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Materialia
journalhomepage:www.elsevier.com/locate/mtla
Synergistic alloy design concept for new high-strength Al–Mg–Si thick plate alloys
Florian Schmid
a,∗, Irmgard Weißensteiner
a, Matheus A. Tunes
b, Thomas Kremmer
b, Thomas Ebner
c, Roland Morak
c, Peter J. Uggowitzer
b, Stefan Pogatscher
a,∗aChristian Doppler Laboratory for Advanced Aluminum Alloys, Chair of Nonferrous Metallurgy, Montanuniversitaet Leoben, Franz-Josef Strasse 18, 8700 Leoben, Austria
bChair of Nonferrous Metallurgy, Montanuniversitaet Leoben, Franz-Josef Strasse 18, 8700 Leoben, Austria
cAMAG rolling GmbH, Lamprechtshausener Strasse 61, Postfach 32, 5282 Ranshofen, Austria
a r t i c le i n f o
Keywords:
Aluminum alloys Precipitation strengthening Alloy design
Wrought alloys
a b s t r a ct
Withtheaimoffullyexploitingtheadvantageousstrength-to-weightratioevidentinAl–Mg–Sialloys,thisstudy presentsmeasuresforincreasingtheyieldstrengthofanENAW-6082typeplatealloy.Inadditiontodescribing thethermodynamicsimulation-basedadjustmentofage-hardenableelements(Si,MgandCu)andamodified artificialageingtreatment,itinvestigatestheeffectsofaddingasmallamountofZr.Thesignificantstrengthen- inginducedbyaddingZriscorrelatedwithsub-grainboundaryhardeninginarecoveredmicrostructureafter solutionannealingat570°C,comparedwiththealmostentirelyrecrystallizedmicrostructureinanunmodified ENAW-6082alloy.Incombinationwithamaximumdissolvablenumberofage-hardenableelementsandinter- ruptedquenching,whichcomprisesanimprovedheattreatmentstrategyforthickplates,itisseenthattheyield strengthcanbeincreasedbymorethan40%to411MPacomparedtoconventionalENAW-6082basematerial asverifiedbytensiletesting.Inthestudyscanningelectronmicroscopyandscanningtransmissionelectronmi- croscopywereperformedformicrostructuralcharacterizationwithafocusonparticleanddeformationanalysis.
Allindividualcontributionswhichgeneratedthesuperiorstrengtharecalculatedanddiscussedinordertoreveal themicrostructure-propertyrelationship.
1. Introduction
Weight-optimized design hasbecomea top priority in the trans- portsector becauseitfacilitates substantialCO2 reduction[1]. Low- densityAlalloys,especiallyage-hardenableAl–Mg–Si(6xxx)plateal- loys,arecurrentlyapreferredchoice.Inadditiontotheirmediumto highstrength,thickplatesmadefrom6xxx-alloyscombinehighfracture toughness,goodcorrosionresistanceandweldability.Theyalsofeature goodrecyclabilityatamoderateprice.Tofurtherexploitthepotential ofthesealloysandpromotetheirselection,anincreaseoftheirstrength- to-weightratiotomorefavorablelevelsseemstobroadentheirfieldof applications[2–4].
Workhardeningplaysakeyroleinsheetmaterials. Deformation- inducedstructuressuchasdislocationsandcellstructureswhichorig- inate from coldrolling contributesignificantlytoenhancedstrength valuesinthinsemi-finishedproducts[2].However,platealloys,which aredeformedathightemperaturesonly,donotnecessarilyexperience
∗Correspondingauthors.
E-mailaddresses:florian.schmid@unileoben.ac.at(F.Schmid),irmgard.weissensteiner@unileoben.ac.at(I.Weißensteiner),matheus.tunes@unileoben.ac.at(M.A.
Tunes),thomas.kremmer@unileoben.ac.at(T.Kremmer),thomas.ebner@amag.at(T.Ebner),roland.morak@amag.at(R.Morak),peter.uggowitzer@unileoben.ac.at (P.J.Uggowitzer),stefan.pogatscher@unileoben.ac.at(S.Pogatscher).
strengtheningfromthisprocessingstepbecausetheirincreaseddisloca- tiondensitywanesduetorecoveryprocesses[5].Inaddition,theso- lutionannealingtreatmentforAl–Mg–Sialloysistypicallyperformed above530°C,whichweakenshardeningviacoldworking[6].Several otherhardeningmechanismsarealsoactiveinthick plates,ofwhich strengtheningviatheformationofmetastableMgandSicontaining𝛽’’- phasesuponartificialagingisthemostimportant[2,7].AddingCuen- hancesthestrengthresponsefurther,mainlyduetotheoccurrenceofL- andQ’-phases[8,9].IncorporationofCuintohardeningphaseswhich containMgandSialsotakesplace.Allofthesealloyingelementsmay alsocontributedistinctlytosolidsolutionhardening[10].
Inthecourseofindustrialheattreatment,highstrengthAl–Mg–Si alloyswithaMgandSisolutecontentexceeding1%sufferfromaloss ofagehardeningpotentialwhichiscausedbynaturalaging(NA)dur- ingroomtemperature(RT)storagebetweensolutionannealingandar- tificial aging[11,12]. Naturalagingin plates maybe suppressedby applying interruptedquenching(IQ)toanelevated temperatureand
https://doi.org/10.1016/j.mtla.2020.100997
Received21December2020;Accepted25December2020 Availableonline2January2021
2589-1529/© 2021ActaMaterialiaInc.PublishedbyElsevierB.V.ThisisanopenaccessarticleundertheCCBYlicense
slowfinalcoolingtoroomtemperaturedirectlyafterwards.Thisgener- atestransformableclustersandnucleiofthemainhardeningphase𝛽’’
[13].InthisrespectIQisthereforecomparabletopre-aging[14],and increasedmechanicalparameterscanbeachievedinthefinalproduct [13,15,16].
In addition to the group of age-hardening elements, dispersoid- formingelementswithlowsolubilitysuchasFe,MnandCrcanbead- vantageouslyadded.Ifprimaryconstituentsarelargelysuppresseddur- ingthecastingprocess,smallanddenselydistributedsecondaryphases willformduringhomogenization[17,18]significantlyinfluencingthe mobilityofboththegrainboundaryandthedislocationmotion[19].Es- peciallyinAl–Mg–Sialloys,effectivelystabilizingthegrainboundaries turnsout tobe quitebeneficial becauseprocessingtemperaturesare comparablyhigh,thusincreasingthekineticsofboundaryrearrange- ments[20–22].
