Development of the as-cast microstructure in magnesium±aluminium.pdf

Development of the as-cast microstructure in magnesiumaluminiumalloys

Arne K. Dahle *, Young C. Lee, Mark D. Nave, Paul L. Schaer, David H. StJohn

?tul=0>CRC for Cast Metals Manufacturing (CAST), Department of Mining, Minerals and Materials Engineering, The University of Queensland,

Brisbane Qld 4072, Australia

Abstract

This paper presents an overview of several projects undertaken at CAST to increase our understanding of the solidication

characteristics of MgAl alloys. With the increased use of magnesium alloys, and with casting dominating as a production route,

there is a need for a more comprehensive understanding of the mechanisms of solidication and defect formation to allow further

optimisation of alloys and casting processes. The paper starts with considering the formation of the primary magnesium dendrites

and the means for grain renement of magnesiumaluminium alloys. The MgAl system is then shown to display a range of eutectic

morphologies for increasing aluminium content, ranging from a divorced structure, through several intermediate structures, to a

fully lamellar structure at the eutectic composition. The eutectic also inuences discontinuous precipitation which occurs in the

aluminium-rich regions of the magnesium phase. The paper concludes with a section on porosity formation as a function of alu-

minium content and an outline of the mechanism responsible for the formation of banded defects in magnesium alloys, particularly

in products made in pressure assisted casting processes. 2001 Elsevier Science Ltd. All rights reserved.

Keywords: Magnesium; Solidication; Casting defects; Eutectic growth; Grain renement; Porosity

1. Introduction

Most commercial magnesium alloys are based on themagnesiumaluminium system and casting is currentlythe most commonly used production process for mag-nesium components [1]. Among the alloys used, AZ91,AM60 and, to a lesser extent, AM50 dominate. The A inthe alloy designation indicates that aluminium is themain alloying element and the rst numeral is the ap-proximate concentration of aluminium in wt%. AZ91therefore contains 9 wt% aluminium, and the Z indicatesit also contains about 1 wt% zinc (actually 0.7 wt% Zn).The range of aluminium contents for the commercialalloys is 39 wt% Al, from AZ31, a wrought alloycomposition cast as billet, to AZ91.

The MgAl alloys are relatively cheap compared withother magnesium alloys available. They are readily ca-stable, particularly by high-pressure die casting, andexhibit good mechanical properties [1]. An importantfeature of these alloys is that they can be cast into longand thin sections by high-pressure die-casting. Typical

applications that utilise this feature include instrumentpanels, steering wheels and seat frames [1]. Althoughmagnesium alloys containing aluminium generally pos-sess good mechanical properties, ternary alloys withzinc, manganese, silicon and rare-earth elements in ad-dition to aluminium are used to obtain improved me-chanical properties. Zinc is added to improve the roomtemperature strength and uidity while silicon is addedto improve the creep strength of the alloys by formingMg2Si particles on the grain boundaries [2]. Addition ofmanganese is required to control the corrosion beha-viour, and magnesium alloys with aluminium andmanganese (AM60, AM50) are commonly used forcomponents where good ductility and impact strengthare required.

Fig. 1 shows the magnesiumaluminium equilibriumphase diagram. The maximum solid solubility at theeutectic temperature is about 13 wt% Al, and a eutecticbetween Mg and the intermetallic Mg17Al12 appears atabout 33 wt% Al. All the aluminium contents used in thecommericial alloys are below the maximum solid solu-bility limit and the alloys therefore solidify with aprimary a-magnesium phase. The equilibrium micro-structure for all the alloys is 100% a-magnesium, butnon-equilibrium, metastable, eutectic normally forms

Journal of Light Metals 1 (2001) 6172

www.elsevier.com/locate/ligandmet

* Corresponding author.

E-mail address: [email protected] (A.K. Dahle).

1471-5317/01/$ - see front matter 2001 Elsevier Science Ltd. All rights reserved.PII: S 1 4 7 1 - 5 3 1 7 ( 0 0 ) 0 0 0 0 7 - 9

during solidication and is present in the as-castmicrostructure in MgAl alloys down to about 2 wt%Al.

Since solidication of MgAl alloys results in astructure consisting of primary dendrites and eutectic, itis tempting to make comparisons between these alloysand the other major commercial light casting alloys, theAlSi alloys. However, there are some very signicantdierences. First, the eutectic is metastable in the MgAlalloys, as discussed above. Heat treatment can thereforeresult in a complete dissolution of the Mg17Al12 inter-metallic phase. In contrast, the Si phase in AlSi alloys isreasonably stable during heat treatment and thereforeremains after heat treatment, although it spheroidisesresulting in some improvement in elongation. Second,the eutectic microstructure in AlSi alloys is an irregulareutectic with Si as the faceted phase and the coarse Simorphology is often modied to reduce the size andrene the eutectic to improve the mechanical properties.In MgAl alloys with more than about 20 wt% Al,the eutectic is a regular eutectic where both phases arenon-faceted [4,5]. However, the volume fraction ofeutectic decreases as the Al-content is decreased, andthe eutectic morphology gradually transforms to a di-vorced eutectic [4,5]. Another dierence is in the volumefraction of eutectic, which is much smaller in the MgAlalloys compared to that in AlSi foundry alloys.

