Philosophical Magazine Letters2012, 1–11, iFirst

Design considerations for achieving simultaneously high-strengthand highly ductile magnesium alloys

A.C. Hanzia, A.S. Sologubenkob, P. Gundea, M. Schinhammera

and P.J. Uggowitzera*

aLaboratory of Metal Physics and Technology, Department of Materials, ETH Zurich,Zurich 8093, Switzerland; bLaboratory for Nanometallurgy, Department of Materials,

ETH Zurich, Zurich 8093, Switzerland

(Received 31 October 2011; final version received 7 January 2012)

A comprehensive scheme of phase configuration optimization in theMg–Zn–Ca(–Zr) system by thermodynamic simulations and microstruc-tural analyses is presented. A composition window of 0.2–0.4wt% Ca and5–6wt% Zn is defined as optimal for establishing a complex heterogeneousmicrostructure allowing for enhanced ductility and simultaneously highstrength of the material. Literature data analysis and our own resultsconfirm the enhanced performance of alloys from this compositionwindow.

Keywords: magnesium alloys; high strength; high ductility; thermodynamicsimulations

Owing to its high strength and simultaneously high ductility, Mg–Zn–Ca(–Zr)wrought alloys are considered as a highly potential structural lightweight materialsfor a number of industrial applications. The possible application field ranges fromthe automotive branch that benefits from fuel economy and the reduction of the CO2

emission arising from vehicle weight reduction, to the medical industry where alloydegradation properties and alloy design freedom are attractive for temporaryimplant solutions. For the former the absence of cost-intensive alloying elements is ofimportance; for the latter it is also advantageous that the alloy system contains nobio-incompatible alloying elements.

As is the case for other multi-component systems, the mechanical performance ofthe alloys is defined by their heterogeneous microstructure, which in turn depends onand can be controlled by thermo-mechanical treatment. It is known that incrystallographically anisotropic Mg-based alloys grain refinement is the mostefficient approach to yielding simultaneous enhancement of both, strength andductility [1–3]. This grain-refinement approach together with the micro-alloyingconcept has proved to be a very effective route for the establishment of a stablemicrostructural configuration. It leads to excellent mechanical characteristicsincluding the reduction of mechanical anisotropy [4–6], which has been

*Corresponding author. Email: peter.uggowitzer@mat.ethz.ch

ISSN 0950–0839 print/ISSN 1362–3036 online

! 2012 Taylor & Francishttp://dx.doi.org/10.1080/09500839.2012.657701http://www.tandfonline.com

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demonstrated also in a number of other works on Mg–Zn–Ca–X alloys [7–9]. Thepinning of matrix grain boundaries by intermetallic phase (IMP) particles ofcontrolled number density ensures the grain-boundary strengthening of the alloy[10], whereas an increase in alloy ductility may be achieved by activation of the non-basal slip modes induced by the compatibility stresses at the matrix grain boundaries[11,12]. Further hardening can be achieved by the development of a heterogeneousintragranular microstructure with small coherent or semicoherent particles alignedparallel or perpendicular to the matrix basal planes [4,6,13].

To assess the influence of alloying elements and their combination on themicrostructural state and to work out the rules for their optimal amounts inconnection with the desired phase constellation for high performance Mg–Zn–Ca(–Zr) alloys, we have performed thorough phase configuration analysis as afunction of Zn and Ca elemental content by thermodynamic simulation, andexamined the microstructure–property relationship of a number of Mg–Zn–Ca-based alloys reported in the literature [4–9,13–19]. The effect of Zr, which isfrequently added to Mg-based industrial alloys to induce grain refinement, on thealloy microstructural state has also been considered. As will be shown below, itsinfluence on the formation of the desired major volume fraction of IMP particles,and consequently on the phase configurations, is however moderate. Wedemonstrate that an appropriate choice of the alloying content for as-extrudedMg–Zn–Ca(–Zr) alloys (i) ensures the formation of ultra-fine grains correspondinglyinducing high strength and ductility of the material, (ii) prevents the formation ofcrack-initiating IMP particles, and (iii) enables considerable age hardening.

