Materials Science and Engineering C 32 (2012) 2396–2403

Contents lists available at SciVerse ScienceDirect

Materials Science and Engineering C

j ourna l homepage: www.e lsev ie r .com/ locate /msec

The role of oleate-functionalized layered double hydroxide in the melt compoundingof polypropylene nanocomposites

Ricardo K. Donato a, Leandro Luza a, Renato F. da Silva a, Celso C. Moro b, Rafael Guzatto c, Dimitrios Samios c,Libor Matějka d, Bojan Dimzoski d, Sandro C. Amico e, Henri S. Schrekker a,⁎a Laboratory of Technological Processes and Catalysis, Institute of Chemistry, Universidade Federal do Rio Grande do Sul, Av. Bento Gonçalves 9500, 91501–970, Porto Alegre, RS, Brazilb Institute of Chemistry, Universidade Federal do Rio Grande do Sul, Av. Bento Gonçalves 9500, 91501–970, Porto Alegre, RS, Brazilc Laboratory of Instrumentation and Molecular Dynamics, Institute of Chemistry, Universidade Federal do Rio Grande do Sul, Av. Bento Gonçalves 9500, 91501–970, Porto Alegre, RS, Brazild Institute of Macromolecular Chemistry, Academy of Sciences of the Czech Republic, Heyrovsky Sq. 2, Prague 6, 162 06, Czech Republice Laboratory of Polymeric Materials, Department of Materials, Universidade Federal do Rio Grande do Sul, Av. Bento Gonçalves 9500, 91501–970, Porto Alegre, RS, Brazil

⁎ Corresponding author. Tel.: +55 51 33086284; fax:E-mail addresses: ricardokeitel@iq.ufrgs.br (R.K. Don

(L. Luza), renato.figueira@ufrgs.br (R.F. da Silva), celso@guzatto@gmail.com (R. Guzatto), dsamios@iq.ufrgs.br (D(L. Matějka), dimzoski@imc.cas.cz (B. Dimzoski), amicohenri.schrekker@ufrgs.br (H.S. Schrekker).

0928-4931/$ – see front matter © 2012 Elsevier B.V. Alldoi:10.1016/j.msec.2012.07.013

a b s t r a c t

a r t i c l e i n f o

Article history:Received 14 July 2011Received in revised form 22 May 2012Accepted 4 July 2012Available online 16 July 2012

Keywords:Renewable feedstock functionalizednanofillerPolymer-matrix nanocompositeMelt compoundingThermoplastic resinLayered double hydroxideFatty acid based compatibilizer

In this research, the oleate-functionalized magnesium and aluminum layered double hydroxide (LDH; Mg:Al=3:1) o-LDH was applied as nanofiller in the melt blending of polypropylene (PP) nanocomposites, inorder to understand its role in this process. o-LDHwas prepared using the memory effect of the calcined car-bonated LDH. Blending of PP and low o-LDH filler contents of 0.45 and 0.90 wt.% afforded thenanocomposites PP0.45 and PP0.90, respectively, which were characterized by transmission electron mi-croscopy, X-ray diffraction, small angle X-ray scattering, thermo-gravimetric analysis, differential scanningcalorimetry and dynamic mechanical analysis. The oleate LDH surface functionalization enhanced the systemcompatibility as a relative regular dispersion of o-LDH tactoids was observed within the matrix, togetherwith partial PP intercalation. This o-LDH incorporation increased the PP relative crystallinity, induced crystal-line orientation and decreased the glass transition temperature. Furthermore, the nanocomposites showedimproved initial resistance to decomposition and stiffness. These results showed that the o-LDH acted asboth nucleating agent and plasticizer, and that the presented approach can be used for the development ofPP nanocomposites with distinguished properties.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

The interest in the application of layered structured clays as nanofillersfor polymer nanocompositeswas triggered by the Toyota report involvingthe preparation of nylon/cationic clay nanocomposites [1,2]. A high de-gree of cationic clay exfoliation and dispersion was necessary for thesenylon nanocomposites to achieve expressive mechanical property im-provements, which did only occur after organo-functionalization of thenanofiller. Until recently, layered double hydroxides (LDH), also knownas anionic clays and hydrotalcites, have attracted comparatively little at-tention for this application, which can be attributed to the natural abun-dance of cationic clays [3].

