Aluminium reinforced by WC and TiC nanoparticles (ex-situ)and aluminide particles (in-situ): Microstructure, wear and corrosionbehaviourA. Lekatoua,, A.E. Karantzalisa, A. Evangeloua, V. Gousiaa, G. Kaptayb, Z. Gcsib, P. Baumlib, A. SimonbaDepartment of Materials Science and Engineering, University of Ioannina, Ioannina 45110, GreecebPhysical Metallurgy Department, Materials Science Institute, The University of Miskolc, H-3515 Miskolc-Egyetemvaros, Hungaryarti cle i nfoArticle history:Received 11 June 2014Accepted 14 August 2014Available online 27 August 2014Keywords:Aluminium Matrix CompositesWC/TiC nanoparticlesMelt inoculationSliding wearCyclic polarizationDilute Harrisons SolutionabstractIn the present effort, Aluminium Matrix Composites (AMCs) were produced by the addition of submicronsized TiC and WC particles of low (up to 1.0 vol%) content into a melt of Al1050. Casting was assisted bytheuseofK2TiF6asawettingagentandmechanicalstirringtolimitparticleclustering. Anextensivepresence of intermetallic phases was observed in the cast products, as a result of both the inoculationby K2TiF6 and the intensive mainly due to the ne carbide particle size reactivity of the carbides withthe molten matrix. Particle distribution was reasonably uniform comprising both clusters and isolatedparticles. Theintermetallicparticledispersionhas changedtheintendednatureof thecomposites.Instead of one type of reinforcement, that of carbide particles, the aluminium matrix contained two maintypes of reinforcement: (a) in-situ intermetallic particles and (b) carbide nanoparticles, as such, or moreoftenas clusters of remainingcarbidenanocores andaluminideparticles. Thereinforcedmaterialsexhibitedanotablyimprovedslidingwearperformanceoverthatofthealloyowingtothebenecialeffectofboththecarbideandtheintermetallicphasedispersion. Awearmechanismwasformulatedbased on microstructural features of the wear surface (repeated hill-valley morphology, surface oxidelayers, crack formation and grooving). Cyclic potentiodynamic polarization in Dilute Harrisons Solution(DHS) revealedthatthecorrosionbehaviourof thereinforcedmaterialswasmainlycontrolledbythecorrosion of the alloy matrix. As such, the predominating form of corrosion was intergranular corrosion(IC) of Al associated with the presence of alloy matrix impurities. Carbide nanoparticles, aluminide phaseassociated with them and their Al-matrix remained essentially intact of corrosion. IC progress was ofteninhibited by the presence of clusters of aluminide and carbide particles. 2014 Elsevier Ltd. All rights reserved.1. IntroductionDue to properties, such as high specic strength and stiffness,lowdensityandlowthermal expansioncoefcient, AluminiumMatrix Composites (AMCs) have attracted great scientic attentionas candidate materials for high-tech, structural andfunctionalapplicationsincludingaerospace, defense, automotive, electronicpackaging, precision instruments, thermal management areas,sports equipment and recreation. Among them, particulatereinforced AMCs (PRAMs) have extensively been investigatedowing to their low production costs and versatility in employingconventional techniquesfortheirproductionandshaping[13].Recent evidence has shown that the mechanical response of AMCscanfurtherbeimprovedifsubmicron-andnano-sizedparticlesare used as reinforcing phase [47]. Within the above framework,themainconcept behindthis workis toexploit theattractivepropertiesofaluminiuminnanotechnologybyfabricatingAMCsreinforcedbysubmicronsizedcarbideparticles andevaluatingthemintermsof theircorrosionandwearbehaviour. However,therearesomemajordrawbacksforsuchconceptapplicability:(a) the high production costs involved in several production routes,(b) theagglomerationof nanoparticleswhenadoptinglowcostcasting techniques and(c) the likely degradationof corrosionresistance due to the introduction of a high number of interfaces.Regarding the rst drawback, amongst the great variety ofmanufacturing methods adopted for AMCs, the conventionalcasting processes are always at the research forefront due to theirrelatively lowcost, ease-to-handle advantages and large scaleproductioncapabilities. Works oncast nano-particlereinforcedAMCshavemainlyinvolvedAl2O3[6,812], SiC[7,13], B4C[14],AlN [15], MgO [16].http://dx.doi.org/10.1016/j.matdes.2014.08.0400261-3069/ 2014 Elsevier Ltd. All rights reserved.Corresponding author. Tel.: +30 26510 07309; fax: +30 26510 07034.E-mail address: alekatou@cc.uoi.gr (A. Lekatou).Materials and Design 65 (2015) 11211135ContentslistsavailableatScienceDirectMaterials and Designj our nal homepage: www. el sevi er . com/ l ocat e/ mat desRegardingtheseconddrawback, inconventional castingpro-cesses, the particle-molten Al wetting behaviour is the most crucialfactor for a successful particle insertion in the melt. Towards thisdirection, ceramicphaseswithastrongmetalliccharacter, suchas TiC and WC, may ensure enhanced particle liquid metal wet-tingcompatibility[17]. Besides, TiCandWCexhibit veryhighhardness and modulus of elasticity, excellent wear and high tem-peraturepropertiesandgoodcorrosionresistance, propertiesofgreat industrial signicance. Nevertheless, research efforts on WCandTiCcastreinforcedAMCsarelimitedandthey, intheirvastmajority, concern microsized TiCp reinforcedAMCs. These workshaveproduced promising results asfar as thewetting behaviourof themelt, thedispersionuniformity, as well as thestrengthand wear resistance are concerned [1825]. Additionally, the parti-cle insertion can further be assisted by the use of uxing agents,suchashalidesaltsthatdissolvetheoxidelayerformedonthesurface of the molten alloy; thus, the involved phases are allowedtoexpresstheirnetwettingcharacteristicsandenhanceparticleincorporation [21,22,24,2629].Regardingthethirddrawback, itiswell establishedthattheaqueouscorrosionbehaviourof AMCscanbeaffectedbymanyAMCfeatures. The matrix-reinforcement interface, the matrix/secondary phase interface, the secondary phase nature and content,the reinforcement/secondary phase interface, the electrolyte activeionconcentration, the productionroute, the heat treatment, the sur-face treatment can signicantly affect the corrosion resistance andmechanisms [3034]. Owing to these too many factors, conictingdata and interpretations exist regarding fundamental issues, suchas corrosion resistance and corrosion initiation sites [35]. Severalstudies have reported lower corrosion resistances for AMCs in com-parison with the respective monolithic alloys. Various reasons havebeenconsideredresponsible for this deteriorationof corrosionresis-tance in AMCs: (a) breakdown of a continuous passive lm at thematrix/reinforcement interfaces [36]; (b) galvanic coupling of alu-miniumandreinforcement[37];(c)voidsatthereinforcement/matrix interface [38]; (d) an increase in the dislocationdensity around particle clusters [39]; (e) interfacial layers aroundparticulatereinforcementsthatpromotegalvaniccorrosion[30];(f) formation of intermetallic phases by reaction of the reinforce-ment with the matrix due to heat treatment [34] or precipitationofintermetallicphasesenhancedbythepresenceofparticulates[33]. More recently, Pardo et al. [40] noted that the corrosion dam-age in AA360/SiCp and AA380/SiCp composites in (13.5) wt% NaClwas caused by pittingattack mainly atthe reinforcement/matrixinterface.However, Grifths and Turnbull [41] did not notice any appar-enteffectofSiCreinforcementontheelectrochemicalbehaviourof Al6061 in aerated 3.5 wt% NaCl. They concluded that the effectsof reinforcement on the corrosion of Al cannot be generalized andare specic functions of the environmental conditions and the pro-cessingroute. Thefollowingworkstendtosupport this claim:Trowsdale et al. [38], whilst nding no signicant galvanic actionbetweenSiCandAl, noticedthat20 wt%incorporationofSiC(ofparticle size of 3 lm) in Al1050 led to a slight reduction of the pit-ting resistance of the alloy. However, large particle sizes (20 lm)led to intensication of pitting due to cracking during the fabrica-tionprocess. Kiourtsidisetal. [42]statedthattheSiCppresencedoes not accelerate failure of the passivation oxide lm, whereaspitting corrosion potentials in aerated 3.5 wt% NaCl are notconsiderably affected by the SiCp content at a given aging condi-tion; nevertheless, alteration in aging kinetics due to SiCp presenceis responsible for the differentiation in the pitting corrosionbehaviour among the composites. Alaneme andBodunrin[43]claimed that unreinforced AA6063 exhibited slightly superior cor-rosion resistance