Investigation of the alumina nanoparticle role in the enhancementof the mechanical properties of polyamide/polycarbonate blends
Fouad Laoutid, Adrian Picard, Olivier Persenaire, Philippe Dubois*
Center of Innovation and Research in Materials & Polymers (CIRMAP), Laboratory of Polymeric and Composite Materials (LPCM), University of MonsUMONS & Materia Nova Research Center, Place du Parc 20, 7000 Mons, Belgium
a r t i c l e i n f o
Article history:Received 6 October 2014Received in revised form26 November 2014Accepted 24 December 2014Available online 3 January 2015
In this work, we have investigated the effect of the incorporation of alumina nanoparticles on the me-chanical properties of polyamide 6/polycarbonate (PA6/PC) blends. Nanocomposites were preparedthrough melt processing using both unreactive and reactive surface-treated alumina nanoparticles.Simple PA6/PC blend comprising 20 wt% of PA6 showed ductile mechanical behaviour with an ultimateelongation close to 50%. Contrarily, the incorporation of 20 wt% of PC into PA6 led to a drastic decrease ofthe strain at break (3%). A thorough study of the morphology, thermal and rheological properties pointedout that these poor mechanical properties were due to an autocatalytic thermal degradation of both PA6and PC phases. Advantageously, the incorporation of 5 wt% of alumina nanoparticles, whatever theirsurface nature, significatively improved the tensile properties of the blends inducing an important in-crease in the elongation at break up to 88% when 5 wt% of hydrophilic alumina is used. Aluminananoparticles proved to behave as protective agents for both PA6 and PC phases within the nano-composite materials.
Polyamide 6 (PA6) and polycarbonate (PC) are widely used invarious industrial fields since both provide useful properties in avariety of applications. Moreover, the blending of both polymerscould allow obtaining an alloy interestingly combining both PA6and PC inherent properties. More precisely, PA6 phase can providegood solvent resistance while PC can ensure good resistance tomoisture. For enhanced performance, the blends need to be com-patibilized since PA6 and PC are non-miscible in the whole range ofcomposition and temperature [1e3].
As a result, the mechanical properties of simple blends of PA6and PC are reported to be greatly inferior to those of the respectiveneat polymers [4]. It has been reported that PA6-b-PC block co-polymers can be generated in situ at the PC/PA6 interface duringmelt blending [5]. The formation of this copolymer, taking place viathe reaction between the eNH2 and eOeCOeOe groups of poly-amide 6 and polycarbonate, respectively, enhances the compati-bility between the two polymeric phases but is not sufficient to
ubois).
21
achieve high mechanical performance, except for blends largelyenriched in PA6, e.g. 90 wt% [5]. However, the formation of thiscopolymer requires long mixing time of ca. 30 min at 240 �C [6e8],which limits the interest of such compatibilization process. The useof compatibilizer agent thus appears necessary to stabilize theblend morphology and prevent its delamination.
The stabilization of immiscible polymer blend morphology canbe achieved by using compatibilizing agents which allow areduction of the interfacial tension between the two immisciblepolymeric phases. This effect is usually reached by using block andgraft copolymers, as well as some functionalized polymers thatpresent a chemical affinity with the two phases. As a result, thecompatibilizing agent can be located at the interface and pro-motes a reduction of the size of the dispersed phase and mayallow some stress transfer between the different phases leading tothe improvement of blend performance [9e12]. As far as PA6/PCblends are concerned, different compatibilizing agents such asmaleated polyolefins [6,7], functionalized and non-functionalizedpoly[styrene-b-(ethylene-co-butylene)-b-styrene] triblock copol-ymer [5], epoxy resin [4] have been investigated. Some of them,such as poly(propylene oxide) (PPO) was shown to affect posi-tively the formation of PA6-b-PC copolymers [1,2]. This effect wasattributed to the plasticizing effect of PPO, which increases the
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PA6 chain mobility, and consequently increases the probability ofthe reaction between the eNH2 and eOeCOeOe groups of thepolyamide 6 and polycarbonate, respectively. The incorporation of5 wt% of PPO in PA6/PC blend proved sufficient to raise the reac-tion speed of PA6-b-PC copolymer formation, but the resultingmaterials were still characterized by limited mechanical proper-ties. Actually, PA6/PC blends have not reached any commercialsignificance because of the lack of an efficient and cost-effectivecompatibilization technique [8].