MgAl alloys display a wide freezing range, Fig. 1,and these alloys are therefore susceptible to a range ofcasting defects including segregation, porosity and hottearing. A distinctive kind of defect often forms when,and because, the alloys are used to cast long, thin sec-tions by high-pressure die casting [68]. This defect hasthe form of bands of segregation, porosity and tears thatusually run parallel to the surface of the casting. Al-though it has been shown that these defects may form in

AlSi alloys in pressure assisted casting processes [8],these defects are much more common in MgAl alloysand they are now regarded as a major impediment toobtaining good mechanical performance from castcomponents [9].

Magnesium components are most commonly pro-duced by high-pressure die casting and, sand and low-pressure die casting are only used to a small extent.Sand-casting is used to some extent for the productionof magnesium alloys not containing aluminium foraeronautical components [1]. There are, however, somecomponents, such as wheels, where a low-pressure orgravity die casting process may be more suitable. Aproblem is that no reliable grain rener addition existsfor MgAl alloys and, thus, the benets of a ne grainsize on the mechanical properties cannot be reliablyachieved [1].

Also, alloy optimisation to minimise porosity has notbeen undertaken to the degree that it has been for AlSialloys. A better understanding of how defects formduring solidication is required to be able to obtain thebest performance from cast MgAl alloys.

This paper presents an overview of a range of projectsthat has been undertaken to improve our understandingof the solidication characteristics of MgAl alloys. Itfollows the solidication process of the alloys, beginningwith the nucleation, grain renement and growth of theprimary a-magnesium phase. Next, the formation of theeutectic is considered showing a range of morphologiesof the eutectic Mg17Al12 phase. The last section consid-ers defect formation in MgAl alloys, rst focussing onthe porosity characteristics as a function of aluminiumcontent and, nally, the mechanism responsible for theformation of shear defects and porosity in pressure as-sisted casting processes commonly used for productionof MgAl components.

2. Nucleation and growth of -magnesium equiaxed den-

drites

The solidication sequence of MgAl alloys startswith nucleation of primary magnesium (a-Mg) in thetemperature range 650600C, ranging from the meltingpoint of pure magnesium to the liquidus temperature ofMg 9 wt% Al, covering the aluminium contents used inmost commercial alloys. Later solidication reactionsinvolve the formation of eutectic phases, with the MgMg17Al12 eutectic reaction occurring at 437C. A typicalmicrostructure of MgAl alloys is shown in Fig. 2,where well-developed primary a-Mg dendrites withsecondary arms (A) showing sixfold symmetry areclearly visible. The eutectic of Mg17Al12 (B) and a-Mgsolid solution (C) is located in the interdendritic regions.According to the MgAl equilibrium phase diagram,Fig. 1, the eutectic phase (Mg17Al12) is expected to ap-Fig. 1. MgAl equilibrium phase diagram, adapted from [3].

62 A.K. Dahle et al. / Journal of Light Metals 1 (2001) 6172

pear when the aluminium content reaches around 13wt%. However, the eutectic phase appears in alloyscontaining as little as 2 wt% Al for non-equilibriumcooling conditions normally encountered in castings[10].

Fig. 3 shows the microstructural changes with in-creasing aluminium content. A small addition of alu-minium to pure magnesium leads to a morphological

change of the primary phase from a cellular to a den-dritic structure. Rosette-like globular equiaxed grainsform with aluminium-rich solid solution between thedendrite arms. As the aluminium content is increasedfurther to 5 wt%, dendrites with pools of eutectic phasebetween the dendrite arms start to develop and, whenthe aluminium content is further increased, a fully de-veloped dendritic structure with sharp tips is observed.

The addition of small amounts of alloying elementssuch as zinc, manganese, silicon and rare-earths to MgAl alloys has little eect on nucleation of the primaryphase since these elements are mostly segregated to formsecondary phases well after the primary phase has nu-cleated [11].

Grain renement is among the important practicesused to improve the properties of castings. It is an es-sential and fundamental approach since grain size sig-nicantly inuences the mechanical properties of thecastings, and the grain size is usually determined at anearly stage of solidication by nucleation of the den-drites.

Several methods of grain renement have been de-veloped. One of the rst methods was grain renementby a simple thermal treatment prior to casting, the so-called `superheating treatment'. This method involvesrapid cooling of the melt to the desired casting tem-perature after a short holding time at an elevated tem-perature, generally between 150C and 260C above theequilibrium liquidus temperature of the alloy [12]. De-spite the successful grain renement achieved by thesuperheating method, alternative techniques weresought due to several practical problems, mainly relatedto the higher operating temperatures involved.

Successful grain renement has been reported by theaddition of ferric chloride (Elnal process) in magne-

Fig. 2. Micrograph of fully developed dendrites in a Mg15 wt% Al

alloy permanent mould casting [3]. The microstructure of the alloy is

similar to 9 wt% aluminium alloy, but the dendritic structure is more

clearly visible than normally observed in a 9 wt% aluminium alloy. The

magnesium dendrites have a characteristic sixfold symmetric shape

(A). The white phase between the dendrites is secondary eutectic phase

Mg17Al12 (B) and the dark regions between the dendrites are alumi-

num-rich solid solution (C). Some of the enriched solid solution forms

within the Mg17Al12 phase (partially divorced eutectic). A more de-

tailed explanation of eutectic phase is presented later in this article.