than the AA6063/Al2O3p composites in NaCl andNaOHmedia;however, thecompositesshowedbettercorrosionresistance in H2SO4 medium. Candan [44] noted that intermetallicsas a result of reaction between an AlMg alloy and SiC reinforce-ment beneted the corrosion resistance of the composites in 3.5%NaCl due to interruption of the continuity of matrix channels.On the other hand, a number of works have reported superiorcorrosionperformanceoftheparticulateAMCsinrelationtotherespectivealloy. ThecorrosionresistanceofLM13-Alto1 MHClwas found by Seah et al. [45] to improve with increasing the garnetparticle content (26 wt% garnet) due to the garnet particles actingas physical barriers to the corrosion process. The positive effect ofzircon (ZrSiO4) particulates on the corrosion resistance of Al6061(HCl of different concentrations) and Al7075 (seawater) has beenreported by Jameel et al. [46] and Nagaswarupa et al. [47], respec-tively. Toptan etal. [48]noted that (1519) vol% addition of B4Cparticles did not signicantly affect the tendency for corrosion ofanAlSiCuMgalloyin0.05 MNaCl;however, itdecreasedthecorrosion tendency and corrosion rate during sliding wear testingin 0.05 M NaCl.The wear performance of AMCs is of major importance, since ithas to satisfy the necessity for long-lasting applications. The slidingwear behaviour of PRAMCs has extensively been investigated dur-ing the last decades [4955]. Key features of the involved degrada-tion mechanisms include: (i) material parameters, such as matrixmicrostructure and hardness, particle size and shape, particle nat-ure, particle volume fraction, particle matrix interface integrity,interfacial bonding, reinforcement wettability by the matrix, sec-ondaryphaseparticles;and(ii)extrinsictribological parameters,such as externally applied normal load, sliding speed, sliding dis-tance, temperature, surfacenish, hardness of counterpart andnominal contactarea[54,56]. Owingtothesetoomanyfactors,no consistent wear behaviour of AMCs has been established [52].Surface oxidation, extensive plastic deformation, debris character-istics and nature can each play a crucial role that can vary the wearmodeformmildtosevere[49,50,53,57]. Thegoverningmecha-nismshavebeendescribedinthreeestablishedtheories:(a)theadhesivewear theorythat considerstheadhesionbetweenthecountersurfacesattheasperitiesandsubsequentdecohesionofasperitiesleadingtomaterial removal[58];(b)thedelaminationwear theory taking place in four steps: cyclic plastic deformationof surface layers, crack or void nucleation, crack growth, formationof debris and debris removal by extension of cracks to the surface[57]; (c) the mechanically mixed layer theory, involving formationof debris owing to oxidation and plastic deformation of the countersurfaces [54,59]; this debris is continuously in a state of comminu-tionandconsolidationand, eventually, formsahardprotectivesurfacelayer that reduces theoverall wear rate. Onnumerousoccasions, theabovetheoriescomplementeachother. Accordingto Al-Qutub [60], erosion wear dominates in mild wear conditions;delamination wear becomes the primary mode in the wear transi-tion state; the severe wear regime is governed by adhesion wear(submicron Al2O3p/Al6061). Sub-micron and nano particulateceramicreinforcements, suchas Al2O3[60,61], B4C[62], MoSi2[63], SiC [64],SiC/graphite [65] and TiC [66],have been found toimprovethewear resistanceof aluminium(AMCs preparedbypowder metallurgy techniques).Overall, the main objective of the present effort is to fabricatePRAMCsbyadoptingfourapproaches:(a)lowcostconventionalcasting assisted by stirring and saltuxing for improved particlewettinganddistribution, (b)additionofsubmicronTiCandWCparticlesastheprimaryreinforcement inordertocombinetheadvantagesof ultranedimensionswiththeexcellentintrinsicproperties of these carbides, (c) employment of lowprimaryreinforcementvolumefractionstolimitsegregation, (d)optimi-zationof surfacepropertyresponsebyattainingfurther in-situreinforcement (whilst at the same time keeping production costslow).1122 A. Lekatou et al. / Materials and Design 65 (2015) 112111352. Experimental procedureAMCs wereprepared bythe addition ofsub-micron sized WCparticles (of approximateparticlesizeof 200400 nm) andTiCparticles (of approximate particle size of 400700 nm) intoAl1050 (Al of 99.5% commercial purity). The compositionsemployedwere: Al-0.7 vol%TiC, Al-1.0 vol%TiC, Al-0.5 vol%WCandAl-1.0 vol%WC. Wettingandhomogenizationwereassistedby two approaches: Fluxing and mechanical stirring. K2TiF6 (10 gK2TiF6/190 gAl1050)wasutilizedasauxingsaltforremovingthe oxide phasefrom the surface of the aluminium melt [26,29].First, mixing of the reinforcement and the salt was carried out. Then,this mixture was added into the alloy melt (830 C). The salt wasallowed to react with Al, a slag was formed, the carbide particlesinltrated into the melt and the slag was removed by a ladle. Rigor-ous stirring was then applied for homogenization and breakage ofany initial particle clusters. Stirring was conducted by an in-housemadeapparatusbasedonanAEGSB2E700Rpowerdrill, withagraphite rod being adapted and a four branch stirring shaft beingassembledat theendof therod. Thestirringspeedwas keptconstant at -3200 rpm. The stirring duration was -20 s. A nal slagcleaning was performedprior tocasting intocylindrical, steelmoulds of 1.5 cm inner diameter and 15 cm height.Specimens were cut fromeachcast bar, mountedandpreparedformetallographicexamination. Standardmetallographicprocedureswerecarried out,which included grindingby SiCpapers followedbypolishingwithdiamond suspensions. Inspectionofallsampleswas performed by Scanning Electron Microscopy (JEOL JSM 6510LVSEM/ Oxford Instruments X- Act EDX).Specimens polished to Ra < 1 lm were subjected to dry slidingwear testing at room temperature. A ball-on-disk tribometer (CSMInstr.)wasemployed. Thefollowingparameterswereemployed:normal loadof 1 N, slidingspeedof 10 cm/s, acquisitionrateof20 Hz, total slidingdistanceof 1000 m. AISI 5210steel ballsof6.0 mm diameter were used as a counterbody material. Each runwas interrupted every 200 m, for measuring the mass loss of thesample(Toledoelectronicbalanceof vedecimal digits). Beforeeach weighing, the specimen was ultrasonically cleaned by acetone.The overall wear rate was calculated from the mass loss vs. slidingdistance data by linear regression analysis (least squares method).Triplicate tests were performed for each material type.Corrosiontestingwasconductedoncylindrical couponsthatwere cut with a diamond saw, ground to 1000 grit, ultrasonicallycleaned and encapsulated in PTFE, leaving a surface area of-1 cm2to be exposed to aerated DHS, at 25 C. DHS (Dilute Harri-sons Solution) is atestingsolutionoftenusedonaeronauticalalloystoapproachatmosphericconditionsoftenencounteredbyairplanes [67]. The solutioncontains ammoniumsulphate andsodiumchloride (0.35 wt% (NH4)2SO4 + 0.05 wt% NaCl) usuallyfound in atmospherically deposited moisture. DHS is an effectiveemulatoroftheeffectsofacidrain. Alltheelectrochemicaltestswere performed using the Gill AC potentiostat/galvanostat byACM Instruments. A standard three electrode cell was employed,with Ag/AgCl (3.5 M KCl) as the reference electrode and a platinumgauge as the counter electrode. Potentiodynamic polarization testswere carried out at a scan rate of 10 mV/min. Polarization scanningstarted after 4 h of recording the Open Circuit Potential in DHS, at25 C. Reverse polarization was conducted to study the susceptibil-ity of the specimens to localized corrosion.3. Results and discussion3.1. Composite microstructure particle incorporationFirstofall, theinterpretationoftheresultshereinpresented,utilizes the AlTi [68], AlW [69], TiW [70] and AlC [71] phasediagrams. TheAlTiandAlWphasediagramsareillustratedinFig. 1.Fig. 2presents themicrostructures of theTiCp/Al materials(Fig. 2a and b) and WCp/Al materials (Fig. 2c and d). It is seen, thatthe carbide phase is present as both isolated particles and particleclusters,which aremainly located atthegrain boundaries. Inallcases, thereisastrongpresenceofintermetallicphases, locatedin the interior and at the boundaries of the Al grains.3.1.1. TiCp reinforced alloyIn the case of the TiCp reinforced alloy, the intermetallic phasepresent has been identied by quantitative EDX as Al3Ti, which isin consistency with the Al-Ti phase diagram (Fig. 1a). Fig. 3 showsthat the Al3Ti phase presents three main morphologies: (a) blocky,longish rectangular plates/rods (e.g. particles from which spectrum10 in Fig. 3a, spectrum 19 and spectrum 23 in Fig. 3b have beenreceived); (b) large, rounded particles of diameter (16) lm (e.g.particlesfromwhichspectra8, 9, 11inFig. 