To begin, we consider the effect of Zn content on the phase state in Mg–Zn–Ca(–Zr) alloys (chemical compositions indicated in Table 1) at a given Ca content. Thealloys possess a range of very attractive mechanical properties in the extruded(extrusion ratio: !30 : 1; Textrusion! 300"C) and completely recrystallized states(Table 2) [4,5,16]. The fine-grained microstructures of the alloys after thermo-mechanical processing are stabilized by evenly distributed IMP particles that exhibitsignificant grain-boundary pinning. Interestingly, the yield stress in the alloyscontaining 5wt% Zn is considerably higher than that in the alloys with 3wt% Zn.This fact cannot be accounted for by considering only the difference in matrix grainsizes. According to [15,20], the Hall–Petch coefficient for fine-grained Mg alloys isabout 5MPamm1/2. Consequently, a reduction in grain size, for example from 6 mm(ZX30) to 4 mm (ZX50), would have resulted in an increase in yield stress ofapproximately 15MPa, and from 3 mm (ZKX30) to 2 mm (ZKX50) in an increase byabout 20MPa. The differences in tensile yield stress between the alloys with 3wt%

Table 1. Nominal chemical composition (in wt%) of Mg–Zn–Ca(–Zr) alloys.

Alloy Mg Zn Ca Zr Mn

ZX30 bal. 3 0.25 – 0.15ZKX30 bal. 3 0.25 0.5 0.15ZX50 bal. 5 0.25 – 0.15ZKX50 bal. 5 0.25 0.3 0.15

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and 5wt% Zn are, however, 60MPa and 70MPa, respectively. Therefore, weconclude that yet another hardening mechanism has to be considered. Moreover, thealloys with 5wt% Zn exhibit a significant aging response of about 45MPa uponaging at 160"C (Table 2, ZKX50 alloy in the recrystallization heat treatment (rht)compared to the rht# aging states), while no aging response is observed in the alloycontaining 3wt% Zn. Logically, the question arises, how high can the Zn content inthe system be for the material still to exhibit high strength and simultaneously highductility without featuring the drawbacks of possible formation of crack-initiatingIMP particles?

The thermodynamic calculations of the phase volume fractions in equilibriumshow that, at the extrusion temperature of approximately 300"C, the major IMPcontributing to grain growth restriction in the Mg–Zn–Ca–Zr system is theCa2Mg6Zn3 phase. All thermodynamic calculations were performed using thePandat software package with the database PanMg8 [21]. Figure 1 presents, as anexample, the result for the ZKX50 alloy. The situation is virtually the same for thealloys without Zr and lower Zn content (i.e., ZX50, ZKX30, ZX30) [4,5], which isalso confirmed by observations of Nie et al. on Mg–Ca(–Zn) alloys [18] and theanalysis by Bamberger et al. for the Mg–Ca–Zn system [14]. The considerablethermal stability of the Ca2Mg6Zn3 particles allows heat treatments close to thesolidus temperature without inducing grain growth. In alloys with a Ca-content of0.25wt%, their phase fraction is approximately 0.8% (Figure 1). According to the3D Monte Carlo simulations performed by Kad and Hazzledine for microstructuralfeatures like the here-discussed [22], stagnation of grain growth is reached atD/d$ (2/f )1/3, where D is the grain size, d the particle size, and f their volumefraction. Our TEM studies [4,16] show that the Ca2Mg6Zn3 particle size d ofextruded alloys is typically 200–500 nm. According to [22] the corresponding matrixgrain size would then range from 1 to 3 mm, which is in good agreement with ourexperimental results.