Fig. 3. Micrographs of magnesiumaluminum alloys with increasing aluminum content. The transition from a globular dendritic structure to a fully

developed dendritic structure with increasing aluminium content is readily noticeable.

A.K. Dahle et al. / Journal of Light Metals 1 (2001) 6172 63

sium alloys containing aluminium and manganese. Theamount of grain renement achieved by the Elnalprocess is somewhat similar to that achieved by super-heating [12,13]. However, due to the detrimental eecton corrosion resistance from the addition of Fe, theElnal process has not attracted industrial attention.

The addition of carbon to the melt (carbon inocula-tion) oers more practical advantages accompanied bylower operating temperatures and less fading. Variouscarbon-containing agents such as organic materials(C2Cl6; CCl4) [12,14], SiC particles [15], or granulargraphite [16] have been reported produce successfulgrain renement in magnesium alloys containing alu-minium.

A number of hypotheses have been proposed to ex-plain the mechanisms by which carbon inoculationmethods cause grain renement. However, none of themprovides an adequate explanation of the mechanismsdue to a lack of understanding of the fundamental fac-tors involved, such as the identication of the activenucleants and/or solute elements with strong segregatingtendency. The formation of Fe and/or Mn containingcompounds is considered to cause grain renement inthe superheating and Elnal treatments, while the for-mation of Al4C3 particles, regarded to be very potentnucleation sites for magnesium, is believed to producerenement in the carbon inoculation methods.

Recent work [1719] on the grain renement of alu-minium alloys has shown that grain renement can befacilitated by two mechanisms. The rst is the formationof crystals in the thermally undercooled region near thewalls of the mould early in the casting process. Thesecrystals are then carried into the bulk of the melt byconvection currents. The other mechanism is a result ofconstitutional undercooling generated by the growth ofa grain adjacent to a nucleant particle suspended in themelt [1719]. In both cases there are two factors whichcan enhance the number of successful nucleation events.The rst is the solute elements present in the melt andthe other is the number and potency of the nucleantparticles. Fig. 4 shows the relative eect of these two

factors where it can be seen that when potent nucleantparticles are present, addition of solute causes the grainsize to decrease rapidly at rst and then more slowly athigher solute contents. As shown below, dierent soluteelements cause this decrease to occur at dierent ratesdepending on their segregating power. Fig. 4 also showsthat as the potency of the nucleant particles decreases,not only is the grain size obtained larger, but the eect ofsolute additions on grain size is decreased.

Through a simplication, the eect of the solute canbe dened by the alloy's growth restriction factor whenthe potency of the nucleant particles is very high [20].The growth restriction factor (GRF) is dened byP

i miC0;iki 1 where m is the slope of the liquidusline, C0 the initial composition, and ki is the equilibriumpartition coecient for element i [21]. A large GRF in-dicates that the growing crystal generates constitutionalundercooling quickly and the liquid around the adjacentnucleants is therefore more quickly undercooled su-ciently to allow a stable nucleus to form on the nucleantparticle compared with an alloy having a small GRF.The eect of nucleant potency is that a higher potency(i.e., a lower amount of undercooling required for nu-cleation) will result in the next nucleation event occur-ring relatively sooner. Therefore, a high-potencynucleant combined with a melt composition with a largegrowth restriction factor is likely to result in a smallgrain size. This is the principle behind the AlTiB grainreners for aluminium alloys where the TiB2 particlesare the potent nucleant particles and the excess titaniumensures that the melt composition has a high-growthrestriction factor. Some alloying elements, such as alu-minium, calcium, silicon and zirconium, have been re-ported to produce signicant grain renement in puremagnesium due to their strong segregating power [22].The closest to satisfying the conditions of the potent AlTiB master alloys in magnesium alloys is the use ofzirconium master alloy to rene alloys that do notcontain Al. In this case a very ne grain size can beobtained. This is because zirconium is the most stronglysegregating element in magnesium [22,23]. The exactnature of the particles present in a MgZr melt is un-known, but they appear to be very potent. It is worthnoting that grain renement by zirconium addition isnot eective in magnesium alloys that contain alumini-um due to an undesirable interaction between alumini-um and zirconium.

The above-described mechanism of nucleation of theprimary phase has the implication that if all alloycompositions are converted to their correspondinggrowth restriction factors, then the plots of grain sizeversus GRF should be the same for all alloys when thenumber and potency of nucleant particles are equal.Fig. 5 shows that this is approximately true for the bi-nary additions of Si, Ca and Zr. However, the result forAl is very dierent. By considering Figs. 4 and 5, this

Fig. 4. Relative grain size as a function of solute level for a range of

nucleant potencies (after [20]).


indicates that the nucleant particles existing in MgAlalloys are much less potent. It has been speculated thatMn causes grain coarsening in these alloys [24]. Also, itwas recently shown that the grain size of AZ91 variessubstantially depending on the purity of the ingots usedto make the alloy [22,23]. It therefore appears that alu-minium in combination with manganese and the impu-rity elements aects the potency of the nucleants in themelt.