Ina2006studyKniplingetal.[23]listedseveralelementsthatform trialuminides,Al3X(X=Sc,Ti,Zr,Hf,Er,Tm,YrandLu),inAlalloys, ofwhichZirconiumandScandiumhavealreadybeenintensivelyinves- tigated.DuetothehighpriceofScandthelowcoarseningresistance ofAl3Sc,Zrmaybeamoreattractivecandidateforindustrialuse.To minimizethedisadvantageoftheprolongedannealingtimenecessitated bythereduceddiffusivityofZrwithinAl,itwasshownthatminorad- ditionsofSi(upto0.18at%)acceleratetheprecipitationkinetics[24]. Generally,Al3X-particlesarenotonlyusedadvantageouslyinunalloyed Al[23,24],butalsoinage-hardenablealloysystems(2xxx[25–27]or 7xxx[28,29])and– thishasbeenquiteprominentlately– inadditively manufacturedparts[30,31].Theseparticleswereshowntohaveasig- nificantimpactondislocationandgrain boundarymotion,whichde- creasesrecrystallizationtendencyand,thus,leadstotheformationof subgrains.IncombinationwithanincreasedamountofOrowanhard- ening,anincreaseinstrengthisthusobservedforScorZrcontaining 2xxxand7xxxseriesalloys[32–34].
TheeffectoftrialuminidesonAl–Mg–Sialloyswasrecentlyinves- tigated[35–38].However,foralloyscontainingSi(inamountsgreater thantheminutenucleationagent additionshownin[24]) itwas re- portedthat Zr,Si-phasescan form [39], thus decreasing theamount ofZravailablefor theformationof finetrialuminides dispersoids.A transformationofthebeneficialmetastableL12-Al3Zrprecipitates(co- herentwithAl-matrix)intoits stableD023-structure(incoherentwith Al-matrix)is alsoobservedatelevated temperatures(roughlyaround 500°C)[40].AhighSicontentsuchasthatofAl–Mg–Sialloyswasre- portedtofavortheoccurrenceofanincoherent(Al,Si)3Zrphaseandto causeacceleratedtransformation[41–43].6xxx-seriesalloysneedcom- parablyhighhomogenizationtemperaturestoobtainthedesiredinter- metalliclandscape(dissolutionofprimary𝛽-Mg2Si,reductionofmicro andmacrosegregation,andtransformationofAl(Fe,Mn)Si-phases)[44–
46], which limitsthepositive effectsof Zron industriallyprocessed Al–Mg–Si alloys due to coarsening. However, investigations demon- stratethatZradditionscanhaveasimilar[47]andevenadditiveeffect [48]onrecrystallizationretardationcomparabletoconventionaldisper- soidscontainingFe,MnorCr.AfewAl–Mg–Sistandards(ENAW-6056, ENAW-6065,ENAW-6110andENAW-6182)allowZruptoamaxi- mumof0.20wt%[49]tomakeuseofthesebeneficialdispersoids.
Inthisstudy,followingthestandardproductionrouteforthickplate alloys(casting, homogenization,hotrolling andartificialaging), we
combinedall ofthestrengtheningmechanismspresented inamodel Al–Mg–Siplatealloy.Thenecessaryprotocolforachievingthedesired highstrengthisasfollows:(i)balancedandoptimizednumberofage- hardenableelements(suchasSi,MgandCu);(ii)applicationofinter- ruptedquenching;and(iii)additionofZr.Itisnotyetfullyunderstood howZrinteractswithMnandFeorhowitreactsat6xxx-typicallyhigh processingtemperatures,andthestrengtheningmechanismduetoZir- coniuminsuchcomplexalloysisalsostillunknown[35].Theseissues areexaminedindetailbelow.Thescientificmethodologyinthisstudy addressesthequestionofwhetheritispossibletocombineallofthese strengtheningmechanismssynergisticallyinonealloy,andwhetherthe yieldstrengthofAl–Mg–Siwroughtthickplatescanbeincreasedtoover 400MPa.
2. Experimentalmethods
Forthisstudy,fivealloysbasedonanENAW-6082wereproduced atalaboratoryscale.Afterre-meltingofprimaryindustrialmaterial(EN AW-6082)inaresistanceheatedfurnace(NaberthermK20/13/S),pure Si,Mg,CuandZirconium(intheformofAlZr10)werealloyedtoobtain thecompositionslistedinTable1.Theseweremeasuredusingoptical emissionspectroscopy(notethatthesimulation-assisteddesignconcept isdescribedinthenextsection).Afterrefiningthemelt,whichwasheld atroughly750°C,viaArgaspurgingandgrainrefineraddition,cast- ingintocustom-madepre-heatedmoldssimilartothatinindustrywas startedaccordingto[50].Acoolingrateof3–5K/swasmaintained.Af- terfilletpiececuttingandhomogenizationtreatment,theresultingslabs werehotrolledfrom40mmdownto6mminseveralrollingsteps.A residualamountof0.02wt%Crispresentinallalloys,duetotheusage ofscrapintheproductionoftheindustrialmaterial.
Solutionheattreatmentoftensiletestsampleswasperformedusinga circulatingairfurnace(NaberthermN60/85SHA)at570°Cfor20min.
Aftereitherwaterquenchingorinterruptedquenchingintoa180°Chot metalbath(Bi57Sn43)withdelayedcoolingwithin24htoRTatade- creasingcoolingrate,thesamplesunderwent14daysofnaturalaging toemulateindustrialprocessing.After7daysofnaturalaging,aplastic pre-deformationof2%tookplace.Afteranother7daysartificialaging wascarriedoutat160°Cfor14hwith10hoframpingup(tosimulate theheatingofthickplates),resultinginatotalagingtimeof24hforthe water-quenchedspecimens(T651state).Startingfromthebasealloy,ar- tificialagingtimesforreachingT651-statewerecheckedforeachalloy variationandnomajordifferencesweredetected.ConcerningIQspeci- mens,2hagingat160°Cwasapplied,i.e.intotal12h(IQstate).The wholeheattreatmentprocessisschematicallyshowninFig.1.Whereas thesolidblacklinerepresentsastandardindustrialprocessingscheme forhotrolledplatealloys,ourhighstrengthvariantsrequireanalter- ationofthehomogenizationstrategy(dashedorangeline)andtuningof thefinalheattreatment(IQ+shortenedartificialaging),asindicated bythedashedblueline(seealsoSections3.3and3.4).
To characterize the mechanical parameters uniaxial tensile tests were performed on a universal testing machine (Zwick-Roell BT1) equippedwitha50kNloadcell.Roundspecimenswerepreparedwith a30mmlongand5mmthickgaugeinaccordancewithEN-ISO6892-1.
As-quenchedspecimensweretestedwithin5minofwaterquenching.
3samplesweretestedtocalculatemeanvalueandstandarddeviation.