3a, spectra17, 18and20inFig. 3bhavebeenreceived); (c) dispersions of neroundedor plate-likeprecipitates(Fig. 3aandc). (Typical EDXanalyses are includedinthe legendof Fig. 3). The ne Al3Tiprecipitatesareusuallyassociatedwiththepresenceof theTiCnanoparticles or their remaining cores (circled cluster in Fig. 3c).a Al (mole fracon)Temperature (K) Al Ti b Al2WAl77W23Al12W Al5W Al7W3Al4W bcc liquid W (mole fracon)AlW Temperature (K) fccFig. 1. Phase diagrams of the systems: (a) TiAl [68], e(l) = Al3Ti; and (b) AlW [69].A. Lekatou et al. / Materials and Design 65 (2015) 11211135 1123The origin of the Al3Ti formation can be twofold: (a) due to reac-tion between K2TiF6 and molten Al [72,73] and (b) due to reactionbetween molten Al and TiC. Taking into account the TiAl (Fig. 1a)and AlC phase diagrams, the following reactions may account forthe presence of Al3Ti particles in the TiCp reinforced materials:At the temperature of the melt (830 C):3K2TiF6(s) 13Al(l) 3Al3Ti(s) 3KAlF4(l) K3AlF6(l) [72[ (1)13Al(l) 3TiC(s) 3Al3Ti(s) Al4C3(s) (2)Uponcooling, attheAl/TiCinterface, theremainingAl(l)reactedwith Al3Ti by a peritectic reaction:Al(l) Al3Ti(s) Al(s)(peritectic; 666 C) (3)a 0.7 vol% TiC b1.0vol% TiC1.0 vol% WC d 0.5 vol% WC c Fig. 2. Microstructures of thecast composites (backscatteredelectron-BSEmode). Isolatedandclusteredcarbideparticles arenotedbyarrows andcircles/ellipses,respectively.cb a1.0 vol% TiCAl3Ti Al3Ti1.0 vol% TiC0.7 vol% TiCAl3Ti Al3TiFig. 3. Intermetallic compound particles observed in the TiCp/Al materials (a & b: Secondary Electron-SE mode, c: BSE mode). Spot EDX analyses in (at%): (a) spectrum 8:76.97 Al-23.03 Ti, spectrum 9: 76.89 Al-23.11 Ti, spectrum 10: 75.97 Al-24.03 Ti, spectrum 11: 76.56 Al-23.44 Ti; (b) spectrum 18: 76.16 Al-23.84 Ti, spectrum 19: 76.06 Al-23.94 Ti, spectrum 20: 76.68 Al-23.32 Ti; (c) Al3Ti formation associated with TiC nanoparticles (in circle).1124 A. Lekatou et al. / Materials and Design 65 (2015) 11211135The peritectic mode of Al3Ti engulfment by the growing aAl grainmay account for the frequent location of Al3Ti inside the Al grains,as observed in Fig. 3.Reaction (1) is expected, as the outcome of a standard Al inocula-tionprocess describedby Mahallawy et al. [72]. Reaction(2) has alsobeen reported to occur at temperature levels that include the melttemperature of the present effort [74] or even at lowertemperatures uponcooling, despitethefact that this reactionis ther-modynamically unlikely [74]. However, previous works with AMCsreinforcedbyTiCpofconventionalparticlesize(45 lm)atthesame processing temperature didnot showany interactionbetweenAl andTiC[18,21,25]. Therefore, it is inferredthat the submicronsizeof the TiC particles has accelerated reactivity phenomena owing tothe high number/specic surface of the interfaces introduced.The non-detection of Al4C3in the nal products can beexplainedbyitsdissolutionandremoval duringmetallographicpreparation, sinceAl4C3ishydrolyzedbywateraccordingtothereaction [75]:Al4C3(s) 12H2O(l) = 4Al(OH)3 3CH4(g) (4)ThedifferentAl3Tiparticle morphologiescanbeexplainedonthebasisof reactions(1)(3): Blocky, longishplatesaremost likelythe product of the reaction between K2TiF6 and Al(l) at the temper-ature of the melt (830 C). These morphologies are usually formedin salt rich regions, i.e. Ti-supersaturated regions [72]. Largerounded particles can also be the product of salt-melt reactions inareas of somewhat lower K2TiF6 concentration, since the morphol-ogy of Al3Ti particles resulting fromthe salt-Al(l) reaction isstronglydependent onboththeTi content inthemelt andthecooling rate; thus, a wide range of different morphologies may beattained [72]. The ne precipitates associated with TiC nanoparti-cles are probably the product of reaction (2).3.1.2. WCp reinforced alloyInthecaseof WCp/Al, twotypesof tungstenaluminideareobserved (Fig. 4): Al12W in the form of coarse polygonal particlesoflargestdiagonalof(318)lm(Fig. 4a)andAl5Wintheformof acicularplates(Fig. 4bandc). (Theirstoichiometryhasbeenidentied by quantitative EDX. Typical EDX analyses are includedinthelegendof Fig. 4). TheirpresenceisfullyjustiedbytheAlW (Fig. 1b) and AlC phase diagrams, as follows: at the temper-ature of the melt (830 C), Al(l) combined with WC to form Al5WandAl4C3. (Al5WneedlesarealwayslocatedbyWCremainingcores,asshown inFig. 4bandc). On cooling,Al12WwasformedasaresultoftheperitecticreactionbetweenAl(l)andAl5W. Onfurther cooling, the remaining Al(l) peritectically reactedwithAl12Wtoform aAl. TheperitecticmodeofAl12Wengulfmentbythegrowing aAlgrainmayaccount forthelocalization ofAl12Win the interior of the Al grains. Therefore, the following sequenceof reactions may account for the presence of Al5W and Al12W par-ticles in the WCp reinforced materials.At the melt temperature(830C) : 19Al(l)3WC(s) 3Al5W(s)Al4C3(s)(5)On cooling(697 C) : Al(l) Al5W(s) Al12W(s)(peritectic) (6)On cooling(661 C) : Al(l) Al12W(s) Al(s)(peritectic) (7)Al3Tiparticlesappearasclustersoragglomeratesofneroundedparticlesorasasystemof coarserectangularplatesforminganincomplete rosette (Fig. 4c, upper right). As aforementionedinSection3.1.1, avarietyof different aluminidemorphologiescanresultfromtheinteractionoftheAl-meltwiththesaltux. Themorphology of incomplete rosette may be associated with salt richregions, i.e. Ti-supersaturated regions [72]. The signicant presenceof W in the Al3Ti composition (see spectrum 13 in Fig. 4a, spectra 2,3, in Fig. 4c) can be explained by the fact that, at the temperature ofthe melt (830 C), Ti and W can form (Ti,W) solid solutions (b1, b2, b)over a wide range of stoichiometry [70]. During casting andsolidication, a fair amount of W has remained trapped in the Al3Tilattice resulting in a metastable Al3(Ti,W) structure. It is known thatAl3Ti of tetragonal D022structure can be transformed to themetastableL12cubicstructurebyalloyingwithtransitionmetalswithmore d-electrons intheir valence band(e.g. W). The L12structureismoreductilethantheD022structureduetocovalentbondingwithenhancedmetallinityascomparedtothecovalentbonding of the D022 structure [76].The great reactivity of WC with molten Al is not only due to itssubmicron particle size (ner than that of TiCp) but also due to itsthermodynamicinstability. {Theenthalpyofformationof WCishigher than that of TiC {DHf(WC): 40.5 kJ/mol, DHf(TiC):184.1 kJ/mol [77])}.3.1.3. Al-Fe based intermetallicsAnotherimportantfeature, observedinFig. 2and, inahighmagnication, inFig5, isthepresenceof eutecticphaseatthegrainboundaries. EDXanalysisrevealedthattheeutecticmicro-constituent, apart from Al, consists of AlFe and AlFeSi interme-tallic phases. Such eutectic presence is commonly encountered incommercial Al-alloys[78]. Feisthemost commonimpurityinaluminium forming a variety of intermetallics, such as Al3Fe, Al6Fe,aAl(Fe, Mn)Si, dAlFeSi, b(Fe,Si), a(Fe,Si), Al12Fe3Si2.3.1.4. Particle distributionDespite the reaction product extent in both types of compositematerials, it could be stated that the particle incorporation withinthe melt was successful and of high rate. Such a successful incorpo-rationisattributedtoboth: (a) thebenecial actionof K2TiF6,whichreactedwithliquidAltoformaKAlFbasedliquidslagthat removed surface oxide phases and allowed the expression oftheparticle-meltnetwettingcharacteristics;and(b)thestirringapplied during processing.Thecarbidenanoparticledistribution, characterized-asafore-mentioned- byclusteredandisolatedparticleslocatedmainlyinthe vicinity of the grain boundaries, could be a result of: (a) initialparticle clustering of the precursor powder that could not be bro-kenbymechanical stirring. Theveryneparticlesizeenhancessuchclusteringendurance;(b)clusterformationduetoparticlepushingbythesolidicationfront. Suchpushingcouldalsoberesponsiblefor the nal particle location at thegrain boundariesand the areas of the lastly solidied liquid. Thermal conductivitydifferencetheories, proposedbyZubkoet al. [79] andSurappaandRohatgi [3], sufcientlydescribethisnal particlelocation.Accordingtothesetheories, duringcooling, thehotterduetotheirlowerthermalconductivityreinforcingparticlespreservethe cooler surrounding liquid. This way, the growth of an advanc-ing grain upon cooling, is obstructed by the reinforcing particles. Asa consequence, aluminium grains become rened and particles arebeing pushed towards the grain boundaries. Thermal conductivitytheories can also