Grain size behaviour with direct particle additions ofSiC, AIN and Al4C3 is shown by the macrographs inFig. 6. All of these particles are expected to be goodnucleants for magnesium due to a small lattice disreg-istry towards magnesium. The macrographs also con-rm that most of the particles possess a relatively goodnucleation potency. These results may support numer-ous hypotheses that these particles, particularly Al4C3,are active nucleants for magnesium in commericial al-loys since MgAl alloys can be rened by the addition ofcarbonaceous substances.

Although our understanding of the principles of grainrenement in magnesium alloys has improved signi-cantly, a reliable and readily added master alloy for grainrenement of MgAl alloys still remains to be developed.Work still remains to determine the eect of Al incombination with Mn, Fe and other impurity elementson grain size and the results may hold a key to under-standing grain renement in the MgAl alloy system.

3. Eutectic growth

As mentioned before, the cooling rates in commercialcasting processes are generally sucient to cause some

eutectic to form during solidication of magnesium al-loys containing more than 2 wt% Al [10]. Die-castings ofthe common commercial magnesium alloys, AZ91,AM50 and AM60, therefore contain a signicant vol-ume fraction of eutectic. Understanding eutectic solidi-cation in these alloys is important for two mainreasons. The rst is that this solidication event controlsthe size, shape and distribution of the more brittleb-Mg17Al12 phase in the nal microstructure, which, inturn, is likely to inuence both the ductility [10] andcreep strength [2527] of the alloys. The second is that,being the nal stage in the solidication process, eutecticgrowth aects feedability at a crucial stage, when feed-ing is interdendritic and large pressure dierentials arerequired to draw liquid through the dendritic network.A dierence in eutectic growth mode could have a largeeect on the ease with which liquid can be drawnthrough the dendritic network, and therefore on theformation of porosity in these alloys.

The eutectic exhibits a wide range of morphologies inhypoeutectic MgAl alloys depending on compositionand cooling rate [4,5]. Alloys with aluminium contentsapproaching the eutectic composition (33 wt% Al) tendto display regular lamellar or brous eutectic micro-structures [28], while those with aluminium contents lessthan about 10 wt% Al (i.e. commercial alloys) exhibiteutectic morphologies that are generally referred to asfully or partially divorced. A fully divorced morphology(Fig. 7(a)) is where the two eutectic phases are com-pletely separate in the microstructure. Each interden-dritic region consists of a single b-Mg17Al12 particlesurrounded by `eutectic' a-Mg, which has grown fromthe primary dendrites. A partially divorced eutecticmorphology (Fig. 7(b)) is characterised by `islands' ofeutectic a-Mg within the b-Mg17Al12 phase, but the bulkof the a-Mg is still outside the Mg17Al12 particle, i.e. thevolume fraction of a-Mg within the Mg17Al12 particle ismuch lower than the proportion predicted by the equi-librium phase diagram.

The eects of aluminium content, zinc content andcooling rate on eutectic morphology in permanentmould cast alloys are shown schematically in Fig. 8. Theeutectic tends to become less divorced with increasingaluminium content [4], but more divorced with increas-ing zinc content [5] and cooling rate [4,5]. The mainmechanisms by which composition and cooling rate mayaect eutectic morphology are discussed in detail byNave et al. [4,5], and are related to the location of thecoupled zone and the undercooling during solidication.An understanding of these mechanisms points the waytowards methods for modication of the MgMg17Al12eutectic. For instance, the addition of a ternary elementwhich does not partition as strongly as zinc to the liquidduring a-Mg growth, and to the Mg17Al12 phase duringeutectic growth, is likely to cause a less divorced eutecticmorphology than the addition of an equivalent amount

Fig. 5. Relationship between grain size and growth restriction factor

(GRF) for binary Mg alloys for the elements Al, Si, Zr and Ca. The

results indicate far less potent nucleant particles in MgAl alloys

compared to other alloys since the saturation level of grain density is

only reached at much higher values of the GRF.


Fig. 7. Fully divorced (a) and partially divorced (b) eutectic morphologies in a hypoeutectic MgAl alloy. The lightest areas are b-Mg17Al12 and thedarkest area are high Al content (`eutectic') a-Mg. The grey areas are primary a-Mg dendrites, showing coring from the low Al content areas near thecentres of their arms (light grey) to the higher Al content areas near the edges of their arms (dark grey).

Fig. 6. Macrographs of (a) Mg1Al, (b) Mg1Al+1vol% AlN-particles, (c) Mg1Al+1vol% Al4C3-particles, and (d) Mg1Al+1vol% SiC-particles.


of zinc, particularly if it also increases the distance be-tween the Mg17Al12 phase boundary and the eutecticpoint.

In terms of feedability during the critical last stages ofsolidication, both eutectic morphology and its mecha-nism of formation are important. Independent nucle-ation and growth of the b-Mg17Al12 phase in theinterdendritic liquid are likely to signicantly increasethe surface area to volume ratio of the feeding channels,providing a much greater resistance to feeding. On theother hand, nucleation of the b-Mg17Al12 phase onthe a-Mg and subsequent growth towards the centre ofthe interdendritic channels should allow the feedingpaths to remain more open for longer and tend to pro-mote a sounder casting. Specimens quenched duringdirectional solidication show quite a good wetting be-tween b-Mg17Al12 and a-Mg phases and growth appearsto occur from the surface of the dendrite into the in-terdendritic liquid see Fig. 9.

