In case of pure PMMA film (black curve) first the removalof adsorbed water occurs, up to �110 �C, which is hardly visible infigure due to the compressed vertical scale. From Tonset � 110 �C upto Tend1 � 213 �C a first mass loss step is observed. This step is as-signed to removal of residual solvent from the casted solution,originally trapped inside the film during its drying out. In supportof the origin of the first mass loss step, an independent measure-ment performed on PMMA powder is also reported as a reference(gray curve). In this case the step at lowest temperature range doesnot appear, as the sample does not contain any solvent. At highertemperatures the PMMA degradation occurs, which can be sepa-rated in two major mass loss steps. Previous observations of otherauthors [27,28] assigned PMMA degradation to scission of weaklinks, whose presence is attributed to the polymerization tech-nique, and to random chain scission. The scission of weak linksshould also depend on the film thickness, as a consequence ofthe change in the diffusion time of the volatile radicals out of thepolymer film [27,28]. The first degradation step observed here

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occurs from Tend1 up to Tend2 � 323 �C, and can be attributedmainly to PMMA depolymerization to form MMA monomer. Anyreaction generating a radical (including b-scission) is expected todegrade a large number of polymer chains by chain-transferprocess with monomer formation. The second step, from Tend2 uptoTend3 � 455 �C, is attributed to random scissions within thepolymer chain and monomer volatilization.

A similar three-step process is also observed for the nanocom-posite samples, colored1 curves in Fig. 1. Again, the first mass lossstep is assigned to residual solvent removal, and the second andthird steps are the actual PMMA degradation reactions. As expected,the final amount of residual mass observed at 500 �C is approxi-mately proportional to the TiO2 NRs concentration, as shown onthe right side of Fig. 1 after the horizontal axis break, where an ex-panded vertical scale is used. Actually, due to the NR surfactantthe residual mass is approximately 61% of the NR concentration, as�39% is estimated to be the mass of the OLAC capping molecules re-moved during the process.

The DTG thermograms, Fig. 2, give more evident indication ofthe change in both position and intensity of the maxima of massloss rate for the different steps, DTGmax i (with i = 1 to 3). In thenanocomposites the position of the first two maxima shifts to-wards higher temperatures, by �25 �C for the first step and�10 �C for the second step, whereas the position of the third max-imum remains approximately the same. The observed shifts to-wards higher temperatures on increasing NRs concentration (seealso Fig. 2 inset) are qualitatively consistent with the increase inglass transition temperature presented in a previous work of ourgroup [14]. The slight shift of the third maximum toward lowertemperature (��3 �C only) can be explained by the presence ofOLAC surfactants, which deteriorate above �380 �C.

On the other hand, the peak intensity (i.e. the maximum rate) isdecreased with respect to PMMA for the first two mass loss steps,whereas it is increased for the third step. This reversed behavior ofthe third step with respect to the previous two is expected, since atvery high temperatures the complete thermal degradation of theorganic content should eventually occur. The shift towards highertemperatures of the peak position and the decrease in peak inten-sity of the first two mass loss steps describe an increased thermalstability in the nanocomposites, which can be partly justified bythe decreasing polymer content in the samples with increasingNRs concentration. However, this is not the only reason for the ob-served effect, since it can be noticed that the changes do not scalewith the NRs concentration. We assume that the higher thermalstability of the nanocomposites with respect to bare PMMA at tem-peratures below 325 �C are due to the presence of nanofillers,which might limit the motion of polymer segments, limiting theinteraction of their end groups with the produced free radicals.In this way chain transfer reactions are suppressed and depolymer-ization is limited. Since oxide TiO2 particles are highly thermallystable, we believe that they act as a barrier to heat, preventingthe nanocomposites to degrade soon, resulting in an increase inthermal stability of the nanocomposites.

3.2. Mechanical behavior

DMTA analysis results are usually expressed through the dy-namic modulus components, i.e. the storage modulus (E0), whichdescribes the elastic response to the deformation, and the lossmodulus (E00), which is related to the viscous response. DMTA anal-ysis was carried out to measure the temperature dependence of

1 For interpretation of color in Figs. 1–4 and 6, the reader is referred to the webversion of this article.

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storage modulus (E0) and loss modulus (E00) of bare PMMA andTiO2 nanocomposites, as shown in Fig. 3.