The LDH crystalline structure consists of layersmade up of octahedralunits with shared borders [4,5]. Each octahedron consists of MII and MIII

cations surrounded by six hydroxyl anions. The presence of MIII gener-ates positive charge excess in the lamellae, which is balanced by

+55 51 33087304.ato), leandro.luza@ufrgs.briq.ufrgs.br (C.C. Moro),. Samios), matejka@imc.cas.cz@ufrgs.br (S.C. Amico),

rights reserved.

exchangeable anions (An−) in the interlamellar region. These clays canbe represented by the general formula [MII

1−x MIIIx. (OH)2]x+Ax

n−/n·mH2O. The most well known LDH is the naturally occurring HydrotalciteMg6All2(OH)16CO3·4H2O.

Polypropylene (PP) is a non-polar and semicrystalline polymer, whichexhibits an attractive combination of low cost and extraordinary versatil-ity of processability, properties, applications and recycling. The physicaland mechanical properties of PP are dependent on the degree of crystal-linity, which influences, i.e., its processability [6]. Incorporation of claynanofillers in the PP matrix has been investigated intensively to obtainmaterialswith improvedmechanical, thermal,fire and gas barrier proper-ties [7]. In the case of nanocomposites, the low PP polarity difficults inter-calation into the polar LDH galleries and homogeneous LDH dispersion.This incompatibility might be overcome by LDH surface modification toturn it similarly hydrophobic, which can be achieved by using the socalled LDH “memory effect”; the ability to recover the original lamellarstructuring after calcination induced complete loss of its original layeredorganization [8]. This allows the incorporation of different anionic speciesin the interlamellar LDH region during the process of structural recovery.The application of organic anions with expressive non-polar groups, suchas alkyl-carboxylates, ‐sulfates and -sulfonates can render the hydrophilicLDH surfaces hydrophobic, which is generally accompanied by expansion

2397R.K. Donato et al. / Materials Science and Engineering C 32 (2012) 2396–2403

of the interlamellar distance [9,10]. Both effects are expected to promotethe preparation of non-polar polyolefin nanocomposites with improvedintercalation and LDH exfoliation and dispersion; however, this has notalways been observed and, especially, exfoliation remains a challenge[11–13].

The application of a renewable source for LDH surface function-alization is especially attractive,which canbe achievedbyusing the oleateorganic anion [14,15]. Previously, oleate-functionalized LDH has been ap-plied in the preparation of polymer nanocomposites, using polyethylene,poly(ethylene-co-butyl acrylate), poly(methylmethacrylate), polypropyl-ene (PP) and PP/poly(propylene-graft-maleic anhydride) mixtures asmatrices [16,17]. Only in the cases of poly(methylmethacrylate) andPP/poly(propylene-graft-maleic anhydride) good LDH dispersion andnanocomposite formation was achieved, which promoted, i.e., im-proved fire properties. The absence of this improvement in the case ofPP nanocomposites was ascribed to poor dispersion of the oleate-functionalized LDH, however, structural, morphological, thermal anddynamic mechanical studies were not provided for these materials.

In this paper, we describe themelt compounding and characterizationof PP/oleate-functionalized LDH nanocomposites. The oleate-modifiedLDH (Mg:Al=3:1) o-LDH was obtained using the clay “memory effect”,forcing the confinement of the oleate anion within the LDH galleries.Wepresent a completeX-ray (fromvery small up towide angles) andmi-croscopy study for a structural and morphological characterization of theobtained o-LDH and PP nanocom-posites, together with thermal and dy-namicmechanical analyses. This provided detailed information about therole of the oleate compatibilizer and the effect of o-LDH on the PP crystal-linity, orientation parameters, thermal resistance to decomposition, glasstransition temperature and stiffness. Fig. 1 shows the general idea of thepursued strategy.

2. Materials and methods

2.1. Materials

A highly isotactic PP (trade name H-503, MFI=3.5 g/10 min,Mw=470,000 g/mol) was used as provided by Braskem S.A. Oleic

Fig. 1. Generalized representation of the pursued strategy: LDH oleate-func

acid (PA), anhydrous sodium carbonate (PA), sodium hydroxide (PA),aluminum nitrate nonahydrate (PA) and magnesium nitrate hexahy-drate (PA) were used as purchased from VETEC Química Fina LTDA.

2.2. Methods

2.2.1. Preparation of oleate-functionalized LDHMethodologies available in the literature were used for the synthesis

of sodium oleate [18] and the carbonated layered double hydroxideLDHwith a Mg:Al ratio of 3:1 [5]. X-ray fluorescence and loss on ignitionanalyses indicated the following LDH contents in wt.%: Al2O3=16.9;MgO=40.0 and CO2+H2O=43.1. This in agreement with a 3:1 ratio ofMg:Al. The carbonated LDH was calcined at 450 °C for 3 h and theresulting oxide was characterized by XRD (Fig. 2; curve b), which con-firmed the loss of the lamellar structure. The memory effect of LDH wasused for the preparation of the oleate-functionalized LDH o-LDH with aMg:Al ratio of 3:1, using calcined LDH and sodium oleate [15].

