Focused Development of Magnesium Alloys Using the Calphad Approach

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Focused Development ofMagnesium Alloys Using theCalphad Approach**By Rainer Schmid-Fetzer,* and Joachim Gröbner

1. IntroductionThe current revival of magnesium as a structural material

is carried by a relatively small number of traditional Mg-alloys. Compared to that, a large number of highly special-ized and sophisticated aluminum alloys, not to speak of steel,was developed in an ongoing effort over the past decades. Itis evident that there is an urgent need for the development ofnew or improved magnesium alloys if we want to fully ex-ploit the potential of this fascinating lightweight material thatalso offers excellent castability, machinability, and bio-com-patibility. Experiments on a technological scale for prepara-tion and testing of new alloys are very expensive and timeconsuming. In view of the huge number of possible alloycomponents, compositions, and processing parameters, onewould like to have at least an ªeducated guessº in whichdirection to go. In this report we want to show that thermody-namic calculations can provide much more than that. Com-putational thermochemistry based on the Calphad method isa modern tool that supplies quantitative data to guide thedevelopment of alloys or the optimization of materialsprocessing.[1]

It enables the calculation of multicomponent phase dia-grams and the tracking of individual alloys during heat treat-ment or solidification by calculation of phase distributionsand phase compositions. It also allows the simulation ofphase transformations during solidification of Mg-alloys,which are responsible for the development of as-cast micro-structures. This can be done either by using the two simpleextreme models of equilibrium solidification or Scheil solidifi-cation, or by more sophisticated models that require addi-tional kinetic material parameters.[2] Other quantities of mul-ticomponent Mg-alloys, important for materials processing,

ADVANCED ENGINEERING MATERIALS 2001, 3, No. 12 947

±[*] Prof. R. Schmid-Fetzer, Dr. J. Gröbner

Technical University of ClausthalInstitute of MetallurgyRobert-Koch-Str. 42D-38678 Clausthal-Zellerfeld (Germany)E-mail: [email protected]

[**] Part of this work is supported in the ªThrust Research ProjectSFB 390: Magnesium Technologyº by the German ResearchCouncil (DFG).

In traditional alloy development, experimental investigations withmany different alloy compositions are performed. The selection criteriafor multicomponent alloying elements and their compositions becomediffuse in a traditional approach. Computational thermochemistry as used in the Calphad approachcan provide a clear guideline for such selections and helps to avoid large scale experiments with lesspromising alloys. Thus, it is a powerful tool to cut down on cost and time during development ofMg-alloys. An overview of the Calphad method is given. As an example of applications, recent devel-opments of new creep resistant magnesium alloys that show about 100 times less creep than the bestcommercial alloys are reported. Also outlined are the methods used in our long-term project of con-struction of the necessary thermodynamic magnesium alloy database for several alloying elements,such as Al, Li, Si, Mn, Ca, Sc, Y, and Zr, and rare earth elements (Ce, Gd, Nd), using the Calphadmethod combined with key experiments.

Calphad ApproachCalphad Approach

DataModels Software

Stable & MetastablePhase Equilibria

ProcessSimulation

DirectApplication

KeyExperiments

Ó WILEY-VCH Verlag GmbH, D-69469 Weinheim, 2001 1438-1656/01/1212-0947 $ 17.50+.50/0

Schmid-Fetzer, Gröbner/Focused Development of Magnesium Alloys Using the Calphad Approach

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ample the heat content, which is important for die casting, orthe chemical potentials of individual alloy components,which are important for selective vaporization or oxidation.

These are the basic data to understand and control thebehavior of any novel or modified Mg-alloy. Large-scaleexperiments for new multicomponent alloys can then befocused on the most promising alloys identified in that ap-proach. Long-term experiments with less promising alloyscan be avoided. Thus, it is a powerful tool to cut down on costand time during the development of magnesium alloys.

Quantitative access to all these basic data in a consistent,numerical, and easily readable form isÐat least for multicom-ponent alloysÐvirtually impossible without the use of com-putational thermochemistry. The fundamental idea goes backto the introduction of the Calphad method by Larry Kaufmanin 1970.[3] In a nutshell it is a method to calculate a myriad ofdata, ranging from stable and metastable phase equilibriadown to the chemical potentials, from a unique set of thermo-dynamic model parameters of the alloy system.

In the following, the foundations, applications, and recentadvancements of the Calphad approach are briefly reviewed,emphasizing the predictive capability to multicomponentalloys. The ongoing long-term effort in developing a thermo-dynamic magnesium alloy database by modeling and key ex-periments is outlined. And, as an example of its application,the focused development of new creep resistant Mg-alloys isillustrated. Two generations of these new alloys, the first with

high and the second with low scandium content, have beendeveloped using computational thermochemistry as a power-ful tool. Their creep properties are superior compared to thebest commercial alloys.

2. The Calphad Approach2.1. Foundations and Current State

Phase diagrams provide the graphical presentation of theequilibrium state of a material as a function of temperature,pressure, and composition of the components. This is whythey are frequently used as roadmaps for alloy design or abetter understanding of the processing of materials. The ther-modynamic properties of materials, such as the heat of solidi-fication or the chemical activity of components, are also fre-quently used to understand, for example, metallurgicalreactions of materials. These two aspects, phase diagramsand thermodynamic properties, have been treated separatelyfor a very long time despite the fact that their fundamental in-terrelations had been established more than a century ago byJ. W. Gibbs, whose work has been summarized by Hertz.[4]

Eventually, the mathematical calculation of phase dia-grams arose; some early examples between 1908 and 1970 aresummarized by Kattner.[5] In 1970, Kaufman et al.[3] initiatedthe Calphad method (calculation of phase diagrams) with adetailed description of procedures together with a listing ofcomputer software. The subsequent meetings of the ever

948 ADVANCED ENGINEERING MATERIALS 2001, 3, No. 12

Rainer Schmid-Fetzer is Professor at the Institute of Metallurgy, Technical University of Clausthal. Hegraduated with a diploma in physics, earned his Dr.-Ing. degree and finished his Habilitation in the Insti-tute of Ferrous Metallurgy and Foundry Technology at the TU Clausthal. Between 1982 and 1984 he wasfirst Visiting Scientist and then Visiting Associate Professor at the University of Wisconsin-Madison,USA, where he became interested in functional materials. After returning to the TU Clausthal he startedhis research on electronic materials. In 1994 he was Visiting Scientist at the Microelectronics and Micro-systems section of the Daimler-Benz Corporate Research Institute, Frankfurt. In 1997 he was Visiting Pro-fessor at the Dept. of Materials Science and Engineering, University of Wisconsin-Madison, contributingto an ongoing joint effort to develop a second generation of software for thermodynamic calculations in mul-ticomponent systems. His research interests comprise the constitution and thermodynamics of materials,alloy development, interface reactions in bulk and thin film electronic materials, metals and composites andcontacts to semiconductors. Professor Schmid-Fetzer has published over 160 papers, holds one patent, andreceived the Werner-Köster award of DGM in 1996. He is the current chairman of the DGM Applied Con-stitution Committee and vice chairman of the Alloy Phase Diagram International Commission.

Joachim Gröbner studied mineralogy at the TU Stuttgart. In the ªPulvermetallurgisches Laboratoriumºhe started working in the field of thermodynamics of ceramic materials and obtained his Ph.D. at theMax Planck Institute for Metal Research in Stuttgart in 1994. He then participated in several interna-tional projects on metal themochemistry at the universities of Vienna and Barcelona. Since 1998 heworks on the development of new magnesium alloys in the frame of the SFB 390 ªMagnesium Technol-ogyº at the Institute of Metallurgy in Clausthal. During this time he contributed to the development ofa multicomponent thermodynamic database for magnesium alloys.


