Non-equilibrium solidification of hyperperitetic Al-Zr alloys

Non-equilibrium solidification of hyperperitetic Al-Zr alloys

NON-EQUILIBRIUM SOLIDlFICATION HYPERPERITETIC Al-2 ALLOYS OF E. SES and H. BILLDAL Central Institute for lndustriat Research, Oslo, Norwag Abstract...

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E. SES and H. BILLDAL Central Institute for lndustriat Research, Oslo, Norwag

Abstract-The transformation products funned during rapid casting of two alloys tvith composition Al-O.8 wt 3, Zr and Al-i.5 %t Y0Zr have been examined by transmission electron microscopy. Thz diffsrent phases to be nucleated during freezing are; {i) the eqliilibrilim Al,Zr phase [ii) the metastabic cubic .41,Zr phase and (iii) supersaturated .-\I-Zr solid solution. The metastable phase forms dendritic configurations up to 4Opm in diameter with the individual dendrite arm segments having a needle form about 10 nm in diameter and oriented in ; 1IO > directions. Intermittent nucleation of supersaturated x-h1 was found to occur during dendritic growth of the metastable phase, resulting in symmetric dendrite-free zones inside the dendritic configuration. In the M-1.5 wt “,,Zr allov. also. large solid particles of the metastablc phase were formed in thz melt. The observations are.inrerpreted within the contect of the stable and meiastable phase diagrams. R&m&-On a observe en microscopic Clectronique par transmission Ies produits de transformations qui se forment au tours de la coulSe wpide de deux alliages dr composition A-O.S”,Zr et AI-i.Y;Zr (en poids). Lcs dilErentes phases qui peuvent germer au tours du refroidissement sont: (i) la phase d’t5quilibre Al,Zr. (ii) la phase cubique mctastabie At3Zr et (iii} la solution solide sursnturtie Al-Zr. La phase m&stable produit des con~~urations dcndrit~ques attei~nant jusqu’~~JOjrm de diamctre: les segments de dendrites ont la forme 2aipuilles de IO nm dr diam&re. orientnties selon les directions < 110 >. On a observ5 une germination %termittente de I’ ;\I 1 surwturi. qui produit des zones symi-triquss sans dendrite. h l’int&ieur dc la confixuration dcndritique. De grandes particules solides de la phase mttastabk se formaient kgalement dansi’alliage Al-I.?,, Zr fondu. On esplique Ies obsenations 3 I’aide des diagrammes de phases stable et metastable. Zusammenfassung-hlittels Durchstrahlungselektronenmikroskopie wirden die Umaandlungsprodukts untersucht, die wlhrend des raschen GieRens zweier Legirrunpen in den Zusammensetzungen AI-O.8 Gea.-“, Zr und M-1.5 Gew.-‘, Zr entstehen. Die verschiedenei wPhrend des Abkiihlcns zu bildenden Phasen sind: lit die GleichPewichtsohass .-\12Zr. (ii) die metastabile kubische Phase Al,Zr und (iii) die iiberslttigt;‘feste Ltisuni AI-Zr.’ Die metastabile Phase biidet dendritische Konfigurationcn aus in einem Durchmesser bis tu 400flm mit einzelnen D~d~tarm~~menten in Nadelform mit etwa IOnm Durchmesser. die in ( 110)-Richtungen orientiert sind. WShrend des dendr~ti~hen Wachsens der me~~stabilen Phase tritt eine aussetzendt Bildung iibers3tigten ~-.~luminiums auf, welchcs zu s~rnmet~~h~n. dendritfreien Zonen innerhaib der Dendritkonfiguration~n ftihrt. In der Legierung Ai-l.5 Gew.-“,Zr wurden aul3erdem fests Partikel der metastabilen Phase in der Schmebe gebildet. Die Beobachtungen werden im Zusammenhang mit Diagrammen stabiler und metastabiler Phases gedcutet.