2.2.2. Melt-compounding of PP/oleate-functionalized LDH nanocompositeso-LDH was dried at 80 °C for at least 5 h prior to processing.

After this time, PP and o-LDH (0.45 (PP0.45) or 0.90 wt.%(PP0.90) based on the inorganic weight of ~30 wt.% as deter-mined by TGA at 800 °C) were melt blended in a pre-heated(180 °C) twin rollers mill (Haake Rheomix 600) at 60 rpm. Therollers were activated, PP (45.0 g) was molten during 5 min, ando-LDH was added. The blend was mixed and heated for another5 min at 180 °C. For comparison reasons, virgin PP was alsoprocessed under the same conditions. The samples were cooledat room temperature. Films for characterization were obtainedby compression molding of PP, PP0.45 and PP0.90 for 5 min at190 °C and a pressure of 70 lb. The samples were removed fromthe mold and cooled at room temperature.

2.2.3. Characterization

2.2.3.1. X-ray diffraction (XRD).XRD diffractogramswere recorded using acopper (λ=1.5406 Å) radiation on a Siemens D500 diffractometer in

tionalization for improved compatibility with the non-polar PP matrix.

Fig. 2. XRD diffractograms of (a) LDH (black line), (b) calcined LDH (blue line) and(c) o-LDH (red line).

2398 R.K. Donato et al. / Materials Science and Engineering C 32 (2012) 2396–2403

Bragg–Brentano geometry, equipped with graphite monochromator andoperating at 17.5 mA and 40 kV. The diffraction patterns were obtainedfor 2θ between 1.0 and 70° (LDH, calcined LDH and o-LDH) or between1.0 and45° (PP,PP0.45 andPP0.90), using a constant step of 0.05° per 2 s.The interplanar distances (d) of LDH and o-LDHwere calculated from thebasal peak by using the Bragg equation: nλ=2dsinθ, where n is the dif-fraction order (n=1), λ is the wavelength of the X-ray, d is the averageinterplanar distance and θ is the incident diffraction angle in relation tothe sample surface.

Orientation parameters (Ahkl/jnm) of PP,PP0.45 andPP0.90: The ori-entation degree ω of the crystalline planes of PP and PP nanocompositeswas determined by XRD by measuring the ratio of total lines per unitarea, LA, at various angles from the orientation vs. the total line per unitarea on the orientation axis. The result is a linear density NL for eachangle. For an oriented system,NLwill have different values at different an-gles away from the orientation axis. The equation used for this measure-ment results in ω=(100·NL −NL )/(NL +0.571·NL ), where NL is thenumber of intercepts per unit length observed in a direction perpendicu-lar to the preferred microelements orientation and NL represents thenumber of intercepts per unit length parallel to that direction [19].

Relative crystallinity index of PP,PP0.45 andPP0.90: The relative crys-tallinity index was determined by XRD. The amount of amorphous phasein the total diffractogram was defined by drawing a baseline. The amor-phous halo was subtracted from the original diffractogram and the crys-talline reflection lines were treated. Each peak of each crystal plane wasintegrated separately; summed and the contribution of the crystallineand amorphous parts was obtained. From these results, the relative crys-tallinity index of the PP and PP nanocomposites was calculated as a ratioof crystalline area to total area, by the following equation: Xc=Ac/(Aa+Ac), where Ac and Aa are the total crystalline and amorphous areasof the diffractogram and Xc is the relative crystallinity index [19].

2.2.3.2. Small-angle X-ray scattering (SAXS). Two different sets of equip-ments were used for a full range of analyses of SAXS spectra; whereFig. 6a is a plot from better-resolved Porod's region and Fig. 6b is a plotfrom better-resolved Guinier's region, acquired on equipments (1) and(2), respectively.

(1) The experiments were performed using a pinhole camera (Molec-ular Metrology SAXS System) attached to a microfocused X-raybeam generator (Osmic MicroMax 002) operating at 45 kV and0.66 mA (30 W). The camera was equipped with a multiwire,gas-filled area detector with an active area diameter of 20 cm(Gabriel design). Two experimental setups were used to coverthe q range of 0.004–1.1 Å−1 where q=(4π /λ)sinθ (λ is thewavelength and 2θ is the scattering angle). The scattering

intensities were put on absolute scale using a glassy carbon stan-dard. o-LDHwas measured in a 1.5 mm diameter sealed capillaryand the nanocomposites as thin films.