explain the Al3Ti localization at grain boundaries.On the other hand, the localization of Al3Ti inside aAl grains can beexplainedby the peritectic mode of Al3Ti engulfment by aAlaccording to reaction (3). In the WCp reinforced alloy, the presenceof Al3Ti at grain boundaries (Fig. 4a and c) is more frequent than inthe TiCp reinforced alloy, because less Al(s) was available to peri-tecticallyengulfAl3Ti;Al(l)-reactant inreaction(3)-hadlargelybeen consumed in reactions (5)(7).Here it should be noted that the termcarbide nanoparticle actu-ally refers to the cores of the original sub-micron carbide particlesremaining after the reaction of their peripheries with the Al-melt.To conclude, the distribution of intermetallic particles, isolatedcarbide nanoparticles andclusters of intermetallic/nanocarbideA. Lekatou et al. / Materials and Design 65 (2015) 11211135 1125particles was reasonable uniform. However, it changed theintended nature of the composites. Instead of one type of reinforce-ment, that of carbide particles, the AMCs contained two main typesof reinforcement: (a) in-situAl3Ti or (Al5W + Al12W + Al3(Ti,W))intermetallicparticlesinthecasesof TiCp-AMCandWCp-AMC,respectively;(b)carbidenanoparticles, assuch, ormoreoftenasclusters of carbide nanoparticles and aluminide particles. A mainquestion is arising: Had the newcomposite improved surface prop-erties in relation to the monolithic alloy?3.2. Sliding wear response3.2.1. Effect of carbide volume fraction and typeFig. 6a, presents the mass loss of the different composites as afunctionof theslidingdistance. It is seenthat, as theslidingdistanceincreased,themasslossalsoincreased, incompatibilitywithpreviousinvestigations[51,52,56,80,81]. Thewearratesofthe produced materials are displayed in Fig. 6b. It is evident thatincreasing the carbide particle volume fraction has led to adecreaseinmassloss, and, consequently, wearrate, whichisinagreement with other experimental efforts [18,49,50,52,53,80,82].The positive effect of the carbide volume fraction on the wearresistance of the composites is both direct and indirect: the directeffect stems from the TiC and WC particles, as such and as clusters;theindirect effectsoriginatesfromthehardintermetalliccom-pound particles formed by the reaction of the carbide phase withtheAl-matrix. Thisdual effect hasledtoanotabledecreaseinthe wear rate with carbide volume fraction increasing despite thelowcarbidevolumefractionsemployed. Inparticular, additionsof just 0.5 vol% WC and 1.0 vol% WC have led to a decrease in thewear rate of Al1050 by a factor of 2.7 and 3.7, respectively; addi-tionsof 1.0 vol%TiChaveledtoadecreaseinthewearrateofAl1050byafactor of 2.2. Suchbenecial behaviour is mainlyattributed to the strengthening effect that the dispersed particles(carbidesandaluminides)inducetothesoftmatrix, delaying, inturn, plastic deformation phenomena which can be mainly respon-sible for the overall degradation sequence. More analytically,carbide particles (intheir majorityas remainingcores of nano-dimensions) andintermetallic particles mayinhibit/retardplasticdeformation-due crack growth in the matrix by: (a) reducing the loadtransfer to the matrix, (b) decreasing the direct matrix-counterfacecontact area, (c) providing thermal stabilitytothe matrix, thuspostponingthermal softeningeffects and(d) inducingAl-grainrenement (see Section 3.1.4) [18,49,50,5456,80]. Regarding con-tribution(a), theparticleclusters(nanocarbideor nanocarbide/intermetallic) andthecoarseintermetallicparticles(e.g. Al12W)maycarrygreat portionsof theappliedload, therebyreducingthe load that ne particles and the soft matrix can carry [49]. Onthe other hand, the probability of reinforcement cracking increaseswith increasing size when size exceeds a critical value [83].Al12W Al3Ti Al3Ti Al3Ti 0.5 vol% WCaAl5W Al3Ti WCb 0.5 vol% WCWC1.0 vol % WCAl5W Al5W cAl3Ti Al3Ti Fig. 4. Intermetallic compound particles observed in the WCp/Al materials (SE mode). Spot EDX analyses in (at%): (a) spectrum 12: 91.73 Al-7.64 W-0.63 Ti, spectrum 13:76.01 Al-18.47 Ti-5.52 W, (b) spectrum 29: 83.88 Al-15.55 W-0.57 Ti, and (c) spectrum 2: 75.44 Al-19.05 Ti-5.51 W, spectrum 3: 75.23 Al-18.95 Ti-5.82 W, spectrum 5: 83.86Al-15.03 W-1.11 Ti.1.0 vol% TiCFig. 5. Intergranular presence of a eutectic (iron aluminide) intermetallic phase (SEmode).1126 A. Lekatou et al. / Materials and Design 65 (2015) 11211135Fig. 6alsodemonstratesthatWCreinforcementhasledtoahigherwearresistanceoftheWCpreinforcedalloyascomparedto the TiCp reinforced alloy. This superiority can be attributed to:(a) the ner particle size of WC resulting in a greater obstructionof the dislocation movement due to the greater number ofinterphaseboundaries, and(b) furtherdensicationof thehardphase dispersion by the more extensive presence of intermetalliccompound particles (Al5W,Al12W);thelatter, asaforementionedin Section 3.1.2, is due to the higher reactivity of WC as comparedto that of TiC (ner particle size and lower thermodynamic stabil-ity in comparison with TiC).3.2.2. Wear track morphologyTheweartracktopographiesof themonolithicalloyandtheAMCs are illustrated in Fig. 7. In all cases, the characteristic hill-valley morphology is observed. The wear track appears wider atthe hill areas and narrower at the valleys; this indicates that,during sliding, signicant material movement towards the hillshasoccurred. SuchlandscapeformationhasbeenexplainedbySarkar[84]intermsofintensiveplasticdeformationof thesoftmatrix infront of the moving counterbody steel ball causingnotablematerial owatdirections, whicharebothparallel andperpendicular to the sliding direction. As a result, a hill is beingbuilt up. When the counterface movement cannot any longer causefurthermaterial owtothehill, thecounterfaceoverpassesitandrepeats thesamecycleonanadjacent area. Eventually, anewhill is beingformed. Thenal outcomeis therepeatedhill-valley morphology observed in Fig. 7.Comparisonoftheweartrackmorphologiesofthereinforcedmaterialswiththat of themonolithicalloyinFig. 7drawsthefollowingobservations: (a) thehill-valley morphologyof themonolithic alloyis moreintensivethanthat of thereinforcedmaterials, asfarasthehill widthtovalley widthratioandthe wear track relief are concerned; (b) the prole of the wear trackedges of the monolithic alloy is rougher than those of the AMCs; (c)the unreinforced alloy shows the widest wear track, whereas the1 vol% WC composite shows the narrowest wear track. Comparisonof theweartrackmorphologiesof aluminiumreinforcedbythesame type of carbide in Fig. 7 reveals that the higher the carbidevolume fraction the more uniform the landscape morphology andthelessroughtheweartrackedges. Itis, thus, evidentthatthemonolithic alloy has been subjected to more severe plasticdeformation than the reinforced alloy. Therefore, it is deduced thatthereinforcingphases(carbidesandaluminides)haverestrictedthematrixplasticowresultinginamoreuniformweartracklandscape [18].EDX analysis on selected hill areas, illustrated in Fig. 8a and b,revealedthepresenceofAl-basedoxidephaseswith Ti(TiCp/Al)andW(WCp/Al)alsobeingpresent. RepresentativeEDXspectraare given in Fig. 8c and d. The formation of oxide layers during slid-ingwearof Al-alloysandtheircompositeshaspreviouslybeenreported[18,50,53,85]. Theirpresenceisowingtoamechanicalmixing process accompanied by oxidation reaction due to the fric-tional heating during dry sliding. Formation of these oxide layersdelaysthetransitionbetweenmildandseverewearregimesinAMCs.Higher magnication micrographs of the wear surfaces, inFig. 9, demonstrateaquitegreaterextentofplasticdeformationfor Al1050 (Fig. 9a), in comparison with the TiC and WC reinforcedalloy (Fig. 9b and c, respectively). In fact, the wear surface of themonolithic alloy exhibits mainly plastic deformation (in terms ofhills/valleys, ridgesand, generally, surfacerelief). Thewearsur-faces of the composites show both the above type of plastic defor-mation, as well as groove formation (aligned in the slidingdirection) with the latter being the main feature. The grooves onthewear surfaces of theAMCs areshallow, neandnarrowlyspaced, atypical patterncausedbytheabrasiveactionof largenumbersof hardparticlesanddebris [50,81]. The0.5 vol%WCcomposite presents slightly more extensive grooving as comparedto the 1.0 vol% TiC composite indicating a sliding action by a highernumberofhardparticles(despitealowervolumefraction). Theabove observation will be discussed in Section 3.2.3.High magnication micrographs of the wear surfaces, in Fig. 10,reveal the presence of cracks/aws within the oxide layer almostperpendicular to the sliding direction. Crack formation during weartesting of AMCs has been reported to be caused by extensive plas-ticdeformation oftheAl-matrix and/orparticlematrixinterfacedebonding, eventually leading to delamination and materialremoval [20,54,56,57]. Another likely reason for such crack forma-tion may be -besides fatigue due to repeated sliding- thermal fati-gueoftheoxidephasesduetothermal cycling, asthetestwasinterrupted every 200 m of sliding distance for specimen weighing[18,82]. Cracking in the surface oxide layer and subsequent delam-inationconstitutecrucialeventsforthewearresponseofanAl-basedmaterial, sincetheyhavebeenconsideredresponsibleforthe transition between mild wear and severe wear regimes[81,86]. Comparisonof Figs. 10ac shows that the monolithicmatrix exhibits the biggest cracks that may lead to delamination(astheoneobservedat thefarleft of Fig. 10a). ComparisonofFig. 