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growing Calphad group, organized by Larry Kaufman, soonattained the level of annual international conferences. Thesemeetings together with the establishment of the internationalCalphad journal helped in the rapid development of sophisti-cated thermodynamic modeling procedures and their cou-pling to phase equilibrium calculations. One important devel-opment concerns the establishment of an acknowledgedstandard for the Gibbs energies of the pure elements in alltheir stable and metastable states,[6] based on the concept oflattice stabilities.[3,7,8] This standard is essential to enable thecombination of separate assessments of binary systems tohigher order systems. More advanced methods to modelthese ªunaryº data for elements and stoichiometric end-mem-bers of solutions have also been developed, addressing theheat capacity of crystalline phases,[9] liquid and amorphousphases,[10] enthalpy estimations for stable and metastablestates,[11] enthalpies and entropies of transition,[12] k-transi-tions,[13] and periodic system effects.[14] Highlights of the de-velopments of models for phases with a range of compositionwill be detailed in Section 2.3.

This progress in thermodynamic modeling together withthe capabilities of modern computational technology led tothe current state of the Calphad approach, which is character-ized by the following highlights.l A predictive capability allows the extrapolation of ther-

modynamic descriptions and phase equilibrium calcula-tions from assessed binary systems to ternary, quaternaryand higher order systems. This feature is indispensablefor an application to real life alloys or processes, it will bedescribed further in section 2.3.

l Identification of key experiments drastically reduces thenecessary experimental effort in multicomponent systems.Key experiments can be selected to give the maximumamount of information, either by verification of extrapo-lated descriptions or in assessing higher order interac-tions. The simultaneous evaluation of all experimentaldata irrespective of their type (phase equilibrium, thermo-dynamic data) and location in the multicomponent systemis only possible with this approach and results in aninternally consistent database of the entire system.

l Stable and metastable phase equilibria can be calculated.By suspending individual phases, which may not formunder specific conditions due to kinetic reasons, the com-plete metastable phase relations can be given quantita-tively.

l The driving forces for all phase transformations are avail-able.

l Local phase equilibria can be calculated, providing a nu-merical input to materials processing software, for exam-ple in solidification simulation or reactor modeling.

l The reading of multicomponent phase diagrams is drasti-cally simplified by the fact that all the interesting two-dimensional sections in temperature, composition or evenchemical potential can be readily calculated, plotted andread directly, instead of trying to envisage or construct

them from an only graphically available set of phase dia-grams and their complex projections. Moreover, for anyselected individual alloy, the amounts of the phasespresent and their compositions can be calculated as func-tion of temperature. It may be rather cumbersome to ex-tract that information if only a printed set of ternary phasediagrams is given, and it is virtually impossible forquaternary or higher order systems.

This powerful tool in materials research goes beyond amere ªcalculation of phase diagramsº. This is why the termªcomputational thermochemistryº is frequently used todescribe the current state of the Calphad approach. It isdepicted in Figure 1 and will be detailed below.

Calphad ApproachCalphad Approach

DataModels Software

Stable & MetastablePhase Equilibria

ProcessSimulation

DirectApplication

KeyExperiments

Fig. 1. Sketch of the Calphad approach, combining thermodynamic modeling with keyexperiments and applications in an iterative database development.

2.2. Applications and Software

Applications of the Calphad approach can be classified infour categories: direct applications, coupling with steady statereactor modeling, coupling with micro-kinetics, and couplingwith macro-kinetics. Only the latter two require the knowl-edge of kinetic data of the materials system.

The direct applications of information read from phasediagrams or thermodynamic equilibrium calculations startedabout a century ago with the first understanding of alloyingbehavior, microstructure development, and metallurgicalreactions. These applications are far from being ªold-fash-ionedº, though. Today we are able to apply quantitativephase equilibrium calculations in multicomponent multi-phase materials to the development of technical alloys andalso to the optimization of near-equilibrium processing condi-tions.[1] This goes as far as setting the electronic carrier con-centration in CdTe semiconductors by annealing at appropri-ate partial pressures of Cd.[15] Many successful applicationsrely on the fact that the actual process conditions are close toequilibrium and a large number of examples covering metals,alloys, intermetallics, ceramics, semiconductors, geologicalmaterials, nuclear core materials, or even waste, and a varietyof processes were given.[1] The calculations of solidificationpaths under equilibrium conditions or even with the popularScheil model are also considered as very useful direct appli-



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blocking the solid state diffusion in the Scheil model.[2] Anapplication to the magnesium alloys AZ91 and AZ62 wasgiven together with different experimental methods.[16]

However, even materials processing far from equilibrium(such as splat cooling, mechanical alloying, and annealingof reactive thin film multilayers) requires a sound knowledgeof thermodynamic data of equilibrium and metastablephases.[17±19] Using the Calphad approach the driving forcesfor metastable phase formation or even complete metastablephase diagrams can be calculated to provide ªroad mapsº forthe engineering of these advanced materials.

A coupling with steady state reactor modeling is done asfollows. A reactor, e.g., an electric arc furnace can be modeledby splitting it up into a certain number of reactor stages.Within each stage a (local) equilibrium calculation can be per-formed using the thermodynamic data set of the materialssystem. In addition to the number of stages one has to specifythe feed phases (amounts, compositions, temperatures, inputstages), the enthalpy supplied/evolved or the temperatureimposed on each stage, pressure of each stage, and optionallydistribution coefficients of phases, which may simulate ki-netic inhibitions. Steady-state equilibrium is attained by equi-librium calculations where the output of one stage is used asinput of the next stage in co- and counter-current flow reac-tors. Convergence is usually achieved with reasonably esti-mated initial values for each stage obtained by separate equi-librium calculations. A nice example is the production ofmetallurgical grade silicon in an electric arc furnace, wherethe ChemSage software was used.[20] It should be noted thatactual kinetic data (diffusion, fluid flow, ...) of the materialssystem are not necessary in this approach. The optionally im-posed phase distribution coefficients are not considered asdetailed kinetic data, for example, the imposed fraction ofNH3 in MOCVD growth of nitride semiconductors in a sim-ple one-stage reactor.[21]

A coupling of the Calphad approach with micro-kineticsconcerns mostly solid state transformations, growth or dis-solution of precipitates, reactions at solid/solid interfaces,and so on. Fluid flow and heat transfer problems do not occuror are irrelevant. Thermal equilibrium is often assumed sincethermal diffusion is much faster compared to species diffu-sion in solids. The framework of multicomponent (multi-phase) diffusion is used, which requires the full data set ofdiffusion coefficients of all species in all phases. Possibly eventhe nucleation may be included with an appropriate model.The phase growth is usually assumed to be diffusion con-trolled. Appropriate software is designed with analyticaland/or numerical methods.[1,2]

Thermodynamic software and data sets are called as aªsubroutineº to provide local equilibrium calculations andthe driving forces for diffusion. These calculations yieldquantitative data at locally prevailing temperature, pressure,composition, or chemical potentials. It is advisable to use theªsubroutineº approach and not to fully integrate the thermo-

dynamic and diffusion software. The thermochemical andphase equilibrium calculations in real systems may becomerather complex and should be thoroughly tested separatelyfor internal consistency and accuracy. A prominent exampleof this approach is the DICTRA software;[22] just a few appli-cations are mentioned here for the multicomponent diffusionin steel.[22,23]