1. INTRODUCTION The properties of aiuminium and aluminium alloys may be profoundly influenced by zirconium additions. A recent example in this regard is that an addition of about 0.5 =x-t“; Zr to several of the commerciat Alalloys make them superplastic [l, 21. Another interesting aspect is that ~l~di~cation of a binary Al-Zr alloy with more than 0.5 wt :< Zr may be associated with a -‘super grain refinement” effect resulting in grain sizes with diameters less than 10 !trn [il. Both of these Zr-induced elects. however, rely on a rapid casting technique. In order to obtain superplasticity. a substantial amount of zirconium in solid solution is needed in the as-cast material. and a fine casting *The terminology suggested by Kerr rt nf. [4] is used in- this work.

grain size requires a high cooling rate in order to nucleate the metastable Al,Zr phase at the expense of the equilibrium phase. Accordingly. while the freezing-out of the metastable At,Zr is a primary objective in terms of grain refinement. this effect may represent a nuisance in acquiring a maximum amount of zirconium in sotid solution. The metastable phase con~gurations that form during rapid cooling has been named “sponge particles” [1] and “petal-like constituents” [3]_ The internal structure and growth mechanism of these configurations have not been established. Further, no attempts have been made to interpret their formation within the context of non-equilibrium reactions in the hyperperitectic* region of the AI-Zr phase diagram. The objective of this work is to clarify these aspects of the solidification behaliour of Zr-bearing aluminium alloys.






2. ESPERDIENTAL The present investigation pertains to two chill-cast alloys; AI-0.S \vt ‘; Zr and X1-1.3 wt “<> Zr. which were prepared from high-purity aluminum and an Al-Swt:,;ZZr master alloy. Melting was carried out under argon cover and the melt was held in the ternperature

range 9%.IOOO’C

for 30min

before it was

cast into a 1 cm dia. and S cm deep copper mould. The alloys have been examined b> optical and transmission electron microscopy. Foils for TEM investigation were prepared by el~tro~I~shing in a methanol-nitric acid solution in a standard way. The foils were examined in a Philips E.&I. 300 electron microscope. The as-cast zirconium distribution was obtained by using energy dispersive microanalysis in association with TEM. 3. EXPERIMENTAL


The low-power optical micrograph in Fig. t(a) represents part of a cross-section of the 0.8 wt yz;Zr-conraining cast ingot, displaying typical features as the outer columnar and the central equiaxed regions. In this cast structure, a salient feature is the frequent presence of either a triangular-. a square-, or a more complicated shaped geometric configuration which usually is situated within the central region of the equiaxed grains (Fig. lb). Ko grains have been found to contain more than one such dendritic configuration. Except for the smaller configuration and grain sizes, the cast structure of the iS%Zr alloy displays the same macroscopic features. In the l.j%Zr alloy, the configurations have a typical diameter in the range 1%2Ojtm vs a mean grain diameter of about 3ltirm compared to 30 pm vs 60 pm configuration/

Fig. 1. Optical micrographs of the Al-Q.S wt *; Zr cast ingot. (a) part d an ingot cross-section showing the outer cofumnar and the central rquiaxed grain regions. fb) the cast-in configurarions in the @axed region.





grain diameters in the O.S”,Zr alloy. On the internal structure of these con~~u~ations. however, the two ahoys exhibited some interesting differences. as described in wparate sections below. 3.1 The Al-0.S wt O0Zr cvrtjkpWiorts Transmission electron microscope examination of foits prepared from the equiaxed pdrt of the as-cast ingots revealed regions of high particle density. These particles which invariably were found to form complicated dendritic patterns. could by diffraction analysis be identified as the metastable Al,Zr phase [j. 6-J. By performing a detailed mapping of larger sections of a foil. composite images like the one shown in Fig. 2 could be constructed. The cast-in config~trations displayed by such composite technique were of the same size and revealed the same macroscopic geometry as those in optical micrographs (Fig. 1b). In their simplest forms. the configurations can be described as having a regular cubic morphology with surfaces paralled to : 100: matrix ptanes. The composite micrograph in Fig. 2 represents a cross-section of such a configuration. No orientation difference between the dendritic region and the non-decomposed matrix could be observed b>- TEM. On the internal configuration structure, each dendritic arm can be described as consisting of needle-