As an alternative to organic compatibilizers, the use of nano-fillers to stabilize the morphology of immiscible polymer blends ismore and more investigated [13e22]. By analogy to “Pickeringemulsions”, in which oil/water emulsion is stabilized by solidcolloidal nanoparticles, several studies have attempted to extrap-olate the findings on colloidal emulsions to immiscible polymerblends such as poly(dimethyl siloxane)/polyisobutylene [23], PP/polystyrene [21], PP/poly[ethylene-co-(vinyl acetate)] [24], PP/PA6[25,26]. In these studies, depending on the interfacial tension be-tween the nanoparticles and polymer phases, hydrophobic pyro-genic silica nanoparticles have been shown to be located at theinterface and to form a rigid mechanical barrier that limits thecoalescence of nodules. In contrast, untreated nanosilica, i.e. hy-drophilic silica nanoparticles, were located in polar matrix forwhich they exhibited the strongest affinity.
As reported, the surface treatment of the nanofiller representsan important parameter that can affect the preferential location ofthe nanoparticles at the equilibrium since it allows controlling theinterfacial tensions at nanoparticleepolymer interface. In this pa-per, we investigated the effect of the incorporation of aluminananoparticles on mechanical and thermal properties of PA6 and PCblends. A thorough study of the effect of surface modification ofalumina nanoparticles on the properties of the formed blends wascarried out. The properties of the resulting nanocomposites werediscussed in the light of their morphology as determined by elec-tronic microscopies.
2. Experimental section
2.1. Materials
The polymers used in this study were commercial products:polyamide-6 (PA6) Akulon F136-c1 from DSM Engineering Plastics,Netherlands and polycarbonate (PC) Lexan Resin 123r from SABIC,Netherlands.
Different pyrogenic alumina nanoparticles were used in thisstudy:
- Hydrophilic alumina nanoparticles (Al-H), from Cabot, Belgium,are untreated alumina nanoparticles (SpectrAl™ 100) withspecific surface area of 95 m2/g.
- Coated alumina nanoparticles (Al-C) from Evonik, Belgium,are hydrophobic alumina nanoparticles (Aeroxide Alu C-805)with a specific surface area of 100 m2/g and surface-treatedwith n-octyl silane. TGA analysis, performed under air at20 �C/min, indicates that Al-C contains around 6.7 wt% of alkylcoating.
- Phosphorylated alumina nanoparticles (Al-P) is obtained aftersurface modification of hydrophilic aluminawith 11-phosphonoundecanoic acid (provided by Specific polymers). A solution of5 g of 11-phosphono undecanoic acid in 100 mL of THF wasadded to a suspension of 20 g of dried (16 h at 80 �C) hydrophilicpyrogenic alumina in 800 mL of THF. After stirring for 24 h atroom temperature, the collected powder was 6 times washed/centrifugated in THF (during 10min at 10,000 rpm), then in THF/water mixture and finally in water. The so-collected modified
nanoparticles were then freeze dried. TGA analysis (under air at20 �C/min) reveals that the final phosphorylated alumina con-tains ca. 6 wt% of 11-phosphono undecanoic acid.
2.2. Processing
Prior processing and injection moulding, PA6, PC and blendsthereof were overnight dried at 80 �C in ventilated oven. Meltcompounding of the PA6/PC blends and nanocomposites was per-formed in a Brabender internal mixer at 240 �C (3 min mixing at30 rpm followed by 7 min at 60 rpm). The formed blends (Table 1)were ground and injection moulded by using a Thermo-Haakemini-injection machine to produce specimens for tensile testing(type ISO 1/2e1BA with dimensions of 1.5 � 60 � 10 mm3). Theinjection moulding parameters are presented on Table 2.
2.3. Characterization
2.3.1. Thermal characterizationThermogravimetric analysis (TGA) was performed on a TA In-
strument Model Q5000IR. The TGA temperature was calibratedwith Curie temperature standards. Conventional TGA were per-formed under nitrogen flow (25 mLmin�1) from room temperatureto 700 �C using a heating rate of 20 �C min�1. Isothermal TGA wereperformed under nitrogen flow (25mLmin�1) at fixed temperature(240 �C) and weight loss recorded versus time. The kinetics of thethermal degradation was characterized by modulated TGA using aheating rate of 5 �C min�1, an amplitude of ±5 �C and a period of200 s.