Table1
Measuredchemicalcompositionofallalloystested,Alinbalance(wt%).
Si Mg Cu Mn Fe Zr Si + Mg Si/Mg Notes
Base alloy – 6082 1.02 0.67 0.08 0.42 0.33 0.00 1.7 1.5 Remolten industrial primary material 6082-Si 1.19 0.70 0.08 0.44 0.39 0.00 1.9 1.7 Addition of Si
6082-Cu 0.98 0.66 0.29 0.42 0.33 0.00 1.6 1.5 Addition of Cu 6082-Zr 1.03 0.66 0.08 0.42 0.37 0.20 1.7 1.6 Addition of Zr
Optimal alloy 1.07 0.81 0.30 0.41 0.36 0.21 1.9 1.3 Co-Add. of Si/Mg, Cu and Zr
Fig.1.Schematicillustrationoftheappliedheattreatmentconsistingofhomogenization,hotrolling,solutionannealing,pre-deformationundartificialaging.
Fig.2. Thermodynamicequilibriumcalculationsof6082,6082-Siand6082-Cu(seeTable1forchemicalcompositions).(a)InfluenceofSiandCucontentonthe onsetofliquidphaseformation.At570°Cannealingtemperature1.25wt%Siisthethresholdvaluefortheoccurrenceofaliquidphaseinthe6082basealloyand 1.16wt%in6082-Cu.(b)PhasefractionofMg2SiandQ-phaseinrelationtothetemperature.Volumetricpercentagesarecalculatedwithadensityof2.91g/cm3 forQ-phase[60]and1.90g/cm3forMg2Si[61].
ThermodynamiccalculationswereperformedwithFactSageTM 7.3 and FactSageTM 8.0 software using the FACT FTLite and FACT PS databases(2020)[51,52].
Tocharacterizethemicrostructureascanningelectronmicroscope (SEM)(JEOL7200FFEG-SEM),equippedwithanEBSD-measurement system (NordlysNano detector, Oxford Instruments), was deployed.
Samples wereprepared viaa standard procedure of grinding, oxide (OPS)andelectrolyticpolishing.Forgrainboundaryanalysisanarea of0.5mm2wasanalyzedoneachsample,withastepsizeof0.6μm.
Therecrystallizedareafractionwasdeterminedfromasurveyedareaof atleast2.35mm2(stepsizesof0.6and3𝜇m)takenatt/2representing thecenteroftheplate.Theminimumgrainboundarymisorientationan- glewassetto5°[53]fortherecrystallizationanalysis.Furtherdetails concerningbasicEBSDdataprocessingbymeansofthemtextoolbox canbefoundelsewhere[54–56].Grainsexhibitingagrainaverageker- nelaveragemisorientation(KAM)below0.5° alongwithagrainaverage bandcontrastof>70%ofthemaximumgrainaveragebandcontrast wereidentifiedasfullyrecrystallized.Toanalyzesub-grains,grainswith aminimumboundarymisorientationof0.8° wererecalculated.
BSEmicrographsweretakenatanaccelerationvoltageof5kVanda workingdistanceofroughly4mmtoevaluateintermetallicdispersoid phases.Eachparticle’sroundnessisderivedbydividingtheactualarea bythecalculatedcircularareausingthemaximumdiameter(thevalueis inverselyproportionaltotheaspectratio).Valuescloseto1correspond toanalmostperfectcircularshape.
Scanning transmissionelectronmicroscope(STEM)measurements werecarriedoutonaThermoScientificTMTalosF200XG2.200kVwas takenastheaccelerationvoltage.Sampleswerecutfromundeformed partsoftensiletestsamples.Aftergrindingdownto100𝜇m,3mmdisc specimens wereprepared viaastandardrouteconsistingof grinding andelectrolyticpolishingwithHNO3inmethanol(1:3)at−20°Cand 15V.ConcerningthedeterminationofSicontainingZr-particlesthefol- lowingmethodologywasapplied.Deployingapictureanalysissoftware (ImageJ1.51f)andusingacolorthresholdfortheSiandZrmapping,bi- naryblackandwhitepictureswithparticlescontainingSiandZrwere made. OverlappingthesetwobyusingtheBooleanoperation“AND” broughtbackathirdpicturewhereonlyparticlesweredepictedthat containedbothZrandSiatthesamespot.
Table2
Overviewofmechanicalparametersofalltestedalloysafterstandardartificialaging(T651)andinter- ruptedquenching+shortenedartificialaging.
As-quenched Standard AA (T651) IQ + shortened AA
R p0.2 R p0.2 R m A R p0.2 R m A
MPa MPa MPa % MPa MPa %
Base alloy – 6082 58 ± 2 289 ± 3 309 ± 2 19 ± 2 349 ± 8 352 ± 9 11 ± 1 6082-Si 66 ± 2 319 ± 8 339 ± 7 15 ± 1 364 ± 4 373 ± 5 12 ± 1 6082-Cu 64 ± 2 313 ± 3 337 ± 3 17 ± 1 355 ± 6 362 ± 6 12 ± 1 6082-Zr 85 ± 0 334 ± 5 362 ± 4 15 ± 1 384 ± 7 398 ± 8 12 ± 1 Optimal alloy 90 ± 3 362 ± 9 392 ± 9 15 ± 1 411 ± 4 425 ± 4 14 ± 0
Fig.3. Isoplethof6082withadditionsofZr.TheoccurrenceofAl3Zrishighly dependentontemperature.Atthesolutionannealingtemperatureof570°CZr hasasolubilityofabout0.1wt%inthesolid𝛼-Al.
3. Designconcept
Inthefollowingwepresentourdesignconcept,whichisbasedonthe exploitationofallscenariosforimprovingstrengthinamodelAl–Mg–Si alloywhichcontainsFeandMn.Theconceptinvolvestuningboththe elementcompositionandtheprocessingparameters.
3.1. Additionandoptimizationofprecipitation-hardeningelements
Asthefirststepinouralloydesign,weadjustedthenumberofage- hardenableelements:SiwasadjustedtoMg,andCuwasaddedasa strength-increasingelement.TheCuamountwaslimitedbecauseofthe decreasingcorrosionresistance(especiallyresistanceagainstintergran- ularcorrosionin6xxx-alloys[57])associatedwithincreasingamounts ofthiselement[2].AfterconsideringtheoptionofincreasingtheCu contenttoacorrosion-irrelevantamountof0.3wt%,andthespecified processparameters(i.e.solutionannealingtemperatureof570°C;see Section2andFig.1forfurtherinformationonthestandardindustrial processingroutefortheproductionofthickplates),wecarriedoutther- modynamiccalculations.TheresultsofwhicharepresentedinFig.2. Fig.2ashowstheliquidphasefractioncoursewithalteredSicontentin the6082basealloyat570°C.Noliquidphaseisallowedtoformatthis temperature,becauseotherwisethematerialwouldbeunusableafter- wards.AthresholdSicontentof1.25wt%canbederivedfromthegraph (notethataddingCuto0.3wt%decreasestheonsetofliquidphasefor- mationtoanSicontentof1.16wt%).TodeterminetheinfluenceofSi onlythiselementwasincreasedfrom1.02to1.19wt%inthe6082-Si alloy,butitwaskeptatthereferencelevelin6082-Cu.MgandSiwere
bothaddedtotheoptimalalloytoapproachthepreferentialSi/Mgratio of1.2inthemainhardeningphase𝛽’’(Mg5Si6[58,59])(seeTable1for thechemicalcompositionsofallalloys).