Fig. 3a and b show the variation with the temperature of thestorage and loss moduli of the nanocomposites measured at a fre-quency of 1 Hz, for different filler concentrations. It can be seenthat all the curves of the nanocomposites (colored lines) lie abovethe curve of the bare PMMA (black line). At a temperature of 30 �C,the storage modulus and the loss modulus increments range be-tween 30% to 45%, and 15% to 40%, respectively, whereas in bothcases the highest values are recorded for the lowest concentration(2 wt.%). These increments in moduli fall within the same order ofmagnitude of similar nanocomposite materials [29]. The significantimprovement in both storage and loss modulus is due to the PMMAchains becoming stiffer due to the incorporation of TiO2 particleson the PMMA, restricting the chain movement. In all the nanocom-posites, the 2 wt.% NRs shows the highest improvement in proper-ties. This is also due to the fact that on increasing the NRsconcentration the aggregates size is also increasing because ofthe mutual interaction of the nanoparticles trying to hinder them-selves inside the polymer matrix. Actually, upon increasing the NRsconcentration the amount of OLAC surfactant used to prevent NRsaggregation is also increasing, which is concurring in decreasingthe Tg of the materials [30]. Actually, both storage and loss modu-lus decrease drastically at around �100 �C, which means the mate-rials undergoes in the Tg region. Furthermore, it is clear from Fig. 3bthat the height of the loss modulus peak decreases and the curvebroadens with increasing TiO2 NRs content close to the glass tran-sition region. A more detailed insight in the observed differencesbetween different NRs concentrations can be given after consider-ing the TEM analysis (see Section 3.3).

In Fig. 3b, the peaks of loss modulus can be associated to the Tg

of the materials. A little change in the peak position as a function ofthe particles concentration is observed for 4 and 8 wt.%, which maybe due to the increase in crosslinking density. Basically, interac-tions between the particles and the matrix are mainly physical,while the interface does not confer mobility to the chains in thestudied concentrations [31,32].

Creep compliance values are shown in Fig. 4. An improvement,i.e. lower creep compliance, of all the nanocomposites with respectto the bare PMMA is observed. In this case the improvement isroughly proportional to the TiO2 NRs concentration, which sug-gests that it is related mainly to the higher viscosity induced bythe fillers. The reduction in compliance ranges from a factor 0.8for 2 wt.% concentration to 0.4 for 8 wt.% in agreement with otherstudies made both in similar stress conditions [29,31] and in high-er stress or longer time conditions [23].

3.3. Transmission electron microscopy

TEM analysis of samples cross-sections can provide importantinsight concerning the quality of the NRs dispersion in the PMMAmatrix. Fig. 5a shows TEM images of bare TiO2 NRs and Fig. 5b–dshowing cross-sections of PMMA–NRs nanocomposites samplescontaining different NRs concentrations.

The images reported in Fig. 5b–d are representative of the mor-phology of the nanocomposites cross-sections. The efficiency of theTiO2 NRs in modifying the properties of the PMMA is primarilydetermined by the degree of its dispersion. The dark lines or spotsin the micrographs are TiO2 NRs aggregates. Aggregates of TiO2 NRsappear already at the lowest NRs concentration of 2 wt.%. Never-theless, the aggregates are uniformly distributed for all the fillerconcentrations, resulting in rather homogeneous nanocompositesamples at the considered length scale.

The analysis of the aggregate domains is reported in Fig. 6. Boththe surface density of the aggregates (‘grains’) and their meandiameter are plot (in blue and red, respectively). Single aggregate

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Fig. 1. Thermogravimetric analysis of PMMA and PMMA–TiO2 nanocomposites inN2 atmosphere.

Fig. 2. Differential thermogravimetric analysis of PMMA and PMMA–TiO2 nano-composites in N2 atmosphere.