Coupling computational thermochemistry with macro-kinetics is at the front end of integration in materials process-ing software. Comprehensive simulation of dynamic, macro-scale processes involving fluid flow constitute the highestlevel of complexity and requirement of additional systemdata. An example is the simulation of technical alloy casting,where three-dimensional simulations of the moving mushysolidification front, void formation at hot spots, and macro-segregation may be covered. Other aspects such as microse-gregation, dendrite shape formation, and ultimately thedevelopment of microstructure is still mostly treated sepa-rately. Both aspects require thermodynamic equilibrium cal-culations at the core of appropriate process simulation soft-ware that tries to solve the simultaneous mass-, heat-, andmomentum-transfer with moving phase boundaries.[2]

Such process simulation software is working in the frame-work of computational fluid dynamics (CFD) or flow-sheet-ing. Depending on the complexity, a full set of data for themacro-kinetic description of the system may be needed, pos-sibly including the micro-kinetic data. The current challengeis to incorporate exact thermodynamic phase equilibrium cal-culations into commercial software packages such as Phoe-nics, Fluent, Aspen, Process, and so on. To this aim a generalsoftware interface TQ (ChemApp) has been developed thatallows to call the most widely used Gibbs energy minimizersas a subroutine.[1,24] An example for the casting process,where the link to thermodynamic software was programmedwith another software interface, is given in the literature.[25] Ifthe simulation of mushy zone is also included in an attemptto couple macro- and micro-kinetics, the computation timebecomes a major issue.[26] Another problem is that the ther-modynamic software subroutine has to be 100 % reliable andconvergent for each individual complex equilibrium calcula-tion even if starting from unrealistic initial values set by theprocess software. Envisaging about 105 calls of the thermody-namic subroutine, a single failure return would obstruct theprocess simulation.

The early versions of software to calculate multicompo-nent phase equilibria, Solgasmix by Eriksson,[27] and to calcu-late (mainly binary) phase diagrams, the Lukas-programs,[28]

were given away free as source code, which substantially fos-tered the development of the field. More general mathemati-cal relations to calculate phase diagrams were laid out byHillert,[29] and were eventually programmed in Thermo-calc.[30] Currently a number of software packages, sometimescombined with databases, are commercially available.Frequently used are ChemSage and Fact, recently merged tobecome FactSage, MTDATA, Thermocalc, and a second genera-



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tion of software, Pandat,[64] aiming at an automatic calculationof multicomponent phase diagrams without the need of userknowledge of the system.

2.3. Thermodynamic Modeling of MulticomponentMultiphase Equilibria

The basic principle to find the equilibrium state of a multi-component alloy (or system) is to minimize its total Gibbsenergy, G, at constant temperature T, pressure P, and overallcomposition:

G �P�

n�G� = min (1)

where Gu is the molar Gibbs energy of any possible phaseu in the system and nu is its amount. Minimization is done bydistributing the given amount of alloy components onto thevarious available phases, subject to the overall material bal-ance. The stable equilibrium is described by a phase assemblyof all phases with nu > 0, whereas the metastable phases(nu = 0) do not materialize. Each phase has its own composi-tion, which is generally different from the fixed overall alloycomposition.

We have thus to supply the functions Gu for all the phases.For the pure element i in the u phase the Gibbs energy

G0;ui (T) = Gu

i (T) ± HSERi (2a)

is described by the equation

G0;�i �T� � a� b � T � c � T � ln T � d � T2 � e � T3 � f � T

ÿ1 � g � T7 � h � Tÿ9 (2b)

where HHSERi is the molar enthalpy of the stable element ref-

erence (SER) at 298.15 K and 1 bar. The values of the parame-ters a, b, c,... are usually taken from the SGTE compilation byDinsdale.[6]

Stoichiometric phases like the binary compound AmBn canbe described by the relation

GAm Bn1:2

� mm�n

� G0;�1� n

m�n� G0;�

2

�amn � bmn � T � ::: (3)

where the last two terms stand for the Gibbs energy of forma-tion of AmBn from the pure components 1(=A) and 2(=B) inthe specified reference state u.

The molar Gibbs energy, Gu, of a binary disordered solu-tion phase u, applicable for most liquids or substitutional sol-id solutions, is generally described as:

G� � x1 G0;�

1 � x2 G0;�2 � RT x1 ln x1 � x2 ln x2

ÿ ��x1 x2 L0;�

12� L1;�

12�x1 ÿ x2� � L2;�

12

��x1 ÿ x2�2 � :::� (4)

Hillert[31] pointed out the advantages of this specific for-mulation[32] of the power series for the excess term to the idealsolution when this equation is used for extrapolation to tern-ary systems.

For a multicomponent solution phase with c componentsthe following equation is used:

G� � Pci�1

xi G0;�i � RT

Pci�1

xi ln xi � E Gbin;�

�E Gtern;� � E Gquat;� � ::: (5)

The excess contributions originating from all the binary in-teractions (EGbin), ternary interactions (EGtern), or quaternaryinteractions (EGquat) are:

E Gbin;� � Pcÿ1

i�1

Pcj>i

xi xjPn

v�0L�;�ij �xi ÿ xj�� (6)

E Gtern;� � Pcÿ2

i�1

Pcÿ1

j>i

Pck>j

xi xj xkfL1;�ijk �xi � �ijk�

�L2;�ijk �xj � �ijk� � L3;�

ijk �xk � �ijk�g (7)

E Gquat;� � Pcÿ3

i�1

Pcÿ2

j>i

Pcÿ1

k>j

Pcl>k

xi xj xkxlfL1;�ijkl �xi � �ijkl�

�L2;�ijkl �xj � �ijkl� � L3;�

ijkl �xk � �ijkl� � L4;�ijkl �xl � �ijkl�g (8)

where

�ijk = (1±xi±xj±xk)/3 (9)

�ijkl = (1±xi±xj±xk±xl)/4 (10)

It is noted that in a ternary system (c = 3) dijk = 0 whereasin a quaternary system with xl ¹ 0 the same term dijk ¹ 0. In-stead of Equations 7,8 one might have used seemingly sim-pler equations where the dijk and dijkl terms would have beenomitted. The main advantage of the given formulation is thatEquation 7 always reduces to the regular solution contribu-tion if all the three L-parameters are identical:

if

L1;�ijk � L2;�

ijk � L3;�ijk � L�ijk (11)



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E Gtern;� � Pcÿ2

i�1

Pcÿ1

j>i

Pck>j

xi xj xk L�ijk (12)

even in a quaternary or higher order system. Without usingthe dijk term in Equation 7 a residual quantity xixjxkLijk(1 ± xl)would be left in quaternary. This occurrence of xl would beincompatible with the regular interaction of the three speciesi-j-k, correctly given in Equation 12. Such a regular interactioncan be derived from thermodynamic principles assuming dif-ferent pairwise i±j bond energies and an ideal entropyterm.[33] The same reasoning is valid for the quaternary inter-action.[65]

This was clearly pointed out by Hillert[31,33] and in fact theEquations 5±10 are just extensions of his equations. The for-mulation is slightly different, the present dijk (and dijkl) vari-ables are invariant to a permutation of indices whereasHillert's original variables (mijk) are not.