2.3.2. Morphology characterizationScanning electron microscopy (SEM) analysis was performed
using a JEOL JSM 6100 apparatus at 10 kV. Samples were preparedby cryogenic fracture and later coated with gold. PC nodules wereselectively extracted with chloroform prior to SEM observations.Transmission electron microscopy (TEM) micrographs were recor-ded with a Philips CM100 apparatus using an acceleration voltageof 100 kV. The specimens for TEM were previously prepared byultra-cryomicrotomy cutting (Leica Ultracut). Tensile test speci-mens were used to prepare the samples and the observed surfaceswere perpendicular to injection direction.
2.3.3. Mechanical analysisTensile testing measurements were performed by using a Lloyd
LR 10 K tensile bench machine at a tensile rate of 10 mm min�1. Allmechanical tests were carried out by using specimens previouslyconditioned for at least 48 h at 20 ± 2 �C under a relative humidityof 50 ± 3% and the values were averaged out over 6 measurements.
2.3.4. Mechanical analysis2.3.4.1. Rheological characterization. Pristine PA6, PC and the cor-responding filled and unfilled PA6/PC (80e20) blends were
specific-polymers
Highlight
Table 2Injection moulding process parameters used to prepare specimens for tensiletesting.
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rheologically characterized in a TA instruments ARES mechanicalspectrometer operating in the parallel-plate dynamic strain mode,at 240 �C in the shear rate range from 1 to 300 rad s�1. For thatpurpose, parts of tensile test specimens were used. The so-obtainedsamples were dried at 80 �C in ventilated oven prior analysis.Measurements were taken at small enough strains (10%), to ensurethat the material response was in the linear viscoelastic regime.Viscosities at 240 �C and 1 rad s�1 were chosen as representative ofthe polymer chain scission during melt blending process.
3. Results and discussion
3.1. Unfilled blends
3.1.1. Morphology characterizationAs aforementioned physical blends of PA6 and PC are considered
thermodynamically immiscible. However, SEM observations (Fig. 1)do not show any macrophase separation but demonstrate thepresence of well dispersed PC droplets in continuous PA6 phase.Most of the dispersed PC domains have sizes within sub-micronregime and the others do not exceed 3 microns. The formation ofsuch morphology highlights good affinity between the two poly-mers (likely resulting from the in situ generation of block co-polymers at the interface, vide supra).
3.1.2. Mechanical characterizationTensile properties of PA6-PC blends were determined and the
recorded values of Young modulus, stress at break and strain at
Fig. 1. SEM images of unfilled (a) and filled PA6/PC (80e20) blends
break are gathered in Table 3 and compared to the pristine poly-mers. As observed, the PC-rich blend, containing 20 wt% of poly-amide, is ductile with an elongation at break similar to that ofpristine PCwhile the incorporation of 20 wt% of PC into PA6 leads toa dramatic decrease of the elongation at break and to the formationof brittle materials.
Thus, the mechanical behaviour of PA6/PC blends is stronglyrelated to the blend composition. Similar observationwas achievedin the case of PP/PA6 blends: the composition containing 20 wt% ofPP showed an elongation at break of 38% [26] while that containing20 wt% of PA6 was very brittle with an elongation at break of 4%[25]. Various parameters could explain the brittleness of dispersedphase on the polymer blend: (i) the poor interfacial adhesion, (ii)the size of dispersed phase nodules, (iii) a cavitation phenomenoninducing a modification of the plastic deformation mechanism ofthe major polymeric phase, (iv) the interface separation betweenthe two phases induced by the contraction of the semi-crystallinenature of the dispersed phase [27], (v) a modification of the skin/core morphology and/or crystallization [28e30].
As far as PA6/PC blends are considered, the decrease of theelongation at break could not be explained by the size of thedispersed phase. In fact, surprisingly, SEM observations of thisblend exhibit finer dispersion of PC nodules into PA6 matrix (Fig. 1)while blending two immiscible polymers generally leads to the
with 5 wt% of Al-C (b), 5 wt% of Al-P (c) and 5 wt% of Al-H (d).