Fig.2bshowsthephasefractionsoftheequilibriumphasescontain- ingSi,MgandCuoverthetemperature,whichcanbeusedasarough estimationofthemainhardeningprecursorphasesformingduringartifi- cialaging.AddingSiincreasestheamountofMg2Siconsiderably,with- outaffectingtheamountofQ-phase(Al4Cu2Mg8Si6[60]).Inthe6082- Cualloy,adrasticriseoftheQ-phasecontainingCuisseen,alongwith areductionofMg2Si.At160°Casignificantlyincreasedtotalhardening phasevolumeispresentinbothalloyvariations(6082-Siand6082-Cu), whichshouldultimatelyleadtoanenhancedagehardeningresponse.
3.2. AdditionofZirconium
Whereas Zr also precipitates from the super-saturated matrix as Al3Zr,itsnucleationandgrowthtakeplaceatmuchhighertemperatures thanthoseassociatedwiththeelementsdiscussedabove.Fig.3shows anisoplethofthe6082basealloywithZrupto0.3wt%.Amaximum amountofroughly0.1wt%isdissolvablein𝛼-aluminumatthetarget solutionannealingtemperature(570°C).Becausebothhomogenization andhotrollingarecarriedoutat6xxx-typicaltemperatureswellabove 500°C,thereisatendencytodissolutionoflargesharesofpreviously formedAl3Zr.Therefore,anincreasedamountof0.2wt%Zirconiumad- dition(comparedtoamaximumamountof0.15wt%inENAW-7050 forexample[49])waschosen.Becausenoprimarynon-hardeningAl3Zr precipitates,whichimpairelongation,shouldbeformed,thetempera- tureduringthecastingproductionprocesswaskeptabove730°C(see redlineinFig.3).
3.3. Adaptionofhomogenization
DuringhomogenizationtreatmentofconventionalAl–Mg–Sialloys, thedissolutionofpreviouslyformedcoarseMg2Sirequiresaratherhigh temperatureofabout550°C(seeFig.2b).Slowheatingtothistemper- ature alsogeneratesthedesiredformation ofdispersoids.Inorder to formdispersoidsinsmallsizesandhighnumberdensities,wesuggest anadaptationofthehomogenizationtreatmentstep.Ontheonehand, itisknownthatAl(Mn,Fe)Si-dispersoidscontainingFeandMnexhibit thehighestnucleationrateat370 °Cuponheatingat0.01K/s [62]. Conversely,anintermediateannealingstepat360°Chasbeenshownto promotetheformationofAl3Zr-nucleiandthusreducethetimeneeded forfullprecipitationat550°C,whichisinfluencedbytheslowdiffusion ofZr[23,63].Ithasalsobeenreportedthatthesetwophase-familiesin- fluenceeachotherbypromotingheterogeneousnucleationsites[64], whichnecessitatesaholdingannealingsteptopreventcoarseningand non-uniformdistributionofdispersoids.Withthisinmind,optimized homogenization(indicatedbythedashedyellowlineinFig.1)wasap- pliedtothealloyscontainingZr.Thishomogenizationconsistedofan intermediateholdingannealingstepat350°Cfor16handasecondstep at550°Cfor2h.
Fig.4.Stress–straincurvesof(a)thebasealloy6082,6082-Siand6082-Cuand(b)6082-ZrinT651state
3.4. Interruptedquenching
BecauseAl–Mg–Sialloyslosetheiragehardeningpotentialdueto naturalaging[11,12],quenchingtoanelevatedtemperaturewasap- pliedasanadditionalmeasuretoalltheexperimentalalloystestedhere [13].Accordingly,allsampleswerequenchedto180°Crightaftersolu- tionannealingandslowlycooledtoroomtemperature(thistakes24hin total,atadecreasingcoolingrate)tosimulatethecoolingofthickplates (indicatedbythedashedbluelineinFig.1).Thisshouldleadtothefor- mationoffavorableclustersandnucleiofthemainhardeningphase𝛽’’, andiscomparabletopre-aging[13,14].Becausethistreatmentalready precipitatesmanyoftheage-hardenableelements,ashortenedfinalar- tificialagingstepwasapplied.
4. Results
ThissectionbuildsonthemeasurespresentedinSection3,whichad- dressedalloyandprocessdesign,andoptimalalloyandmicrostructure characterization.Asanoverview,Table2liststhemostimportantstress–
straincurvevaluesfortheagedsamplespresentedbelow.Allalloyswere alsotestedrightaftersolutionannealingintheas-quenchedstate.These valuesarealsolistedinTable2andarediscussedinSection5. 4.1. Alloydesign
4.1.1. Additionandoptimizationofprecipitation-hardeningelements Fig.4ashowsthestressstraincurvesofalloysoptimizedbyadding SiorCuincomparisonwiththebasealloy6082inthefullystrength- enedT651state.Regardingmechanicalproperties,bothmeasurements (Si/Cuaddition)generateaconsiderableincreaseinstrength.Whereas 6082-Cushowsatotalyieldstrengthincreaseof24MPa,addingSire- sultsinanincrementof28MPa.Elongationtofractureissignificantly reducedin6082-Si.6082-Cushowshardlyanychangesinthisrespect.
4.1.2. AdditionofZirconium
Fig.4bshowsthestress–straincurvesofthealloywithaZraddition comparedtothe6082basealloyinT651state(6082-Zrwashomog- enizedby2-steptreatment). Theachievable propertychanges gener- atedbyaddingZrarequiteremarkable.Asignificantincreaseinyield
strengthofroughly45MPaisobserved.Theelongationtofracturestays highataround15%.Notethatasimilareffectonmechanicalparame- tersisachievedwhenusingpureAlinsteadofindustriallyproducedEN AW-6082asbasematerial(seeSupplementaryMaterial:TableS1and Fig.S1).