The second largest fraction IMP in these alloys is a binary MgZn phase thatprecipitates in the solid state within the matrix grains at temperatures below 250"C

Table 2. Engineering mechanical properties of extruded ZX30, ZKX30, ZX50 and ZKX50 atroom temperature (relative scattering of the data 510%). Data from [4,16]; unpublishedresults for ZX30 sht# aging.

Alloy State!0.2C(MPa) Ratio !T/!C

!0.2T(MPa)

UTS(MPa)

"u(%)

"f(%)

Grainsize (mm)

ZX30 a-e 125 1.2 150 255 20 27 6ZX30 sht# aging n.m. n.m. 145 255 22 28 6ZKX30 rht 165 1.12 190 285 17 26 3ZX50 a-e 190 1.11 210 295 18 26 4ZKX50 rht 220 1.18 260 320 13.5 24 2ZKX50 rht# aging 240 1.27 305 330 10 23 2

Notes: !0.2C,T: yield stress in compression/tension; UTS: ultimate tensile strength; "u: uniformelongation; "f: elongation to fracture.a-e: as-extruded; sht: solution heat treatment at 380"C for 4 h; rht: recrystallization heattreatment at 360"C for 4 h; aging: at 160"C for 24 h.

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(indicated by an arrow in Figure 1). Its significant phase fraction is assumed toinduce a high precipitation tendency and a correspondingly considerable agingresponse, which has in fact been observed as an increase in yield stress of the ZKX50alloy upon aging treatment (Table 2). The Mg–Zn phase particles were found insmall amounts even in as-extruded ZKX50 (Figure 2a and b), forming mostprobably during cooling from the extrusion temperature. In the bright-field TEMmicrograph in Figure 2a, the precipitates are marked with arrows. They are verysmall (%5 nm) and aligned parallel and perpendicular to the basal plane of the matrix(see corresponding electron diffraction pattern acquired from the same area). Theparticles do not cause any additional features in the selected-area diffraction (SAD)pattern (the [2!1!10] zone axis) (Figure 2a, inset). The atomic-number-sensitive high-angle annular dark-field mode of scanning TEM (HAADF STEM) image acquiredin the [01!10] zone axis reveals the chemical contrast of the particles (Figure 2b).According to our EDX-STEM analysis, the matrix contains about 2–3 at.% Zn. Theprecipitates are composed almost exclusively of Mg and Zn and evaluation of anEDX spectrum acquired from a particle yields a value within about 5–6 at.% of Zn.Since the size of a particle is much smaller than the thickness of a TEM specimen, thematrix necessarily contributes to the EDX spectrum and the evaluation result cannotbe considered as the particle composition. In very few cases, up to 1 at.% Zr wasdetected in some particles of the as-extruded ZKX50. High-resolution TEM

Figure 1. Equilibrium phase fractions in the ZKX50 alloy. The extrusion temperature of300"C (downward-facing arrow at 300"C) indicates the major IMP present during thermo-mechanical processing. The framed inset on the right illustrates the roughly sphericalCa2Mg6Zn3 particles located at the grain boundary [6]. The downward-facing arrow at 160"Cmarks the phase fraction of binary MgZn (volume fraction of about 1.6% at 160"C); theframed inset on the left shows the intragranular rod-like precipitates that were formed uponT6 treatment [4].

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(HRTEM) with fast Fourier transform (FFT) of the areas of interest was used toreveal the particle–matrix crystallographic relationship. Figure 2c shows an HRTEMimage of a grain area containing a few precipitates (arrowed). The FFT of the framedarea in Figure 2c is presented in Figure 2d and shows the presence of additional