(2) These analyses were performed on a high-energy synchro-tron facility of the Brazilian Synchrotron Light LaboratoryLNLS (Campinas-SP, Brazil) in transmission geometry withλ=1.608 Å. The exposure time for each sample was 300 s andthe distance between the sample and the plane detection was1581 mm. o-LDH was measured as powder and PP and thenanocomposites as thin films. A FIT2D software was used to nor-malize and convert the images area of the detector in scattering in-tensity curves versus wave vector (q in nm−1), i.e., I(q) versus q.The scattering SAXS curves analyses were normalized by Guinierapproximation [20]. The lamellar long period (L) may be obtainedfrom the scattering angle (θ) at maximum scattered intensity(qmax) using the equation 2Lsinθ0=nλ, where λ is the wavelengthof the X-ray and n is the diffraction order (n=1).

2.2.3.3. Thermo-gravimetric analysis (TGA). TGA was carried out on a TAInstruments Q50 thermo-gravimetric analyzer. An average sampleweight of 8–12 mg was placed in a platinum pan and heated from ca.30 to 800 °C at 20 °C/min under a flow of nitrogen.

2.2.3.4. Differential scanning calorimetry (DSC). Equilibrium thermody-namic parameters were determined using a TA Instruments DSC 2010differential scanning calorimeter under nitrogen atmosphere. An averagesample weight of 5–10 mgwas sealed in an aluminium pan. The sampleswere heated and cooled between 30 and 200 °C at a heating and coolingrate of 10 °C/min. Crystallization temperature (Tc, determined at themaximum of the exothermic peak) and melting point (Tm, determinedat themaximum of the endothermic peak) were estimated using the sec-ond heating and the first cooling ramp, respectively.

2.2.3.5. Scanning electron microscopy (SEM). SEM analyses of LDH ando-LDH were performed on a microscope Vega Plus TS 5135 (Tescan).The powder samples were fixed to a copper support with a conductivecarbon double-adhesive tape and sputtered with 4 nm of platinum in avacuum sputter coater SCD 050 (Balzers) prior to analysis. The speci-mens were observed using accelerating voltage 30 kV and secondaryelectron detector.

2.2.3.6. Transmission electronmicroscopy (TEM). TEMwas performedwitha microscope Tecnai G2 Spirit Twin (FEI). Ultrathin sections for TEM(thickness ~50 nm) were prepared with an ultramicrotome UltracutUCT (Leica) under cryo-conditions: The temperatures during the cuttingprocedure were −80 and −50 °C for the sample and the knife, respec-tively. The ultrathin sections (50 nm thickness) were transferred ontomicroscopic support grids and covered with thin carbon layer (thicknessof the layer: ~4 nm; preparation: vacuum evaporation device JEE-4C,Jeol) in order to limit sample damage under the electron beam. TEMmi-crographs were obtained at accelerating voltage 120 kV, using brightfield imaging.

2.2.3.7. Dynamic mechanical analysis (DMA). These analyses were carriedout on a TA instruments Q 800. The storage modulus, stiffness and tan δwere studied between−30 and 130 °C, using a heating rate of 3 °C/min.

2.2.3.8. X-ray fluorescence (XRF). Themagnesium and aluminum contentsof LDH were determined by XRF analysis that was performed on aRIGAKU RIX 2000 X-ray fluorescence spectrometer, equippedwith a rho-dium tube with refrigerated anode, a LiF 200 analyzing crystal, six ana-lyzing crystals for wavelength dispersion of the secondary X-rayspectrum (LiF200, LiF 220, Ge, TAP, PET and RX 35) and both flow-proportional and scintillation counter tubes, operating at 50 kV and50 mA. LDH (1.000 g) was melt with 7.000 g Flux (Li2B4O7) in

Fig. 3. SEM micrographs of (a) o-LDH and (b) LDH: Scale bar=2 μm.

Fig. 4. TGA trace of o-LDH.

2399R.K. Donato et al. / Materials Science and Engineering C 32 (2012) 2396–2403

Pt-crucible at 1.200 °C for 5 min. The fused beads were analyzedaccording to the measurement conditions listed above.