4.2. Processdesign
4.2.1. Adaptionofhomogenization– formationofdispersoids
Fig.5showsSEM-backscatteredelectron(BSE)micrographs,where all dispersoids containing Fe, Mn, Cr and Zr appear as bright dots (Fig.5a–d)afterone-step andtwo-step homogenizationtreatmentin theundeformedstatebeforehotrollingfor6082and6082-Zr.Thecor- respondingparticledistributionstatistics areshown inFig.5eandf.
InFig.5a-d.Thenumberdensitiesofallparticlesafterhomogenization showthefollowingtrends(seeFig.5f):(i)thenumberdensityincreases whenZrisadded;(ii)forbothalloysthenumberdensityincreaseswhen switchingfromone-totwo-stephomogenization;(iii)two-stephomoge- nizationappliedtothebasealloy6082generatesahighernumberden- sitythanstandardone-stephomogenizationofthe6082-Zralloy.Com- paredtothematerialcontainingZr,anincreasednumberofelongated particles(0–0.33roundness;seeSection2Experimentalmethods,foran explanationofthisterm)formsinthebasealloy6082,especiallyafter one-step homogenization.Withtwo-stephomogenizationthenumber ofroundparticles(withanevensmallerdiameter;seeFig.5e)ismore thandoubled.Generally,two-stephomogenizationleadstoanoverall decreaseinthediameterofallparticlesin6082(from116nmto85nm).
In6082-Zrthediameteroftheparticlesstaysonalevelcomparableto anyhomogenization(between70and80nm),especiallyinthemost importantfractionofroundish,smallparticles(0.66–1.00roundness).
However,thenumberdensityincreasessignificantlywithtwo-stepho- mogenization,asalsoreportedinRef.[63].Fig.5gandhshowthedis- tributionofparticlesafterhotrollingandsolutionannealingat570°C inT4state.Thenumberdensitiesarereducedinbothalloys.Inparticu- lar,thecompletedisappearanceofstronglyelongatedparticles(0–0.33 roundness)andastrongreductioninthesecondshareofparticles(0.33–
0.66roundness)isalsoapparentinboth.Whereasthemeandiameterof allparticlesincreasesfrom85to101nmin6082,noalterationinthe particlesizeisseenfor6082-Zr(Fig.5i).
Fig.5.SEM-BSEmicrographsofthemicrostructureindifferentstates(afterhomogenization– a–e;inT4-state– g–i)afterone-stephomogenization(a,c,g)and two-stephomogenization(b,d,h)ofthebasealloy6082(a,b,g)and6082-Zr(c,d,h).Meandiameter(e,i)andnumberdensities(f,i)ofalldetecteddispersoids arecategorizedinthreegroupsaccordingtotheirshape.
Fig.6. Stress–straincurvesfor6082and6082-ZrinT651stateandafterin- terruptedquenching.Inbothcasesanincreaseinyieldstrengthofabout40to 50MPaisobserved.
4.2.2. Interruptedquenching
Fig.6showsstressstraincurvesfor6082(onestephomogenization) and6082-Zr(two-stephomogenization)afterstandardAA(T651state) andinterruptedquenching+shortenedAA(160°C/2h).Thistreatment generatedanincreaseinyieldstrengthofabout40to50MPaforboth alloys.Obviouslythisprocessvariationnotonlystrengthensthebase alloy,butalso– almostequally– 6082-Zr,whichdemonstratesthatIQ andtheeffectofZradditionworksynergistically.
4.3. Optimalalloy
Afterdemonstratingthegaininstrengthcausedbyapplyingthein- dividual measurespresented,we now investigate whetherthese can be combinedtocreate anoptimalalloy. Themeasures consistof (i) amaximumdissolvableamountof SiadjustedtoMg,combinedwith theadditionof0.3wt%Cu;(ii)theadditionof0.2wt%Zr;and(iii) theapplicationofIQtosuppressboththeformationofnon-hardening phasesand,subsequently,thelossofhardeningpotential.Fig.7shows theresultsofthislastalloydesignstep.Comparedtothereferenceal- loy,thesumofmeasures(i)and(ii)generatesayieldstrengthincrease of73MPa.Interruptedquenchingfurtherincreasesstrengthby49MPa (lightgreendashedcurveinFig.7),resultinginasignificantlyincreased yieldstrengthof411MPa,ascomparedtocommonlyprocessedENAW- 6082referenceplatematerialwith289MPayieldstrength.Notethat
Fig.7. Demonstrationofanoptimalstrategyforhighstrengthplatematerial viastress–straincurves.Theminuteoptimizationofchemicalcompositionand modifiedindustrialprocessingcanboosttheyieldstrengthoftoday’sstate-of- the-artENAW-6082platematerialbymorethan40%.
theseeffectscanalsobeachievedusingpureAlasthebasematerial(see supplementarymaterial:TableS1andFig.S1).
4.4. Microstructureanalysis
4.4.1. 6082-Zr
Fig.8showstheresultsofSEM-EBSDmeasurementsofthebasealloy 6082and6082-Zr.Adepictionofthe3rdorderKAM(stepsize0.6μm, threshold5°)inFig.8aandbhighlightsthenoticeabledifferencebe- tweenthesetwoalloysaftersolutionannealingat570°C.Whereasonly littledeformationisstoredin6082aftersolutionannealing,misorien- tationisstillhighin6082-Zr.ThiscanalsobeseenfromFig.8cand d,whichshowsthebandcontrastforeachspecimen.Afurtherdepic- tionisfoundinthesupplementaryFig.S2aandb,whererecrystallized areasareshadedin blue.Thedistributionofallboundaries forthese twoalloysinFig.8eandffurtherconfirmsthesignificantdifferencein themicrostructurespresented.Nosharptransitionofsmallanglegrain boundaries(SAB)andhighanglegrainboundariesisevident.However, McQueendefinesSABuptoathresholdvalueof8°[5],Gottsteinuntil 15°[65].Exemplarilytakingthelattervalueof15°,thereisaroughly uniformdistributionin6082,whereasin6082-Zrthevastmajorityof grainboundaries(88%)aremeasuredatunder15°.