Feeding during eutectic solidication is also likely tobe aected by the solidication range of the eutectic andthe smoothness of the solid/liquid interface during eu-tectic growth (i.e. if any phase is a leading phase,growing ahead of the other phase during coupledgrowth), or how isothermal the eutectic growth interfaceis. An isothermal, smooth, interface allows easier feed-ing, while an interface in which one of the phases growswith a considerable lead over the other phase wouldrequire feeding along increasingly narrow and convo-luted paths in order to avoid porosity formation. Zinc,which is present in the common commercial alloy AZ91,segregates strongly to the Mg17Al12 phase during eu-tectic growth [5], increasing the lead of this phase overthe a-Mg phase and causing the eutectic to solidify witha less isothermal interface. The addition of zinc couldtherefore be expected to promote porosity formation

based on the above considerations. While there is someevidence that the addition of 2 wt% Zn increases micr-oporosity in sand-cast magnesium alloys containing 2, 4,8 and 10 wt% Al [29], more detailed work is required toestablish the eect of zinc on porosity formation in MgAl alloys.

4. Precipitation reactions

Completion of eutectic solidication does not neces-sarily mark the end of phase transformations in a castmagnesiumaluminium alloy. When the cooling rate ofthe casting is suciently slow (typical of sand-casting),precipitation may occur in the supersaturated areas ofthe a-Mg. This precipitation may take two forms, con-tinuous (Fig. 10(a)) and discontinuous (Fig. 10(b)) pre-cipitation. The most obvious form, and the form bywhich the bulk of the precipitation occurs, is discon-tinuous precipitation. This involves the growth oflamellar precipitates of Mg17Al12 into the a-Mg grains ina similar manner to the way pearlite colonies grow intoaustenite grains during the cooling of steel. The alu-minium partitions to the Mg17Al12 lamellae as theygrow, leaving the a-Mg between the lamellae muchleaner in aluminium than before discontinuous precipi-tation commenced. The discontinuous precipitationappears to grow from near the eutectic Mg17Al12 into thea-Mg grains, but whether the precipitates actually havethe same orientation as the Mg17Al12 phase, or whetherthey nucleate separately in the supersaturated a-Mgphase (e.g. on a dislocation) has not been conrmed.Discontinuous precipitation occurs mostly in the a-Mgregions near the Mg17Al12 phase, since these regionshave higher aluminium contents (approx. 1013 wt% Al)

Fig. 9. A section of a directionally solidied Mg9.1Al0.4Zn alloy

that has been quenched during eutectic growth. There appears to be

good wetting between the a-Mg and b-Mg17Al12 phases.

Fig. 8. The eect of aluminium content, zinc content and cooling rate

on eutectic morphology in permanent mould cast hypoeutectic MgAl

alloys.


than the centres of the dendrites, where the aluminiumconcentrations may be as low as 2 wt% Al.

5. Defect formation

5.1. Porosity characteristics

The solidication sequence of MgAl alloys plays animportant role in the formation of defects. The causes ofmicroporosity in MgAl alloys has been a contentiousissue since early work conducted by Baker [30] in the1940s and much debate has occurred concerning whichof the variables, solidication shrinkage or dissolvedgas, contribute to the formation of microporosity.

Initial investigations led to the belief that the forma-tion of microporosity was not signicantly aected bydissolved hydrogen. The dierence in solubility of hy-drogen between the solid and liquid phase is relativelysmall compared to aluminium alloys and it was believedthat solidication shrinkage was responsible for porosityformation. However, this theory was soon dispelled andinvestigations measuring the hydrogen content of themelt before solidication have suggested that dissolvedhydrogen does contribute to the incidence of micropo-rosity [3032].

It is unlikely that the formation of microporosity inMgAl alloys is isolated to one of the two mechanismsoutlined above. Moreover, solidication shrinkage andevolution of dissolved gas occur in unison and act col-laboratively to form microporosity just as in other alloysystems [33]. Magnesium alloys solidify relatively slowerthan aluminium alloys because of their low thermalconductivity. Therefore progressive feeding is dicultand interdendritic feeding is a very important stageduring solidication of magnesium alloys as a result ofthe wide freezing range [30]. vrelid et al. [34] reportedthat increased aluminium content decreases the solu-bility of hydrogen in the liquid.

Recent work has shown that the aluminium contentof MgAl alloys strongly aects the volume of porosity[35]. Cylindrical, unfed, castings were produced fordierent aluminium contents in the hypoeutectic MgAlregion. The results are reproduced in Fig. 11. The resultsshow that the peak in porosity occurs at about 9 wt%Al, slightly below the maximum equilibrium freezingrange as also plotted in the gure. This result can becompared to a reported maximum in hot tearing sus-ceptibility at about 1 wt% Al in the MgAl system [36].Maximum hot tearing is normally related to the occur-rence of the rst eutectic liquid and a maximum, non-equilibrium, freezing range. Furthermore Campbell [37]has argued that there is a maximum of hydrostaticpressure when the rst eutectic liquid appears and thatthis point also should correlate with a maximum inporosity. It is therefore clear from Fig. 11 that porosityin MgAl alloys does not follow the same trend andtherefore requires a dierent explanation. The low levelsof porosity in pure magnesium and the eutectic alloy (33wt% Al) are related to the isothermal, or near-isother-

Fig. 11. Percentage porosity versus aluminum content and freezing

range versus aluminium content. Note that the large freezing range

correlates with a greater incidence of porosity.