2.2.3.9. Loss on ignition (LOI). The water and carbon dioxide content ofLDHwas determined by LOI analysis, which was carried out in an EDG3P-S muffle furnace with digital temperature display and thermostat-ic temperature control at 1000 °C for 2 h. Empty crucibles and LDH(1.000 g) were dried at 105 °C for 2 h and cooled to room tempera-ture in a desiccator before any measurements were made. To avoidoverheating, crucibles were put into the furnace only after a constanttemperature was reached.

3. Results and discussion

3.1. Structural, morphological and thermal analysis of LDH and o-LDH

TheMg:Al ratio of 3:1was confirmed for the carbonated LDH byXRFanalysis. Fig. 2 shows the XRD patterns of LDH, calcined LDH ando-LDH. LDH exhibited the characteristic pattern of a layered solidwith sharp, symmetrical peaks at low angles, corresponding to thebasal reflection and higher order reflections, together with some asym-metrical peaks at higher angles. All these peaks are characteristic ofrather well crystallized layered material with the hydrotalcite-typestructure, showing the characteristic planes (003), (006) and (009) at2θ 11.3°, 22.8° and34.1°, respectively. The basal peak (003) correspondsto an interlamellar distance of 0.78 nm, which is expected for LDHwiththe CO3

2− anion [10]. Intercalation of the oleate anion within the LDHgalleries by the ‘memory effect’ procedure dramatically changed theXRD diffractogram. Now, the XRD pattern of o-LDH shows a basalpeak (003) with maximum intensity at 2θ=1.9°, which indicates aninterlamellar distance of 4.7 nm, together with another novel peak at2θ=4.5º (d-spacing=2.10 nm). These observations are in agreementwith the successful intercalation of the oleate anion [15]. Assuming athickness of 0.48 nm for the LDH layers [10], the o-LDH gallery heightwas 4.2 nm, which suggests that the oleate ions were intercalated in abilayer- or micellar-type organization [14,15].

The SEM micrographs of LDH and o-LDH are shown in Fig. 3. Sur-face irregularities with secondary growth of small layers are presentin LDH (Fig. 3b). The particle surface of o-LDH displayed a defined la-mellar structure and platelet surface texture (Fig. 3a), which con-firmed that the original lamellar organization was recovered by theoleate induced restructuring memory of LDH.

The TGA trace ofo-LDH is shown in Fig. 4, which reveals similar stagesof weight loss as previously reported for MgZnAl oleate-functionalizedLDH [16]: (1) The events below 250 °C were relative to water adsorbedat the LDH surface and water entrapped within the LDH galleries. (2) Ex-pressive weight loss between 250 and 600 °C due to principally decom-position of the oleate anion and simultaneous dehydroxylation of theLDH layer structure. Neglecting the loss due to LDH dehydroxylation,this corresponds to an oleate content of ~50 wt.%. Continuousdehydroxylation resulted in the o-LDH residual inorganic mass of~30 wt.% at 800 °C, which was used to define the filler content of thePP nanocomposites (0.45 and 0.90 wt.%o-LDH forPP0.45 andPP0.90, re-spectively). The same criterion was used byWilkie et al. for the prepara-tion of PP nanocomposites with 1, 2 and 4 wt.% of oleate-functionalizedLDH [16], showing that higher nanofiller contents were applied in theirstudy.

3.2. Melt compounding of PP/o-LDH nanocomposites

3.2.1. Structural and morphological analysisTEM micrographs of the PP nanocomposites are shown in Fig. 5. The

areas of low contrast represent PP and the areas of strong contrast,black spots, are due to o-LDH. The image of PP0.45 shows amoderate de-gree of o-LDH dispersionwithin the PPmatrix and the presence of o-LDHtactoids with different sizes and compactness (Fig. 5a). The presence of

tactoidswith lower compactness indicates a better compatibility betweenthe PP matrix and these o-LDH particles. A better permeation of the PPchains can be observed in these lower compactness tactoids, which is ex-emplified by a magnification of an individual tactoid of PP0.90 in Fig. 5b.These observations confirm the statement ofWilkie et al. that the low fireresistance performance of their PP/oleate-LDH composite systems wasconsistent with a poor exfoliation and dispersion of the LDH nanofillerin the PP matrix [16]. In general, LDH surface functionalization with hy-drophobic organic groups like, i.e., oleate, undecenoate and palmitate ap-pears to be insufficient for goodnanocomposite formationwith non-polarpolymers and additional strategies are required to increase the matrix-filler compatibility [12,13,16,17].