F. Laoutid et al. / Polymer Degradation and Stability 112 (2015) 137e144140
formation of unstable morphologies, which tend to macro-phaseseparate. SEM images of PA6/PC (80e20) blend show the pres-ence of small cavities, corresponding to PC nodules previouslyextracted by selective solubilization in chloroform. As observed,these particles exhibit spherical and elongated structure withdiameter smaller than 5 mm. The formation of such structures re-sults from the competition between the coalescence and theelongation of the minor phase, induced by the action of shearduringmelt mixing. This process occurs generally in the case of noncompatibilized polymer blends but generally leads to the formationof larger domains. The formation of PA6-b-PC copolymers (videsupra), during melt mixing may explain the formation of suchsmaller dispersed domains. Actually, the in situ generated PA6-b-PCcopolymers may increase the compatibility between PA6 and PCphases and therefore stabilize the blend morphology, which pre-vents/refrains coalescence. Therefore, the regular dispersion ofrather low-sized PC nodules is not likely responsible for theobserved decrease in strain at break. Another explanation needs tofound out. It is worth pointing out that an important release ofvolatile compounds occurs during the melt processing of the brittlecompositions, i.e., the PA6/PC (80e20) blend. In contrast, no evi-dence of volatile formation during the processing of PC-rich blend,pristine PC and PA6, has been observed. The formation of thesevolatile compounds is likely due to a reaction between the twopolymers and/or between one polymer and the additives/stabi-lizers/catalyst residues of the other polymer. It is noteworthy thatthe formation of volatiles is more pronounced when the PA6 is thecontinuous polymer phase (i.e., for PA6-rich blends). This reactionreadily occurs since the formation of volatiles is also observedwhen the blend mixing time is reduced to only 2 min by using atwin-screw extruder instead of Brabender internal kneader. Furtherinvestigations of this phenomenon are discussed in the forth-coming sections.
3.1.3. Rheological characterizationThe generation of volatile compounds during the melt pro-
cessing of PA6-rich blend results from a chemical reaction betweenPA6 and PC. Therefore, it can be expected that this reaction affectsthe blend viscosity. Usually, the blend viscosity decreases when themajor polymeric phase undergoes a thermal decomposition whileit increases if coupling or grafting reactions take place.
Fig. 2. Effect of the incorporation of 5 wt% of alumina nanoparticles on the evolution ofthe dynamic viscosity of PA6/PC (80e20) blends in comparison with pristine PA6,pristine PC and unfilled PA6/PC (80e20) blend.
Fig. 2 shows the plot of dynamic viscosity (h) versus angularfrequency (u) for PA6/PC (80e20) blends and neat polymers aftermelt processing. The complex viscosity, determined at small fre-quency of oscillation (1 rad s�1), is lower in the case of the blend(140 Pa s) with respect to pristine polymers (2400 Pa s for PC and600 Pa s for PA6). This result gives credits to our speculations andindicates that the thermal degradation, during melt compounding,is more important in the case of PA6/PC (80e20) blend than whenboth polymers are processed separately. On the other hand, it isalso observed that the dynamic viscosity decreases slowly with theangular frequency in the case of pristine PA6 and PC while thedecrease appears more prominent for the blend. This result likelyindicates that the thermal degradation occurs also during the vis-cosity measurement at 240 �C.
3.1.4. Thermogravimetric analysisTo study the thermal degradation observed by rheological
analysis, the PA6/PC (80e20) and PA6/PC (20e80) blends werecharacterized by thermogravimetry (TGA). Results are presented onFig. 3 and Table 4.
First, the theoretical TGA response has been calculated andcompared to the experimental curve. For so doing, the TGA curvesof PA6 and PC (taking into account the blend composition and asrecorded under Nitrogen) have been arithmetically added. Suchapproach allowed evidencing the actual modification of the ther-mal resistance of PA6 by the addition of PC. Fig. 3 showed individualTGA curves of PA6, PC and experimental curve of the PA6/PC(80e20) blend in comparison with the theoretical TGA curve of thecorresponding blends.
Surprisingly, even if PC has a better thermal stability than PA6,the addition of only 20 wt% of PC to PA6 leads to an importantdecrease of the thermogravimetric resistance of the blend. Indeed,it appears that the theoretical curve of PA6/PC blend is different tothe experimental one (Fig. 3). These results are in accordance withprevious results reported by Costa et al. [6,7] who highlighted thepresence of chemical reactions between the two polymers that leadto the appearance of a first decomposition step (300e400 �C)corresponding to a weigh loss of about 20%. Therefore, it can beassumed that such decrease of thermal stability of the blend ex-plains the decrease of the molten viscosity of the blend and theformation of the volatile compounds during processing.
PC-rich composition is also concerned by a decrease of thethermal stability of the blend. In fact, comparison of experimental
Fig. 3. Experimental and predicted TGA curves of PA6/PC (80e20) and PA6/PC (20e80)under nitrogen flow (for sake of comparison, PA6 and PC TGA curves are also shown).