10b and c reveals a smoother wear surface, fewer cracks/sur-face damage and ner, shallower grooves for the 0.5 vol% WC com-posite in relation to the 0.7 vol% TiC composite.Overall, the wear mode of the produced AMCs is characterizedas mild,as it lacks any signicantseizure,material delaminationor extensive material deformation. The mild wear regime is attrib-uted to the benecial action of two main factors: (a) the extensivehard phase (carbides and intermetallics) dispersion has restrictedmatrixowanddelayedcrackgrowthand(b)thesurfaceoxide05101520253035400 200 400 600 800 1000 1200Sliding distance (m)Mass loss (g x10-3)Al-1.0 vol% WCAl-0.5 vol% WCAl-1.0 vol% TiCAl-0.7 vol% TiCAl1050a b Fig. 6. (a) Mass loss versus sliding distance during dry ball-on-disk testing, for Al1050, monolithic and reinforced by TiC and WC submicron particles; (b) the wearrate of the different materials.A. Lekatou et al. / Materials and Design 65 (2015) 11211135 1127hillvalleyb 0.7 vol% TiC0.5vol% WCvalleyhilldmonolithichillvalleyavalley1.0 vol% TiChillchillvalley1.0 vol% WC eFig. 7. Panoramic views of the wear track morphologies of the monolithic alloy and the different composites produced (SE mode), illustrating the hill-valley landscape.AlOTi*c1.0 vol% TiC*aWAlO+d+b 1.0 vol% WCFig. 8. Wear track morphologies in BSE mode and EDX spectra from hill (dark contrast) areas. (a) and (c) Al-1.0 vol% TiC and EDX spectrum from a hill, respectively; (b) and(d) Al-1.0 vol% WC and EDX spectrum from a hill, respectively.1128 A. Lekatou et al. / Materials and Design 65 (2015) 11211135layerhasprotectedunderlyingmatrixareasfromdirectcontactwith the counterbody material.3.2.3. Wear debrisComparison of the micrographs of the WC reinforced alloy withthose of the TiC reinforced alloy and the monolithic alloy, showsthattheweartracksandtheneighbouringexternalzonesoftheWC-AMCspresent notablymoredebristhanthecorrespondingsurface zones of the TiC-AMCs and the monolithic alloy (compareFigs. 7dandewithFigs. 7ac, Fig. 8bwithFig. 8a, Fig. 9cwithFigs. 9a and b,Fig. 10cwith Figs. 10a and b). Thisobservation iscompatible with the relatively extensive (ne and dense) groovingalong the wear track of the WC-reinforced alloy (compare Fig. 9cwith Figs. 9b and a).EDXanalysisofdebrisparticlesfromthewearsurfaceoftheWCreinforcedalloy(Fig. 11a) revealedtwotypesof materials:(a) bright contrast particlesof highWcontent (e.g. spectra13and15)anddarkercontrastoxideparticlesof relativelylowWcontent (spectra 14, 16). It is indicatedthat the former onesoriginated from oxidized broken particles of W-aluminides (mostmonolithic ac 0.5 vol% WCb 1.0 vol% TiCFig. 9. Wear track morphologies of different materials produced (SE mode), illustrating a high extent of plastic deformation for the monolithic alloy and the formation of neand dense grooves on the wear surface of the AMCs.AldOWa monolithic b 0.7 vol% TiC0.5 vol% WCcFig. 10. (a)(c): Crack formation in oxide layer areas (SE mode); (d) EDX analysis in the vicinity of the cracks revealing Al-based oxide presence (Al-0.5 vol% WC).A. Lekatou et al. / Materials and Design 65 (2015) 11211135 1129likely Al12W), whilst the latter ones originated fromoxidizedmixtures of Al-matrix and fragments of W-, Ti- aluminides.Al12Wparticlesarehighlysusceptibletofragmentationnotonlybecause of their large size but also because of their intrinsic brit-tleness; the latter is due to an ordered ve-fold icosahedralstructure, which is commonly met in quasicrystals and bulkmetallicglasses[87]. Theacicularintermetalliccompoundparti-cles(Al5W) alsopresent ahightendencytofracture, especiallytheones lyingnormal tothesurface; as such, theyarelikelyto participate in the relatively dark contrast debris particlesshowninFig. 11a.EDXanalysisofdebrisfromthewearsurfaceoftheTiCrein-forced alloy(Fig. 11b)revealedonlyonetypeofmaterialthat ofalumina without any or with little Ti (spectrum 8); thus, it is indi-cated that debris from these composites mostly derived from theAl-matrix mixed to a low extent with Al3Ti.Based on the EDX analysis of debris, as aforementioned, and themicrostructural examination, aspresentedinSections3.1.1and3.1.2, the higher amount of wear debris in the case of theWC-AMCsascomparedtotheTiC-AMCs, maybeexplainedbythemoreextensivepresenceof intermetallicparticles that arehighly susceptible to fragmentation and can cause third body abra-sion either as monolithic fragments or as mixtures with the alloymatrix. In the case of WC-AMCs, third body abrasion has been con-ductedbyAl12W(monolithicfragmentswithoxidizedsurfaces)andacicular Al5W(mixedwithalumina fromthe matrix andAl3(Ti,W)).Althoughthenumber of abrasiveparticles producedduringsliding wear is higher in the case of the WC-AMCs than in the caseof the TiC-AMCs, the relatively shallow and ne grooves along thewear tracks inthe case of WC-AMCs (compare Fig. 10bwithFig. 10c) are evident of the relatively high resistance of WC-AMCsto third body abrasion.3.2.4. Mechanism of wearBased on the aforementioned observations, a likely wear mech-anism for the composites of the present effort can be summarizedas follows:(i)Ontheonset of slidingwear testing, thesoft aluminiummatrixwas subjectedtointensiveplasticdeformationinfrontofthecounterfaceball. Asaconsequence, signicantmaterial owoccurredthat, asslidingprogressed, ledtothe repeated hill-valley morphology. Carbide and interme-tallic compound particles restricted the matrix ow result-ing in smoother wear surfaces.(ii)Atthesametime, thefrictionalheatingduringdryslidingwear induced the formation of surface Al-based oxide layers,enriched by Ti- or W- oxides.(iii)Thefrictional forces, thebrittlenessof thealumina-basedsurface layer and wear fatigue enhanced by thermal fatigue(duetotherepeatedinterruptionofthetestformasslossmeasurements) caused the formation of cracks in the oxidelayer.(iv)Crackpropagationandgrowtheventuallyledtomaterialremoval. Crackpropagationwas delayedbythecarbide/intermetallic phase dispersionandthe renedaAl grainboundaries.b 1.0 vol% TiCSpectrum 8AlTiO 1.0 vol% WC aAlWOWSpectrum 13OAlWTiSpectrum 16Fig. 11. Debris particles on the wear surface of the AMCs. (a) 1.0 vol% WC reinforced alloy and EDX spectra of brighter contrast particles (Spextrum13) and darker contrastparticles (Spectrum16); and (b) 1.0 vol% TiC reinforced alloy and EDX analysis from a debris particle (spectrum 8).1130 A. Lekatou et al. / Materials and Design 65 (2015) 11211135(v)The abrasive action of debris (TiCp/Al: oxides mostly origi-natingfromthematrix; WCp/Al: oxidesoriginatingfrommixturesof Al matrixandW-, Ti-aluminidesandoxidesoriginating from fragments of large Al12W particles) causedthe formation of shallow, ne and dense grooves along thewear tracks.3.3. Corrosion behaviour3.3.1. Potentiodynamic polarizationThe cyclic voltammograms of the tested materials are presentedin Fig. 12(ad). The negative hysteresis loops of the anodic polariza-tion curves (i.e. higher current densities upon reverse polarizationas compared to the forward polarization) suggest that all materials(monolithic alloy and AMCs) have been subjectedto localized corro-sion processes. The great similarity of the polarization curve shapes,the nearly same areas of the hysteresis loops, as well as the similarcorrosion potential (Ecor) values and anodic-to-cathodic transitionpotential (Ea/c tr) values (despite the different volume fractions andparticle reinforcements),indicate that the corrosion behaviour ofthe tested materials was mainly controlled by the corrosion of themonolithic alloy.The only differences that are worthwhile to mention concern