2400 R.K. Donato et al. / Materials Science and Engineering C 32 (2012) 2396–2403

A detailed X-ray study was undertaken in an attempt obtaining fur-ther information about the role of o-LDH in the preparation of PPnanocomposites. The SAXS and XRD patterns of o-LDH, PP and the PPnanocomposites are shown in Figs. 6 and 7, respectively. The o-LDHpeaks observed in Fig. 6a (curve a) at q~0.35 and 0.15 (1/Å) correspondto interlamellar distances of 2.0 and 4.8 nm, respectively, relating to thealready described XRD peaks at 2θ~4.5 and 1.9° (Fig. 2, curve c). Bothpeaks were observed for the nanocomposites, where the intensity ofthose peaks increased with increasing o-LDH content (Fig. 6a, curves band c), although the intensity of the peak at q~0.35 was much less in-tense. The correlating XRD peak at 2θ~4.5 was not visible for the PPnanocomposites (Fig. 7b, curves b and c). Both SAXS and XRD observa-tions imply loss of structural organization of the o-LDH filler tactoids,which was similarly observed for the oleate-LDH filler applied in thepreparation of polyethylene nanocomposites [17]. A third peak, not pres-ent in the o-LDH curve, that was also more intense with a higher o-LDHcontent canbeobserved at q~0.04, indicating a15.4 nm interlamellar dis-tance. This peak correlates precisely with the one from the PP lamellarlong period presented in Fig. 6b, implying the intercalation of PP chains

Fig. 5. TEM micrographs of (a) PP0.45 and (b) PP0.90: scale bar=1 μm.

within the o-LDH galleries and supporting the role of the oleate anionas compatibilizer.

The patterns presented in Fig. 6b show symmetrical peak shapesthat indicate a narrow distribution of the lamellar long period for PPand the PP nanocomposites. Furthermore, the obtained values of the la-mellar long period, given in Table 1, are approximately the same,suggesting no strong change in the crystallinity due to the addition ofo-LDH. On the other hand, evaluation of the XRD patterns, presented in

Fig. 6. a. SAXS diffractograms of (a) o-LDH (red line), (b) PP0.45 (green line) and (c)PP0.90(black line). b. SAXS diffractograms of (a) o-LDH, (b) PP0.90, (c) PP0.45 and (d) PP.

Fig. 8. TGA and DTA traces of (a) PP (red line), (b) PP0.45 (green line) and (c) PP0.90(black line).

Fig. 7. XRD diffractograms of (a) PP (red line), (b) PP0.45 (green line), (c) PP0.90(black line) and (d) o-LDH (blue line): (Fig. 7a) 2θ: 1.0–45°; (Fig. 7b) 2θ: 1.0–6.0°.

2401R.K. Donato et al. / Materials Science and Engineering C 32 (2012) 2396–2403

Fig. 7, shows strong changes in the intensity values of the different reflec-tive planes. The orientation parameters Ahkl/jnm defined as Ahkl/jnm=[Ihkl]/[Ihkl+Ijnm] permits evaluation of induced orientation by the pro-cessing (Table 1) [6,21,22]. It is clearly observed that the Ahkl/jnm

values change drastically from original material to the modifiedones. This fact indicates an effective interaction of the PP matrixwith the o-LDH filler and an o-LDH induced preferential growth ofPP crystallites (orientation of the crystalline phase). Furthermore,the o-LDH interacted with the amorphous PP phase as its meltcompounding with PP afforded nanocomposites with increased rela-tive crystallinity index (Table 1), indicating that o-LDH acted as nu-cleating agent. Both effects were previously observed with other claynanofillers, including the C16–C18 fatty acid modified hydrotalciteF100 [7,23]. Furthermore, the o-LDH incorporation induced the pres-ence of some β-PP (2θ~16°).

Table 1Lamellar long period, relative crystallinity index and orientation parameters (Ahkl/jnm)of PP and the PP nanocomposites.

Entry Sample L (Å) a A040/110b A130/110

b A110/111b Ic

c

1 PP 152 0.18 0.14 0.89 0.732 PP0.45 163 0.38 0.25 0.79 0.783 PP0.90 148 0.43 0.28 0.76 0.81

a Lamellar long period as determined by SAXS.b Orientation parameters as determined by XRD.c Relative crystallinity index as determined by XRD.