Fig. 10. (a) Continuous and (b) discontinuous precipitation in alloy AZ91E.


mal, solidication of these alloys. In the other alloys, thevolume fraction of eutectic will increase with increasedaluminium content, which can be expected to reduceporosity when the volume fraction is sucient. Thereason for the peak in porosity at about 9 wt% musttherefore be related to the worst combination of mushyzone size, interdendritic feeding, permeability and eu-tectic volume fraction. When these alloys solidify thereis insucient liquid to feed the shrinkage of the eutectic,and porosity therefore initially increase with eutecticvolume fraction. The equilibrium solidication range forthe alloys is also indicated in the diagram for the sake ofcomparison and if can be observed that the maximumamount of porosity is located closer to the maximumequilibrium solidication range rather than the maxi-mum for non-equilibrium conditions predicted to bearound 12 wt% Al. As discussed in the section on eu-tectic growth there are signicant changes in the eutecticmorphology with increased aluminium content thatcertainly may inuence porosity formation and inter-dendritic ow. Further discussion is provided in [35].This is an area of research that deserves further study.

5.2. Defect formation in pressure assisted casting pro-cesses

It has recently been documented that casting of longthin sections of magnesium alloys in pressurised castingprocesses can result in long, continuous, bands of de-fects with an outline that follows the contour of thecasting [68]. The appearance of the defect varies frombeing highly segregated, through porous, to torn. Fig. 12

shows an example of a shear band in an AM60 alloycasting where the band contains segregated eutectic andporosity.

A model has been developed which can predict theoccurrence, location and appearance of the bands basedon understanding the solidication characteristics andthe mushy zone mechanical properties [7,8]. This modelcan be simplied to describe the mushy material as aseries of thermal contours that divides the solid fractionand the mechanical behaviour of the mush into fourregions. At an early stage of lling the thermal contoursextend from a temperature near the liquidus temperaturein the centre of the casting to near or below the solidustemperature at the walls of the die cavity. The toptemperature is dened by the thermal characteristics ofthe shot sleeve of the casting machine. Often the liquidentering the die cavity will have a solid fraction of about0.2 due to premature solidication in the shot sleeve [38].The contours are delineated by critical temperatures thatdene a signicant change in the mechanical response toshear stresses generated during mould lling. The criticaltemperatures are the liquidus temperature, the coher-ency point where the dendrites rst touch each othercausing the rst measured resistance to motion, themaximum packing point where the dendrites becomefully interlocked, the eutectic temperature where themicrostructure becomes more rigid due to eutecticbridging joining the dendrites together, and the solidustemperature [7,39,40]. Table 1 summarises the dier-ences in mechanical behaviour of the regions betweenthese contours. During solidication the range ofproperties of the mush in the die cavity can be dened by

Fig. 12. (a) A defect band containing segregation of eutectic and porosity in a circular section of a high-pressure die cast AM60. The ow direction is

into the page. Note the very ne grain size on the outside of the defect band compared with the coarser microstructure inside. Higher magnication in

(b) with the region outside the band on the left hand side of the micrograph. (Courtesy: A. Bowles, CAST.)


the locations of these critical contour boundaries. Ad-ditionally, the feeding mechanism within each of thesezones can be estimated (also listed in Table 1). Fig. 13 isa representation of the change in shear strength withsolid fraction. The values of coherency and maximumpacking solid fractions depend on the dendrite size andmorphology, and small spherical dendrites have largervalues than large irregular dendrites [7,39,40].

When signicant ow has to occur while partiallysolidied mush is present, the mush will deform, and thelow-strength regions are the most likely to yield. Due tothe extraction of heat through the mould walls, themush near the walls is likely to reach a low temperaturealmost immediately, forming a rigid skin. Shear defectstherefore occur at the edge of the skin where the solidfraction is less than the maximum packing fraction andthis is the reason for the formation of the bands. Thedierent appearances of the band are caused by theamount of deformation and the solid fraction whendeformation occurs. Deformation at low solid fractionsresults in a highly segregated band, as liquid is concen-

trated to lubricate the ow of the central semisolidmush. This behaviour is similar to `plug ow' of slurrieswhere the solid material moves towards the centre of theow channel and liquid towards the edge. It is oftenobserved that the larger coarse dendrites coming fromthe shot sleeve segregate towards the centre of a thinsection [8]. With increased solid fraction and deforma-tion there is less liquid available and the amount ofporosity in the band increases, until at the extreme itappears completely torn [8].