thecathodiccurrent densityvaluesrecorded. Thesedifferences aremore clearly seen in Fig. 13, which includes only the cathodic partof selected forward polarization curves. The WC reinforced mate-rials present higher cathodic current densities compared with theTiC reinforced materials (Figs. 12a and d and 13a). This can be asso-ciated with (a) the coarse Al12W particles, which have large enoughsurfaces tosustain cathodic reactions, and (b) theincreased areafraction of noble intermetallic particles (Al12W, Al5Wand Al3(Ti,W))due to the relatively high reactivity of WC particles, as aforemen-tioned in Sections 3.1.2 and 3.2.1. Furthermore, the alloy reinforcedwiththehighvolumefractionof carbidephasepresentshighercathodic current densities than the alloy reinforced with the lowvolume fraction of carbide phase, for the same carbide type(Figs. 12b and c and 13b). This trend may also be associated withthe increased area fraction of noble intermetallic particles that havesufcient surface areato sustain thecathodic reactions [41]. Theabove postulations are going to be investigated by SEMexaminationin Section 3.3.2.3.3.2. Microstructure of corrosion and correlation with polarizationbehaviourSEM micrographs of the corroded materials (cross sections) aregiven in Fig. 14(ad). The main degradation morphology is inter-granular corrosion (IC) associated with the presence of the AlFe/Al or AlFeSi/Al eutectic microconstituent. In order to clarify thepredominant mechanism of corrosion, one has to initially considerthe generally accepted four steps involved in the localized corro-sion of aluminium [88]:(1)The adsorption of the reactive anions on localized sites of thesurface lm of aluminium (i.e. sites where the lm presentsinhomogeneities [89]). In the presence of incoherent precip-itates, such as Al3Fe [90] and aAlFe(Mn)Si [91], the precipi-tate/Al interfaces would be the preferred sites for adsorption.(2)Thechemicalreaction oftheadsorbedanionwiththealu-minium ion in the aluminium oxide/hydroxide lattice.(3)The thinning of the oxide lm by dissolution. This dissolu-tionisaawassisted/awcenteredprocess. (Thepassivelm on Al-alloys exhibits semi-conductive properties owingto the non-stoichiometry of composition and local structuralinhomogeneities [89]).(4)The direct attack of the exposed metal by the anion possiblyassisted by an anodic potential.Al-FebasedintermetalliccompoundsarenoblerthantheirAlmatrix[89]. AtthepHof theelectrolyte(pH ~ 5.0), bothAlFeintermetallicsandAl displayregionsofpassivity[92]. However,the lm over an AlFe based intermetallic phase is thin and elec-tronically conductive; thus, when in galvanic couple with Al, it stillhas a proven ability to efciently sustain cathodic reactions [93].Basedontheabovepostulations, thefollowingmechanismisconsidered to have taken place:(1)Onceaggressiveanionadsorptiononthealuminiumoxidesurface lm at the Al/AlFe intermetallic interface occurred,an active centre was developed. The active centre was thenthe site for accelerated lm thinning [88].-2000-1500-1000-500050010001500Potenal (mV, Ag/AgCl)-2000-1500-1000-5000500100015000.00001 0.0001 0.001 0.01 0.1 1 10 100Potenal (mV, Ag/AgCl)Current density (mA/cm2)0.00001 0.0001 0.001 0.01 0.1 1 10 100Current density (mA/cm2)0.00001 0.0001 0.001 0.01 0.1 1 10 100Current density (mA/cm2)0.00001 0.0001 0.001 0.01 0.1 1 10 100Current density (mA/cm2)-2000-1500-1000-500050010001500Potenal (mV, Ag/AgCl)-2000-1500-1000-500050010001500Potenal (mV, Ag/AgCl)Fig. 12. Cyclicpotentiodynamicpolarizationcurvesofmonolithicandreinforcedmaterials at various combinations of carbide reinforcement in terms of carbide typeandvolume fraction(DHS, 25 C). (a) Variationof carbide type (1.0 vol%), (b)variationof TiCvolumefraction, (c) variationof WCvolumefraction, and(d)variation of carbide type (0.7 vol% TiC, 0.5 vol% WC).A. Lekatou et al. / Materials and Design 65 (2015) 11211135 1131(2)Once the lm was sufciently thinned,direct attack of theexposed metallic Al occurred. Because the lm was thinnedlocally, the attack on the metal was also concentrated. In thisstage, theelectrochemicalbehaviouroftheironaluminidebecame of major importance. Consequently, due to theelectrochemical potential difference between Al andAlFe or Al and AlFeSi (even with the lm on the interme-tallicphase), localizeddissolutionof theanodicAl (inoradjacent to the eutectic microconstituent) occurred andsmall pits were formed.(3)Asthepitsweregettingdeeper, differential aerationcellswereformedbetweenthebottomof thepitsandthepitwalls.(4)Pitting was evolved to intergranular corrosion at the bound-aries where the Al-Fe intermetallics exist.On closer inspection of Fig. 14,the following observations aremade: (a) features associated with the carbide reinforcement (clus-tersofaluminidesandcarbides)alongwiththeirAl-matrix haveremained intact of corrosion; (b) often, aluminide and carbide clus-tershave inhibited ICprogressby actingas physicalbarriers; (c)furthermore, in Fig. 14c, the voids around Al12W particles indicatethat the latter have acted as large cathodic sites.The microstructural observations correlate well with the poten-tiodynamicpolarizationperformanceofthereinforcedmaterials,with respect to the following aspects:(a)The fact that corrosion was mainly associated with the inter-granular iron aluminide phase (afeature of the monolithicalloy), whilst features associated with the carbide reinforce-ment have remained essentially free of corrosion traces sup-port the claimderiving fromthe similar polarizationbehaviourinFig. 12andformulatedinSection3.3.1: thecorrosionbehaviour of the tested materials was mainlycontrolled by the corrosion of the monolithic alloy.(b)The only notable evidence of corrosion associated with thereinforcement istheanodicdissolutionofAlaroundA12Windicatingacathodic rolefor thelargesurfaceof Al12Wparticles. Thisevidencecorrelateswell withtherelativelyhighcathodiccurrentsrecordedduringpolarizationoftheWCreinforcedAMCs (incomparisonwiththeTiC-AMCs,-1700-1500-1300-1100-900-700-500-3000.0001 0.001 0.01 0.1 1 10Potential (mV, Ag/AgCl)Current density (mA/cm2)Forward Al 1050Forward Al-1.0 vol%TiCForward Al-1.0 vol%WC-1700-1500-1300-1100-900-700-500-3000.0001 0.001 0.01 0.1 1 10Potential (mV, Ag/AgCl)Current density (mA/cm2)Forward Al 1050Forward Al-0.5 vol%WCForwardAl-1.0 vol%WCabFig. 13. Cathodic polarization curves of monolithic and reinforced materials (DHS,25 C). (a) Variationof carbidetype(1.0 vol%), and(b) variationof WCvolumefraction.a1.0 vol% TiC b1.0 vol% WCd 1.0 vol% WC c 1.0 vol% WCFig. 14. SEM micrographs of AMCs after cyclic polarization in DHS, 25 C (cross-sections). (a) and (b): extensive intergranular corrosion associated with the AlFe/Al andAlFeSi/Al eutectic microconstituent; (c) evidence that Al12W particles have acted as cathodic surfaces. (d) WC & aluminide clusters and their matrix appear corrosion-free.Black outlined ellipses: intergranular clusters of aluminides (Al5W, Al3Ti)/carbides and their matrix have remained free of corrosion signs. White arrows: aluminide (Al5W,Al3Ti) and carbide clusters have acted as physical barriers to IC progress. White outlined ellipses: clusters of aluminides/carbides and their matrix have remained free ofcorrosion signs.1132 A. Lekatou et al. / Materials and Design 65 (2015) 11211135Fig. 12aanddandFig. 13a)andthe1.0 vol%WC-AMC(incomparison with the 0.5 vol% WC-AMC, Figs. 12c and 13b).According to Birbilis and Buchheit [94], the size of interme-tallicparticlesplaysanimportant roleonthekineticsofcathodic reactions, as this will govern the amount of currentthe intermetallic can support.It should also be noted, that although clusters of aluminide par-ticles and carbide nanocores do not appear to have acted as catho-dicsites(even inthecaseoflargesurface areaoccupation, asinFig. 