ArticlePDF Available

Effect of Ultrasonic Vibration and Interpass Temperature on Microstructure and Mechanical Properties of Cu-8Al-2Ni-2Fe-2Mn Alloy Fabricated by Wire Arc Additive Manufacturing

Authors:

Abstract and Figures

A novel ultrasonic vibration assisted (UVA) wire arc additive manufacturing (WAAM) was used to fabricate Cu-8Al-2Ni-2Fe-2Mn alloy in this study. The effect of different interpass temperatures with and without ultrasonic vibration on the microstructural evolution and mechanical properties of the fabricated part were investigated by optical microscope (OM), scanning electron microscope (SEM), transmission electron microscope (TEM), nanoindentation, and mechanical tensile testing. The results showed that reduction of the interpass temperature without UVA treatment cannot prevent the columnar dendrites directionally growing along the deposition direction. Under the UVA treatment, the coarse columnar dendrites were broken at the interpass temperature of 400 °C, and formed a fine cellular structure with an interpass temperature of 100 °C, owing to the acoustic streaming effect and cavitation effect. In addition, globular κII phase was based on Fe3Al and lamellar κIII phase was based on NiAl distributed in the interdendritic region, whereas κIV phase (rich-Fe) were precipitated in the α-Cu matrix. The improvement of microstructural characteristics caused by UVA treatment further improved the tensile properties and nano-hardness of WAAM fabricated parts. Eventually, it is experimentally demonstrated that WAAM fabricated Cu-8Al-2Ni-2Mn-2Fe alloy can obtain high-performance at UVA process under an interpass temperature of 100 °C.
Content may be subject to copyright.
Metals2020,10,215;doi:10.3390/met10020215www.mdpi.com/journal/metals
Article
EffectofUltrasonicVibrationandInterpass
TemperatureonMicrostructureandMechanical
PropertiesofCu8Al2Ni2Fe2MnAlloyFabricated
byWireArcAdditiveManufacturing
WeiChen
1,†
,YuhuaChen
1,
*
,†
,TimingZhang
1
,TaotaoWen
1
,ZuozhuYin
1
andXiaosongFeng
2
1
JiangxiKeyLaboratoryofFormingandJoiningTechnologyforAerospaceComponents,Nanchang
HangkongUniversity,Nanchang330063,China;waam96@163.com(W.C.);
timingzhang1987@163.com(T.Z.);18702529802@163.com(T.W.);yinzuozhu@163.com(Z.Y.)
2
ShanghaiAerospaceEquipmentsManufacturer,Shanghai200245,China;fxsupc@163.com
*
Correspondence:ch.yu.hu@163.com;Tel.:+8613330067995
Theseauthorscontributedequallytothiswork.
Received:27December2019;Accepted:31January2020;Published:3February2020
Abstract:Anovelultrasonicvibrationassisted(UVA)wirearcadditivemanufacturing(WAAM)
wasusedtofabricateCu8Al2Ni2Fe2Mnalloyinthisstudy.Theeffectofdifferentinterpass
temperatureswithandwithoutultrasonicvibrationonthemicrostructuralevolutionand
mechanicalpropertiesofthefabricatedpartwereinvestigatedbyopticalmicroscope(OM),
scanningelectronmicroscope(SEM),transmissionelectronmicroscope(TEM),nanoindentation,
andmechanicaltensiletesting.Theresultsshowedthatreductionoftheinterpasstemperature
withoutUVAtreatmentcannotpreventthecolumnardendritesdirectionallygrowingalongthe
depositiondirection.UndertheUVAtreatment,thecoarsecolumnardendriteswerebrokenatthe
interpasstemperatureof400°C,andformedafinecellularstructurewithaninterpasstemperature
of100°C,owingtotheacousticstreamingeffectandcavitationeffect.Inaddition,globularκ
II
phase
wasbasedonFe
3
Alandlamellarκ
III
phasewasbasedonNiAldistributedintheinterdendritic
region,whereasκ
IV
phase(richFe)wereprecipitatedintheαCumatrix.Theimprovementof
microstructuralcharacteristicscausedbyUVAtreatmentfurtherimprovedthetensileproperties
andnanohardnessofWAAMfabricatedparts.Eventually,itisexperimentallydemonstratedthat
WAAMfabricatedCu8Al2Ni2Mn2FealloycanobtainhighperformanceatUVAprocessunder
aninterpasstemperatureof100°C.
Keywords:Cu8Al2Ni2Fe2Mnalloy;wirearcadditivemanufacturing;ultrasonicvibration;
interpasstemperature;microstructure;mechanicalproperties
1.Introduction
Nickelaluminumbronze(NAB)alloyshavehighstrengthandoxidation,corrosion,andwear
resistance,andhasbeenwidelyusedinmarineindustries[1,2].Furthermore,NABalloysdonotshow
abrittletemperaturerangeandareusefulforweldingandrepairingstructuralpartswithasimilar
chemicalcomposition[3].NABalloysareusuallycomposedofcopper(Cu),aluminum(Al),nickel
(Ni),iron(Fe),andmanganese(Mn)elements[4].Theadditionofalloyingelementsimprovesthe
mechanicalpropertiesofthealloyandinhibitstheformationofγ
2
phaseandAl
4
Cu
9
,whichare
deleterioustocorrosionresistanceduetothehighAlcontent[5].Inthecasting,NABalloysincludes
α‐Cumatrixphase,martensiticβphase(βʹphase),andfourintermetallicκphases(Ferichκ
I
,κ
II
,κ
IV
andNirichκ
III
phases)[2,6].Themicrostructureevolutionunderequilibriumstatecanbe
approximatelydescribedasthefollowingfourstages[2,6,7].Firstly,theprimaryα‐Cuphasestartsto
Metals2020,10,2152of21
formwithintheβphasearoundthetemperature1030°C.Secondly,aroundthetemperatureof930
°C,iftheweightpercentofFeisgreaterthan5%,thelargeprecipitates(κI)carryoutintheβphase,
andanothersmallintermetallicglobularκIIphaseformswithalowlevelofweightFe.Thirdly,atthe
temperatureof800°C,theremainingβphaseistransformedintotheintermetallicκIIIphasethrough
theeutectoidreaction.Finally,thesaturationsolubilityofFe(κIVphase)willprecipitateintheα‐Cu
matrixwhenthetemperaturefallsat860°C.Further,theretainedβʹphasemayappearintheNAB
alloysunderquenchedfrom1020°C[8].ThecomplexmicrostructureofthecastNABalloysincreases
thepossibilityofgalvaniccorrosionwhichisheldinseawater[9].Defectssuchasshrinkageporosity
andcoarsergrainsarepronetoappearinthecastalloysowingtotheirpoorcastingproperties.These
weaknessesareharmfultothemechanicalandcorrosionpropertiesofcomponents.
Inrecentyears,wirearcadditivemanufacturing(WAAM)hasgainedconsiderableinterestdue
toitshighdepositionrate,efficiency,fulldensity,andlowequipmentcost[10].InWAAMprocess,
anelectricarcisusedasaheatsourceandemployseithergasmetalarcwelding(GMAW[11]),gas
tungstenarcwelding(GTAW[12])orplasmaarcwelding(PAW[13])tomeltthewireasthefeed
stock.Ascomparedtotheconventionalcasting,thecoolingratesinWAAMaresignificantlyhigher
(WAAM:102K/s[14],cast:10K/s[6])toobtainfinegrains.Todata,somestudieshavereportedthe
propertiesofNABalloysfabricatedbytheWAAM.Dingetal.[15]foundthatthepostweldheat
treatmentfurtherrefinedthemicrostructure.Therewasanimprovementintensilestrengthanda
decreaseinelongationofthedepositedNiAlbronzealloy.Shenetal.[16]obtainedthemechanical
propertiesinlongitudinal,transverseandnormaldirectionsofthedepositedCu9Al4.5Ni3.5Fe
1.3Mnalloy.Theresultsshowedthedepositexhibitshigherstrengthinthelongitudinaland
transversedirectionthaninthenormaldirection.TheanisotropyofWAAMNABalloyscanbe
modifiedbyquenchingafter900°Cfor2handthentemperingat650°Cfor6h.Dharmendraetal.