It should be also noted that the ternary and quaternary in-teraction terms in Equations 7 and 8 provide a subregularbehavior with a linear composition dependence inside the{}-bracket. No higher terms have yet been used in practice. Ina predictive approach from the known binary systems to ac-component system one would not use any ternary or qua-ternary interaction terms. If those are used at all it certainlydoes not make sense to envisage a higher than the subregularorder given in Equations 7 and 8.

Reviewing Equations 5±10 it becomes evident what isreally meant by the predictive capability of this approach. Letus, for simplicity, assume that we deal with a c = 4 compo-nent magnesium alloy system, which comprises, in additionto stoichiometric compounds, just two solution phases, theliquid phase L and the (Mg) solid solution. Once EGbin of thesix binary edge system have been assessed, GL and G(Mg) canbe calculated from Equation 5, setting EGtern = EGquat = 0. Thisenables equilibrium calculations in the entire quaternaryalloy system. It is a general experience that the binary interac-tions are most prominent, that is, the higher order interac-tions like EGtern are usually small.

In the next level of refinement these calculations can beused to identify key experiments at decisive conditions in thefour ternary edge systems, either to verify that EGtern ^ 0 orto determine quantitative values. In addition, solid solutionsof binary compounds into the ternary systems and true tern-ary compounds have to be modeled.

Using both EGbin and EGtern in Equation 5 results not only inexperimentally supported calculations in the ternary edgesystems, it also improves the reliability of calculations in thequaternary alloy system. In the next level key experiments inthe quaternary system can be identified. Experience showsthat generally EGquat ^ 0 if EGbin and EGtern are describedproperly.

This interactive approach ªextrapolative calculationsº±ªkey experimentsº±ªsupported modelingº is also useful dur-ing assessment of the EGbin of individual binaries. In ternary

or higher order systems, however, it is indispensable consid-ering the huge number of possible experimental conditionsthat cannot be handled by a plain trial and error approach,due to limited resources.

More complex solution phases require models that takethe specific physical nature of the interactions between thecomponents and/or the crystalline structure of the phase intoaccount. This reduces the number of adjustable parametersand can provide realistic functional behaviors, which are outof the scope of Equation 4, such as a sharp bend in the Gu±xcurve. A classical example is the Wagner±Schottky model[34]

for ordered intermetallic phases, which derives the composi-tion dependence of the Gibbs energy from the generation andinteraction of point defects in the perfect crystal. Substantialprogress had been made on the thermodynamic modeling ofsolutions,[35] specifically for liquids,[36] selected intermetalliccompounds,[37] order±disorder phase transitions,[38] oxides,[39]

electronic materials,[40] and more advanced models for thecomposition dependence of intermediate phases.[41]

3. Magnesium Database Development3.1. Thermodynamic Modeling

In order to use the Calphad approach as a tool in magne-sium alloy design, a high quality thermodynamic databasefor Mg-alloys is needed. Development of a reliable thermody-namic database for multicomponent alloys requires a combi-nation of experiments and computational thermochemistrywith a feedback of data from alloy application. Since numer-ous binary and ternary subsystems have to be treated and va-lidated with key experiments before multicomponent alloyscan be calculated reliably, this development becomes a long-term project. In our group a thermodynamic database for sev-eral alloying elements, e.g., Al, Li, Si, Mn, Ca, Sc, Y, Zr andrare earth elements has been under construction for morethan six years and is still ongoing. The base of this 13-compo-nent system are the binary systems. Figure 2 gives an over-view of the necessary systems and the current state of assess-ment. Highlighted in the first column are the includedelements. Also shown are the current quaternary aim systemsAl-Mg-Li-Si and Mg-Mn-Sc-RE (with RE = Ce, Gd, Sc, and


Mg-RE-Mn-ScRE = Ce, Gd, Y

AlCaCeGdLiMgMnNdScSiYZn Zr

-Ca Al-Ce Al-Gd Al-Li Al-Mg Al-Mn Al-Nd Al-Sc Al-Si Al-Y Al-Zn Al-Zr-Ce Ca-Gd Ca-Li Ca-Mg Ca-Mn Ca-Nd Ca-Sc Ca-Si Ca-Y Ca-Zn Ca-Zr-Gd Ce-Li Ce-Mg Ce-Mn Ce-Nd Ce-Sc Ce-Si Ce-Y Ce-Zn Ce-Zr-Li Gd-Mg Gd-Mn Gd-Nd Gd-Sc Gd-Si Gd-Y Gd-Zn Gd-Zr-Mg Li-Mn Li-Nd Li-Sc Li-Si Li-Y Li-Zn Li-Zr-Mn Mg-Nd Mg-Sc Mg-Si Mg-Y Mg-Zn Mg-Zr-Nd Mn-Sc Mn-Si Mn-Y Mn-Zn Mn-Zr-Sc Nd-Si Nd-Y Nd-Zn Nd-Zr-Si Sc-Y Sc-Zn Sc-Zr-Y Si-Zn Si-Zr-Zn Y-Zr- Mg-Mn-Sc

Mg-Mn-REMg-Sc-REMn-Sc-RE

Al-Mg-Li-Si

Al-Li-SiAl-Mg-LiAl-Mg-SiMg-Li-Si

Gray: donelight: missing

Fig. 2. The binary systems of the currently developed 13-component Mg-alloy databaseand some selected multicomponent systems.


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Y).78 independent binary systems are necessary as basis for a13-component system. 26 of them could be taken from litera-ture and included in our database after checking the consis-tency. For 23 binary systems the thermodynamic descriptionswere created by our group. These systems with existing con-sistent thermodynamic data sets are marked with a graybackground in Figure 2. Two main paths for the current Mg-alloy development are indicated by arrows. The first group ofalloys are Li-containing, aiming at a further density reductionand improved ductility. The second group with Sc and RE-additions aims at increased creep resistance. In this paper thesecond group is discussed as an example of focused alloy de-velopment.

To create the thermodynamic phase descriptions the Cal-phad method is used as given above and depicted in Fig-ure 1. After critically reviewing the experimental data and se-lecting the appropriate models for all phases, the actualdetermination of the thermodynamic parameters (e.g., theamn, bmn, or Lm,u

ij ) is done by multidimensional least square op-timization. This optimization is, depending on the complexityof the system, stepwise until finally all parameters of that sys-tem are optimized simultaneously taking all the experimentalinformation (phase equilibria and thermodynamic quantities)into account.

3.2. Key Experiments Coupled with Modeling andApplications

In Mg-systems the experimental literature data were verysparse in both phase diagram data and thermodynamic prop-erties. Only a few ternary system were investigated in moredetail and calculated thermodynamically.[42,43,44] An attemptto apply these calculations for alloy development was pub-lished.[45]

Several experimental methods are being applied in ourgroup to produce a sufficient experimental basis for a multi-component Mg-database. For preparation of highly reactiveand evaporating alloys containing Mg, Li, and Ca the levita-tion melting technique was adapted. The main advantagesare crucible free melting, suppressing of evaporation withprotective gas overpressure, and visual control of energy sup-ply to avoid overheating. Other preparation methods includearc and electron beam melting for extremely high melting in-termetallic phases, see Table 1.

During thermal analysis (DTA) the alloys were protectedin a sealed tantalum crucible, Figure 3. The development oftechniques for producing these crucibles and sealing the lidsfree of micro-cracks were found to be essential for reproduc-ible DTA measurements of these highly reactive and easilyvaporizing alloys. For phase and crystal structure analysis X-ray diffractometry at room and high temperature, electronmicroscopy, and metallography are also used in a combinedapproach.