Table 4Temperature values corresponding to 5% (Te5 wt%), 10% (Te10 wt%) and 20% (Te20 wt%)weight loss. Comparison of theoretic and experimental values.
F. Laoutid et al. / Polymer Degradation and Stability 112 (2015) 137e144 141
and theoretical TGA curves of PA6/PC (20e80) composition high-lights that the blend is less thermally stable than it should be.However, in contrast to PA6/PC (80/20) blend, the decrease of thethermal stability of the PC-rich blend does not result in any loss ofmechanical properties since its thermal degradation starts athigher temperature with respect to PA6-rich blend. Isothermal TGAanalysis, performed at 240 �C under air (Fig. 4), also indicates thatthe thermal stability of PA6/PC (80e20) is lower to that of pristinePA6 and PA6/PC (20/80) blend. Indeed, after 30min, the weight lossof PA6-rich blend is about 12% while it is only of 2% when the blendis largely enriched in PC. This would explain why no volatile com-pounds production has been observed during melt processing ofPA6/PC (20/80) blend.
It is now clearly evidenced that the melt blending of PA6 and PCleads to a low thermally stable alloy. Taken separately, PA6 and PCare relatively thermally stable, but once blended, chemical re-actions between the two phases take place during melt processing(<300 �C). These chemical reactions that were evidenced by theimportant release of volatile compounds during processing lead tothe formation of a blend having lower thermal stability withrespect to pristine polymers, as showed by TGA analysis. It can beassumed that the thermal degradation pathways of PA6 and PCdepend whether the two polymers are separately or together meltprocessed. One polymer seems to interact with the second duringthe thermal degradation and modify the thermal degradationmechanism. This is not surprising insofar as the mode of degra-dation of both PA6 and PC is strongly dependent on several pa-rameters. As far as PA6 is concerned, its thermal degradationdepends on the environment; i.e. the temperature and the presenceof nucleophiles [31,32]. Below 300 �C, PA6 thermal degradationoccurs primally via 3-caprolactame (monomer) formation in theabsence of a nucleophile. Monomer formation often begins attemperatures slightly above 200 �C and occurs mainly by intra-molecular cyclization. In the presence of a nucleophile (specificallywater), PA6 degradation leads to chain cleavage with the formationof acid and amine end-groups. The thermal degradation of PC isalso complex and controversy still exists as to the true chemicalreactions occurring during thermal degradation even though PChas been studied for over 50 years. Under nitrogen, the chemical
Fig. 4. Isothermal thermograms of pristine PA6 and PC, PA6/PC (80e20) and PA6/PC(20e80) at 240 �C under nitrogen flow.
degradation process of PC dependsmainly on the elimination or notof the volatile product that are mainly composed by carbon dioxide,bisphenol-A, with lesser amounts of carbon monoxide, methane,phenol, diphenyl carbonate, and 2-(4-hydroxyphenyl)-2-phenylpropane [33]. The subsequent breakdown of bisphenol-A leads toethyl phenol, isopropenylphenol, isopropylphenol and cresol. Thematerial undergoes branching and eventual crosslinking to form aninsoluble gel when the system is continuously evacuated and thevolatile products are removed. If not, chain scission is observed.These two modes of degradation are the result of competition be-tween condensation and hydrolysis reactions. Davis and Goldenconcluded that PC degradation is a random chain scission processoccurring predominantly at the carbonate linkages and the initialdegradation begins at the end-group (phenyl or hydroxyl), whichreacts with any free proton present, such as from water or freehydroxyl groups. The chains with PhCO3- end-groups give offquantitative amounts of phenol and diphenyl carbonate whilethose with hydroxyl end-groups evolve only trace amounts ofphenol. Under air, it was found that PC first undergoes an oxidationstep at temperatures from 300 to 320 �C followed by a depoly-merization that occurs in the range 340e380 �C and consisting ofhydrolysis and alcoholysis of the carbonic ester, which is a form ofester exchange where chains break somewhere near the middle.The hydrolysis produces CO2 and either bisphenol-A (if it occursnear the chain end) or two shorter chains [33]. At higher temper-ature (480e600 �C), a complex random chain scission dominatesand consists of decarboxylation, hydrolysis, hydrogen abstraction,chain scission, ether cleavage, and, under certain conditions,crosslinking.
