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:
[email protected] (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
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