These techniques have been applied in our work on tern-ary and quaternary alloy systems for improvement of creepresistance,[46] thermal stability, spray forming (Mg-Mn-Sc-RE,Mg-Al-Ca-RE) and density reduction (Mg-Li-Al-Si, Mg-Al-Ca-Si).[47,48] Thermodynamic descriptions supported by ex-perimental work of several subsystems were published forMg-Sc,[49] Mn-Sc,[50] Al-Sc,[51] Al-Gd-Mg,[52] Gd-Li-Mg,[53]

Ca-Li-Mg,[54] Al-Li-Si,[55,56] and Al-Ca.[57] The thermodynamicand technical application of the ternary Al-Mg-Sc system wasalso discussed.[58,59]

In the following the new cerium containing quaternarysystem Mg-Mn-Sc-Ce together with Mg-Mn-Sc-Gd, Mg-Mn-Sc-Y, and Mg-Mn-Sc-Zr are shown as example for the selec-tion of new creep resistant alloys using computational ther-mochemistry. More details of alloy selection in the quaternarysystems Mg-Mn-(Sc, Gd, Y, Zr) will be published soon.[60]

4. First Generation of Creep Resistant TernaryMg-Mn-Sc-(Ce) Alloys

4.1. Ternary Mg-Mn-Sc Alloys

Investigations started with binary Mg-Sc alloys. Scandiumwas chosen for precipitation hardening because of its largesolubility in (Mg) and the retrograde solubility at lower tem-peratures. The binary phase diagram had to be re-investi-gated,[61] it is shown in Figure 4. However, binary MgSc pre-cipitates form very slowly during aging and improve themechanical properties only slightly because of their incoher-ent interface. Therefore, Mn was added as second alloyingelement. The precipitation of Mn2Sc was predicted by ther-modynamic calculations. Mn2Sc precipitations form coher-


Fig. 3. Cross section of a sealed tantalum crucible used for experiments with liquid Mgalloys.

Table 1. Methods for key experiments in Mg-database development


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ently and were found to be very efficient for improving creepresistance and hardness. New MgSc15Mn1 or MgSc6Mn1 al-loys show about 100 times better creep resistance than thebest commercial WE43 alloy at 350 �C and 30 MPa.[46]

In spite of the good properties of this first generation ofMg-Mn-Sc alloys, the high cost of Sc addition (6 wt.-% Sc ormore) initiated a search for a second generation by investigat-ing quaternary systems.

Additional alloying elements Ce, Gd, Y, and Zr were con-sidered for this purpose to achieve a sufficient quantity ofsuitable precipitates to improve mechanical properties usinga minimum of expensive alloy element addition. These ele-ment combinations Mg-Mn-(Sc, Ce, Gd, Y, Zr) form a varietyof quaternary systems and within those there is a hugeamount of possibilities to select alloy compositions. Thereforephase diagram and other calculations were performed toidentify promising candidates.

4.2. Alloy Selection in the Mg-Mn-Ce-Sc System

For a first alloy selection several vertical sections ofphase diagrams in various ranges of Mn and Ce were cal-culated. For practical reasons the Mn content was limited at1.5 wt.-%. First, alloys were cast with 3 to 5 wt.-% Ce andup to 15 wt.-% Sc. The reason for the addition of these highalloying elements was the lack of knowledge about theamount and the mechanical effect of the precipitations thatare possibly formed during heat treatment. The calculatedquaternary phase equilibria are presented in two-dimen-sional sections with constant Mn and Ce content. Figure 5shows a T±x section with constant Mn and Ce content(1.5 wt.-% Mn, 4.2 wt.-% Ce) for Sc contents varying from0±15 wt.-%. Primary crystallization of Mn2Sc phase can beobserved for a wide range of higher Sc contents. TheL + Mn2Sc phase field and several different solid phasesstable at lower temperatures can be seen.

A large variety of phase diagram sections can be calculatedusing the thermodynamic database. At this point the questionof how to identify promising alloy candidates from all thesecalculated diagrams becomes important. Which phase dia-gram features are related to which alloy processing steps?What is needed is a list of beneficial phase diagram features,derived from the relevant alloy processing steps. The mostimportant of these points are given in Table 2, they areapplied in the following discussion.

Figure 6 shows the phase amounts during equilibriumsolidification for an alloy with 1.5 wt.-% Mn, 4.2 wt.-% Ce, and6.3 wt.-% Sc, indicated by the dotted line in Figure 5. Smallamounts of primary Mn2Sc and significantly larger amounts ofCeMg12 are expected in this alloy. In fact, in the micrograph ofthis alloy after T5 heat treatment (aged after casting) fine-grained Mn2Sc as well as larger crystals of CeMg12 can be seenin Figure 7.[62] It will be shown later that the finely distributedMn2Sc particles have a very positive effect. One alloy withhigher Sc content was studied to promote the favored Mn2Scprecipitations. The calculated phase amounts in Figure 8 for analloy with 1.5 wt.-% Mn, 3.2 wt.-% Ce, and 9.0 wt.-% Sc show


0

200

400

600

800

1000

1200

1400

1600

Temperature[°C]

0 20 40 60 80 100

at. % Sc

[97Vos]

[69Bea]

DTA on Heating

DTA on Cooling

Single Phase

Two-Phase limits

Liquid

Mg Sc

MgSc

( Sc)

( Sc)

(Mg)

β

α

Fig. 4. The binary Mg-Sc phase diagram.[61]

5 10

wt.% Sc4.2 Ce1.5 Mn94.3 Mg0.0 Sc

4.2 Ce1.5 Mn79.3 Mg15.0 Sc

0

200

400

600

800

Temperature[°C]

(Mg) + Mn Sc +

CeMg + MgSc2

12

(Mg) + Mn Sc +

CeMg2

12

L + (Mg) + Mn Sc2

L + Mn Sc2

LL + Mn Sc

+ S c2

β

Fig. 5. Phase diagram section with 4.2 wt.-% Ce, 1.5 wt.-% Mn, and Sc content from0±15 wt.-%. Dotted line indicates the MgCe4.2Mn1.5Sc6.3 alloy.

Table 2. Beneficial phase equilibrium features and their relevance for (Mg)-alloy proces-seing


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no remarkable differences to the alloy with 6.3 wt.-% Sc in Fig-ure 6. The low-temperature formation of MgSc should be disre-garded because of the very slow kinetics and the incoherenceof this precipitate. In other words the higher Sc-content doesnot substantially improve the calculated amount of mechani-cally efficient precipitates. In fact, the mechanical properties ofthis alloy with 9 wt.-% Sc show no remarkable difference tothat with 6.3 wt.-% Sc.[62]

The amount of the favored Mn2Sc particles is mainlycontrolled by the Mn-content. Therefore the Sc-contentcould be decreased to economically and technically suitableamounts and possibly be supported by precipitates stem-ming from cheaper rare earth element additions. In a sec-ond generation of creep resistant alloys, promising quater-nary candidates were selected primarily by computationalthermochemistry.