In this study, TGA experiments coupled with mass spectrometryfor evolved gas analysis (results not presented here) were per-formed on the PA6/PC (80e20) blend and did not show the for-mation of 3-caprolactam. This result may indicate that PA6decomposes under the action of a nucleophile agent. Taking intoaccount the PC decomposition process, It can be assume that asmall part of PA6 degrades during the melt processing and pro-duces acid and amine end-groups that induce a premature thermaldegradation of PC. Thermal degradation of PA6 could also becatalyzed by PC end-groups. The thermal degradation of PC leads tothe formation of nucleophiles that in turn could accelerate the PA6thermal degradation. Polyamide phase seems to be the mostaffected since the PC-rich blend remains ductile while the PA6-richcomposition becomes brittle. The continuous phase is supposed toensure themechanical resistance of the blend. This assumption is infull agreement with the results of isothermal TGA that showed thatthe weight loss recorded for PA6/PC (20/80) is only of about 2 wt%after 30 min degradation at 240 �C. For the sake of comparison,PA6/PC (80/20) blend losses 12 wt% in the same conditions.
In the rest of our work, we have limited our study to the PA6/PC(80e20) blend actually characterized by low mechanical perfor-mance and thermal resistance. Accordingly, surface-functionalizedalumina nanoparticles have been studied in order to minimize thethermal degradation of the blend and to enhance its mechanicalproperties.
3.2. PA6/PC (80e20) blends filled with 5 wt% Al2O3
The use of polymeric compatibilizing agents, having chemicalaffinity with both polymeric phases, is generally used to improvemechanical and morphological properties of immiscible polymerblends. Several recent studies have demonstrated the interest ofusing surface-modified nanoparticles as compatibilizing agents.Indeed, under certain conditions, these nanoparticles could belocalized at the interface forming a rigid barrier that preventscoalescence of the minority phase nodules [25,26]. In these
Fig. 5. TEM images of PA6/PC (80e20) filled with 5 wt% of Al-H (a), 5 wt% of Al-C (b) and 5 wt% of Al-P (c).
Fig. 6. TGA curves under nitrogen flow at 20 �C min�1.
F. Laoutid et al. / Polymer Degradation and Stability 112 (2015) 137e144142
studies, no chemical interaction between the nanoparticles andthe polymeric phases was reported and the localization of thenanoparticles at the interface was governed only by interfacialtension. As far as PA6/PC (80e20) blends are concerned, we usedtwo different surface-treated alumina nanoparticles in order tobetter control the blend morphology and to limit the degradingchemical interactions between the two polymeric phases. Indeed,it is expected that hydrophobic alumina nanoparticles (Al-C), thatare supposed to have no affinity with the two polar polymericphases, i.e., PA6 and PC, will be located at the interface. Conse-quently, the rigid barrier that could be formed around PC nodulescould limit the chemical reactions between PA6 and PC. Reactivealumina (Al-P), thus bearing carboxyl groups at its surface, is ex-pected to act as a coupling agent between the two polymer pha-ses. Indeed, carboxylic acid functions could react with both amineand phenol functions, formed during the thermal degradation ofPA6 and PC, respectively, or simply present as chain end-groups.For sake of comparison, PA6/PC blends containing 5 wt% of un-treated alumina nanoparticles (Al-H) has been also prepared andcharacterized.
3.2.1. Morphology characterizationThe incorporation of alumina nanoparticles seems to affect
positively the blend morphology as only sub-micronic nodulescould be observed. Nomore elongated domain is present in the caseof PA6/PC-based nanocomposites. This observation is more likelydue to some limitation of droplets coalescence that may result fromthe increase of the polyamide viscosity induced by the specificlocation of the alumina nanoparticles within the PA6 phase (seeTEM picture, Fig. 5). Such an increase in PA6 viscosity can affect thebalance between breakup and coalescence of dispersed polymerdomains during melt mixing. Such a behaviour has been alreadyobserved in case of PA6/PP blends filled with hydrophilic silicananoparticles and again mainly composed by PA6 [26]. Even in thecase of functionalized alumina (Al-P), the nanoparticles proved tobe located only into the continuous polyamide phase and neitherinto PC nodules nor at PA6/PC interface.