14ac), itisreasonabletoassumethattheentirevolumeofhigh conductivity may have also contributed a cathodic inuenceon the matrix in compatibility with Deuis et al. [95].Theaboveobservationssuggestthatthematrix/primaryrein-forcement interface was not the mainfactor inthe corrosionbehaviourofthesecomposites. ThefactthattheAl/WCandtheAl/TiCinterfaceshaveremainedunaffectedinconjunctionwiththeslidingwear response, indicatetheexistenceof cleanandstrong reinforcement-matrix interfaces.4. ConclusionsAluminiumMatrixComposites(AMCs)wereproducedbytheaddition of sub-micron sized WC and TiC particles (61.0 vol%) intoa melt of Al1050. Casting was assisted by the employment of K2TiF6as a wetting agent and mechanical stirring. The main conclusionsdrawnfromthestudyof themicrostructure, wearperformanceand corrosion performance of the cast materials are:v The AMC contained two main types of reinforcement: (a) in-situAl3Ti and(Al5W + Al12W + Al3(Ti,W))intermetallicparticlesinthe cases of TiCp-AMC and WCp-AMC, respectively; (b) carbidenanoparticles, as such, or moreoftenas clusters of carbidenanocores and aluminide particles.v Particle distribution was considered as reasonably uniformcomprising both clusters and isolated particles. Carbide nano-cores, inclusters withaluminides or isolated were mainlylocated at grain boundaries and the areas of the lastly solidiedliquid.v The sliding wear performance of the alloy was markedlyimproved by the addition of the carbide phase through a directeffect stemming from the TiC and WC particles, as such and asclusterswithintermetallicphase, andanindirecteffectorigi-nating from the intermetallic hard particles.v WCp-AMCs presented higher wear resistance than TiCp-AMCs.v Awearmechanismwasformulatedincludingfoursteps: (a)plastic deformation of the Al-matrix leading to a repeatedhill-valley morphologyof thewear surface; (b) formationof surfacealuminabasedoxidelayersduetofrictional heat-ing; (c) crack formation due to friction, fatigue and brittlenessof theoxidesurfacelayers; (d) negrooveformationalongthe sliding direction. The wear mode was characterizedasmild.v Thecorrosionbehaviour of thereinforcedmaterials inDHS,at 25 C, was mainly controlled by the corrosion of thealloymatrix. As such, the predominatingformof corrosionwas intergranular corrosion(IC) of Al intheAl/ironalumi-nide eutectic microconstituent or adjacent to the grainboundaries.v Features associated with carbide reinforcement (clusters ofaluminide and carbide nanoparticles) along with their Al-matrixhave remainedintact of corrosion, whilst inseveralcases, theyhaveactedasphysicalbarrierstotheICprogress.However, a cathodic action of large Al12Wparticles wasobserved.v Overall, the addition of submicron reinforcing carbide particleshad a benecial effect on the wear response of the monolithicmatrix, whilstitdidnotworsenthecorrosionresponseirre-spective of the particle volume fraction.AcknowledgementsTo the Greek General Secretariat for Research & Technology andtheEuropeanCommission(NSRF2007-2013) withintheframe-work of Joint Research & Technology Programs /Hellas-Hungary.References[1] KainerKU. Basicsof metal matrixcomposites. In: KainerKU, editor. Metalmatrixcomposites: custom-madematerials for automotiveandaerospaceengineering. Weinheim: Wiley-VCH Verlag GmbH & Co; 2006. p. 152.[2] SurappaMK. Aluminiummatrixcomposites: challenges andopportunities.Sadhana 2003;28:31934.[3] Surrapa MK, Rohatgi PK. Preparation and properties of cast aluminium ceramicparticle composites. J Mater Sci 1981;16:98393.[4] Karbalaei Akbari M, MirzaeeO, Baharvandi HR. Fabricationandstudyonmechanical properties andfracturebehavior of nanometricAl2O3particle-reinforcedA356compositesfocusingontheparametersof Vortexmethod.Mater Des 2013;46:199205.[5] AhamedH, SenthilkumarV. Experimentalinvestigationonnewlydevelopedultrane-grained aluminium based nano-composites with improvedmechanical properties. Mater Des 2012;37:18292.[6] Mula S, Padhi P, Panigrahi SC, Pabi SK, Ghosh S. On structure and mechanicalpropertiesofultrasonicallycastAl-2%A1203nanocomposite. MaterResBull2009;44:115460.[7] Haoze Liu, Luhong Wang, Aimin Wang, Taiping Lou, Bingzhe Ding, Zuangqi Hu.Studies of SiC/Al nanocomposites under highpressure. Nanostruct Mater1997;9:2258.[8] Ezatpour HR, Sajjadi SA, Sabzevar MH, Huang Yizhong. Investigation ofmicrostructure and mechanical properties of Al6061-nanocompositefabricated by stir casting. Mater Des 2014;55:9218.[9] TahamtanS, HalvaeeA, EmamyM, ZabihiMS. FabricationofAl/A206-Al2O3nano/micro composite by combining ball milling and stir casting technology.Mater Des 2013;49:34759.[10] Su H, Gao W, Feng Z, Lu Z. Processing, microstructure and tensile properties ofnano-sized Al2O3 particle reinforced aluminum matrix composites. Mater Des2012;36:5906.[11] Sharitabar M, Sarani A, KhorshahianS, ShaeeAfarani M. Fabricationof5052Al/Al2O3nanoceramic particle reinforced composite via friction stirprocessing route. Mater Des 2011;32:416472.[12] Mazahery A, Abdizadeh H, Baharvandi HR. Development of high-performanceA356/nano-Al2O3 composites. Mater Sci Eng A 2009;518:614.[13] Dehghan Hamedan A, Shahmiri M. Production of A356-1 wt% SiCnanocompositebythemodiedstircastingmethod. MaterSciEngA2012;556:9216.[14] AlizadehA, Taheri-Nassaj E, Hajizamani M. Hotextrusionprocesseffectonmechanical behavior of stir cast Al based composites reinforced withmechanically milled B4C nanoparticles. J Mater Sci Technol 2011;27:11139.[15] WangJ, Yi D, SuX, YinF, Li H. Properties of submicronA1Nparticulatereinforced aluminum matrix composite. Mater Des 2009;30:7881.[16] AnsaryYarA, MontazerianM, AbdizadehH. Microstructureandmechanicalproperties of aluminum alloy matrix composite reinforced with nano-particleMgO. J Alloys Compd 2009;484:4004.[17] NaidichJV. Thewettabilityofsolidsbyliquidmetals. ProgSurfMembrSci1998;14:353492.[18] MavrosH, KarantzalisAE, LekatouA. Solidicationobservationsandslidingwear behaviour of cast TiC particulate reinforced AlMgSi matrix composites. JCompos Mater 2013;47:214962.[19] GopalakrishnanS, MuruganN. Productionandwearcharacterizationof AA6061 matrixtitanium carbide particulatereinforced compositebyenhancedstir casting method. Composites Part B 2012;43:3028.[20] KaftelenH, nlN, Gller G, LtveogluM, HeneinH. Comparativeprocessing-structureproperty studies of AlCu matrix composites reinforcedwith TiC particulates. Composites Part A 2011;42:81224.[21] KarantzalisAE, LekatouA, GeorgatisE, PoulasV, MavrosH. MicrostructuralobservationsofcastAlSiCuTiCcompositematerial. JMaterEngPerform2010;19:58590.[22] PeijieLi, KandalovaEG, NikitinVI. InsitusynthesisofAlTiCinaluminummelt. Mater Lett 2005;59:25458.[23] Contreras A, Bedolla E, Prez R. Interfacial phenomena in wettability of TiC byAlMg alloys. Acta Mater 2004;52:98594.[24] LopezVH, KennedyAR. Flux-assistedwettingandspreadingofAl onTiC. JColloid Interface Sci 2006;298:35662.[25] Karantzalis AE, Lekatou A, Georgatis E, Tsiligiannis Th, Mavros H. Solidicationobservations in dendritic cast Al-alloys reinforced with TiC particles. J MaterEng Perform 2010;19:126875.A. Lekatou et al. / Materials and Design 65 (2015) 11211135 1133[26] Baumli P, SychevJ, Budai I, SzaboJT, KaptayG. Fabricationofcarbonberreinforcedaluminummatrix composites via titanium-ioncontaining ux.Composites Part A 2013;44:4750.[27] Rajan HB, Ramabalan S, Dinaharan I, Vijay SJ. Synthesis and characterization ofinsituformedtitaniumdiborideparticulatereinforcedAA7075aluminumalloy cast composites. Mater Des 2013;44:43845.[28] Kalaiselvan K, Murugan N, Parameswaran S. Production and characterizationof AA6061-B4C stir cast composite. Mater Des 2011;32:40049.[29] Toptan F, Kilicarslan A, Karaaslan A, CigdemM, Kerti I. Processing andmicrostructural characterization of AA 1070 and AA 6063 matrix B4Cpreinforced composites. Mater Des 2010;31:S8791.[30] DeSalazarJMG, UrenaA, ManzanedoS, BarrenaMI. CorrosionbehaviorofAA6061 and AA7005 reinforced with Al2O3 particles in aerated 3.5% chloridesolutions: potentiodynamic measurements and microstructure evaluation.Corros Sci 1999;41:52945.[31] Shimizu Y, Nishimura T, Matsushima I. Corrosion resistance of Al-based metalmatrix composites. Mater Sci Eng A 1995;198:1138.[32] Buarzaiga MM, Thorpe SJ. Corrosionbehaviour of as-cast silicon carbideparticulate aluminium alloy metal-matrix composites. Corrosion Metals Alloys1994;50:17685.[33] Trzaskoma PP. Pit morphology of aluminumalloy and silicon carbide/aluminumalloymetalmatrixcomposites. CorrosionMetalsAlloys1990;46:4029.[34] Mclntyre JF, Conrad RK, Golledge SL. Technical note: the effect of heattreatment on the pitting behavior of SiCW/AA2124. Corrosion 1990;46:9025.[35] Zhu J, Hihara LH. Corrosion of continuous alumina-bre reinforced Al-2 wt.%CuT6 metalmatrix composite in 3.15 wt.% NaCl solution. Corros Sci 2010;52:40615.[36] Nath D, Namboodhirt TKG. Some corrosion characteristics of aluminium-micaparticulate composites. Corros Sci 1989;29:121521.[37] Hihara LH, Latanision RM. Galvanic corrosion of aluminium-matrixcomposites. Corrosion 1992;48:54652.[38] TrowsdaleAJ,Noble B,Harris SJ, GibbinsISR, ThompsonGE,Woods GC. Theinuence of silicon carbide reinforcement on the pitting behaviour ofaluminium. Corros Sci 1996;38:17791.[39] Paciej RC, AgarwalaVS. Inuenceof processingvariablesonthecorrosionsusceptibility of metal-matrix composites. Corrosion 1998;44:6804.