Fig. 2. TEM




like segments branching out in different ( 1103 directions. As illustrated by the micrograph in Fig. 2. the dendrites which are oriented in < 110 > directions are Iong and straight, while branches in other directions appear to be due to a zig-zag growth pattern altemating between nearby ( 110 > directions. The dendrites have a uniform thickness of about 1Onm and a tip radius of about 2-t nm. The arm spacing decreases from 0.2-M itrn in the outer part of the dendritic configurations to a nearly unresolvable distance in the central region. 3.2 The AI-L.5 Zr corrjigtrmtiorts The cast-in configurations of the more highly alloyed material display the same macroscopic morphologies as described above. Also, the internal structure may be identical, with the important difference that the Al-l.5 Zr alloy configurations frequently contain core regions which are completely transformed to the metastable A1,Zr phase. An example of such a duplex con~g~~r~tion is presented in Fig. 3. where the two dark-field micrographs are slightly tilted with respect to the [Xl] particle,‘matrix orientation to bring the dendritic region (Fig. 3a) and the central completely transformed region (Fig. 3b) into maximum reflecting conditions. The outer dendritic region have a similar structure to that described above, with

dark-fefd composite image of a mstastablc Srl-0.Ywt “,, Zr ingot.




in an






Fig. 3. Two dark-field images slightI> tilted with respect to the [Xl] particle matrix or~entawn to bring the dendritic region faj atld the comptstely transformed region (b) into masimum reii,z-ring conditions. From an Al-l.5 wt “,Zr ailoy.

the dendrites of Fig. 3 oriented in three principal directions, i.e. on the AB and BC sides of the core particle. the dendrites are strongly elongated in the [lOI] and [01 I] directions. respectively, while the dendrites branching out from the bottom AC side are alternating between [Oli] and [lOi] directions. The diffraction patterns obtained from the transformed core region are indistinguishable from those of the dendritic region, i.e. they can be indexed according to the metastable Al,Zr phase. An interesting structural detail in Fig. 3(b) is the planar boundaries dividing the triangular particle into three sections D. E and F. These boundaries which define : 110; crys~llographi~ planes [i.e. (01il. (1 i0) and (lOl)] are intersecting each other along the common [I I I] direction in the center of the triangular partide. A similar boundary pattern has been found in all the investigated core region precipitates, and accordingly appears to be a general structual characteristic. If it is assumed that these metastable A1,Zr particles are not fully ordered according to the f.c.c. Ct~,~ttu structure. but have a small tetragonal component towards the D02,-structure of the equilib rium A1,Zr phase [j]. then several possibilities exist for the interpretation of these boundaries (i.e. stacking faults, twin boundaries. antiphase domain boundaries). However. to distinguish between these possibilities. a more detailed contrast and diffraction pattern analysis is required.

The dark-field mi~rogrdpll in Fig. 4 shows a more complicated cast-in configuration imaged in a [OOI] lattice orientation. The dendritic network. which in this orientation clearly exhibits the preferential ( 1IO > growth directions. appears to have orisnated from the three particles marked A, B and C. The particle marked A has been identified as the equilibrium tetragonal Al,Zr phase [5] while the B and C particles are of the metastable type similar to that described above. On the origin of the equilib~um phase particle, A, this may be an undissolved remnant from the master alloy. However. the long straight needles in the region marked D are also of the equiIibrium phase, ill~lstrating that this phase also ma>- grow dendritically during casting.

Some dendritic con~gurations display a peculiar structural irregularity in the form of about 1 blrn wide particle free zones which were oriented parallel to the configuration-matrix interface. 3s shown in Fig. 5. Such zones have been found in con~gurations in both of the investigated alloys. By combining TEM with energy dispersive microanalysis. it was demonstrated that these zones were associated with a zirconium deficiency. For instance. the scan itcross ti;t configuration in Fig. j revealed a zirconium concentration of about 3 wt y.0 in the central region decreasing to about 1 wt O,,in the outer. while the particle-free zone


Fig. 5. TEM dark-field image al a dendritic configuration having n qmttric 1 ,ttm wide dendrite-k zone. From an AI-&Y w ?;Zr alloy. The zirconium wncrntration profile has been obtained by combining TEM aith energy dispersive microan&sis.




Fig. 6. Detail from a dendrite-free zone in a Al-l.5 wt T0Zr configuration. The central region of the configuration is located to the left of the zone.

contains less than 1 wt 0:. As can be seen from Fig. 5. the central region contains. in addition to the metastable phase. a small volume fraction of plate-shaped particles (marked by arrow and seen edge on in Fig. 5). Due to the high particle concentration in this region, we were not able to identify these small particles, but they are believed to be of the equilibrium phase. Another interesting detail is that the individual dendrite particles appear to be much coarser inside the particle-free zone. as illustrated by the dark-field micrograph in Fig. 6. Both the assumed presence of the stable phase and the coarser particle structure within the symmetrical particle-free zones indicates that these are frozen out at a considerably higher temperature than their surroundings.