Figure 2. TEM micrographs of the ZKX50 alloy; (a)–(d): as-extruded state; (e), (f): aged at160"C for 24 h. (a) Bright-field TEM micrograph, acquired in the [2!1!10] zone axis. The arrowspoint to small regular-shaped precipitates within the matrix. The corresponding electrondiffraction pattern (shown in the inset) demonstrates that the Mg–Zn particles are alignedparallel or perpendicular to the (0002) plane. (b) HAADF STEM micrograph acquiredfrom another grain oriented in the [01!10] zone axis on the same specimen. The brightspecks in the image are Zn-enriched precipitates located predominantly at defects. Thecorrespondent diffraction pattern is shown in the inset. (c) High-resolution image acquiredin the [2!1!10] zone axis. The arrows point to areas of the Mg–Zn precipitates in the matrix.Owing to the coherent nature of the precipitates and the fact that they are located at adislocation, the image of the particles is partially distorted by lattice strain. (d) FFT acquiredfrom the particles area clearly confirms their presence by demonstrating additional reflections.(e) HAADF STEM image of a specimen aged at 160"C for 24 h; Zn-enriched particles appearin very large number and are very bright against the Mg-matrix background (f) EDX-STEMspectrum acquired from one of the areas marked with rectangles in (e).

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reflections (arrowed) lying equidistant between the fundamental reflections of thematrix and as such revealing the coherent nature of the precipitates. For example, thedistance between the incident spot and the [01!12] reflection is divided into three equalparts by two additional reflections.

In the aged ZKX50 specimens (160"C for 24 h), the number density of theparticles was found to be drastically increasing (Figure 2e). In the HAADF STEMimage, Zn-enriched particles appear in very large number and are very bright againstthe Mg-matrix background. The EDX-STEM analyses (Figure 2f shows a spectrumacquired from one of the areas marked in Figure 2e) give values of about 11–12 at.%Zn in the particles. In aged ZKX50, no Zr was found in the particles. As it wasmentioned earlier, keeping in mind the limited accuracy of EDX-STEM analysis ofsmall precipitates, we do not discuss the particle composition in either of the alloystates. Nevertheless, since the EDX spectra acquisition conditions were kept identicalfor TEM-analyses of specimens of both heat-treatment states, a definite increase ofZn content in the particles upon aging is certainly observed. Detailed structuralanalysis of the precipitates in the ZKX50 alloy is a subject of current studies and willbe reported elsewhere.

Literature analysis on the phase occurrence in wrought and aged Mg–Zn(–Ca)(–Zr) alloys does not allow for a definite identification of the Mg–Zn phase. TheMg–Zn phase particles of prominently similar morphology were found in a numberof Mg–Zn(–Ca)(–Zr) alloys [4,8,13,17]. It is still debatable, though, what structurethe Mg–Zn phase actually adopts [8,17,23–25]. Gao et al. presented convincingevidences of small precipitates of monoclinic Mg4Zn7 and hexagonal MgZn2 phasesforming in a binary Mg–8Zn alloy (wt%) upon aging at 200"C [17]. Mendis et al.reported the occurrence of intragranular metastable rod-like "01 (MgZn2) and plate-like "02 (MgZn2) precipitates in age-hardened Zr-containing Mg–6.1Zn–0.4Ag–0.2Ca(–0.6Zr) (wt%) alloys [8]. According to Mendis et al. the number density ofrod-like MgZn2 precipitates was markedly higher in the Zr-containing as-extrudedand also subsequently aged alloys [8]. Nevertheless, the complex studies of Oh-ishi

Figure 2. Continued.

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et al. by three-dimensional atomic probe (3DAP) and SAD TEM analyses on thesame as in [8] alloy system revealed the occurrence of the rod-like Mg–Znintragranular particles (containing equal amounts of Mg and Zn) and having thesame crystal structure as MgZn2 (hP12) [13]. The particles in the as-extruded alloycontained Zr in contrast to the T6-treated alloy [13]. The presence of Zr in theprecipitates of an as-extruded Zr-containing alloy and the large increase of theprecipitate number density upon aging accompanied with a simultaneous reductionof Zr content in the particles allowed Oh-ishi et al. [13] to conclude that Zrcontributes to the precipitate formation partially substituting Zn in Mg–Zn phaseparticles. Accordingly, the authors suggested that these particles dissolve uponsolution treatment and MgZn2 phase precipitates during aging at 160"C [13]. TheMg–(Zn, Zr) phase was also reported in a ZK60 alloy by Watanabe et al. [19]. Such avariety of experimental structural and composition data on the Mg–Zn phasesuggests a rather shallow Gibbs free-energy profile for the Mg–Zn formation path,yielding a number of possible but energetically similar metastable polymorphs ofMg–Zn in Mg–Zn–Ca(–Zr) systems.