5. Second Generation of Creep Resistant, LowScandium Alloys

5.1. Alloy Selection in the Mg-Mn-Sc-Gd System

Several vertical sections in the quaternary Mg-Mn-Sc-Gdsystem were studied in the ranges of 0±1.5 wt.-% Mn,0±10 wt.-% Sc, and 0±10 wt.-% Gd. The calculated quaternaryphase equilibria are presented in two-dimensional sectionswith constant Mn and Gd content. Figure 9 shows a T±x sec-tion with constant Mn and Gd content (1 wt.-% Mn, 5 wt.-%


0

300

600

900

MnSc

2MgSc

CeMg12

(Mg)

(Mg) liquid

Temperature[°C]

Phase amounts [mol]

0 0.2 0.4 0.6 8 1.0

Fig. 6. Calculated equilibrium phase amounts for the MgCe4.2Mn1.5Sc6.3 alloy.

Fig. 7. Micrograph of MgCe4.2Mn1.5Sc6.3 alloy (T5 heat treated).[62]

(Mg)

(Mg) liquid

0

300

600

900

Mn

Sc

2

MgS

cC

eM

g1

2

Tem

pera

ture

[°C

]

Phase amounts [mol]

0 0.2 0.4 0.6 0.8 1.0

Fig. 8. Calculated equilibrium phase amounts for the MgCe3.2Mn1.5Sc9.0 alloy.

0

200

400

600

800

0.2 0.4 0.6 0.8

Te

mp

era

ture

[°C

]

wt.% Sc

(Mg) +Mn Sc +

GdMg2

5

(Mg) +Mn Sc +

GdMg +

Mn Sc

2

5

23 6

(Mg) +Mn Gd

+GdMg

+Mn Sc

12

5

23 6

(Mg) +Mn Gd12

(Mg) + Mn Sc +

GdMg + MgSc2

5

(Mg) +Mn Sc2

L +Mn Sc2

(Mg)

L + (Mg)

L (liquid)

(Mg) +Mn Sc23 6

5.0 Gd1.0 Mn

94.0 Mg0.0 Sc

5.0 Gd1.0 Mn

93.0 Mg1.0 Sc

Fig. 9. Phase diagram section with 5 wt.-% Gd, 1 wt.-% Mn, and Sc content from0±1 wt.-%. Dotted lines indicate the MgGd5Mn1Sc0.3 and the MgGd5Mn1Sc0.8 alloy.


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mary solidification of (Mg), a large one-phase field of (Mg)and several different solid phases stable at lower tempera-tures can be seen.

For an alloy with 1 wt.-% Mn, 5 wt.-% Gd, and 0.8 wt.-%Sc (indicated by the right dotted line in Fig. 9) equilibriumphase amounts during solidification and heat treatment aregiven in Figure 10. At the liquefaction temperature of 650 �C,primary (Mg) is formed and consumes the melt totally downto the solidification point at 619 �C. At 590 �C the first precipi-tate Mn2Sc starts forming, which can be seen in the inset ofFigure 10. Large amounts of the second precipitate GdMg5

start forming at 425 �C. Similar positive features are observedfor an alloy with only 0.3 wt.-% Sc.

The equilibrium phase amounts during solidification andheat treatment for the alloy with 1 wt.-% Mn, 5 wt.-% Gd,and 0.3 wt.-% Sc (indicated by the left dotted line in Fig. 9)are given in Figure 11. The Mn23Sc6 phase is formed as firstprecipitate in the solid state. Both alloys fulfill all features il-lustrated in Table 2 and were classified as very promising forfurther large scale creep experiments.

In fact, first results of these alloys show a creep resistancesimilar to the previous ternary high-scandium Mg-Mn-Scalloys, i.e., about 100 times better than commercial WE43alloys at 350 �C and 30 MPa, see further.

5.2. Alloy Selection in the Mg-Mn-Sc-Y System

In the Mg-Mn-Sc-Y system several vertical sections in theranges of 0±1.5 wt.-% Mn, 0±10 wt.-% Sc, and 0±10 wt.-% Ywere studied. Figure 12 shows a T±x section of a calculatedquaternary phase diagram with constant Mn and Y content(1 wt.-% Mn, 5 wt.-% Y) for Sc contents varying from0±1 wt.-%. A large one-phase field of (Mg) and several differ-

ent solid phases stable at lower temperatures can be seen, e.g.,in the corresponding diagram for Gd. Differences compared tothe Gd system can be observed concerning the stable solidphases and the formation temperatures. In the whole range ofprimary crystallization of (Mg) the secondary phase is Mn12Y.The desired Mn2Sc precipitate forms only at temperaturesbelow 500 �C or even lower with decreasing Sc content.

For an alloy with 1 wt.-% Mn, 5 wt.-% Y, and 0.8 wt.-% Sc,indicated by the dotted line in Figure 12, phase amounts dur-ing equilibrium solidification are given in Figure 13. At theliquefaction temperature 644 �C, primary (Mg) is formed andconsumes the melt totally down to the solidification point at613 �C. The Mn12Y phase is only stable in the solid state in a


Te

mp

era

ture

[°C

]

Mn Sc

MgSc (Mg)

(Mg)

(Mg)

liquid

liquid

GdMg

Gd

Mg

5

0

0

200

200

400

400

600

600

800

800

0

0

0.2

0.02

0.4

0.04

0.6

0.06

0.8

0.08

1.0

0.1

5

2

Phase amount [mol]

Fig. 10. Calculated equilibrium phase amounts for the MgGd5Mn1Sc0.8 alloy.

Te

mp

era

ture

[°C

]

Phase amounts [mol]

Mn Sc2

(Mg)

(Mg)

(Mg)

liquid

liquid

GdMg5

0

0

200200

400

400

600

600

800

800

0

0

0.2

0.02

0.4

0.04

0.6

0.06

0.8

0.08

1.0

0.1

Mn Sc23 6

Gd

Mg

5

Fig. 11. Calculated equilibrium phase amounts for the MgGd5Mn1Sc0.3 alloy.

0.2 0.4 0.6 0.8

wt.% Sc5.0 Y1.0 Mn

94.0 Mg0.0 Sc

5.0 Y1.0 Mn

93.0 Mg1.0 Sc

0

200

400

600

800Tem

pera

ture

[°C

]L +Mn Sc2L + (Mg)

L (liquid)

(Mg)

(Mg) +

Mn Sc

+ Mn Y2

12

(Mg) +Mn Sc +

Mn Y +

Mg Y

2

12

24 5

(Mg) +Mn Y +

Mg Y12

24 5

(Mg) +Mn Sc

+Mg Y

2

24 5

(Mg) + Mn Sc +

Mg Y +MgSc2

24 5

(Mg) +Mn Y12

Fig. 12. Phase diagram section with 5 wt.-% Y, 1 wt.-% Mn, and Sc content from0±1 wt.-%. Dotted line indicates the MgY5Mn1Sc0.8 alloy.


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high temperature range from 605 to 322 �C and may not format all during fast cooling. The established Mn2Sc and the newMg24Y5, even in very large amount, could form by aging in afavorable temperature range. This makes MgMn1Y5Sc0.8 alsoa promising alloy.

In order to check if the amount of Mn2Sc could be raisedby increasing the manganese content to 1.5 wt.-%, we can ex-amine Figure 14 for the alloy MgMn1.5Y5Sc0.8. This alloy isdisqualified since Mn2Sc forms as the primary phase from theliquid, clearly seen in the inset of Figure 14. (Mg) forms onlysecondary, and even Mn12Y forms as a tertiary phase duringsolidification. Such a microstructure cannot be ªrepairedº byheat treatment. It would be difficult to envisage such a drasticchange caused by raising the Mn-content from 1.0 to1.5 wt.-% without the thermodynamic calculations.