[17]analyzedCu9Al4Fe4Ni1MnalloywereformedκIIandκIIIphaseintheinterdendriticregions
andprecipitatedκIVparticleswithintheα‐matrix.Theresearchalsopredictedtheheattreatmentsin
therangeof500–600°CmayincreasethestrengthlevelofWAAMNABalloys.Aboveall,the
propertiesofWAAMNABalloysmaybeimprovedbythepostheattreatment,butitalsoincreased
theenergyconsumptionandworkingperiodofthealloypreparation.
Basedonissuesrelatedtotheinhomogeneityofthestructureandtheanisotropypropertiesof
WAAMNABalloys,itisessentialtoexploreatypeofsupplementarymeansthatintegratesinthe
WAAMsystemtoimprovethealloysstrengthandminimizeanisotropy.Theacousticnonlinear
effects(includingspecificcavitationandacousticstreaming)arisinginultrasonicvibrationcan
influencetheconvectionandsolidificationbehaviorsofthemoltenpool[18].Hence,ultrasonic
vibrationhasbeeneffectivelyusedinmeltingmetalsolidificationprocessessuchascasting[19,20],
welding[21],cladding[22],additivemanufacturing[23–26].Wangetal.[25]discoveredthatlaser
engineerednetshapingofInconel718withultrasonicvibrationcouldrefinethemicrostructure,
decreasemicropore,improvehomogeneityofchemicalcontentsandshortenthesizeofLavesphases.
UltrasonicvibrationwasusedtoassistlaserengineerednetshapingtofabricateFeCrstainlesssteel
[23]and4047aluminum[26],whichhadanalogousphenomenainreducingmicroporeandrefining
grains.However,therearenoreportedinvestigationsonultrasonicvibrationassisted(UVA)
treatmentduringtheWAAMprocessoffabricationtheNABalloysparts.
Inthispaper,theNABalloy(Cu8Al2Ni2Fe2Mn)fabricatedbyWAAMwithandwithout
UVAtreatmentunderdifferentinterpasstemperatureswerestudied.Thechangesinmicrostructure,
elementaldistribution,phaseconstituent,nanohardness,andtensilepropertieswereinvestigated.
Moreover,theformationofmicrostructureandfracturecharacteristicwerealsodiscussedindetail.
2.ExperimentalProcedures
2.1.ExperimentalSetup
Aweldingmachineassociatedwithasystemofultrasonicvibrationwasutilizedtoconductthe
experimentalinvestigations.Figure1aillustratedtheexperimentalsetupofUVAWAAM,whichwas
composedofaweldsystem(FroniusTransPulsSynergic2700weldingmachine,Fronius,Pettenbach,
Metals2020,10,2153of21
Austria)withamotioncontrolsystem,asystemofultrasonicvibration.Ultrasonicvibrationwas
producedbyanultrasonicgenerator(JianPaiUltrasonicTechnologyCo.,Ltd.,Dongguan,China)
withamaximum2000Woutputpowerand20kHzvibrationfrequencies.Thevibrationdirection
wasperpendiculartothesubstratemetal.AninfraredthermometerandKtypethermocoupleswere
usedtomeasuretheinsituinterpasstemperature,therebycontrollingthedwelltimebetween
depositionoflayers.ThesamplingrateoftheKtypethermocouplewas1Hzandthemeasuring
rangeoftemperatureswasbetween0and1350°C.Inthisstudy,theultrasonicvibrationsystem
functionedatapowerof2000W,anddifferentinterpasstemperaturesof100°Cand400°Cwithand
withoutultrasonicvibrationwerecompared(Comparedwithahighertemperature,abetter
macroscopicmorphologycanbeobtainedattheinterpasstemperatureof400°C.Theinterpass
temperatureof100°Cisselectedwiththepurposeofavoidingtheinfluenceofheataccumulationon
themicrostructureandmechanicalproperties).
Figure1.(a)SchematicdiagramoftheexperimentalapparatusofultrasonicvibrationduringWAAM,
(b)typicalWAAMfabricatedsamples,(c)schematicconfigurationofthesamplepreparation,(d)
dimensionsoftensilespecimensinmm,(e)schematicdiagramofmicrostructureobservationposition.
2.2.Materials
CommerciallyavailableCu8Al2Ni2Fe2Mnalloywirewithadiameterof1.0mmwas
employedasafillermaterial,ofwhichthechemicalcompositionwas86Cu,8Al,2Ni,2Fe,and2Mn
(inwt.%)(madebyNewland(Tianjin)WeldingMaterialCo.,Ltd.,Tianjin,China.Thesubstratein
thepresentworkwasalowcarbonsteelplatewithdimensionsof180×100×6mm
3
.Beforethe
WAAMprocess,thesubstratesurfacewascleanedtoremovegreaseandoilwithacetone.Basedon
thepreliminaryexperimentalresults,somewallsampleswereproducedwiththesamewelding
parameter(arccurrent:97A,arcvoltage:10.3V,weldingrate:0.48m/min,wirefeedrate:4.0m/min,
nominalincrementofZaxis:1.8–2.2mm,nominalincrementofYaxis:3mmandargonflowrate15
L/min).Thedepositiontorchtiptoworkdistancewaskeptto15mm,andcontactanglewas90°for
alltheexperiments.Theweldingpathwasalternatedbackandforth,andthetypicalfabricatedwall
sampleswerepresentedinFigure1b.
ThesinglebeadwallswereusedtostudytheeffectofUVAandinterpasstemperatureonthe
microstructureandmechanicalproperties.Inordertoensuretheintegrityofthesamplingandthe
accuracyoftheresults,alltestedsamplesaredepositedwithalengthof120mmandaheightofnot
lessthan50mm.ThenomenclatureandprocessingparametersusedaregiveninTable1.
Metals2020,10,2154of21
Table1.Thenomenclatureandprocessingparametersofthesample.
NumbersInterpass
Temperature/°C
Ultrasonic
Power/W
Numberof
Deposition
Layers/Cycles
Height
/mm
Ratioof
Effective
Area/%
C1730–750‐1839.3665.5
C2Below400‐2353.7470.0
C3Below100‐2253.6579.5
U1710–73520002852.7653.9
U2Below40020002653.6875.9
U3Below10020002553.1984.8
2.3.CharacterizationofMicrostructureandProperties
Therequiredspecimensweremachinedbywireelectricdischargemachine(Maidenvan
MachineryCo.,Ltd.,Nanchang,China).Thespecimensusedformetallographicobservationand
nanohardnesstestwerecutfromthecenterarea(asshownbythedottedlineinFigure1c).Amixed
solutionof5vol.%ferricchloride,20vol.%hydrochloricacid,and75vol.%distilledwaterwasused
torevealthemicrostructureofthesamplesaftermounting,grindingandpolishing.The
microstructuresofthesampleswereexaminedusingaZeissAxioScopeA1opticalmicroscope(OM)
(CarlZeissMicroscopyGmbH,Gottingen,Germany)andHitachiSU1510scanningelectron
microscope(SEM)(Hitachi,Tokyo,Japan).FEITalosf200xtransmissionelectronmicroscope(TEM)
(FEI,Hillsboro,OR,USA)wasusedtoidentifytheelementaldistributionandphasesconstituentin
themicrostructure.Thesampleswerepreparedbymechanicalpolishingtoathicknessof50μm
followedbyfurtherthinningviaionbeammillingusingaGatanPrecisionIonPolishingSystem
(GatanModel691)(GatanInc.,Sarasota,FL,USA)operatingatanacceleratingvoltageof200kVfor
TEMstudy.Highangleannulardarkfield(HAADF)detectors(FEI,Hillsboro,OR,USA)wereused
incombinationwithEDStogenerateelementalmaps.Selectiveareadiractionwasusedtoobtain
patternsfromtheα‐Cumatrixandinterdendriticregion,respectively.