Fig. 4. Typical creep tests of bare and loaded PMMA at 80 �C.

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domains at 2 wt.% NRs concentration (Fig. 5b) have mean diameterof �100 nm, and undergo a significant increase in size withincreasing NRs concentration from 2 to 4 wt.% (mean domaindiameter �170 nm in Fig. 5c), tending to a saturation at 8% concen-tration (diameter �200 nm in Fig. 5d). The large error bars for theaggregate diameter represent the widespread distributions in therespective aggregate populations. In the same graph, the wholehistograms are also shown, in red as the respective (mean) datapoints. The histogram widths are at the scale with the axis values,and have all single bin width of 15 nm. From the histograms it isclear that the distribution of aggregates diameter changes frommonomodal at 2% to bimodal already at 4% and even more at 8%.However, the increase in overall mean diameter is still significant.

Concurrently, the density of aggregate domains is decreasing,starting from �3.4 lm�2 in Fig. 5b, and changing to �1.4 lm�2 inFig. 5c and �1.3 lm�2 in Fig. 5d. Obviously, both grain density

Fig. 3. Typical DMTA curves of (a) storage modulus E0 and (b) loss modulus E0 0 of PMMA–

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and particles diameter tend to saturation at 8%, and this trend isespecially clear for the grain density, which remains constant be-tween 4% and 8% NRs concentration. Associating this analysis withthe mechanical characterization results, it can be concluded thatthe aggregate domains density, rather than the nanoparticles con-centration, is the key parameter for the improvement of both elas-tic and damping properties of the nanocomposites, whereasimprovements in the short term creep compliance is a bulk effectclosely related to the filler concentration.

The uniform mesoscale aggregate distribution resulting fromthe TEM images is also in agreement with the improvement inthermal stability of nanocomposites observed previously with re-spect to the pure polymer [33]. Actually, the uniformly distributedinorganic aggregates may prevent quick heat diffusion into thepolymer matrix limiting further material degradation [34].

In a previous work it has been demonstrated that the in-depthNRs distribution is not uniform and the NRs rather tend to accumu-late close to the film surface when the sample preparation is doneby the spin coating technique. Instead, the use of the drop castingtechnique herein seems to lead to a homogeneous distribution ofaggregates of NRs throughout the sample’s depth, probably dueto a different dynamic of PMMA solution drying.

4. Conclusions

Nanocomposites of PMMA incorporating TiO2 NRs prepared bydrop casting were investigated in terms of thermal and dynamicmechanical behavior. Both the thermal stability and the elasticmodulus appear to be increased in the nanocomposites, up to

TiO2 nanocomposites with different concentration as a function of the temperature.

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Fig. 5. TEM images of (a) bare TiO2 NRs (magnification 35,000X), and cross-sections of PMMA–NRs nanocomposites with different concentration: (b) 2, (c) 4, and (d) 8 wt.%,(magnification 750�).

Fig. 6. Results of the aggregate domains analysis performed on the TEM imagesfrom Fig. 5. The plot shows the number of aggregates per unit area and theaggregate mean diameter along with the respective size distributions. The splinelines are just guides to the eye.

N. Patra et al. / Composites: Part B xxx (2012) xxx–xxx 5

8 wt.% NRs concentration compared to the pure polymer. Electronmicroscopy analysis of the films cross-sections showed occur-rences of uniformly distributed NRs aggregates in the polymer ma-trix, whose size scales approximately with the NRs concentration.The observed improvement of the thermo-mechanical propertiesof the materials can be ascribed to the good dispersion of theseaggregates. In particular, the enhancement in storage (elastic)and loss (viscous) modulus is attributed to chains interlocking,which depends mainly on the aggregates density, i.e. on the distri-bution, whereas the instantaneous viscosity, measured on shorttime creep tests, is roughly proportional to the filler concentration.The thermal properties enhancement could arise from the limit inthe motion of polymer segments due to the nanofillers. Therefore,the end groups of the polymer chains become less reactive with thefree radicals suppressing the depolymerization.

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cterization of poly(methyl methacrylate) nanocomposites filled with TiO2

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