3.2.2. Rheological characterizationFig. 2 shows the plot between dynamic viscosity (h) and angular
frequency (u) for pristine polymers, PA6/PC (80e20) blend andrelated nanocomposites. It can be observed that the incorporationof nanoparticles into the blend leads to some increase of thecomplex viscosity, determined at 1 rad s�1. These values corre-spond to the melt viscosity of the blend after melt processing. Theuse of non-reactive alumina; i.e., hydrophobic (Al-C) and hydro-philic (Al-H), seems to limits the thermal degradation during pro-cessing since the viscosity of these blend is similar to that ofpristine processed PA6, i.e., around 680 Pa s. In contrast, the use ofreactive alumina (Al-P) in the blend leads to an important decrease
of the melt viscosity (230 Pa s) but remains higher than that of theunfilled blend (140 Pa s). Thus, the use of nanoparticles seems to bebenefit to the blend since it allows reduction of the thermaldegradation during melt processing.
However, during rheological testing the viscosity decrease ismore pronounced for nanofiller-based blends than pristine PA6.Since the viscosity of the blend is, at best, equivalent to that ofprocessed PA6, this result indicates that the three filled PA6/PCblends undergo some thermal degradation during melt processing.Actually, the observed viscosity increase is also attributed to thepresence of well dispersed nanoparticles (see Fig. 5) that are wellknown to increase the viscosity of polymer nanocomposites.Nevertheless, since the viscosity of the blend filled with Al-P is thelowest,we can assume thatAl-H andAl-C are those allowing a bettercontrol over the thermaldegradationof theblendduringprocessing.In fact, the increase of the blend viscosity, due to the presence ofnanoparticles, could be considered to be equal since the dispersionstate of the three alumina nanoparticles is similar (Fig. 5).
3.2.3. Thermogravimetric analysisIn comparison to unfilled PA6/PC (80e20) blend, the incorpo-
ration of 5 wt% of the different alumina nanoparticles leads to verylimited changes (Fig. 6). In fact, the first degradation step, below400 �C, is only slightly shifted to higher temperatures showing thatalumina nanoparticles do not significantly improve the thermog-ravimetric resistance of PA6/PC (80e20) blends. Furthermore, TGAcurves of the blends filled with alumina nanoparticles are similarwhatever their surface nature. The presence of carboxylic acidfunctions does not allow any increase of the blend thermal stabilityand only the presence of the mineral fraction of the nanoparticlesseems to be responsible of the slight increase of the thermalresistance. In fact, the increase of both T�5 wt%, T�10 wt% and T�20 wt%,corresponding to temperatures of�5%,�10% and�20%weight loss,
Table 6Activation energy (Ea) related to the first step of PA6 thermal degradation for thedifferent nanocomposites (filled with 5 wt% Al) in comparison to unfilled PA6/PC(80e20) blend and pristine PA6.
PA6 PA6/PC(80/20)
PA6/PC (80/20)eAl-C
PA6/PC (80/20)eAl-H
PA6/PC (80/20)eAl-P
Ea (kJ mol�1) 207 86 101 104 102
Table 5Temperature values corresponding to 5% (T�5 wt%), 10% (T�10 wt%) and 20% (T�20 wt%)weight loss.
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respectively, is similar whatever the nature of alumina nano-particles used. The general shape of TGA curves remains similar tothat of unfilled PA6/PC (80e20) blend except the amount of thefinal residue that corresponds only to the amount of nanofillersinitially in the blend (Table 5).
The fact that no more residue is formed during the thermaldegradation of these blends, even when alumina nanoparticles areused, indicates that the presence of alumina does not change thedegradation pathway of PA6/PC blend, and the antagonist thermo-degrading effect of the blend is slightly reduced but remainseffective.
3.2.4. Tensile propertiesTensile properties of PA6/PC-based nanocomposites filled with
5 wt% of alumina nanoparticles were determined. The recordedvalues of Young modulus, stress at break and strain at break aregathered in Table 3 and compared to the pristine polyamide andunfilled blend. It can be seen that the addition of 5% of aluminananoparticles, whatever their surface nature, increased strongly thestrain at break and led to the formation of ductile materials. Thisresult may look surprising since the thermal resistance of theseblends did not appear much higher than that of the unfilled blend,which demonstrates a brittle behaviour. Indeed, the strain at breakincreases from 3% to values higher than 50% when 5% of alumina isincorporated to PA6/PC (80/20) blend whatever the surface treat-ment of the nanoparticles.
The incorporation of 5 wt% of alumina nanoparticles in the PA6/PC (80e20) blends does not significantly affect the blendmorphology as attested by SEM images of the blends (see Fig. 1).Only a small decrease of the size of PC domains could be observedwhen alumina nanoparticles are used. Al-H seems to lead to theformation of smaller PC domains. Consequently, the increase of theblend ductility could not be related to the morphology of theblends.