3.2.2. Thermal analysisThe resistance to thermal degradation of the PP samples was studied

by TGA and DTA (Fig. 8) and the results are summarized in Table 2. Incor-poration of o-LDH in the PP matrix resulted in improved resistance tothermal degradation when comparing Td5% and Td10%, showing thebest results for PP0.45 (+37 °C and+22 °C, respectively). This behaviorwas also reported for polyethylene/oleate-LDH nanocomposites [17].These results might be due to the presence of o-LDH residues at the PPsurface, which act as isolator and protect the bulk polymer. Besides,these residues might retard the diffusion of decomposition volatiles[24]. In comparison, the temperature at maximum degradation of bothPP nanocomposites reduced slightly,whichmight be due to a catalytic ac-tion of the nanofiller like in the catalytic thermal cracking of alkanes [23].The residual weight contents at 600 °C for PP0.45 and PP0.90 are 1.5%and 2.0%, respectively. This could reflect the higher thermal stability ofPP0.45 that might have resulted in a higher residual charcoal content be-sides the added 0.45 wt.% of o-LDH inorganic mass.

The effect of PP/o-LDHmelt blending on the PP melting and crystalli-zation properties are detailed in Table 2. Fig. 9 shows the heating (Fig. 9b)and cooling (Fig. 9a) curves for the neat PP and the PP nanocomposites.The crystallization temperature (Tc) increased slightly with increasingo-LDH content, suggesting the existence of a tendency. This trend wasfollowed by the crystallization enthalpy, showing higher values with in-creasing o-LDH content. Such a tendency was not observed for the melt-ing temperature. These changes in the crystallization temperature andenthalpy could be explained by the role of o-LDH as nucleating agentwithin the PP macromolecular segments [23].

3.2.3. Dynamic mechanical analysisFig. 10a shows the dynamic mechanical curves (dynamic storage

modulus and loss factor tan δ) as a function of temperature for PP,

Table 2Thermal properties of PP and the PP nanocomposites.

Entry Sample Td 5%(°C)a

Td 10% (°C)b Td max (°C)c Tc(°C)d

Tm(°C)e

Tg(°C)f

ΔHc

(J/g)g

1 PP 334 373 461 112 166 13.4 80.92 PP0.45 371 395 458 114 164 11.0 91.53 PP0.90 358 378 450 115 165 9.71 95.8

a Temperature at decomposition of 5 wt.% as determined by TGA.b Temperature at decomposition of 10 wt.% as determined by TGA.c Temperature at maximum decomposition as determined by TGA.d Crystallization temperature as determined by DSC.e Melting point as determined by DSC.f Glass transition temperature as determined by DMA.g Crystallization enthalpy as determined by DSC.

Table 3Dynamic mechanical properties of PP and the PP nanocomposites.

Entry Sample Storage modulus (MPa) Stiffness (103N/m)

−20 °C 20 °C 80 °C −20 °C 20 °C 80 °C

1 PP 3111 1838 594 157 93 302 PP0.45 2725 1457 411 234 124 333 PP0.90 2605 1492 376 304 174 48

Fig. 10. DMA analyses of (a) PP (red line), (b) PP0.45 (green line) and (c) PP0.90(black line): (Fig. 10a) Storage modulus and tan delta; (Fig. 10b) stiffness.

Fig. 9. DSC analyses showing the Tc (Fig. 9a) and Tm (Fig. 9b) of (a) PP (red line), (b) PP0.45(green line) and (c) PP0.90 (black line).

2402 R.K. Donato et al. / Materials Science and Engineering C 32 (2012) 2396–2403

PP0.45 and PP0.90, while Fig. 10b shows the stiffness results for thesematerials. The dominant relaxation at 13.4 °C corresponds to theglass-rubber relaxation of the amorphous PP portion. The broadpeak around 100 °C is associated with the crystalline PP regions. In-creasing o-LDH contents afforded nanocomposites with decreasingTg (Table 2). This indicates that o-LDH addition affords amorphousdomains with more flexible polymer chains and that o-LDH acts asplasticizer. Such a PP plasticizing effect was also observed for theC16–C18 fatty acid modified hydrotalcite F100 [23]. Altogether, the in-crease of both crystallinity (higher Ic and Tc values) and plasticization(lower Tg) suggests the formation of nanocomposites with crystallineand amorphous parts that are more brittle and ductile, respectively.The thermo-viscoelastic behavior of both nanocomposites was charac-terized by storage modulus decrease and stiffness increase (Table 3).This increase in stiffness could be related to the increased crystallinityof the PP nanocomposites.

