(7.9)--Precipitation and Hardening in M机械工程材料机械工程材料.pdf

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1、Precipitation and Hardening in Magnesium AlloysJIAN-FENG NIEMagnesium alloys have received an increasing interest in the past 12 years for potential appli-cations in the automotive,aircraft,aerospace,and electronic industries.Many of these alloys arestrong because of solid-state precipitates that ar

2、e produced by an age-hardening process.Although some strength improvements of existing magnesium alloys have been made and somenovel alloys with improved strength have been developed,the strength level that has beenachieved so far is still substantially lower than that obtained in counterpart alumin

3、um alloys.Further improvements in the alloy strength require a better understanding of the structure,morphology,orientation of precipitates,effects of precipitate morphology,and orientation onthe strengthening and microstructural factors that are important in controlling the nucleationand growth of

4、these precipitates.In this review,precipitation in most precipitation-hardenablemagnesium alloys is reviewed,and its relationship with strengthening is examined.It is dem-onstrated that the precipitation phenomena in these alloys,especially in the very early stage of theprecipitation process,are sti

5、ll far from being well understood,and many fundamental issuesremain unsolved even after some extensive and concerted efforts made in the past 12 years.Thechallenges associated with precipitation hardening and age hardening are identified and dis-cussed,and guidelines are outlined for the rational de

6、sign and development of higher strength,and ultimately ultrahigh strength,magnesium alloys via precipitation hardening.DOI:10.1007/s11661-012-1217-2?The Minerals,Metals&Materials Society and ASM International 2012I.INTRODUCTIONMAGNESIUMis the lightest of all commonly usedstructural metals,with a den

7、sity approximately twothirds that of aluminum and one quarter that of steels.Magnesium is an abundant element,comprising 2.7 pctof the Earths crust,and it is available commerciallywith purity exceeding99.8 pct.Magnesium has arelatively low melting temperature and high specificheat.Hence,magnesium an

8、d its alloys may,thus,bereadily cast to near-net shape by conventional castingmethods.Because of such attractive features,magne-sium alloys have received considerable research over thelast decade for potentially wider and larger applicationsin the automotive,aircraft,aerospace,and 3C(com-puter,commu

9、nication,and consumer electronic prod-uct)industries.Theannualproductionrateofmagnesium metal was approximately 450,000 tons in2001 and reached 720,000 tons in 2008.Despite theconsiderable efforts made thus far,the adoption ofmagnesium alloys in engineering applications remainslimited compared with

10、that achieved for aluminumalloys.One important technical reason is that there arelimited magnesium alloys for designers to select from forspecific applications,and within these limited choices,the most cost-effective magnesium alloys have inade-quate properties such as yield strength,creep-resistanc

11、e,formability,and corrosion resistance.The accumulatedempirical experience,rather than basic understanding,provides the tools for practical design and developmentof magnesium alloys with better mechanical and chem-ical properties.Many magnesium casting and wrought alloys achievetheir useful mechanic

12、al properties via age hardening,which involves(1)solution treatment at a relatively hightemperaturewithinthea-Mgsingle-phaseregion,(2)water quenching to obtain a supersaturated solidsolutionofalloyingelementsinmagnesium,and(3)subsequent aging at a relatively low temperature toachieve a controlled de

13、composition of the supersatu-rated solid solution into a fine distribution of precipi-tates in the magnesium matrix.The decomposition ofthe supersaturated solid solution often involves theformation of a series of metastable or equilibriumprecipitate phases that have a different resistance todislocat

14、ion shearing.Therefore,the control of theprecipitation is important if the maximum precipitationstrengthening effect is to be achieved.Attempts to improve the age-hardening response ofmagnesium alloys inevitably requires an in-depth under-standing of precipitation,precipitation hardening,andmicrostr

15、uctural factors that are most important in con-trolling the precipitation of strengthening phases and thestrength of precipitation-hardenable alloys.For precip-itation-hardenedmagnesiumalloys,theirmicrostructuresoften contain a distribution of plate-shaped or lath-rod-shapedprecipitatesofintermediat

16、eorequilibrium phasesformed parallel or normal to the basal plane of themagnesium matrix phase.In the last century,the crystalstructure,composition,and orientation relationship oftheseprecipitateshavebeencharacterizedprimarilyusingconventional transmission electron microscopy(TEM)JIAN-FENG NIE,Profe

17、ssor,is with the Department of MaterialsEngineering,Monash University,Clayton,VIC 3800,Australia.Contact e-mail:Jianfeng.niemonash.eduManuscript submitted February 1,2012.Article published online July 21,2012METALLURGICAL AND MATERIALS TRANSACTIONS AVOLUME 43A,NOVEMBER 20123891and electron diffracti