4.4.2. Optimalalloy
Fig.9showsthemicrostructureoftheoptimalalloyafterstandard AAcontainingseveraldifferenttypesofdispersoids.Twodistinctphase familiesarepresent:(i)phasescontaining Fe,MnandCr(Fig.9b–e) and(ii)phasescontainingZr(Fig.9g–i).Fig.9b–eshowsthatFe,Mn andCrareallmixedanddonotprecipitateontheirown,whichcon- firmstheseparticlesasAl(Fe,Mn,Cr)Si-phases.Incontrast,Zr(seeFig.9g andh)eitherformspureAl3Zr-precipitatesorasubstitutionofacertain amountofAlbySitakesplacetoform(Al,Si)3Zr(afewexamplesare indicatedbyredarrows).OfalldetectedphasescontainingZr,62%had adetectableamount ofSi(seeFig.9g–i).Fig.9ialsoshows thatZr- particlesareoftenattachedtoAl(Fe,Mn,Cr)Si-particles,whichindicates heterogeneousnucleation(forbettervisibilityonlyMnwasselectedin multi-elementEDXmapping).Notethatusingvarioustiltinganglesin theSTEMconfirmedthatthesetwophasesarelocatedoneachother (seesupplementarymaterialFig.S3).ThesizedistributioninFig.9fin- dicatesthattheZrparticlesaremuchsmaller(74±32nm)thanthe othertypeofdispersoid(134±72nm).Althoughthecumulativesum oftheFe,MnandCrcontent(0.38at%)ismorethansixtimeshigher
thanZr(0.06at%),thearealdensityof(Al,Si)3Zrismorethantwiceas high(7.52×1012m−2comparedto2.97×1012m−2).
Fig.10ashowsadarkfieldSTEMmicrographwithbrightparticles alongdislocationsandentangleddislocationsthattendtobestuckto theseparticles.Themulti-elementmappinginFig.10bshowsMncon- tainingaswellasZrcontainingdispersoidsandMgandSirichharden- ingphases.Heterogeneousprecipitationofthetwodispersoid-phasesis alsoobservedinthismicrograph(asinFig.9h),aswellasatendency towards theformationofthemainhardeningparticlesondispersoid phases(indicatedwithredarrowsinFig.10b).Besidesthenucleation onthephaseboundariesofsuchparticles,coarsenedphasescontaining MgandSiareobservedalongdislocations(indicatedbydashedyellow linesinFig.10b).ThisisdepictedinmoredetailinFig.10c–f.
Fig.11showsseveralSTEMmicrographsofasub-grainwithanori- entationclosetothe<001>-axisoftheoptimalalloy(indicatedbythe redcircleinFig.11a).InFig.11aandbmanybrightconstituents,identi- fiedaseitherZr-orMn-rich(Fig.11b),arelocatedexactlyattheborder ofthesub-graininthemiddleofthepicture(indicatedbytheyellowand redarrowsalongthegrainboundary).Thisindicatestheinterferenceof theseparticleswithsub-grainboundaries.Thedark-fieldmicrographof theoptimalalloyinFig.11crevealsthatmanydislocationsarepinned insidesub-grainsbydispersoids(seeredcircle).
5. Discussion
Astheabove datashows,alteringthechemicalcomposition(age- hardenableelementsMg,SiandCu;Zrasdispersoidformer)andadapt- ing the heat treatment (optimized homogenization and interrupted quenching) generated an exceptionally high yield strength of over 400MPainanoptimal6082-typeplatealloy.Inthefollowingwedis- cusstheunderlyingmechanismswhichcausedthisgaininstrength.We alsoprovideadetailedanalysisoftheroleplayedbyminorZraddition inmicrostructureevolution.
5.1. InfluenceofSiandCuaddition
AddingSiandCuisknowntocauseanincreaseinstrengthviasolid solutionstrengthening andprecipitationstrengthening[66]. Table 3 showsthemechanicalparametersofthealloystestedhereindifferent stateswhileundergoingstandardartificialaging.Rightafterquenching (“As-quenched”)onlyaminordifferenceinyieldstrengthismeasured.
Intheabsenceofanyhardeningphasesthisdifferencecanbeattributed solelytosolidsolutionhardeningcausedbydifferentSiandCucontent [10].However,accordingtothemodelofsolidsolutionhardeningby Leysonetal.[67]wecalculatearelativestrengthdifferenceoflessthan 1MPa tobecausedbythealterationsin6082-Siand6082-Cu[10]. Thereforethisincreaseintheas-quenchedstrengthisattributedtoclus- ters,whichformeitherduringquenchingorduringtheveryshortperiod oflessthan5minbetweenquenchingandtensiletesting[10,68].Con- cerningartificialaging,bothalloyvariationsshowanenhancedstrength responseofroughly20MPa.FromFig.2bitisqualitativelyseenthat addingSiandCugeneratesanincreasedvolumefractionofstrengthen- ingphasesat160°C(1.56vol%inthereference,1.63vol%in6082-Si and1.87vol%in6082-Cu),fromwhichanincreasedprecipitationpres- sureandthusanincreasedhardeningeffectcanbederived.
5.2. InfluenceofZraddition
AddingZrtothebasealloy6082results inastrengthincreaseof 28MPaintheas-quenchedstateafterhotrolling(6082-Zr).Artificial agingcausesafurtherincreaseinyieldstrength(17MPa),whichgen- eratesatotalincreaseof45MPainthefullyhardenedT651statecom- paredto6082.FurtheralloyingwithSi,MgandCuintheoptimalalloy showsanadditiveeffect(increasedyieldstrengthintheT651stateof 73 MPacomparedto6082).Toattributethese findingstothecorre- sponding mechanisms,adetailedinvestigation ofthemicrostructure,
Fig.8.ComparisonofthemicrostructureinT4temperaftersolutionannealing(570°Cfor20min)of6082(a,c,e)and6082-Zr(b,d,f).a,b)3rdorderkernel averagemisorientationuptoaKAMthresholdof5°.Boundariesabove5° areshowninblue,aswellaslargernon-indexedareas.(c,d)Depictionofthebandcontrast.
(e,f)Misorientationdistributionofallmeasuredboundaries;infosonthethresholdusedcanbefoundinthetext.
Table3
OverviewofstrengthvaluesofallalloystestedindifferentstatesforstandardAA.Datawereacquiredrightafter solutionannealingandquenching(as-quenched)andafterfinalartificialagingtreatment(allvaluesinMPa).
Solidsolutionhardeningiscalculatedaccordingto[67].Thecontributionduetoartificialagingiscalculatedas thedifferencebetweenas-quenchedstrengthandstrengthintheT651state.
As-quenched Artificial aging T651 (conv. AA)
R p0.2 Gain to the base alloy Strength gain Gain to the base alloy R p 0.2
Base alloy – 6082 58 ± 2 + 232 ± 3 289 ± 3
6082-Si 66 ± 2 8 ± 3 + 253 ± 8 + 21 ± 9 319 ± 8
6082-Cu 64 ± 2 7 ± 3 + 249 ± 3 + 17 ± 5 313 ± 3
6082-Zr 85 ± 0 28 ± 2 + 249 ± 5 + 17 ± 6 334 ± 5
Optimal alloy 90 ± 3 32 ± 3 + 272 ± 10 + 40 ± 11 362 ± 9
Fig.9. STEM-EDXmicrographsofthedispersoidlandscapeoftheoptimalalloyinconditionT651.(a)HAADF-overviewofthemicrostructure.(b-e)EDXmapping ofFeMnCr(b),Fe(c),Mn(d)andCr(e).(f)ParticledistributionofFeMnCr-phasesand(Al,Si)3Zr-phases.Thedash-dottedlinerepresentsthemeandiameter.(g–i) EDXmappingofZr(g),Si(h)andMnZr(i).TheredarrowsindicateparticlesthatcontainbothSiandZr.
whichconsideredthedispersoidsinparticular,wascarriedout.Bothof thealloyscontainingZrarediscussedbelow.