Tostudythetensilebehavior,cylindricaltensilesamples3mmindiameterwereextractedfrom
thedepositedwall,alongtheverticaldirectionandthehorizontaldirection.Hereinafter,thetwo
orientationswerereferredtoasVandH,respectively.Thelocationandsizeofthetensilespecimens
wereshowninFigure1c,d.ThecomparisonoftensilepropertiesbetweentheHandVdirectionwas
investigated.Consideringtheheterogeneousmechanicalpropertieswithinthefabricatedcomponent
[27],fivespecimensweremachinedindifferenthorizontalregionstorevealtheeffectsoflocationon
thetensileproperties(asshowninFigure1c).Inaddition,areferencecastNABsample,whichwas
closetothecompositionoftheweldingwire,wasusedtotestthetensileproperties.ThecastNAB
specificationcorrespondedtoUNSC95220andthenominalcomposition(inwt.%)ofthismaterial
was:Cu:81%minimum,Al:8.0–10.5%,Ni:2.5%,Fe:2.0–4.0%,Mn:0.5–2.0%.Thetensiletestswere
performedatroomtemperaturewithanextensionrateof0.5mm/minusinganInstron8872universal
testingmachine(InstronCorporation,Norwood,MA,USA).Thestrainwasmeasuredbythe
electronicextensometer.ThefracturesurfaceswereexaminedbySEM.Nanoindentationwasusedto
measurenanohardnessandmodulusofdifferentsamples.Thenanohardnesstestwasperformed
usingaNanomechanics(NanomechanicsInc.,OakRidge,TN,USA)andaloadingforceof50mN
atroomtemperature.Theindenterwasgraduallyloadedonthesurfaceofsamplewithin20s,and
thenkeptsteadyfor5s,finallyunloadedgraduallywithin2s.ThemeltingpointsofWAAM
fabricatedsamplewasdeterminedbyusingNetzschSta449F3differentialscanningcalorimetry
(DSC)(Netzsch,Selb,Germany)withaheatingrateof10°C/min.
Metals2020,10,2155of21
3.ResultsandDiscussion
3.1.Macrostructure
Theeffectofultrasonicvibrationandinterpasstemperatureonmacrostructureisshownin
Figure2.TheNABcomponentsfabricatedusingtheWAAMprocesswithorwithoutUVAtreatment
reachedfulldensityandnoporositydefectwasfoundinthecrosssections(C1andU1samplesare
continuousdepositionwithhigherinterpasstemperatureandpoorforming,whicharenotcompared
inthispaper).Itcanbefoundthatcolumnargrainsgrowepitaxiallyalongthedepositiondirection,
andtheinterlayerbands(ILB)areclearlyvisibleintheinterpasstemperature400°Csample(Figure
2a).Thesimilarresultsareshownintheinterpasstemperature100°Csamples,asthesizeofthe
columnargrainsdecreased,resultinginthereductionofthecontrast,whichmadeitdifficultto
characterizetheILB(Figure2b).IntheUVAtreatmentofWAAMspecimens,thegrainsizeisfurther
refinedandthedepositedlayerbecomeswider(Figure2c,d).Thewidthofthethinwallsincreased
from7.5–8.0mmto8.5–9.0mmaftertreatmentUVA.ThesizeoftheWAAMcomponentisdecided
bytheshapeofthemoltenpool[12].Theultrasonicvibrationisabletoacceleratetherelative
movementbetweenthesolidphaseandsurroundingliquidphaseinthemoltenpool,leadingtoa
betterspreadingofthemoltenpoolonthepreviouslydepositedlayers.Thisisbeneficialfor
increasingtheeffectiveareaofdeposition.Aschematicdiagramoftheeffectiveareaofdepositionis
shownontherightsideofthemacrostructureinFigure2,andtheresultsarelistedinTable1.The
experimentalresultsshowthattheU3sampleenablesthemaximumdepositedefficiency,reaching
84.8%.
Figure2.MorphologyofCu8Al2NiFe2Mnalloywithdifferentparameters(YZplane)and
effectivearea:(a)sampleC2,(b)sampleC3,(c)sampleU2,(d)sampleU3.(Themacroimageonthe
leftisanenlargedviewofthepositionintheblackdottedbox).
3.2.MicrostructureandComposition
Changesofthealloy’sgrainsizeareusuallyrelatedtotheweldprocessingparameters.
Interestingly,theinterpasstemperatureandultrasonicvibrationcouldalsoeffectivelychangethe
microstructureoftheNABalloyduringWAAMwhentheweldprocessingparametersremain
unchanged.Figure3showstheOMmicrographsofthemicrostructurebetweentwoadjacentILB
fabricatedbyWAAMindifferentinterpasstemperatureswithandwithoutUVAtreatment(YZ
plane,theobservationpositionisshowninFigure1e).ThemicrostructureoftheCu8Al2Ni2Fe
2MnalloyswithoutUVAtreatmentismainlycomposedofcolumnardendrites,asshowninFigure
3a1–a4,b1–b4.Thesizeandmorphologyofthemicrostructurearequitedifferent.Theregionof
elongatedcolumnargrainswithoutsecondarydendritearms(SDA)growthalongthedeposition
directionduetoahightemperaturegradient(G)andaslowgrowthrate(R)atthebottomofthe
moltenpooldelineatedtheILBandismarkedwithawhitearrow(Figure3a1,b1).Withthedecrease
ofthetemperaturegradientfromthebottomtothetopofthemoltenpool,theliquidmetalbecomes
Metals2020,10,2156of21
undercooledandleadstothedestabilizationofthesolidificationfront[28],thecolumnargrainsturn
tobecoarserandgraduallyformSDA(Figure3a3,b3).ThemorphologyofSDAisrelatedtothe
coolingrate[29].MoreSDAsarefoundinthesampleswhichundergohigherinterpasstemperatures.
Infact,atthetopofeachdepositedlayer,theepitaxialgrowthofthecolumnardendritesstructureis
blockedoffbyarandomorientedequiaxeddendritesstructure,theequiaxeddendriteswitha
thicknessofabout200–300μm(InsertpictureinFigure4a).Thetransitioncharacterofthetopregion
isrevealedinahighermagnification(Figure4a).Itisbelievablethatthecolumnartoequiaxed
transitionoccursatthetopofthemoltenpoolduetothedecreaseofthetemperaturegradientand
theincreaseoftheinterfacemovingvelocity[30].Thereisnosubsequentlayersremeltingit,hence,
onlythetopregionshowsthistransitioncharacteristic.
WhentheUVAtreatmentWAAMisundertheinterpasstemperatureof400°C,partofthe
columnardendritesstructureisfracturedandthegrainsizeisfurtherdecreased(asshowninFigure
3c1–c3).AtthetopofthemoltenpoolthatisneartheILB,abundantcolumnardendriteswithSDAs
arealsofound(Figure3c4).Hence,amixedmicrostructurecontainsthecoarsecellulardendrites,
cellulargrainsandcolumnardendrites,whichareformedinUVAtreatmentWAAMsampleswitha
higherinterpasstemperature.Withfurtherdecreaseofinterpasstemperatureat100°C,the
microstructureisrefinedtoformacellularstructureandthegrainboundariesareshortened(Figure
3d1–d4).Similarmicrostructureevolutionhasbeenreportedforfabricationusinglaseradditive
manufacturing[23,25]andweldingprocess[21,31].Ultrasonicvibrationinducesacousticstreaming
andcavitationcausedthestirringandmixingofliquidmaterialsinthemoltenpool[18,32,33].