18、on.As a consequence of the resolu-tion and limitation of these techniques,the characteristicfeatures of some precipitate phases and the precipitationsequence in many alloys were not clearly established.Inthefirstdecadeofthiscentury,withtheassistanceofhigh-resolutiontransmissionelectronmicroscopy,par

19、ticularlyatomic-resolutionhigh-angleannulardark-fieldscanningtransmissionelectronmicroscopy(HAADF-STEM),andthree-dimensional atom probe(3DAP),some puzzles onthe structure and composition of precipitate phases insome existing magnesium alloys have been solved.Thesemodern characterization facilities a

20、lso greatly facilitatethe identificationofprecipitates inmagnesiumalloysthatare developed in recent years.Such knowledge on thecrystallography of precipitate phases provides the basisfortheunderstandingoftheformationandstrengtheningmechanisms of the precipitate phases and,more impor-tantly,for the r

21、ational alloy design in practice.The purpose of this article is to provide a compre-hensive review of the literature on precipitation andhardening in most,if not all,age-hardenable magnesiumalloys.Because a few books on magnesium alloys14and some review articles on precipitation in magnesiumalloys58

22、and particle hardening912are already in theliterature,the emphasis of this article will be focused on(1)the structure,morphology,and orientation ofprecipitates,precipitationsequenceandhardeningresponse in each of the major alloy systems;(2)theeffects of precipitate shapes on strengthening;and(3)the

23、rational design of microstructures for larger age-hardening response and therefore higher strength.Someunsolved issues that require additional research are alsohighlighted and discussed.II.PRECIPITATION AND AGE-HARDENINGRESPONSEA.Mg-Al-Based Alloys1.PrecipitationThe magnesium-rich side of the Mg-Al

24、binary phasediagram includes equilibrium solid phases a-Mg and b-Mg17Al12,as well as a eutectic temperature of 710 K(437?C).The b phase has a body-centered cubicstructure(space group I?43m)with the lattice parametera 1.06 nm.13The equilibrium solid solubility of Al ina-Mg is 11.8 at.pct(12.9 wt pct)

25、at the eutectic tem-perature,and it decreases to approximately 3.3 at.pct at473 K(200?C).14The equilibrium volume fraction ofprecipitates achievable in the Mg-Al alloys aged at473 K(200?C)can reach a substantially large value of11.4 pct.This thermodynamic feature provides a uniqueopportunity for gen

26、erating a large volume fraction ofprecipitates by using conventional aging treatments,i.e.,solution treatment at approximately 692 K(420?C),followed by water quench and subsequent aging at atemperature in the range of 373 K to 573 K(100?C to300?C).Unfortunately,during the isothermal agingtreatment i

27、n the temperature range 373 K to 573 K(100?C to 300?C),the precipitation process seems toinvolve solely the formation of the equilibrium b phase(Table I).Although the b precipitates are resistant todislocation shearing,15,16their distribution is relativelycoarse,presumably because of the relatively

28、high diffu-sion rate of Al atoms in the solid matrix of magnesiumand a possibly high concentration of vacancies in the a-Mg matrix.Consequently,the age-hardening responseof Mg-Al alloys15,1720is not as appreciable as expected(Figure 1(a).19,20Previous studies16,21revealed that the precipitation ofth

29、e equilibrium b phase occurs both discontinuously andcontinuously.The discontinuous precipitation is alsoknown as cellular precipitation,and in this reaction,thesupersaturated solid solution a phase decomposes intothe b phase and an a phase that is structurally identical tothe a phase but has a less

30、 saturated concentration ofaluminum.The discontinuous precipitation initiates ingrain boundaries and expands toward the grain center ina cellular form.22The cell comprises a lamellar structureof b and a phases,and the cell interface separating a anda is a high angle boundary.The continuous precipita

31、tionoccurs inside the grains.The continuous and discontin-uous precipitations occur simultaneously and competewith each other during isothermal aging of Mg-Al alloys.Duly et al.2325reported that the continuous precipita-tion is favored at both high and low aging temperaturesand that discontinuous pr

32、ecipitation dominates themicrostructure at intermediate temperatures.They pro-posed that the disappearance of discontinuous precipi-tation at high aging temperatures is caused by thevolume diffusion of solute that prevents the nucleationand growth of the cellular colonies and that the absenceof disc

33、ontinuous precipitation at low aging temperaturesis the result of a lower driving force,which is caused bythe occurrence of continuous precipitation in the earlystage of aging treatment.A more recent study of a binary Mg-9 wt pct Al alloyand alloy AZ9126indicates that in the binary Mg-9wt pct Al all

34、oy samples aged at 423 K(150?C)or cooledfrom the solution temperature to room temperature,only discontinuous precipitates are observed,whereasthat only continuous precipitates form when the binaryand the AZ91 alloys are aged at 623 K(350?C).It is alsofound that both discontinuous and continuous prec