Most of the measurements of the mechanical beha-viour of partially solidied material have been under-taken on aluminium alloys [7,39,40]. However,observation of banded defects is far more common inmagnesium alloys. Also, it is observed that magnesiumalloys can ll large distances along thin sections, muchlarger and thinner than that achievable with aluminiumcasting alloys. This can be explained by considering thedierences in the main solidication parameters be-tween, for example, the common Al7 wt% Si and AZ91casting alloys. There is a signicant dierence in thefreezing range, approximately ve times, between thetwo alloys; about 30C for Al7 wt% Si and 160C forAZ91. The volume fraction of eutectic is about 50% forAl7 wt% Si and about 25% for AZ91. Fluidity mea-surements show that AZ91 is much better than Al7wt% Si and the thermal conductivity of magnesium isalso much lower than that for aluminium. By consid-eration of the dierences in both the amount of solidfraction formed and the freezing range before eutecticsolidication begins, it can be realised that AZ91 canow signicant distances before the microstructure locksup when eutectic solidication occurs in the segregatedband. Also, the time to coherency and maximumpacking is greater. Therefore, once a segregated band ofliquid forms, much more ow is possible before thisband solidies. Eventually the temperature drops,leading to the formation of pores, and then tearing, dueto a lack of liquid to adjust for shrinkage.

Fig. 13. Representation of the change in shear strength with solid

fraction. The two plots show the dierence in behaviour between ne

globular grains and coarse dendritic grains. The strength increases

rapidly once eutectic soidication begins.

Table 1

A summary of behaviour of the partially solidied casting alloy within the zones dened by the critical thermal contoursa

Thermal zone Microstructure Feeding mechanisms Mechanical behaviour

Tl Tch Dendrites surrounded by liquid Mass feeding No strength and liquid-likebehaviour

Tch Tpk Dendrites touching each other withliquid lms

Mass feeding with an increasing

proportion of interdendritic feeding

as solid fraction increases

Low strength with slurry-like or

thixotropic behaviour. Packing

occurs. Shear causes liquid to

segregate to shear plane

Tpk Te Dendrites packed together andmechanically interlocked

Interdendritic and burst feeding Strength increases at a faster rate.

Solid-like behaviour with dendrites

deforming and network fracture

Te Ts Dendrites surrounded by partiallysolidied eutectic

Interdendritic and solid feeding while

eutectic solidication

Rapid increase in strength to Ts.Behaves as a solid

a The liquidus Tl, coherency Tch, maximum packing, Tpk, eutectic Te, and the solidus Ts.


Whether or not these banded defects form depends onthe volume of ow and the thermal state of the mushwhen ow stops. By controlling the casting parametersand die design it is possible to minimise the degree ofdamage caused by these defects [8]. How all these factorsinteract is dicult to determine, as there are very limiteddata available on the mechanical and physical propertiesof magnesium alloys. This is an area of research thatrequires signicant attention.

6. Summary

The solidication of magnesiumaluminium alloysbegins with the nucleation of a-Mg dendrites exhibitingsixfold symmetry. The grain size is set by a combinationof the cooling conditions, alloy composition and thetype of nucleant particles present. The eect of compo-sition on grain size can be estimated by the value of thegrowth restriction factor for the alloy. However, alloyscontaining aluminium do not exhibit the ne grain sizesachievable in other magnesium alloys. Also, the grainsize is aected by the purity of the base magnesium usedto make the alloys. It is speculated that aluminium incombination with manganese and the impurity elementsaects the potency of the nucleants presents in the melt.The types of nucleant particles naturally occurring in themelt have not been identied and work has shown thatthere are many particles that can be deliberately addedwhich facilitate nucleation to some extent. However, areliable, easy to use, commercial grain rener still needsto be developed.

After solidication of the a-Mg dendrites, eutecticsolidication occurs as divorced or partially divorcedb-Mg17Al12 in the interdendritic and grain boundaryregions, surrounded by eutectic a-Mg which is enrichedin aluminium. The degree of divorced growth is aectedby the zinc content and cooling rate, with an increase ineither leading to a more divorced microstructure. Thesensitivity of the eutectic microstructure to ternary ele-ments suggests that further research could be under-taken to manipulate the distribution and morphology ofthe b-phase to gain an improvement in as-cast proper-ties. At slow cooling rates, the a-Mg that forms late inthe solidication process decomposes to an a and blamellar structure by discontinuous precipitation. Somecontinuous precipitation may also occur. Heat treat-ment is able to completely dissolve the b-phase.

Due to the large freezing range and low eutectic vol-ume fraction of MgAl casting alloys, they can be castinto large thin sections. However, this advantage alsoincreases their susceptibility to banded defects of seg-regated eutectic, porosity or tears. By controlling thecasting parameters the degree of damage caused by thesedefects can be minimised. However, much more researchneeds to be carried out to gain data on the mushy zone

mechanical behaviour and solidication properties ofMgAl alloys in order to develop a capability to predict,and then avoid, the formation of these defects.

Acknowledgements

CAST was established under the Australian Govern-ment's Cooperative Research Centres Scheme.

References

[1] A.K. Dahle, D.H. StJohn, G.L. Dunlop, Developments and

challenges in the utilisation of magnesium alloys, Mater. Forum

24 (2000) 167182.

[2] D.J. Sakkinen, SAE Trans. 103 (1994) 558569.

[3] ASM Handbook Committee, Alloy Phase Diagrams ASM

International, Ohio, USA, 1992.

[4] M.D. Nave, A.K. Dahle, D.H. StJohn, Magnesium technology

2000, in: H.I. Kaplan, J.N. Hryn, B.B. Clow (Eds.), The Minerals,

Metals and Materials Society (TMS), Warrendale, PA, USA,

2000, pp. 233242.

[5] M.D. Nave, A.K. Dahle, D.H. StJohn, Magnesium technology


Metals and Materials Society (TMS), Warrendale, PA, USA,

2000, pp. 243250.