Meanwhile,theeffectofheatingaccumulationcanbeeliminatedbycontrollingtheinterpass
temperature.Theactionsofstirringandmixingleadtoahomogenizingdensityandheatinthe
moltenpool,whichcontributetotheuniformdistributionofgrainsandtheformationofacellular
structurewithuniformgrainboundaries.
Figure3.ThemicrostructurebetweentwoadjacentILBindifferenttreatment(Y
Zplane):(a1
a4)
sampleC2,(b1
b4)sampleC3,(c1
c4)sampleU2,(d1
d4)sampleU3.
Metals2020,10,2157of21
ThemicrostructureofsomecharacteristicregionsisshowninFigure4.Figure4b,cshowsthe
comparisonsofthedepositionlayersnearthesubstrate.TheC2samplewithoutUVAtreatment
presentstheplanargrowthinthebondinginterfacesduetothehightemperaturegradientandslow
growthrateatthebottomofthemoltenpoolnearthesubstrateandismarkedbyablackarrow.With
thedecreaseofG/R,randomdistributionofdendritesformintoαCumatrixandthenchangeto
cellularstructure(Figure4c).Thiscellularstructuregrowsintheoppositedirectiontothe
temperaturegradientandeventuallyformscolumnardendrites.Similarphenomenaareobtainedfor
theC3sample.
IncomparisontotheWAAMwithoutUVAtreatment(Figure4b),theUVAtreatmentWAAM
producesaweakerbondinginterface,asshowninFigure4d.Further,thepressureintheacoustic
fieldmadethesolidliquidinterfacemorphologicallyunstableandfragmentedthedendritesand
epitaxialcellulartips[33].Thesedendritetipsareconveyedintotheundercooledmeltviaconvection.
Inthemeantime,alargeamountofgrainnucleationisformed.Hence,thealloyfabricatedwithUVA
treatmentgeneratedequiaxedgraininasmallergrainsize(about5–12μm)atthebottomregion.It
canbediscernedthatthereexisttwokindsofmorphologyofILB,whichisanimportantfeature.One
isthattheILBexhibittypicalcolumnargrainsorientupwards,whichoftenshowsalittledarker
contrastcompareswiththesurroundingsunderOM(theOMinsertinFigure4eandmacrostructure
inFigure2a).Thehighermagnificationobservationofthemicrostructure(Figure4e,f)showsthatthe
belowILBexpressestherelativelylargerprimarydendritearms(PDA)spacingwithSDAs.In
contrast,thesecondaryarmsarepartiallyremeltingintheILB.ThecharacteristicsofILBintheC2,
C3andU2samplesallconformtothisformationrule.OtherkindsofcamberedbandsinU3sample
showsthedifferentsizeofcellularstructureinILB(Figure4g).WiththemicrographinFigure3d1–
d4,themicrostructureoftheU3samplecanbedividedintothreeparts:ILB,finecellularzone,and
coarsecellularzone.TheformationofILBcanbeattributedtothepartialremeltingofthepreviously
depositedlayersandcomplexthermalhistoryexperiencedbytheWAAMprocess.Zhuetal.[34]also
foundthattheILBwerespaceduniformlyinthewholesample,whichdemonstratedthatitwasjust
aresultofthedepositinglayer.Furthermore,thetwoadjacentILBare2.0–2.2mminseparation,
whichcorrespondstotheaverageaddedlayerheight.Itiseasytomeasurethatthedepositionheight
ofthelastlayerismorethan2.4mmfortheconvenienceofexplainingtheILBareremeltingstructure.
Figure4hshowsahighermagnificationofinterdendriticstructurefromU3sample.The
microstructureconsistsofαCumatrixandperishableetchinterdendriticκphase.Besides,no
evidenceofretainβoracicularβmartensiteisobservedintheUVAtreatmentWAAMsample,a
similarresultisreportedinDharmendra’sstudies[17].However,SEMandOMaredifficultto
qualitativelyevaluatethetypeoftheprecipitationphaseintheUVAtreatmentWAAMsamples.The
TEMexaminationisdiscussedindetailbelow.
Figure4.Themicrostructureofsomecharacteristicregions(Y
Zplane):(a)thetransitionof
microstructurefromC2sample,(b)thebottomregionofC2sample,(c)thedendritestransformedto
cellularstructureinC2sample,(d)thebottomregionofU3sample,(e)interlayerbands(ILBregion)
inC2sample,(f)ILBofC2sample,(g)ILBofU3sample,(h)interdendriticregioninmagnification
fromU3sample.
Metals2020,10,2158of21
Asweknow,themicrostructureisverysensitivetotemperature,whilethedifferentlocationsof
thedepositedwallduringtheWAAMprocessundergorapidmeltingandsolidificationwitha
complexthermalhistory,soitisnecessarytoobservethemicrostructureindifferentdeposited
heights.Figure5showsthemicrostructureintheXZplaneofeachcomponentisdescribedinlower
andupperregions(thelocationisshowninFigure1e).AllofthesampleswithoutUVAtreatment
havecoarsecolumnardendritesgrownepitaxiallyalongthedepositiondirection,asshowninFigure
5a–d,whichaccordwiththeobservationintheYZplane.Continuouscolumnardendritesare
blockedbyUVAtreatmentintheU2sampleandthebrokencharactercanbeseenmoreclearlyinthe
highmagnificationSEMimage(Figure5e,f).AsimilarcellularstructureisfoundintheUVA
treatmentataninterpasstemperatureof100°Csample(Figure5g,h).Itisimportanttopointoutthat
differentverticallocationsstronglyeffectPDAspacing.IntheC2,C3,andU2samples,thePDA
spacingiscoarsenedwithincreasingdepositionheight.Afteranalysisbythequantitative
metallographymethod,theaveragePDAspaceoftheC2,C3,U2,andU3samplesisaround28.5±
4.2μm,21.6±3.6μm,26.8±3.1μm,and18.6±2.6μm,respectively.
Figure5.Effectofdepositedheightonmicrostructureindifferenttreatment(XZplane):(a,b)sample
C2,(c,d)sampleC3,(e,f)sampleU2,(g,h)sampleU3.
ThetypicalmicrostructureoftheXYplaneisshowninFigure6(theobservationlocationis
markedwithabluearrowinFigure1e).Sincethemicrostructuregrowsepitaxialalongthedeposition
direction(i.e.,Zdirection)inC2,C3andU2samples,itappearsascoarsedendriticstructureinC2
duetodevelopSDA(Figure6a,b).TypicalequiaxedgrainsarefoundinC3sample(Figure6c,d)in
thissection,whichresultinfewerSDAsatlowerinterpasstemperatures.Inthecrosssectionofthe
UVAtreatmentsamples,thesizeofdendriticstructureisfurtherreduced(i.e.,U2sampleinFigure
6e,f),andthemicrostructureoftheU3sampleissimilartothatoftheothertwocrosssections(Figure
6g,h),whichalsoprovesthatthemicrostructureformintheU3sampleismoreuniform.
Metals2020,10,2159of21
Figure6.Effectofdepositedheightonmicrostructureindifferenttreatment(XYplane):(a,b)sample
C2,(c,d)sampleC3,(e,f)sampleU2,(g,h)sampleU3.