35、ipi-tates form when the alloys are aged at intermediatetemperatures of 473 K or 523 K(200?C or 250?C).Itwas proposed26that whether discontinuous precipita-tion occurs also depends on the concentration ofvacancies in addition to the aging temperature.Formostcontinuousprecipitatesanddiscon-tinuous pre

36、cipitates of the b phase in the lamellarstructure,theywereinitiallyreportedtoadopttheexactBurgersorientationrelationship,i.e.,011b=0001a;1?11?b=2?1?10?a.27Subsequenttransmission electron microscopy studies28,29indicatethat the orientation relationship is actually near theBurgers.The b precipitates i

37、n this orientation relation-ship have a plate morphology,with their broad surfaceparallel to(0001)a(Figures 1(b)and(c).Although the bplates in this orientation relationship are often describedas incoherent,7,8,30ample experimental evidence dem-onstrates that the equilibrium b phase is in fact notinc

38、oherent.Apart from the apparent lattice matchingbetween the b phase and surrounding matrix phase in3892VOLUME 43A,NOVEMBER 2012METALLURGICAL AND MATERIALS TRANSACTIONS ATable I.Part of the Whole Precipitation Sequence in Individual Magnesium Alloy SystemsPrecipitation process is not well studied;*d

39、is separation distance of columns of RE atoms;#low Gd:Zn weight ratio and low Gd content.METALLURGICAL AND MATERIALS TRANSACTIONS AVOLUME 43A,NOVEMBER 20123893the plate broad surface,or habit plane,the latticematching is also found in interfaces defining the majorand minor side facets of individual

40、b plates.28,29,31Despite the irrational orientation of these side facetswith respect to both precipitate and matrix lattices,themajor and minor side facets(Figure 2(a)are invariablyparallel to the moire fringes defined by the intersectionof1?100?aand0?33?b,and of10?10?aand4?11?b,respectively.Figure

41、2(a)shows the major side facet of athin b plate that is embedded in the matrix phase.Thismajor interface is parallel to the moire fringes resultingFig.1(a)Isothermal aging curves of magnesium alloy AZ91 at373 K and 473 K(100?C and 200?C)(adapted and reproducedfrom Refs.19 and 20).(b and c)Transmissi

42、on electron micro-graphs showing the distribution and morphology of b precipitates insamples aged for 8 h at 473 K(200?C).Electron beam is parallel to2?1?10?ain(b)and 0001ain(c).(b)is reproduced from Refs.29and 20,and(c)is from Ref.20.Fig.2(a)0001aTransmission electron micrograph showing theparallel

43、ogram shape of b precipitates in AZ91(adapted and repro-duced from Refs.29 and 20).(b)High-resolution transmission elec-tron micrograph showing the major planar interface of b precipitatesuch as that shown in(a);the planar interface is indicated by moire planes(reproduced from Ref.33).(c)Schematic d

44、iagram showingthe orientation relationship between two sets of lattice planes,e.g.,1?100?aand0?33?b;and their resultant moire plane,which corre-sponds to shown in(b).This diagram demonstrates the coherentmatching of the two sets of lattice planes within the moire plane.343894VOLUME 43A,NOVEMBER 2012

45、METALLURGICAL AND MATERIALS TRANSACTIONS Afrom the overlapping of the1?100?aand0?33?bplanes(Figure 2(c),and it contains some ledges whose unitheight is defined by the interplanar spacing of the moire fringes.The migration of this interface in its normaldirection seems to involve the formation and la

46、tergliding of moire ledges within the interface plane.32,33These observations suggest the existence of commensu-rate matching of1?100?aand0?33?bplanes29,33,34inthe major facet interface,and of 10?10?aand 4?11?binthe minor facet interface.This commensurate matching,together with the fact that0?33?ban

47、d4?11?bare theclosest-packed planes in the b lattice and1?100?ais thenear closest-packed plane in the magnesium lattice,suggests that the major and minor side facets of each bplate have relatively low interfacial energies.Apart from the near Burgers orientation relationship,two other orientation rel

48、ationships have also beenreported for the b phase,19,27,29,35,36namely1?10?b=1?100?a,111?b=0001?aand1?10?b?=1?100?a,11?5?b?=0001?a.For the former orientation relation-ship,the b precipitates have a rod shape with their longaxes parallel to0001?a.The0001?arods have ahexagonal cross section,with the b

49、ounding facetsparallel to1?100?a=?330?b.For the latter orientationrelationship,the b precipitates develop a rod shape withtheir long axes inclined with respect to the0001?adirection.Although these rods are more effective thanthe(0001)ain impeding dislocation gliding on the basalplane,only a small fr

50、action of them exists in themicrostructure.It is currently unclear how to promotethe rod-shape precipitates at the expenses of the(0001)aplates.Such an effort inevitably requires an in-depthunderstanding of the transformation strains associatedwith each of the orientation relationships and theactiva