In our simulations, we rely on the thermodynamically stable MgZn phase.Taking into account the abovementioned involvement of Zr on the formation andevolution of Mg–Zn particles, we consider the effect of Zr on the establishment of astable phase configuration in the Mg–Zn–Ca(–Zr)-based alloys as plausible, butdefinitely not decisive for their evolution. However, an important influence of Zr onthe microstructure evolution is known to take place predominantly duringsolidification owing to its large grain growth restriction coefficient [26]. The smallvolume fraction Zr-containing IMPs appearing in equilibrium are found as tinyparticles at sites of high dislocation densities (i.e., in non-recrystallized areas), butnot at the grain boundaries of recrystallized grains [16].

Since Zn and Ca are necessary elements participating in the formation of bothtypes of IMPs, Ca2Mg6Zn3 and MgZn, one has to consider the balance of elementcontents within two limits. On one side, the Zn and Ca amounts have to be sufficientfor the necessary volume fractions of the Ca2Mg6Zn3 and MgZn phases in the alloyswithout inducing the formation of undesired, crack-initiating IMP particles. On theother side, the reduction in the Ca and Zn amounts must not exceed the lowest limit soas not to diminish their beneficial grain growth restricting influence duringsolidification (Q-factor, [4,5,26]) and hot extrusion (grain-boundary pinning particles[5,10]). Figure 3a shows the equilibrium phase fractions in an Mg–5Zn (wt%) alloy at160"C as a function of the Ca content. The Ca2Mg6Zn3 phase fraction increasesproportionally to the Ca content. Thereby, on account of a very small solubility of Cain the Mg-matrix, particularly at the aging temperature of 160"C, Ca is entirelyconsumed for the formation of Ca2Mg6Zn3, as indicated by the element content insolid solution depicted in Figure 3a. Moreover, it is important to consider that, uponincreasing the Ca content, as manifested by an increase in Ca2Mg6Zn3 volumefraction, the fraction of the MgZn phase decreases, and consequently its precipitationtendency also. An increase of Ca from 0.25wt% to 0.6wt% would reduce the MgZnphase fraction by 50%. Owing to the preferential formation of the Ca2Mg6Zn3 phaseover the MgZn phase, too high Ca content (40.6wt%) would therefore negativelyinfluence the aging response. Since the Zn content must be kept below 5–6wt%, asdiscussed in a later section, the consumption of Zn by the formation of Ca2Mg6Zn3

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and thus the reduced precipitation tendency of the hardening MgZn phase cannot becompensated by an increase in the Zn content. Moreover, too high a fraction ofCa2Mg6Zn3 particles may increase the risk of crack initiation and thus reduce theductility. On the other hand, Ca content below 0.2wt% would reduce the volumecontent of Ca2Mg6Zn3 below 0.5% (Figure 3a) and consequently reduce the grain-boundary pinning effect in an undesired manner. Referring to these considerations,the optimal lower limit of the Ca content is assumed to be in the range of 0.2wt% andthe upper limit in the range of 0.4–0.6wt%.

With respect to the influence of the Zn addition and its optimal content, Figure 3bindicates an approximate lower Zn limit required for the formation of MgZnprecipitates of approximately 3wt% at 160"C. Because the Ca2Mg6Zn3 phasepredominantly consumes Zn, the lower Zn limit increases upon increasing the Cacontent. As can be seen in Figure 3c, Ca exhibits nearly no solid solubility while themaximum solubility of Zn in matrix is approximately 2.2wt% at 160"C. This is alsothe amount of Zn that contributes to solid-solution strengthening. However, choosinga Zn content of only 3wt% will not result in any precipitation-hardening response.