TheWAAMprocessinvolvesalargermeltpooldiameterandlayerheightthanpowerbed
technologies,whichgenerallyleadtoagreatermicrostructuralheterogeneityandmoresevereILB
[35].Hence,themicrostructurebetweentwoadjacentILBchangesregularlywiththetemperature
gradientandgrowthrate,andthephenomenonismoreobviousunderthecontrolofinterpass
temperature.TheCu8Al2Ni2Fe2MnsamplesfabricatedbyWAAMwithoutUVAtreatmentnot
onlycontaincolumnardendrites,butalsoexhibitcolumnartoequiaxeddendritestransition.Kurz
andFisher[36]establishedasolidliquidinterfacemodelforinstabilitycriterionduringsolidification,
whichpointedouttheratioofG/RdeterminedthegrowthmorphologyandGVcontrolledthescale
ofthemicrostructureformed.Bermingham[37]reportedthethermalenvironmentwasdynamic
duringadditivemanufacturingandequiaxedgrainformationwasonlyachievablewhentemperature
gradientsdecreasesufficientlytopermitconstitutionalsupercooling.WiththedecreaseofG/Rfrom
thebottomtothetopofthemoltenpool,thegrainmorphologymightchangefromcolumnar
dendritestoequiaxialdendritesduetoseriousconstitutionalundercooling.Forcedcontrolofthe
interpasstemperaturecausesthedepositedlayeractasamainlyheatsink,whichmadetheamounts
ofheatfluxdissipatesdownwardduringthesolidificationprocessandresultsinregularchangesin
themicrostructurebetweenadjacentILB.
AsshowninFigure7a,theendothermicpeaksincalorimetriccurverevealthemeltingpointof
theWAAMCu8Al2Ni2Fe2Mnalloyisabout1020°C.Thetypicalthermalcyclemeasuresinthe
WAAMbuiltforasequenceofsixaddedlayersisshowninFigure7b.Thefirstcurveisthethermal
cycleofthermocoupleplungeintothemeltpool.Thepeaktemperatureofthesecondcurveisvery
closetothemeltingpointmeasuredbyDSC.Hence,wecanpredictthatthetemperatureofthearc
meltingpoolishighenoughtoremelttheequiaxeddendritesregion(Figure4a),whiletheSDAs
helptomaintainthegrowthdirectionofprimarydendritesinsubsequentlayers(Figures3a1–b4and
4f),andfinallyformstheepitaxialgrowth.
Figure7.(a)DSCcurveoftheWAAMNABsample,(b)
thethermalcyclefordepositedlayers.
Previousstudiesrevealthatthemicrostructureofmaterialsdependsonboththenucleation
conditionandsubsequentgrowthstage[19,26].TheUVAtreatmentWAAMatlowerinterpass
temperaturecanhinderthegrowthofepitaxialcolumnarmoreeffectivelyandpromotetheformation
ofnewgrains.Therefinementofthemicrostructuresmainlyresultsfromtheactionsofacoustic
streamingandcavitationinducedbyultrasonicvibration[18,32,33].Thepressureintheacoustic
streamingmakesthesolidliquidinterfacemorphologicallyunstableandfragmentsthecolumnar
graintips.Withtheassistanceofinstanthighacousticstreaming,theprimarycoarsecolumnargrains
arebrokenandfragmentedbecauseoftheforce.Thesedendritetipsareconveyedintothe
undercooledmeltviaconvectionandstirring.Inthemeantime,alargeamountofgrainnucleationis
formed.Withalowerinterpasstemperatureof100°C,theeffectofheataccumulationontheWAAM
processiseliminated.Eventually,WAAMpartswithrefinedgrainsareobtainedUVAtreatmentata
lowerinterpasstemperatureprocessduringthenucleationstage.
Metals2020,10,21510of21
Inordertorevealthevariousintermetallicphasesintheinterdendriticregionsandα‐Cumatrix,
detailedTEMobservationsoftheU3samplewerestudied.Figure8andFigure9showthe
distributionfeaturesofelementsattheinterdendriticregionsandαCumatrix.Theinterdendiritic
regionsdisplaytwotypesofintermetallicphases,κ
II
phaseinthedifferentglobularsizes(Rangefrom
somenanometerstotensnanometers)andκ
III
phaseinthelamellarmorphology,andEDSelemental
mappingdistinctlyrevealsthemorphology(Figure8).Theκ
II
andκ
III
phasearefoundtobeFe
3
Aland
NiAl,respectively,basedontheselectedareadiffractionpatternanalysis,asshowninFigure
10a,b,d,e.Someamountsareveryfinetothesizesofκ
IV
phase(Ferich,5to10nminsize)precipitate
isobservedintheαCumatrix(Figure9).TheobservationoffineFerichprecipitationand
precipitationfreeareasisdistributingnonuniformlyintheα‐Cumatrix,whichmayberesultedby
thedifferenteffectsofthermalcyclingoftheWAAMondifferentlocations.Selectedareadiffraction
patternanalysisrevealstheκ
IV
precipitateshaveacompositionandcrystalstructureissimilartothat
ofκ
II
basedonFe
3
Al(Figure10c,f).Itispossiblethattheκ
II
andκ
IV
arebeingprecipitatedatthesame
temperaturefromthesamephase[5].Thesmallanddisperseκ
IV
phasesmayinhibitgrainboundary
mobility,whichbeneficialforimprovingstrength[38],andalsopreventinggraingrowth[4].
Figure8.(a)HAADFimageoftheUVA
treatment
WAAMintheinterdendriticregion,(bf)
correspondingEDSelementalmapsforCu,Al,Fe,Ni,andMn,respectively.Theglobularprecipitates
areFe
3
Al(κ
II
)andthelamellarprecipitatesareNiAl(κ
III
)intheinterdendriticregion.
Metals2020,10,21511of21
Figure9.(a)HAADFimageoftheUVA
treatment
WAAMintheαCumatrixregion,(bf)
correspondingEDSelementalmapsforCu,Al,Fe,Ni,andMn,respectively.TheprecipitatesareFe
rich(κ
IV
)intheα‐Cumatrix.
Figure10.TEMbrightfieldmicrographsimagesandselectedareadiffractionpattern:(a)κ
II
precipitates(Fe
3
Al)intheinterdendritic,(b)κ
III
precipitates(NiAl)intheinterdendritic,(c)κ
IV
precipitates(Ferich)inα‐Cumatrix,(df)correspondingselectedareadiffractionpatternFe
3
Al,NiAl
andFerich,respectively.
Metals2020,10,21512of21
TheresultsofTEMobservationshowsthatthetwinsanddislocationscanbefoundnearthe
interdendriticregionofthesamplewithUVAtreatment.InthespecimenprocesswithUVA
treatmentatinterpasstemperatureof100°C,thecoexistenceoftwinsandinterdendriticisshownin
Figure11a.ThehighresolutionHRTEMmicrographandcorrespondingselectedareaelectron
diffractionareinsertedinFigure11a,whichfurtherverifiesthestructureoftwins.Theareasbetween
thetwinbands,whicharetensofnanometersinthickness,seemtobeparalleltoeachother,andEDS
elementalmappingindicatesthatthisparalleltwinningmayrelyonκ
III
precipitates(NiAl)inthe
interdendriticregion(Figure12).Figure11bpresentsadensityofdislocationsintheU3sample.The
dislocationsformrectangularblocksof50–120nminsize.EDSelementalmappingrevealsthatthere
isnoobviousprecipitatedphaseinthedislocations(Figure13).Meanwhile,themicrostructure,that
iswithoutUVAtreatment,specimensissimpler,asnoobvioustwinsanddislocationsoccursinthe
C3specimen,asshowninFigure11c.Hence,combiningpreviousstudies[39,40]andtheresults
obtainedinthisstudy,itcanbeconcludedthattheactionofUVAtreatmentWAAMsolidification
processcanhelptocreatetwinsanddislocations.Thetwinsanddislocationsmaycontributeto
strengtheningoftheWAAMNABalloybyactingasagrainboundaryandblockingthelattice
dislocationmotion[41,42].