5.3. Alloy Selection in the Mg-Mn-Y-Zr System

For complete substitution of expensive scandium, zirco-nium was checked as a possible alloying element. The impactof such a substitution can be seen in Figure 15 for scandium-free alloys with 1 wt.-% Mn, 4.5 wt.-% Y, and 0±1 wt.-% Zr.This phase diagram section shows a very steeply rising lique-faction line, 1000 �C are reached for less than 0.1 wt.-% Zr.Moreover, and actually the reason for the steep liquefactionline, a huge primary crystallization field, L + Mn2Zr, stretchesover the entire composition range in Figure 14. Only forextremely small Zr-additions, not discernible in Figure 15, isa primary (Mg) solidification expected. The reason for this de-structive phase diagram feature is the extremely high thermo-dynamic stability of Mn2Zr in comparison to the otherphases. Since yttrium does not play a significant role in thatpart, the only way to diminish the L+Mn2Zr primary fieldwould be a drastic reduction of the manganese content. Butthis would also drastically reduce the amount of beneficialMn-containing precipitates. As a result, the entire quaternaryMg-Mn-Y-Zr alloy system is disqualified.

5.4. Defining the Commercial Creep Resistant BenchmarkAlloy

In order to assess the results on the creep resistance mea-surements of the selected new alloys a commercial alloy hasto be chosen as a benchmark. The history of commerciallydeveloped creep resistant alloys is briefly summarized in Ta-ble 3. These alloys usually contain about 0.7 wt.-% Zr as agrain refiner. The influence of the rare earth elements (RE) onthe creep resistance generally decreases in the series Nd (Pr)> Ce > La.[63]


Te

mp

era

ture

[°C

]

MgSc

(Mg)

(Mg)

(Mg)

liquid

liquid

Mg

Y

0

0

200200

400

400

600

600

800

800

0

0

0.2

0.02

0.4

0.04

0.6

0.06

0.8

0.08

1.0

0.1

24

5

Mn Y

Mg Y

12

24 5

Mn Sc2

Phase amount [mol]

Fig. 13. Calculated equilibrium phase amounts for the MgY5Mn1Sc0.8 alloy.

Te

mp

era

ture

[°C

]

Mn Sc2

Mn Y

MgSc (Mg)

(Mg)

(Mg)

liquid

liquid

0

0

200200

400

400

600

600

800

800

0

0

0.2

0.02

0.4

0.04

0.6

0.06

0.8

0.08

1.0

0.1

Mg Y

Mg

Y24

5

12

24 5

Mn Sc2

Phase amount [mol]

Fig. 14. Calculated equilibrium phase amounts for the MgY5Mn1.5Sc0.8 alloy.This alloy contains only 0.5 wt.-% Mn more than that in Figure 13 but shows unfeasibleprimary solidification of Mn2Sc.

0.2 0.4 0.6 0.8

wt.% Zr4.5 Y1.0 Mn

94.5 Mg0.0 Zr

4.5 Y1.0 Mn

93.5 Mg1.0 Zr

Te

mp

era

ture

[°C

]

(Mg) + Mn Zr +

Mn Y + Mg Y2

12 24 5

(Mg) +Mn Zr +

( Zr) +Mg Y

2

24 5

a

(Mg) + Mn Zr

+ Mn Y2

12

(Mg) + Mn Zr2

L + (Mg) + Mn Zr2

L + Mn Zr2

0

400

200

600

800

1000

L (liquid)

Fig. 15. The phase diagram section with 5 wt.-% Y, 1 wt.-% Mn, and Zr content from0±1 wt.-% shows detrimental primary solidification of Mn2Zr in the entire composi-tion range These alloys cannot be completely molten under technical conditions.


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The current end point of commercial alloys is given by theWE series containing Y and rare earth elements. In fact, themaximum stress that can be tolerated at 200 �C for 100 h and0.2 % elongation is highest for the WE series as shown in Fig-ure 16. The alloy WE43 with T6 heat treatment was chosen asthe benchmark alloy. Die casting alloys such as AE42 orAS21, known to be more creep resistant than standard AZ orAM alloys, are not included in the comparison since theyrank even below the ZE alloys in Figure 16.

6. Alloy Preparation and Creep ResistanceMeasurements

6.1. First Generation of Mg-Mn-Sc-(Ce) Alloys

The most promising alloy compositions identified by thethermochemical calculations were prepared by squeeze cast-ing by our project partners within the Thrust Research ProjectSFB 390, ªMagnesium Technologyº, at the TU Clausthal.Yield strength and creep rates of these alloys were measuredin the as-cast condition and also after heat treatment.[62] It isclearly seen from Figure 17 that the secondary creep rate at350 �C of our first generation (high-scandium) alloys is betterby a factor of 100 compared to the WE43 benchmark alloy.

0

0 20 40 60 80

time [10 s]3

WE43

MgCe3Mn1Sc9

MgMn1Sc6

MgMn1 5Sc1

MgCe4Mn1Sc6

ε 1 ·10-8s-1

strain[%]

T=350°C, =30 MPaσ

0.2

0.4

0.6

0.8

1.0

ε 1·10-6s-1∼∼

∼∼

Fig. 17. Creep curves of first generation (high scandium) alloys at 350 �C after T5treatment[62] compared to the WE43 benchmark.

The first generation of alloys showed a strong annealingresponse due to the formation of the Mn2Sc precipitates. Theexistence of the Mn2Sc precipitates was confirmed by X-raydiffraction, SEM and TEM investigations and energy disper-sive X-ray microanalysis. A micrograph showing finely dis-persed Mn2Sc precipitates in a MgSc6Mn1 alloy after T5 treat-ment is shown in Figure 18.[62]

At 400 �C an increase of the secondary creep rate can beobserved for the Ce-containing alloys, Figure 19. This disad-vantage is partly compensated by higher creep ductilitycaused by significant grain boundary gliding in the Ce-con-taining alloys.[62] Comparing Ce-containing alloys with differ-ent Sc-content, no significant change in the creep rate can beobserved, Figure 20. The result is in line with the calculationsin section 4.2 and the minimized Sc-contents of the secondgeneration alloys.


Table 3. Commercial creep resistand Mg-alloys

WE54 T6 WE43 T6 QE22 T6 ZE41 T5 AZ81 T4

165 161

88

52

12

0

20

40

60

80

100

120

140

160

180

Yie

ldstr

ength

,R

[MP

a]

p0

,2,

100

,200

WE54 T6 WE43 T6 QE22 T6 ZE41 T5 AZ81 T4

Alloy

Load 100 h, 200°C

T4 solution heat treatedT5 agedT6 solution h.tr. + aged

Fig. 16. Creep properties of commercial magnesium alloys. Fig. 18. Micrograph of MgSc6Mn1 (T5) alloy.[62]


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6.2. Second Generation of Creep Resistant, Low ScandiumAlloys

The second generation (low scandium) alloys wereselected from the most promising candidates identified bythe thermodynamic calculations given in Section 5 from thecerium-free alloys. The creep curves for these low scan-dium quaternary alloys are shown in Figure 21.[62] Again,an almost 100 times lower creep rate is achieved at 350 �Cand 30 MPa compared to the best commercial alloy WE43.The best alloy MgGd5Mn1Sc0.3 with the smallest elonga-tion after 15´104 s contains only 0.3 wt.-% of expensive Sc. Itis exactly one of the few most promising alloys indicatedby the thermodynamic calculations. The calculated equilib-rium phase amounts during solidification and heat treat-ment for this alloy was shown in Figure 11. This is consid-ered to be a substantial progress compared to the firstgeneration of alloys with much higher scandium contentsof 6 or 15 wt.-% Sc.