[40] Pardo A, Merino MC, Merino S, Viejo F, Carboneras M, Arrabal R. Inuence ofreinforcement proportion and matrix composition on pitting corrosionbehaviour of cast aluminiummatrixcomposites (A3xx.x/SiCp). Corros Sci2005;47:175064.[41] GrifthsAJ, TurnbullA. Aninvestigationoftheelectrochemicalpolarizationbehaviour of 6061 aluminum metal matrix composites. Corros Sci1994;36:2335.[42] KiourtsidisGE, SkolianosSM, PavlidouEG. AstudyonpittingbehaviourofAA1913: SiCp composites using the double cycle polarization technique.Corros Sci 1999;41:1183203.[43] Alaneme KK, Bodunrin MO. Corrosion behavior of alumina reinforcedaluminum(6063) metal matrixcomposites. J Miner Mater Character Eng2011;10:115365.[44] Candan S. An investigation on corrosion behaviour of pressure inltrated AlMg alloy/SiCp composites. Corros Sci 2009;51:13928.[45] Seah KHW, Krishna M, Vijayalakshmi VT, Uchil J. Corrosion behaviour of garnetparticulate reinforced LM13 Al alloy MMCs. Corros Sci 2002;44:91725.[46] Jameel AA, Nagaswarupa HP, Krupakara PV, Vijayamma KC. CorrosioncharacterizationofAl6061/Zirconmetalmatrixcompositesinacidchloridemediums by open circuit potential studies. Int J Appl Chem 2009;5:110.[47] Nagaswarupa HP, Bheemanna HG, Banuprakash G. Electrochemical studies ofAl 7075/zirconmetal matrixcompositesinnatural seawater. UltraChem2012;8:31928.[48] ToptanF, AlvesAC, Kerti I, ArizaE, RochaLA. Corrosionandtribocorrosionbehaviour of AlSiCuMg alloy and its composites reinforced withB4Cparticles in 0.05 M NaCl solution. Wear 2013;306:2735.[49] Bindumadhavan PN, Wah HK, Prabhaka O. Dual particle size (DPS) composites:effect on wear and mechanical properties of particulate metal matrixcomposites. Wear 2001;248:11220.[50] Mandal A, MurtyBS, ChakrabortyM. Slidingwearbehaviourof T6treatedA356-TiB2 in-situ composites. Wear 2009;266:86572.[51] Yang LJ. The effect of nominal specimen contact area on the wear coefcient ofA6061 aluminium matrix composite reinforced with alumina particles. Wear2007;263:93948.[52] KkM, zdinK. Wear resistanceof aluminiumalloyandits compositesreinforced by Al2O3 particles. J Mater Process Technol 2007;83:3019.[53] Yalcin Y, Akbulut H. Dry wear properties of A356-SiC particle reinforced MMCsproduced by two melting routes. Mater Des 2006;27:87281.[54] Deuis RL, Subramanian C, Yellup JM. Dry sliding wear of aluminiumcomposites a review. Compos Sci Technol 1997;57:41535.[55] Sannino AP, Rack HJ. Dry sliding wear of discontinuously reinforcedaluminium composites: review and discussion. Wear 1995;189:119.[56] Veeresh Kumar GB, Rao CSP, Selvaraj N. Mechanical and tribological behaviorof particulatereinforcedaluminummetal matrixcompositesareview. JMiner Mater Character Eng 2011;10:5991.[57] Suh NP. The delamination theory of wear. Wear 1973;25:11124.[58] Archard JF. Contact and rubbing of at surfaces. J Appl Phys 1953;24:9818.[59] HeilmannP, DonJ, SunTC, Rigney DA. Sliding wear andtransfer. Wear1983;91: 17190.[60] Al-Qutub AM, Allam IM, Qureshi TW. Effect of sub-micron Al2O3 concentrationon dry wear properties of 6061 aluminum based composite. J Mater ProcessTechnol 2006;172:32731.[61] Karbalaei Akbari M, Baharvandi HR, Mirzaee O. Nano-sized aluminumoxidereinforcedcommercial castingA356alloymatrix: evaluationof hardness,wear resistance and compressive strength focusing on particle distribution inaluminum matrix. Composites Part B 2013;52:2628.[62] Mohammad Shari E, Karimzadeh F, Enayati MH. Fabrication and evaluation ofmechanical and tribological properties of boron carbide reinforced aluminummatrix nanocomposites. Mater Des 2011;32:326371.[63] SameezadehM, EmamyM, Farhangi H. EffectsofparticulatereinforcementandheattreatmentonthehardnessandwearpropertiesofAA2024-MoSi2nanocomposites. Mater Des 2011;32:215764.[64] AnyaCC. Wet erosivewearof aluminaanditscompositeswithSiCnano-particles. Ceram Int 1998;24:53342.[65] Ravindran P, Manisekar K, Vinoth Kumar S, Rathikac P. Investigation ofmicrostructure and mechanical properties of aluminum hybrid nanocomposites with the additions of solid lubricant. Mater Des 2013;51:44856.[66] Nemati N, Khosroshahi R, Emamy M, Zolriasatein A. Investigation ofmicrostructure, hardness and wear properties of Al-4.5 wt.% CuTiCnanocomposites produced by mechanical milling. Mater Des 2011;32:371829.[67] Battocchi D, SimesAM, TallmanDE, BierwagenGP. Comparisonoftestingsolutionsontheprotectionof Al-alloysusingaMg-richprimer. CorrosSci2006;48:222640.[68] Witusiewicz VT, Bonder AA, Hecht U, Rex S, Velikanova TYa. The AlBNbTisystem III. Thermodynamic re-evaluation of the constituent binary system AlTi. J Alloys Compd 2008;465:6477.[69] Franke P, Neuschtz D, editors. Elements and binary systemsfrom AgAItoAuTI, series: Landolt-Brnstein: numerical data and functional relationshipsin science and technology Newseries, Part 19B1, Subseries: PhysicalChemistry. Berlin, Heidelberg: Springer; 2002. p. 2146.[70] NPL (National Physical Laboratory), MTDATA Phase Diagram Software, Ti-Wphase diagram, calculated 15-07-2003, .[71] NPL (National Physical Laboratory), MTDATA Phase Diagram Software, AlCphase diagram, calculated 09-03-1999, .[72] El-Mahallawy N, Taha MA, Jarfors AEW, Fredriksson H. On the reactionbetween aluminium, K2TiF6 and KBF4. J Alloys Compd 1999;292:2219.[73] Birol Y. Production of AlTiB master alloys from Ti sponge and KBF4. J AlloysCompd 2007;440:10812.[74] ContrerasA, Angeles-ChvezC, FloresO, PerezR. Structural, morphologicalandinterfacial characterizationof AlMg/TiCcomposites. Mater Character2007;58:68593.[75] House JE. Inorganic chemistry. Academic Press; 2012. p. 353.[76] JahntekM, Kraj M, HafnerJ. Interatomicbonding, elasticproperties, andideal strength of transition metal aluminides: a case study for Al3(V,Ti). PhysRev B 2005;71. 024101(1-16).[77] Lide DR, editor. CRC Handbook of Chemistry and Physics. 72nd ed. CRC Press;19911992.[78] GeorgatisE, LekatouA, KarantzalisAE, PetropoulosH, KatsamakisS, PouliaA. Development of a cast AlMg2SiSi in-situcomposite: microstructure,heat treatment andmechanical properties. J Mater EngPerform2013;22:72941.[79] ZubkoAM, LobanovVG, Nikonova VV. Reactionof foreignparticles withcrystallization front. Sov Phys Crystallogr 1978;18:23941.[80] Tyagi R. Synthesis andtribological characterizationof insitucast AlTiCcomposites. Wear 2005;259:56976.[81] RanganathG, SharmaSC, KrishnaM. Dryslidingwearof garnetreinforcedzinc/aluminium metal matrix composites. Wear 2001;251:140813.[82] Venkataraman B, Sundararajan G. The sliding wear behaviour of AlSiCparticulate composites I. Macrobehaviour. Acta Mater 1996;44:45160.[83] Sannino AP,Rack HJ. Tribological investigation of 2009 Al-20 vol% SiCp/17-4PH Part I: composite performance. Wear 1996;197:1519.[84] Sarkar AD. Friction and wear. London: Academic Press; 1980. p. 2059.[85] Urena A, Rams J, Campo M, Snchez M. Effect of reinforcement coatings on thedryslidingwearbehaviourofaluminium/SiCparticles/carbonbreshybridcomposites. Wear 2009;266:112836.[86] Las L, Rodrigez-Ibabe JM. Wear behaviour of eutectic and hypereutectic AlSiCuMg casting alloys tested against a composite brake pad. Mater Sci Eng A2003;363:193200.[87] Niu Haiyang, Chen Xing-Qiu, Liu Peitao, Xing Weiwei, Cheng Xiyue, LiDianzhong, et al. Extra-electron induced covalent strengthening andgeneralizationof intrinsic ductile-to-brittle criterion. Sci Rep2012;2:718.10.1038/srep0071.[88] Foley RT. Localized corrosion of aluminum alloys a review. Corrosion 1986;42:27788.[89] Szklarska-SmialowskaZ. Pittingcorrosionofaluminum. CorrosSci1999;41:174367.[90] HolmK, Hornbogen E. Annealing of supersaturated and deformed Al-0.042 wt% Fe solid solutions. J Mater Sci 1970;5:65562.[91] Stickels CA, Bush RH. Precipitation in the System Al0.05 wt pct Fe. Metall Trans1971;2:203142.1134 A. Lekatou et al. / Materials and Design 65 (2015) 11211135[92] Birbilis N, Buchheit RG. Investigationanddiscussionof characteristics forintermetallic phases common to aluminum alloys as a function of solution pH.J Electrochem Soc 2008;155(3):C11726.[93] Stansbury EE, Buchanan RA. Fundamentals of electrochemical corrosion.Materials Park Ohio: ASM Int.; 2000. p. 326.[94] Birbilis N, Buchheit RG. Electrochemical characteristics of intermetallic phasesin aluminum alloys: an experimental survey and discussion. J Electrochem Soc2005;152(4):B14051.[95] Deuis RL, Green L, Subramanian C, Yellup JM. Corrosion behaviour ofaluminium composite coatings. Corrosion 1997;53:88090.A. Lekatou et al. / Materials and Design 65 (2015) 11211135 1135