Figure11.TEMBrightfieldmicrographsimages:(a)twinsfromU3sample,(b)dislocationsfromU3
sample,(c)α‐CumatrixwithcolumnargrainboundaryfromC3sample.
Figure12.(a)HAADFimageoftheUVA
treatment
WAAMinthetwinningregion,(bf)
correspondingEDSelementalmapsforCu,Al,Fe,Ni,andMn,respectively.Thelamellarprecipitates
areNiAl(κ
III
),andglobularFe
3
Al(κ
II
).
Metals2020,10,21513of21
Figure13.(a)HAADFimageoftheUVA
treatment
WAAMinthedislocationsregion,(bf)
correspondingEDSelementalmapsforCu,Al,Fe,Ni,andMn,respectively.Thelamellarprecipitates
areNiAl(κ
III
),andglobularFe
3
Al(κ
II
).
3.3.NanoHardnessandTensileProperties
Toacknowledgethevariationofnanohardnessandmodulusbytheultrasonicvibrationand
interpasstemperature,thenanoindentationtestsareconductedontheinterdendriticregion,ILB,α‐
CumatrixloweranduppertheILB,respectively.TypicalloaddepthcuresforfourregionsinC2,C3,
U2andU3samplesareplottedinFigure14a–d.Itcanbeseenthatthecontractdepth,maximum
depthandfinaldepthintheinterdendriticregionareshallowerthanthoseintheα‐Cumatrixand
ILB,whichcouldpartlyreflectagreaternanohardnessoftheinterdendriticregion.Besides,theILB
exhibitsthemaximumdepth,whichmeansthatthenanohardnessvalueoftheareaistheminimum.
Inaddition,theindentationtestcurvesindicateagoodrepeatabilitywithUVAtreatmentunder
interpasstemperatureof100°Csamples(Figure14d).
ThevariationtrendsofbothnanohardnessandmodulusinfoursamplesareshowninFigure
14e,f.Eachpositionnanohardnessvaluerepresentsanaverageofatleast8measurements.Withthe
interpasstemperaturedecreasing(C3sample)andUVAtreatmentunderinterpasstemperatureof
100°C(U3sample),thenanohardnessofα‐CumatrixandILBisincrease,andthenanohardnessin
theinterdendriticregionisdecreases(Figure14b,d).Apartfromtheenhancementinnanohardness,
themodulusalsoslightlyincreasesinα‐CumatrixandILB,anddecreasesintheinterdendriticregion.
Itisoftendesirabletohavemoreuniformnanohardnessinthematerial.WithUVAtreatmentunder
interpasstemperatureof100°C,thenanohardnessininterdendriticregion,ILBzone,theαCu
matrixuppertheILB,theα‐CumatrixlowertheILBare2.28±0.09GPa,2.11±0.02GPa,2.24±0.03
GPaand2.22±0.02GPa,whilethefourmodulusare129.4±3.80GPa,122.4±1.48GPa,126.4±2.22
GPaand126.1±2.06GPa,respectively.ThenanohardnessandmodulusoftheWAAMCu8Al2Ni
2Fe2MnalloyinU3sampleisrelativelyclose,whichindicatesthattheCu8Al2Ni2Fe2Mnalloy
fabricatedbyUVAtreatmentunderinterpasstemperatureof100°Cismorestable.Theuniformand
increaseoftheα‐CumatrixandILBinU3sampleismainlyattributedtotworeasons:(i)basedon
theHallPetchlaw,thenanohardnessincreaseswiththedecreasesofthegrainsizewhichiscaused
byUVAtreatment[43].(ii)accordingtotheTEMresults,theprecipitationofκ
IV
phaseintheα‐Cu
matrixmayincreasethemeannanohardness[17,44].
Metals2020,10,21514of21
Figure14.(ad)LoaddepthcurvesofC2,C3,U2andU3,respectively,(e)nanohardnessofC2,C3,
U2andU3,respectively,(f)modulusofC2,C3,U2andU3,respectively.
TensilepropertiesofspecimenU2,C3andU3aretestedaccordingtotheeffectiveareaof
deposition(Table1)andthefinergrainsize(Figure3).Thetypicalstressstraincurvesforthethree
conditionsoftheWAAMsamplesandcastC95220alloys,aswellasthetensilepropertiesdetermine
fromthesecurves,arecomparedinFigure15.ItcanbeseenthatalloftheWAAMsamplesarefound
toexhibitconsiderablyhigheryieldstrength(YS)andelongation(EL)ascomparedtothecast
samples.ThecastC95220alloyshashigherultimatetensilestrength(UTS)thanthatofWAAM
samplesduetomorecontentsofAl,NiandFeelements(themicrostructureofC95220alloysisshown
intheinsertedpictureatthetopleft).Thefiguresdemonstratethatthoughthenatureoftheflow
curvesoftheU2,C3andU3samplesaresimilar,thereisamarkeddifferenceintheirproperties.The
yieldstrengthoftheU3samples(224–227MPa)arehigherthanthatoftheC3(200–204MPa)andU2
samples(189–193MPa).IntheperspectiveofNABcomponentengineeringdesign,yieldstrengthis
morerelevantthanultimatetensilestrength.Ontheotherhand,theultimatetensilestrengthofU3
(515–518MPa)samplesarealsohigherthanC3(434–458MPa)andU2samples(387–438MPa).The
anisotropyofUTSexistsinC3andU2samples.Regardingductility,theelongationofU3samples
(40.56–41.42%)arecomparabletothatoftheC3samples(39.95–42.8%),whilespecimensprepared
fromU2samplesexhibithigherductility(42.2–50.87%).Thedetailedresultsoftheaveragetensile
propertiesintheverticaldirection(V)andhorizontaldirection(H)withintheUTS,YSandELare
listedinTable2.
Metals2020,10,21515of21
Figure15.ComparisonofstressstraincurvesofcastC95220alloys,U2,C3andU3samples(Inserted
picturesrepresenttherelationshipbetweenapplyingtensionandgrainsdirectioninadditive
manufacturingsamples,Fmeanstensileforcedirection.Insertedpictureatthetopleftisthe
microstructureofC95520alloys).
Table2.TensilepropertiesoftheWAAMsamples.
DirectionUTS/MPaYS/MPaEL/%
U2V390.3±3.8190.7±1.048.9±1.8
U2H430.1±6.9191.3±0.843.1±0.8
C3V436.5±2.1199.3±0.942.1±0.8
C3H454.9±3.1202.0±1.340.5±0.7
U3V516.1±0.8225.3±0.940.9±0.2
U3H517.5±0.5226.8±0.541.2±0.4
InordertostudytheinterpasstemperatureandUVAtreatmenteffectsondifferentlocations
alongthedepositiondirectiononthetensilepropertiesofthefabricatedwall,thetensileproperties
atdifferentheightsalongthedepositiondirectionaretested.Experimentalresultsareshownin
Figure16.IntheU2andC3samples,theUTSismoresensitivetothelocationswithinthebuildand
decreasingalongthebuildheight.TheneartosubstratesampleshowsthehighestUTS,whichis
causedbythefinermicrostructureinthebottomregion.Inaddition,althoughtheUTSofthesamples
atdifferentheightsdiffer,nocleartrendcanbeidentifiedinYSandductilityintheU2andC3
samples.Inaddition,whenalowinterpasstemperatureof100°Cisused,lessdecreaseofUTSis
generatedalongthedepositeddirectioninC3samples.AfterUVAtreatmentunderaninterpass
temperatureof100°C,thetensilepropertiesofthesamplesatdifferentheightstendtobestable.