5.2.1. Formationofdispersoids
AddingZrleadstotheformationofAl3Zr.Herealmosttwo-thirds ofallprecipitatesshowthatasubstitutionofAlbySiforms(Al,Si)3Zr (seeFig.9i)intheoptimalalloy.ForboththeZr-freebasealloy6082 and6082-Zr,theoptimized2-stephomogenizationhasprovedeffective increatinga dispersoidlandscapewithanincreasednumberdensity anddecreasedaveragediameterofallparticles(seeFig.5).STEMmea- surements(seeFig.9h)haveconfirmedpreviousobservationsofhet- erogeneousnucleationofMn-andZr-richparticlesoneachotherinthe optimalalloy[64,69].
5.2.2. EffectofZronthemicrostructure
SEM-EBSDmeasurementsinFig.8revealcompletelydifferentmi- crostructuresinthebasealloyand6082-Zr.Thecauseliesinquitediffer- entdispersoidlandscapesandtheincreasednumberdensityofsmaller particlesin6082-Zr(seeFig.5)duetoboththeZradditionandanop- timizedhomogenizationstrategy.Itisknownthatsuchsecondphase particlesinterferewithmovinggrainboundariesandblocktheirmove- mentduringrecrystallization,andotherprocessesthatincludeabound- aryrearrangement[70].Dispersoids,suchasAl3Zrthathaveacoherent interfacewiththematrixarepostulatedtohavearetardingpressureon boundarieswhichistwiceashighasincoherentones.Neverthelessit isquiteremarkablethattheincoherentMn-dispersoidsdonotinhibit recrystallizationmoreeffectively(Fig.8c),eventhoughthereference alloyactuallycontainssufficientlevelsofMn (0.42wt%) toactwell againstgrainboundaryanddislocationmovements[20,71].Thestan- dardhomogenizationone-steptreatment,however,causestheforma- tionofcomparativelycoarse particleswithlownumberdensity.This
resultsinalesseffectiveretardingpressureagainstrecrystallizationin thebasealloy,thusrenderingtheMnadditionfairlyineffective.
5.2.3. Influenceonmechanicalparameters
Intheas-quenchedstate,6082-Zr,whichismodifiedbyZr-addition only,showsa strengthincrease of28 MPacompared tothebase al- loy.DispersoidstypicallycontributetoastrengthincreaseviaOrowan hardening,duetotypicalsizes,whichforeclosesthemtobecutbydislo- cations[10,35].StrengthcontributionsofZr-particlesaccordingtothe Orowanmechanismcalculatedaccordingto[9]aregiveninTable4, andareonlyofminorimportance.In6082-Zronly10MPaandinthe optimalalloyonly14MPaarecalculatedinadditiontothebasealloy.
AssumingthatparticlescontainingZrthatprecipitateheterogeneously atpresentMn-richdispersoids(countedinthecalculationfortheopti- malalloyseparately)donotcontributemuchtoincreasedstrength,this valueisreducedevenfurther.Consequently,theOrowanmechanism alonecannotexplaintheincreaseinstrengthcausedbyaddingZr.
Followingthestandardproductionrouteofthickplates,hotrollingat 540°Cwasappliedforallalloystested.Duetothehighstackingfaulten- ergyofAl,sub-grainscaneasilyformduringthisproductionstep.More severeconditions(increaseofdeformationspeeḋεanddecreaseoftem- peratureT)generateadecreaseinthesub-grainsize[5,73].Afterhot rolling,similarelongatedmicrostructures,indicatingnon-recrystallized areas,arepresentinboththereferencealloyandin6082-Zr(Fig.S4).
Asalreadydescribedabove,recrystallizationin6082-Zrissuppressed duringsolutionannealingduetothebeneficialdispersoidlandscapes, resultinginalargelyrecoveredmicrostructure.Wecanthereforeexpect sub-grain boundaryhardeningasanadditionalstrengtheningmecha- nisminthealloyscontainingZr.
Anevaluationofthemicrostructureintheoptimalalloyshowsan averagesub-grainsize(𝛿)ofroughly5𝜇minrollingdirection(equiv-
Fig.10. STEMmicrographsoftheoptimalalloyintheT651-state.(a)Darkfieldmicrographshowingdispersoidsalongwithentangleddislocations.(b)EDX-mapping ofMg,Si,MnandZrwhichshowstheheterogeneousnucleationofseveralphases(hardeningphasecontainingMg,Siondispersoidsindicatedbyredarrows)anda tendencytowardsprecipitationofthehardeningphaseondislocations(seedashedyellowline).Thiscanalsobeseeninmoredetailinc–f.
Fig.11. STEMmicrographsoftheoptimalalloyintheT651-state.Thegraininthemiddle(lightgreyishgrain,indicatedbyaredcircle)in(a)isnearthe<001>- axis.(a)HAADFmicrographshowingsmallsub-grainswithasizeofroughly3–5𝜇manddispersoids.(b)Multi-elementEDXmappingofthesub-grain(grainfrom (a),circledinred)inthe<001>-axis.Manyparticles(indicatedbyarrows)arelocatedrightontheboundary.(c)Dark-fieldmicrographshowingthedislocation distribution.Anentanglementofdislocationscanbeobservedinthevicinityofbrightparticles(redcircle).Sub-grainboundariesaremadeupofahighnumberof stackeddislocations.
alenttothetestingdirection).PuttingthisintoEqs.(1)and(2),which weredevelopedbyGinterandFarghalli[74],withaburgersvector(b)of 2.86×10−10m,ashearmodulus(G)of25.4GPaandaTaylorfactor(M) of3givesastrengthincrease(Δ𝜎)of44MPaduetosub-grainboundary hardening[74].Takingintoaccountthat– dependingonthethreshold value(see4.4.16082-Zr)– 74%or88%ofthetotalboundariesaremade
upofSAB(Fig.8f),thisincrementisreducedtoavaluebetween33and 39MPa.Heterogeneousnucleationof50%ofallZr-containingparticles (Fig.9i)leadstoaconsiderablereductionoftheirhardeningcontribu- tion.ReducingOrowanhardeningtoabouthalftheamountwouldleave 33MPafortheoptimalalloy,whichisonly3MPahigherthaninthe basealloy.Thereforethiscontributionjudgestobenotsignificantand
Table4
SimulatedstrengthcontributionsduetoOrowanhardeningofthereference,6082-Zrandoptimalalloy(allvaluesinMPa).Calculatedaccordingto [10].