7. ConclusionsFocused magnesium alloy development is possible using

the powerful tool of thermodynamic calculations. Alloy com-positions with promising possibilities of alloy microstructuredesign can be selected by means of thermodynamically calcu-lated phase diagrams, phase amount charts, and solidificationcurves. Most importantly, element combinations and compo-sitions with unwanted properties can be recognized beforestarting large-style experiments, thus reducing the experi-mental effort to a reasonable volume. The next step, the ex-perimental study of mechanical properties of identifiedpromising alloys, has shown excellent results both in the firstgeneration (ternary and high scandium) and second genera-tion (quaternary, low scandium) of new creep resistant alloys.Obviously, these experiments cannot be replaced by thermo-dynamic calculations. However, considering the huge num-ber of less promising alloy combinations that could have beenselected by trial and error from multicomponent systems, thefocused alloy development following this approach avoids awaste of time and effort.

Received: May 18, 2001

±[1] U. R. Kattner, G. Eriksson, I. Hahn, R. Schmid-Fetzer,

B. Sundman, V. Swamy, A. Kussmaul, P. J. Spencer,T. J. Anderson, T. G. Chart, A. Costa e Silva, B. Jansson,B. J. Lee, M. Schalin, Calphad 2000, 24, 55±94.

[2] J. �gren, F. H. Hayes, L. Höglund, U. R. Kattner, B. Le-gendre, R. Schmid-Fetzer, Z. Metallkde. 2001, 92, inpress.

[3] L. Kaufman, H. Bernstein, Computer Calculation of PhaseDiagrams with Special Reference to Refractory Materials,Academic Press, New York 1970.


0

0 5 10 15 20

time [10 s]4

WE43

MgCe3Mn1Sc9

MgMn1Sc6

MgMn1 5Sc1

MgCe4Mn1Sc6ε 3.3·10-6 s-1

ε

ε

1.9 ·10

2.4 ·10

-7

-6

s

s

-1

-1

strain[%]

T=400°C, =30 MPaσ

2

4

6

8

10

ε 1.6·10-5 s-1

ε 9.7·10-5s-1∼∼

∼∼

∼∼

∼∼

∼∼

Fig. 19. Creep curves of first generation (high scandium) alloys at 400 �C[62] comparedto the WE43 benchmark.

0

0 20 40 60 80

time [10 s]3

WE43

MgCe3Mn1Sc9

MgCe3Mn1Sc1

MgCe4Mn1Sc6

ε 1 ·10-8s-1

strain[%]

T=350°C, =30 MPaσ

0.2

0.4

0.6

0.8

1.0

ε 1·10 -6 s-1∼∼

∼∼

Fig. 20. Creep curves of Mg-Ce-Mn-Sc alloys at 350 �C after T5 treatment.[62]

0

5

10

15

0 5 10 15time [10 s]

4

WE43 T6

MgGd5Mn1 FSc0.3

MgGd10Mn1 FSc0.8

MgY4Mn1Sc1 F

ε 1·10-6 s-1

ε 3 to 6·10 -8 s-1

strain[%]

T=350°C, =30 MPaσ

∼∼

∼∼

Fig. 21. Creep curves[62] of second generation (low scandium) alloys, identified by theCalphad approach, compared to the WE43 benchmark. The alloy with only 0.3 wt.-%Sc shows best properties, see also Figure 11.


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WS [4] J. Hertz, J. Phase Equilibria 1992, 13, 450.

[5] U. R. Kattner, J. Metals 1997, 12, 14.[6] A. T. Dinsdale, Calphad 1991, 15, 317.[7] L. Kaufman, Calphad 1977, 1, 7.[8] L. Kaufman, Calphad 1978, 2, 55.[9] M. W. Chase, I. Ansara, A. Dinsdale, G. Eriksson,

G. Grimvall, L. Hoglund, H. Yokokawa, Calphad 1995,19, 437.

[10] J. �gren, B. Cheynet, M. T. Clavaguera-Mora, K. Hack,J. Hertz, F. Sommer, U. Kattner, Calphad 1995, 19, 449.

[11] A. Chang, C. Colinet, M. Hillert, Z Moser, J. M Sanchez,N. Saunders, R. E Watson, A. Kussmaul, Calphad 1995,19, 481.

[12] D. de Fontaine, S. G. Fries, G. Inden, P. Miodownik, R.Schmid-Fetzer, C. Shuang-Lin, Calphad 1995, 19, 499.

[13] B. Burton, T. G. Chart, H. L. Lukas, A. D. Pelton, H. Sei-fert, P. Spencer, Calphad 1995, 19, 537.

[14] F. Aldinger, A. F. Guillermet, V. S Iorich, L. Kaufman,W. A. Oates, H. Ohtani, M. Rand, M. Schalin, Calphad1995, 19, 555.

[15] Q. Chen, M. Hillert, B. Sundman, W. A. Oates,S. G. Fries, R. Schmid-Fetzer, J. Electron. Mater. 1998, 27,961.

[16] D. Mirkovic, J. Gröbner, R. Schmid-Fetzer, in MagnesiumAlloys and their Applications (Ed: K. U. Kainer), DGM,Wiley-VCH, Weinheim 2000, pp. 783±788.

[17] C. Gente, M. Oehring, R. Bormann, Phys. Rev. B 1993,48, 13 244.

[18] R. Bormann, in Polycrystalline Thin Films (Eds: K. Bar-mak, M. A. Parker, J. A. Floro, R. Sinclair, D. A. Smith),Mat. Res. Soc. Symp. Proc. 1994, 343, 169.

[19] R. Busch, F. Gaertner, C. Borchers, P. Haasen, R. Bor-mann, Acta Metall. Mater. 1995, 43, 3467.

[20] G. Eriksson, K. Hack, in The SGTE Casebook, (Ed: K. Hack),The Institute of Materials, London 1996, p. 200.

[21] A. Koukitu, N. Takahashi, T. Taki, H. Seki, Jpn. J. Appl.Phys. 1996, 35, L673.

[22] J. �gren, ISIJ International 1992, 32, 291.[23] J. �gren, in The SGTE Casebook (Ed: K. Hack), The Insti-

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REVIE

WS

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[64] More and updated information can be found at the cor-responding websites:ChemSage at http://gttserv.lth.rwth-aachen.de/gtt/,Fact at http://www.crct.polymtl.ca/fact/fact.htm, Fact-Sage at http://www.factsage.com/, MTDATA athttp://www.npl.co.uk/npl/cmmt/mtdata/, Pandat athttp://www.computherm.com/, Thermocalc at http://www.thermocalc.se/

[65] It should be noted that the binary interaction of Eq. 6could equivalently be written in the same style as Eqs. 7and 8:

E Gbin;� � Pcÿ1

i�1

Pcj>i

xi xjPnm�0

Am�1;�ij �xi � �ij�nÿv�xj � �ij�m (13)

where

�ij = (1±xi±xj)/2 (14)

The subregular behavior is given for the order n = 1.The conversion between the parameters Am�1;�

ij andLm,u

ij has been given by Hillert. [31] Since virtually allbinary assessments are published in the form of Eq. 6instead of Eq. 13 it is reasonable to keep this form in thegeneral Gibbs energy equation for a c-component solu-tion phase.


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Focused Development of Magnesium Alloys Using the Calphad Approach

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