Hence,frombothhorizontalandverticaltensilepropertiesconsequences,itisconsideredthatthe
tensilepropertiesaresymmetricalwithintheU3samples.TheUTS,YS,andELoftheU3samplesare
517.0±0.89MPa,226.2±0.95MPa,and41.01±0.29%,respectively.Amongallthetestsamples,
anisotropyiswellcontrolledinthesampleswithUVAtreatmentataninterpasstemperatureof100
°C,andtheincreaseofUTSandYSisnotaccompaniedbyanydropintheductility.
Metals2020,10,21516of21
Figure16.Effectoflocationontensilepropertiesandelongationofdepositedmetalindifferent
regionsofthecorrespondingFigure1c.
Figure17showsthatSEMmicrographsofthetensilefracturesurfacesfortheCu8Al2Ni2Fe
2MnalloyfabricatedwithandwithoutUVAtreatmentunderdifferentinterpasstemperature.The
fracturesurfacemorphologyoftheU2andC3samplesexhibitsatransgranularductilefailuremode
(Figure17a–f).InU2Hdirectionsamples,theslipseparation(markedbythewhitearrowinFigure
17a,b)anddimplesmixintheboundariessothatthePDAisclearlyobserved.TheU2Vdirectionis
mainlydominatedbyslipseparationinfracturecharacteristics(Figure17c).However,therearea
largenumberofsmallanddeepdimplesalignedalongthePADboundarieswhentheUVAtreatment
isnotapplied(Figure17d–f),implyingaductiletypeoffracture.WiththeUVAtreatmentunderthe
interpasstemperatureof100°C,thegrainssizesaregreatlyreduced,asgrainboundariesbecome
ambiguousinthefracturesurfaceoftheU3samplesandabundantfinerdimplescanbeobserved,as
showninFigure17g–i,whichalsoindicatesaductilefailuremode.Inaddition,inthehorizontal
tensilefracturemorphologiesoftheU2andC3samples,thePDAspacinginthefractureincreases
withdifferentheightsalongthedepositiondirection,whichisconsistentwiththeresultsof
microstructureobservation(Figure5)andalsoexplainsthetensilestrengthdecreaseswithdeposited
height(Figure16).Theconsistencyofthesetensiletestresultsisinagreementwiththechangesin
grainsizeobservedinthemicrostructure,suggestingthatUVAtreatmentunderaninterpass
temperatureof100°CisbeneficialtothematerialpropertiesofWAAMfabricatedCu8Al2Ni2Fe
2Mnalloycomponents.
Metals2020,10,21517of21
Figure17.Fracturesurfacesoftensilesamples:(a)U2H2,(b)U2H4,(c)U2V2,(d)C3H2,(e)C3H4,
(f)C3V2,(g)U3H2,(h)U3H4,(i)U3V2.(Thewhitearrowindicatesthecharacteristicsofslip
separation).
Itwasfoundthatthetensilepropertiesareanisotropicbetweenthehorizontalandvertical
directionsinU2andC3samples.Previousstudieshaveshownthatthestrengthisrelatedtothe
growthofcolumnargrainsalongthedepositiondirection[45].Duetothedifferentorientationsof
thetensilespecimens,horizontalsamplescontainmoregrainboundaries,whiletheverticalsamples
containmorecolumnargrains(asshownthetestsamplespicturesareinsertedinFigure15).Since
thestrengthresultingfromthegrainboundariesishigherthanthestrengthofthegrainsthemselves,
thestrengthofthehorizontaltensilespecimenishigherthantheverticalone,andtheelongationof
theverticalsampleishigherthanthatofthehorizontalsample[46].Inaddition,astheinsertedtest
samplespicturesareshowninFigure15,theverticalsamplesshowmoreobviousuniformelongation
andneckingwhencomparedwiththehorizontalsamples,whichmeansverticalsampleshavegreater
plasticdeformationandleadingtoalargeamountofslipseparationsonthefracture(Figure17c).As
thestrengthincreasesandelongationdecreases,fractureschangetoamixtureofslipseparationand
dimples(Figure17a,b,d–f).IntheU3samples,themicrostructureiscellular,andtherelationship
betweenthetensileforceandgrainboundariesisshowninFigure15(inserttestsamplespictures),
wherethenumberofdimplesisincreasingandthesizeissmallerthanothersamples.
Asidefromtheanisotropyofmechanicalproperties,whencomparedtheU2withC3samples,
theU3samplefabricatewithUVAtreatmentunderaninterpasstemperatureof100°Chave
remarkablylargertensilestrengthandyieldstrength.Themeanvaluesofthesepropertiesincrease
by19.7%and13.3%inUTS,and15.8%and11.1%inYS,respectively.Incontrast,theelongation
decreasesby9.4%inU2samples,whichisalmostthesameaselongationofC3samples.Themain
factorscontributetohigherstrengthareasfollows:thegrainrefinement,theformationoftwinsand
dislocationsstructure,andthepresenceofprecipitatingphaseinαCumatrix.Firstly,itisclearly
seenthatU3sampleshavefinegrainsthanotherparameterssamples(Figures3,5).Thisfinecellular
structureisthedirectoutcomeofUVAtreatmentunderaninterpasstemperatureof100°Candis
observedinmanyotherUVAtreatmentsformanufacturingofprocessedmaterials[25].Thegrain
sizereducedsignificantly,andthematerialisstrengthenedaccordingtotheHallPetcheffect[47].
Further,Wangetal.[48]alsomadeitclearthatthefinecellularstructuresignificantlyenhancedthe
strengthofadditivemanufacturingsamples.Secondly,duetotheUVAtreatmentWAAMprocess,
thetwinsanddislocationsform,coherenttwinshaveastrengtheningeffectthatissimilartothe
influenceofgrainboundaries[41].Thegeneratedtwinsactasadditionalgrainboundariesand
Metals2020,10,21518of21
interruptdislocationgliding,resultinginthedynamicHallPetcheffect[49].Thirdly,theresultsof
TEMconfirmtheexistenceoftheprecipitateκIVphaseintheα‐Cumatrix(Figure9).Theprecipitates
arebypassedthroughthedislocationloopingmechanism,whichincreasesinstrength[41].Itmaybe
notedthattheWAAMprocessunderUVAataninterpasstemperatureof100°Cprobablyrestricts
thesolidstatetransformationofκIVphaseduetomoresuperiorcoolingrates.However,infact,this
processmayreducetheformationofinterdendritic(moreκ
IIphase,composedofFerich),and
promotethedistributionofFeintotheα‐Cumatrix.DuringtheWAAMprocess,thepreviouslayers
aresubjectedtoreheatingbythethermalcycleofthesubsequentdeposition.Thetemperaturerange
of500to600°Cwouldresultinhomogeneousκ
IVprecipitationintheαCumatrix[17].Inthis
temperaturerange,theWAAMprocesshasasignificantreductionincoolingrateandthiscondition
canberetainedforalongtime(Figure7b),andwhenFeissufficient,moreκIVphasewillprecipitate
inαCumatrix.Moreover,itcanmaintaingoodductilityinU3samplethatisfabricatedbyUVA
treatmentWAAMunderinterpasstemperatureof100°Cmethodscouldbemainlyduetothe
formationofcellularstructure[42].Furthermore,thefactthatthefracturemorphologydevelopsfrom
theslipseparationofhighinterpasstemperaturetofinedimplesatthelowerinterpasstemperature,
whichalsoverifythegeneraltrendofincreasingstrengthandkeepductility.Inaddition,compared
withtheU2samples,thetensilestrengthoftheC3samplesincreaseslightly,whichmainlyresultsin
therefinementofthecolumnardendrites.