T651 (conv. AA) Total dispersoids strengthening (calculated) Input parameters
R p0.2 Al(FeMnCr)Si (Al,Si) 3Zr Diameter [nm] Volume fraction [m 3/m 3] a
Base alloy – 6082 289 ± 3 30 – SEM ( Fig. 5 g)
101 ± 59
0.0090
6082-Zr 334 ± 5 40 SEM ( Fig. 5 h)
75 ± 47 0.0115 b
Optimal alloy 362 ± 9 23 21 STEM ( Fig. 9 b and d)
134 ± 72 (Mn) 74 ± 32 (Zr)
0.0093 (Mn) 0.0024 (Zr)
aThevolumefractioniscalculatedbyusingtheat%ofMnandZrandassumingthatallofitprecipitatesasAl12Mn2FeSi[72]orasAl3Zr[35]. Thereforethisvaluerepresentsthemaximumamount,becausenotallMnorZrwillprecipitateentirelyduringprocessing.
b Thevolumefractionofthedispersoidsin6082-ZristhesumoftheparticlescontainingMnandZr.Thereforethediameterusedheredoesnot differentiatebetweenthesetwodifferentphases.
Table5
Overviewofmechanicalparametersofallalloystestedforinterruptedquenching(allvaluesinMPa).
T651 (conv. AA) Interrupted quenching + shortened AA (2 h/160 °C)
R p0.2 R p0.2 Strength gain
Base alloy – 6082 289 ± 3 349 ± 8 60 ± 9
6082-Si 319 ± 8 364 ± 4 45 ± 9
6082-Cu 313 ± 3 355 ± 6 42 ± 7
6082-Zr 334 ± 5 384 ± 7 50 ± 8
Optimal alloy 362 ± 9 411 ± 4 49 ± 10
themeasuredincreaseinyieldstrengthlistedinTable3,of28MPain 6082-Zrand32MPaintheoptimalalloy,ismostlyattributedtosub- grainboundaryhardening(as-quenchedstrengthgainwithoutcontri- butionofartificialaging).AdditionalalloyingwithSi/MgandCudoes notsignificantlyinfluencetheas-quenchedstrength;seeSection5.1.
𝛿𝑏=10⋅(𝜏 𝐺
)−1
(1)
Δ𝜎=𝑀 ⋅𝜏 (2)
Nowweconsiderthestrengthgaincausedbyartificialagingofthe alloyscontainingZr,6082-Zrandtheoptimalalloy.Table3showsen- hancedvaluesduetoartificialaging(+17MPa)for6082-Zrcompared tothebase alloy.Becausethenumberofage-hardenable elementsis notalteredin6082-Zr,anadditionalmechanismisproposedhere.We notethatduringthe14daysofnaturalagingaplasticpre-deformation of2%wasapplied,whichgeneratedastrain(dislocation)hardeningof approximately39MPa.TheresearchofGruberetal.[75]showedthat duringagingat185°Casignificantreductionofthedislocationdensity takesplaceinleanAl–Mg–Sialloys.Suchsofteningcanalsobeassumed tooccurduringartificialageing(160°C/14h)of thebase alloy.Be- causethemodifieddispersoidlandscapepresentin6082-Zriscapable ofpreventingdislocationmotion(seeFig.11),thereforepreventingan- nihilationofdislocations,astrengthdecreaseduetoareductioninthe dislocationdensityisatleastpartlysuppressed.Thesameappliestothe optimalalloy.
Summingup theindividualstrength gainsachievedbyadding Si (6082-Si,21MPa),Cu(6082-Cu,17MPa)andZr(6082-Zr,17MPa), amaximumincrement of roughly55MPa shouldbe reachedduring artificialagingintheoptimalalloy,providedthesemechanismsactad- ditively.However,itisseenthatbothdispersoids(Fig.10b)anddisloca- tions(Fig.10d–f)actasheterogeneousprecipitationsitesintheoptimal alloy.Ontheonehand,hardeningparticlesprecipitatedondispersoids donotcontributemuchtotheoverallstrength,butdecreasethehard- eningpotential.Ontheotherhand,itcanbeassumedthathardening particlesnucleatingondislocationstendtocoarsenfaster,probablydue topipediffusionalongdislocations.Itisseenthatthesizeofparticles differsdistinctlyaccordingtowhethertheylieonornexttodisloca-
tionsinFig.10d–f.Consequently,thehardeningpotentialisweakened asparticlesgrowbigger.Someof theCucontentisalsoincorporated into(Al,Si)3Zr[26,76],whichlowerstheCucontentavailableforpre- cipitationformationaswell.Thisexplainswhythemeasuredgaininthe optimalalloycausedbyartificialagingcomparedtothebasealloysis actuallyabitlowerthan55MPa,showinganincreaseinyieldstrength of40MPa.
5.3. Interruptedquenching
ThethirdmeasureforincreasingthestrengthofAl–Mg–Sialloys,in- terruptedquenchingduringsolutionannealing,alsoprovedsuccessful (formechanicalparametersseeTable5).Whenquenchingthealloysto anelevatedtemperature(180°C)andslowlycoolingthemtoRT,which simulatesthecoolingofthickplates(roughly100mmthickness)inam- bientair,aminimumstrengthincreaseof42MPacouldbeseeninall alloyscomparedtotheconventionalT651state.Previousinvestigations showedthatunfavorableclustersformduringRTstorage,causedbythe highlevelofsuper-saturatedsolutesandquenched-invacanciesatthe endofsolutionannealing.Thisdecreasestheartificialagingresponse [77,78].Byapplying interruptedquenchingweexploited theadvan- tageousstateofthealloyrightattheendofsolutionannealingtopro- motetheformationofdenselydistributednucleifor𝛽’’ordirectlytrans- formableprecursors[13].Thisprocessingmethodworksforanyalloy variationandpromptedafurtherincreaseinstrength,finallyreaching ayieldstrengthof411MPaintheoptimalalloy.
6. Conclusions
This study investigates measures for significantly increasing the strengthofanENAW-6082alloy,whichisprocessedaccordingtoan industrialschemefortheproductionofthickplates.Themostimportant findingsaresummarizedasfollows:
• Athermodynamic simulation-basedadjustmentof age-hardenable elementsgeneratesastrengthincreaseof~ 40MPa.
• Interrupted quenching (IQ)can increase thestrength by at least 42MPa.
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