[6] S. Sannes, H. Westengen, in: B.L. Mordike, K.U. Kainer (Eds.),

Magnesium alloys and their applications, Werksto-Information-

sgesellschaft mbH, Frankfurt, Germany, 1998, pp. 223228.

[7] A.K. Dahle, D.H. StJohn, Acta Mater. 47 (1999) 3141.

[8] A.K. Dahle, D.H. StJohn, Transactions of 20th International Die

Casting Congress & Exposition North American Die Casting

Association (NADCA), November 14, 1999, Cleveland (OH),

Paper #T99062, pp. 203210.

[9] A.L. Bowles, J.R. Griths, C.J. Davidson, Magnesium Technol-

ogy 2001, The Minerals, Metals and Materials Society (TMS),

Warrendale, PA, USA, 2001, in press.

[10] I.J. Polmear, Light Alloys, Chapman and Hall, London, 1989.

[11] A. Luo, in: G.W. Lorimer (Ed.), Third International Magnesium

Conference, Manchester, UK, 1996, pp. 449464.

[12] C.E. Nelson, AFS Trans. 56 (1948) 123.

[13] R.T. Wood, The Foundryman (1953) 98.

[14] R. Hultgren, D.W. Mitchell, Trans. AIME 161 (1945) 323327.

[15] A. Luo, Scr. Metall. Mater. 31 (1994) 12531258.

[16] D.O. Karlsen, D. Oymo, H. Westengen, P.M.D. Pinfold, S.I.

Stromhaug, Light Metals Processing and Applications Canadian

Institute of Mining, Metallurgy and Petroleum, Quebec City,

Canada, 1993, pp. 397408.

[17] J.E.C. Hutt, D.H. StJohn, L. Hogan, A.K. Dahle, Mater. Sci.

Technol. 15 (5) (1999) 495500.

[18] J.E.C. Hutt, A.K. Dahle, Y.C. Lee, D.H. StJohn, Light Metals

1999, in: C.E. Eckert (Ed.), The Minerals, Metals and Materials

Society (TMS), Warrendale, PA, USA, 1999, pp. 685692.

[19] M. Easton, D. StJohn, Metall. Mater. Trans. A 30A (1999) 1625

1633.

[20] M. Easton, D. StJohn, Acta. Mater. in press.

[21] J. Hunt, Mater. Sci. Eng. 65 (1984) 7583.

[22] Y.C. Lee, A.K. Dahle, D.H. StJohn, Metall. Mater. Trans. A. in

press.

[23] Y.C. Lee, A.K. Dahle, D.H. StJohn, Magnesium Technology


Metals and Materials Society, 2000, pp. 211218.


[24] Y. Tamura, T. Motegi, N. Kono, E. Sato, Mater. Sci. Forum

350351 (2000) 199204.

[25] M.S. Dargusch, G.L. Dunlop, K. Pettersen, in: B.L. Mordike,

K.U. Kainer (Eds.), Magnesium Alloys and Their Applications,

Werksto-Informationsgesellschaft mbH, Wolfsburg, Germany,

1998, pp. 277282.

[26] M. Regev, E. Aghion, A. Rosen, M. Bamberger, Mater. Sci. Eng.

A 252 (1998) 616.

[27] E.M. Gutman, Y. Unigovski, M. Levkovich, Z. Koren,

E. Aghion, M. Dangur, Mater. Sci. Eng. A 234236 (1997)

880883.

[28] S. Guldberg, Ph.D. Thesis, The Norwegian University of Science

and Technology, Trondheium, Norway, 1997.

[29] F.A. Fox, J. Inst. Metals 71 (1945) 415439.

[30] W.A. Baker, J. Inst. Metals 71 (1945) 165204.

[31] E.A.G. Liddard, W.A. Baker, Trans. Amer. Foundrymen's Soc.

53 (1945) 5465.

[32] B.A. Mikucki, J.D. Shearouse, Magnesium Properties and

Applications for Automobiles (1993) 107115.

[33] E.J. Whittenberger, F.N. Rhines, J. Metals (1952) 409420.

[34] E. vrelid, P. Bakke, T.A. Engh, Light Metals (1997) 141154.

[35] P.L. Schaer, Y.C. Lee, A.K. Dahle, Magnesium Technology

2001, in: J. Hryn (Ed.), The Minerals, Metals and Materials

Society, in press.

[36] R.A. Dodd, W.A. Pollard, J.W. Meier, AFS Trans. 65 (1957)

100117.

[37] J. Campbell, Trans. Met. Soc. AIME 245 (1969) 23252334.

[38] W.P. Sequeira, M.T. Murray, G.L. Dunlop, D.H. StJohn, in:

S.K. Das (Ed.), Automotive Alloys 1997, TMS, Warrendale, PA,

1997, pp. 169183.

[39] A.K. Dahle, L. Arnberg, Acta. Mater. 45 (2) (1997) 547559.

[40] T. Sumitomo, D.H. StJohn, T. Steinberg, Mat Sci. Eng. A 289

(2000) 1829.


Development of the as-cast microstructure in magnesium±aluminium.pdf

Documents

Transcript of Development of the as-cast microstructure in magnesium±aluminium.pdf