Figure 3. (a) Dependence of phase fraction and element content in solid solution in theMg-matrix on Ca content at 160"C in Mg–5Zn–xCa. (b) Dependence of phase fraction on Zncontent at 160"C at 0.25wt% Ca and 0.5wt% Ca. (c) Dependence of element content in theMg matrix at 160"C on the Zn content at 0.25wt% Ca and 0.5wt% Ca. (d) Dependence ofelement content of Ca and Zn in solid solution at 350"C (compared to 160"C) on the Zncontent in the alloy at 0.25wt% Ca.

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Thus, in order to benefit from both, solid solution and age hardening, a lower Zn limitof at least 4wt% seems reasonable.

Besides grain-boundary strengthening and precipitation hardening (at sufficientlyhigh Zn contents) solid-solution hardening contributes to alloy strengthening. Sincethe formation of the MgZn hardening phase consumes Zn, an increase of its volumefraction results in a decrease of the solid-solution hardening effect of Zn. Still, anoverall strength benefit will remain. For the ZKX50 alloy, the amount of Zn in solidsolution at 350"C (just below the solidus temperature) is approximately 4wt%(Figure 3d). After an aging treatment at 160"C, 2wt% Zn would still remain in solidsolution. In other words, at least 2wt% Zn does not contribute to solid-solutionhardening, but to precipitation hardening. This is advantageous because a net yieldstress benefit of approximately 45MPa results, as demonstrated by the experimentaldata presented in Table 2.

For the evaluation of the upper limit of the Zn content, one has to consider thephase diagram of an Mg–Zn alloy containing 0.25wt% Ca (Figure 4). For 0.25wt%Ca, the maximum Zn content in solid solution is approximately 6.4wt% (note: it is6.1wt% without Ca and 6.9wt% with 0.5wt% Ca). Exceeding the maximum Znsolubility is considered counterproductive, because it would result in the formationof eutectic phases that cannot be dissolved anymore and may dramatically decreasethe ductility. Moreover, with increasing Zn content, the alloy’s solidus temperaturedecreases, which in turn narrows the window for solution heat treatments andconsequently would increase the treatment duration. Considering the variouspositive impacts of Zn on the mechanical performance of the alloy, we suggest the Zncontent to lie optimally between 5.0wt% and 6.0wt%.

Figure 4. Isopleth for 0.25wt% Ca in the Mg–Zn–Ca system.

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In conclusion, systematic thermodynamic simulations and detailed studies of themicrostructure–property relationships in the Mg–Zn–Ca(–Zr) system, performed ona number of our own and literature data, confirmed the necessity of both mainalloying elements, Ca and Zn, for the establishment of a complex heterogeneousmicrostructure consisting of inter- and intra-granular IMP particles of controlleddensity and morphology. Since both alloying elements participate in main IMPs andZn also contributes to solid-solution strengthening, a minimal Ca and Zn amountand their balance was worked out to ensure solid-solution and precipitationhardening as well as grain-boundary strengthening effects. Furthermore, it wasfound that a loss of ductility sets the upper limit on the element amounts. Tosummarize, it was found that the Mg–Zn–Ca(–Zr) system possesses a compositionwindow (0.2–0.4wt% Ca, 5–6wt% Zn) in which the most optimally performinghigh-strength and simultaneously highly ductile alloys can be obtained. Typically,so-designed alloys in the as-extruded state exhibit a !Y of 300–350MPa combinedwith "f of 20–25%, and a tension–compression yield stress asymmetry of 51.3.

Acknowledgments

This work was supported by the Austrian Institute of Technology within the framework of the‘‘Biocompatible Materials and Applications’’ project, and the Staub/Kaiser Foundation(Switzerland).

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