4. Conclusions

A detailed study allowed understanding the role of oleate-functionalized o-LDH in themelt compounding of PP nanocomposites.o-LDH was successfully prepared using the memory effect of calcinedLDH and sodium oleate. Melt blending of PP and o-LDH resulted in amoderate dispersion degree of clay tactoids of different sizes and com-pactness, together with PP intercalation in the o-LDH galleries. This in-dicated that the oleate anion acted as compatibilizer, although otherstrategies will be required to promote full compatibility between the

filler and matrix. The o-LDH nanofiller induced orientation of the crys-talline PP phase under the processing conditions and increased the PPrelative crystallinity, demonstrating that it functioned as nucleatingagent. o-LDH affected both the crystalline and amorphous PP domainsas it also showed to be a plasticizer. Both PP0.45 and PP0.90 showedimproved resistance to thermal degradation and stiffness. As such, thestudied strategy allowed the preparation of materials with distinguish-able properties using a renewable feedstock for LDH functionalization.

Acknowledgments

The authors are grateful to the Brazilian Synchrotron Light Labora-tory LNLS for the SAXS analyses, to CNPq and FAPERGS for the finan-cial support, and to Braskem S.A. for the PP.

References

[1] A. Usuki, Y. Kojima, M. Kawasumi, A. Okada, Y. Fukushima, T. Kurauchi, O. Kamigaito,J. Mater. Res. 8 (1993) 1179.

[2] Y. Kojima, A. Usuki, M. Kawasumi, Y. Fukushima, A. Okada, T. Kurauchi, O. Kamigaito,J. Mater. Res. 8 (1993) 1185.

2403R.K. Donato et al. / Materials Science and Engineering C 32 (2012) 2396–2403

[3] A. Vaccari, Catal. Today 41 (1998) 53.[4] V. Rives, M.A. Ulibarri, Coord. Chem. Rev. 181 (1999) 61.[5] S.K. Yun, T.J. Pinnavaia, Chem. Mater. 7 (1995) 348.[6] M.F.S. Lima, M.A.Z. Vasconcellos, D. Samios, J. Polym. Sci., Part B: Polym. Phys. 40

(2002) 896.[7] D.R. Paul, L.M. Robeson, Polymer 49 (2008) 3187.[8] J.C. Rodrigues, T.M.H. Costa, M.R. Gallas, C.C. Moro, Phys. Chem. Miner. 36 (2009)

439.[9] M.Z.B. Hussein, Z. Zainal, C.Y. Ming, J. Mater. Sci. Lett. 19 (2000) 879.

[10] F.R. Costa, A. Leuteritz, U. Wagenknecht, D. Jehnichen, L. Häußler, G. Heinrich,Appl. Clay Sci. 38 (2008) 153.

[11] D.G. Evans, X. Duan, Chem. Commun. 5 (2006) 485.[12] C. Nyambo, D. Wang, C.A. Wilkie, Polym. Adv. Technol. 20 (2009) 332.[13] C. Nyambo, P. Songtipya, E. Manias, M.M. Jimenez-Gasco, C.A. Wilkie, J. Mater. Chem.

18 (2008) 4827.[14] Z.P. Xu, P.S. Braterman, K. Yu, H. Xu, Y. Wang, C.J. Brinker, Chem. Mater. 16 (2004)

2750.

[15] Y. Kameshima, H. Yoshizaki, A. Nakajima, K. Okada, J. Colloid Interface Sci. 298(2006) 624.

[16] C. Manzi-Nshuti, P. Songtipya, E. Manias, M.M. Jimenez-Gasco, J.M. Hossenlopp,C.A. Wilkie, Polym. Degrad. Stab. 94 (2009) 2042.

[17] C. Manzi-Nshuti, P. Songtipya, E. Manias, M.M.J. Gasco, J.M. Hossenlopp, C.A. Wilkie,Polymer 50 (2009) 3564.

[18] O.D.D. Jr, N.T. Prevost, G.D. Strahan, Clean-Soil Air Water 36 (2008) 687.[19] G. Machado, E.L.G. Denardin, E.J. Kinast, M.C. Gonçalves, M.A. de Luca, S.R. Teixeira,

D. Samios, Eur. Polym. J. 41 (2005) 129.[20] E.L.G. Denardin, Structural rearrangement of PET during full compression. PhD thesis.

Postgraduate course in Chemistry, UFRGS, Porto Alegre, Brazil, 2004.[21] D. Samios, S. Tokumoto, E.L.G. Denardin, Macromol. Symp. 229 (2005) 179.[22] G. Machado, E.L.G. Denardin, E.J. Kinast, M.C. Gonçalves, M.A. de Luca, S.R. Teixeira,

D. Samios, Eur. Polym. J. 41 (2005) 129.[23] D. Tabuani, S. Ceccia, G. Camino, Macromol. Symp. 301 (2011) 114.[24] N. Greesh, P.C. Hartmann, R.D. Sanderson, Macromol. Mater. Eng. 294 (2009) 787.