We cannot attribute the improvement of the mechanical prop-erties of PA6/PC blend to a compatibilizing effect of nanofillers.Indeed, the use of either reactive (Al-P) or non-reactive nano-particles (Al-C and Al-H) allows obtaining a ductile mechanical
Fig. 7. TGA isothermal curves of pristine PA6 and PC, unfilled and filled PA6/PC(80e20) blend at 240 �C under nitrogen flow.
behaviour. This observation indicates the carboxylic functionspresent at the surface of Al-P cannot be considered as responsiblefor improving the mechanical behaviour of the blend. Indeed theother two alumina nanofillers provide the same improvement inmaterials ductility.
The reason behind the difference in mechanical properties ofthe alumina-filled nanocomposites with respect to the unfilledblend is not clear. One explanation could be that the presence ofnanoparticles increases the melt viscosity of PA6 phase, whichdecreases the PA6 chain mobility and consequently reduces theprobability of the reaction between PA6 and PC. This behaviourwould be at the opposite of what was reported and attributed toPPO, which acts as plasticizer that increases the PA6 chain mobility,and consequently increases the probability of the reaction betweenthe eNH2 and eOeCOeOe groups of polyamide 6 and poly-carbonate, respectively [1,2].
In order to further evidence the degradation reaction occurringduring melt processing, thermogravimetric analyses have beencarried out. For that purpose, isotherms at processing tempera-tures (240 �C) have been advantageously performed. TGAisothermal curves are reported on Fig. 7 and show that neatpolymers are relatively thermally stable at 240 �C with limitedweight loss, i.e. less than 1% in the case of PC and around 6% forPA6. However, the incorporation of 20 wt% of PC to PA6 induces animportant weight loss, higher than 10%. The first degradation stepof this blend is similar to that of neat PA6. According to theseresults, we can suppose that, in the case of the brittle composition,the PA6 phase starts to degrade first, catalyzing the thermaldegradation of PC that could also induce in return an accelerationof the thermal degradation of polyamide phase. The thermaldegradation of PA6/PC (80e20) blend seems therefore to be anautocatalytic reaction inducing thermal degradation of both PA6and PC phases.
This catalytic effect of PC on the PA6 thermal degradation hasbeen further investigated by assessment of the activation energy(Ea) of the PA6 thermal degradation at 240 �C. For that purpose,experiments have been run using modulated thermogravimetry,which allow determining the kinetic parameters of thermal deg-radations bymeans of sinusoidal heating ramps. The so-obtained Eavalues are reported in Table 6. As observed, high activation energyof 207 kJ mol�1 was measured for pristine PA6 while unfilled PA6/PC (80e20) blend presented a much lower value of 86 kJ mol�1.Such a decrease of activation energy supported the fact that PCbehaves as catalyst of the thermal degradation of PA6. Interestingly,both unreactive and reactive alumina nanoparticle-based compo-sitions exhibited higher Ea values (ca. 100 kJ mol�1) with respect tounfilled blends. These observations clearly evidenced that aluminananoparticles could be considered as protective agents reducingthe thermal degradation of both polymers during the blendprocessing.
4. Conclusion
In the present study, the effect of alumina nanoparticles on themorphology and properties of PA6/PC blends was investigated. Thein-depth thermal and rheological study clearly evidenced that
F. Laoutid et al. / Polymer Degradation and Stability 112 (2015) 137e144144
nanoparticles act as protective agents reducing the thermaldegradation of the polymer pair during melt processing. This pro-tective effect has been further evidenced by assessment of theactivation energy of the PA6 thermal degradation by modulatedthermogravimetry. Actually, both unreactive and reactive aluminananoparticles showed efficiency to lower the PA6 degradation rate.As a consequence, the mechanical properties of the blends provedto be strongly dependent upon the presence of nanoparticleswhatever their surface treatment.
As demonstrated, an opportune choice of nanoparticles allowedtailoring the final morphology and properties of polymer blends.This concept can be applied more generally to other polymers as anew route for enhancing polymer blend performances such asmechanical but also thermal and fire properties.
Acknowledgements
The authors thank the Wallonia Region, Nord-Pas de Calais Re-gion and European Community for the financial support in theframe of the INTERREG IV e NAVARE project. This work was alsosupported by the European Commission and R�egion WallonneFEDER program (Materia Nova) and OPTI2MAT program of excel-lence, by the Interuniversity Attraction Pole program of the BelgianFederal Science Policy Office (PAI 7/05) and by FNRS-FRFC.
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