4.Conclusions
Inthisstudy,theeffectsofinterpasstemperatureandultrasonicvibrationonmicrostructural
characteristics(includingthemacrostructure,microstructure,elementaldistribution,andphase
composition)andmechanicalproperties(includingnanohardnessandtensilestrength)ofWAAM
fabricatedCu8Al2Ni2Fe2Mnpartsareinvestigated.Theresultsareasfollows:
(1)Interpasstemperaturesof100°Cand400°Chavenosignificanteffectonthemicrostructural
feature.ThealloyfabricatedunderdifferentinterpasstemperatureswithoutUVAtreatmentismainly
composedofthecolumnardendritesgrowingepitaxiallyalongthedepositiondirection.
(2)ThemicrostructureoftheCu8Al2Ni2Fe2Mnalloyisrefinedbyultrasonicvibration.With
theUVAtreatmentprocess,continuousdirectionalgrowthofcolumnardendritesisbrokenatthe
interpasstemperatureof400°C,andformscellularstructureattheinterpasstemperatureof100°C.
AcomparisonofpartsnottreatedbyUVAandpartstreatedbyUVAat400°Cindicatesthatparts
treatedbyUVAat100°Chaveasmallergrainsizeandbettermaterialdispersion.
(3)WithUVAtreatmentattheinterpasstemperatureof100°C,thealloyhasaninterdendritic
microstructurecontainingglobularκIIphasebasedonFe3AlandlamellarκIIIphasebasedonNiAl,
whereasprecipitatesκIVphase(richFe)isnucleatedintheα‐Cumatrix.Thetwinninganddislocation
havebeenobservedintheUVAtreatmentalloy.
(4)Thenanohardnessandmodulusofinterdendriticandα‐Cumatrixaredifferentatinterpass
temperatureof400°C,andthenanohardnessandmodulusofinterdendriticarehigher.Thenano
hardnessandmodulusoftheαCumatrixincreasesastheinterpasstemperaturedecreases.Inthe
UVAtreatmentattheinterpasstemperatureof100°C,thenanohardnessandmodulusaremore
uniform,thenanohardnessrangesfrom2.11to2.28GPawhilethemodulusrangesfrom122.4to
129.4GPa.
(5)Tensilepropertiesareanisotropicatinterpasstemperatureof100°CandUVAtreatmentat
aninterpasstemperatureof400°C.IntheUVAtreatmentataninterpasstemperatureof100°C,it
displaysalmostisotropictensilepropertiesandobtainesthebeststrengthductilitycombinations.The
UTS,YS,andELare517.0±0.89MPa,226.2±0.95MPaand41.01±0.29%,respectively.Twotypical
fracturecharacteristics,i.e.,slipseparationoccursathighinterpasstemperature,anditismainlya
dimplefeatureatalowinterpasstemperature.
AuthorContributions:methodology,W.C.andT.W.;writing—originaldraftpreparation,W.C.;writing—
reviewandediting,Y.C.andT.M.;fundingacquisition,Y.C.;datacuration,Z.Y.;conceptualization,Y.C.and
X.F.Allauthorshavereadandagreedtothepublishedversionofthemanuscript.
Metals2020,10,21519of21
Funding:ThispaperwassupportedbytheNationalDefenseBasicResearchProgramgrantnumber
JCKY2018401C003;KeyResearchandDevelopmentProjectofJiangxiProvincegrantnumber20192BBH80018;
JiangxiDistinguishedYoungScholarsgrantnumber2018ACB21016;JiangxiAdvantageousScientificand
TechnologicalInnovationTeamgrantnumber20171BCB24007.
ConflictsofInterest:Theauthorsdeclarenoconflictofinterest.
References
1. Ding,Y.;Zhao,R.;Qin,Z.;Wu,Z.;Wang,L.;Liu,L.;Lu,W.Evolutionofthecorrosionproductfilmon
nickelaluminumbronzeanditscorrosionbehaviorin3.5wt%NaClsolution.Materials2019,12,209.
2. Fonlupt,S.;Bayle,B.;Delafosse,D.;Heuze,J.Roleofsecondphasesinthestresscorrosioncrackingofa
nickelaluminiumbronzeinsalinewater.Corros.Sci.2005,47,2792–2806.
3. Rahni,M.R.M.;Beidokhti,B.;Haddadsabzevar,M.Effectoffillermetalonmicrostructureandmechanical
propertiesofmanganese–aluminumbronzerepairwelds.Trans.NonferrousMet.Soc.China2017,27,507–
513.
4. Lv,Y.;Wang,L.;Han,Y.;Xu,X.;Lu,W.Investigationofmicrostructureandmechanicalpropertiesofhot
workedNiAlbronzealloywithdifferentdeformationdegree.Mater.Sci.Eng.,A2015,643,17–24.
5. Culpan,E.A.;Rose,G.Microstructuralcharacterizationofcastnickelaluminiumbronze.J.Mater.Sci.1978,
13,1647–1657.
6. Hasan,F.;Jahanafrooz,A.;Lorimer,G.W.;Ridley,N.Themorphology,crystallography,andchemistryof
phasesinascastnickelaluminumbronze.Metall.Trans.A1982,13,1337–1345.
7. Nelson,E.Microstructuraleffectsofmultiplepassesduringfrictionstirprocessingofnickelaluminum
bronze.M.S.thesis,NavalPostgraduateSchool,Monterey,CA,USA,December2009.
8. Hasan,F.;Lorimer,G.;Ridley,N.CrystallographyofmartensiteinaCu10Al5Ni5Fealloy.J.Phys.1982,
43,653–658.
9. Wharton,J.A.;Barik,R.C.;Kear,G.;Wood,R.J.K.;Stokes,K.R.;Walsh,F.C.Thecorrosionofnickel
aluminiumbronzeinseawater.Corros.Sci.2005,47,3336–3367.
10. Rodrigues,T.;Duarte,V.;Miranda,R.M.;Santos,T.;Oloveira,J.P.Currentstatusandperspectivesonwire
andarcadditivemanufacturing(WAAM).Materials2019,12,1121.
11. Köhler,M.;Fiebig,S.;Hensel,J.;Dilger,K.Wireandarcadditivemanufacturingofaluminumcomponents.
Metals2019,9,608.
12. Li,Z.;Cui,Y.;Wang,J.;Liu,C.;Wang,J.;Xu,T.;Lu,T.;Zhang,H.;Lu,J.;Ma,S.;etal.Characterizationof
microstructureandmechanicalpropertiesofstellite6partfabricatedbywirearcadditivemanufacturing.
Metals2019,9,474.
13. Bai,X.;Colegrove,P.;Ding,J.;Zhou,X.;Diao,C.;Bridgeman,P.;Hönnige,J.R.;Zhang,H.;Williams,S.
NumericalanalysisofheattransferandfluidflowinmultilayerdepositionofPAWbasedwireandarc
additivemanufacturing.Int.J.Heat.Mass.Transfer2018,124,504–516.
14. Fachinotti,V.D.;Cardona,A.;Baufeld.B.;Biest,O.V.Finiteelementmodellingofheattransferinshaped
metaldepositionandexperimentalvalidation.ActaMater.2012,60,6621–6630.
15. Ding,D.;Pan,Z.;Duin,S.;Li,H.;Shen,C.FabricatingsuperiorNiAlbronzecomponentsthroughwirearc
additivemanufacturing.Materials2016,9,652.
16. Shen,C.;Pan,Z.;Ding,D.;Yuan,L.;Nie,N.;Wang,Y.;Luo,D.;Cuiuri,D.;Duin,S.;Li,H.Theinfluenceof