the equilibrium A1,Zr phase will be nucleated at temperatures below T,. An increase in the cooling rate. however, will suppress the formation of the equilib rium phase. and at temperatures below Tr the melt will also be supersaturated with respect to the metastable A1,Zr phase. Althou_gh this phase has a higher free energy than p, its formation may be favoured due to a lower surface energy. By a further increase in cooling rate. the formation of p may also be suppressed, causing the melt eventually. at temperatures below T,.. to be supersaturated with respect to the Al-Zr solid solution. causing the solidification to follow the dotted lines in the phase diagram in Fig. 7. The cooling rate of the present chill casting has not been determined experimentally. but can. on the basis of the work by Ohashi and Ichikawa [3]. be estimated to be in the range IO-50’C,‘sec, which in this contest is considered an intermediate cooling rate. A salient result, however, is that this casting rate produced all of the above-mentioned reaction products in both of the investigated hyperperitectic alloys The formation of the metastable AlsZr phase represents the dominating transformation in the present casting experiments. This phase appears to be easily nucleated in the melt, growing dendritically at a rate of about 10j~m/sec. With a sufficiently high Zr content, the metastable AI,Zr phase may nucleate and grow as solid particles in the melt. As these particles grow, the Zr concentration in the melt ahead of the reaction front will decrease and eventually- the dendritic morphology again is favoured. The result is dendritic configurations which contain large core region particles (Fig. 3). The puzzling dendrite-free zones can be looked upon as a recording of a short period in the freezing history where the formation of the metastable phase has ceased at the expense of the supersaturated x-Al. Although the nucleation of r-Al above the petitectic temperature, with the metastable phase particles acting as nucleation sites, was reported also in the work by Kerr et al. [4], the present observation that this

4. DISCUSSION A convincing analysis of the possible reactions involved during peritectic transformations has been given by Kerr er al. [4]. Their work pertains to the Al-l-i system, which is known to exhibit a solidification pattern similar to Al-Zr [3]. Following Kerr et al.. the possible transformations are: (i) the kquilibrium petitectic reaction. (ii) non-equilibrium peritectic reactions or (iii) suppression of the peritectic transformation. with increasing cooling rates favouring the latter reactions. Now, by adapting the Kerr et (II. analysis to the present casting experiments. the possible hyperperitectic transformations can. in terms of the hypothetical phase diagram presented in Fig. 7. be summarized as follows: with a slow cooling rate.







Fig. 7. Hypothetical equilibrium (fully drawn tines) and non-equilibrium (broken and dotted lines) phase diagram for an Xl-Zr alloy.

solidification of r-A1 occurs only as an intermission between @’dendritic growth periods may seem difficult to u~ders~~ However, by returning to the phase dipgram in Fig. 7. this weird Freezing pattern can be interpreted as follows, By again assuming that the ~~i~~~cat~on of Ai-Zr is analogous to that of Al--%, then to a specific coofing rate corresponds ;i critical zirconium concentmti~n, C*, above which the nucleation of supersaturated u-Al willl not occur [4]. With the present intermediate cooling rate, this concentration C* should be considerably less theta the alby compositions [3,4]. and the freezing out of g-:-AI restricted to the shaded area in the phase diagram (Fig. 7). Now, upon tooting of the melt from higher tem~ratur~s~ the me~stabl~ phase will nucleate below 7’,*and the subsequent high dendtitic growth rate should cause a rapid zirconium depletion of the meLt ahead of the advancing dendritic growth front, causing these front regions to enter the temperatureconcentration region in the phase diagram where the nucleation of x-Al is favoure~ During the freezing of x-Al, the zirconium concentration ahead of the reaction front may again approach COand a renucleation of 13”occur before the melt reaches the non-equilibrium peritetic tcrn~mt~r~ TP..

It is realized that some of the interpretations above are specukttive; they are. however, consistent with reasonable interpre~tions of the data that can be found in the literature on non-equilibrium transformations in peritectic systems. Continued research should now focus on establishing the detailed thermodynamics of the various tr~nsfo~ations involved.


authors wish to express their gratitude to the Royal Norwegian Council for Industrial and S&r&k Research for financial support.


2. B. M. Watts, M. J. Stow& B. L. Baike and D. G. E, Owen, Metal Sci. 1% 189 (19763. 3. T. Ohashi and R. fchikawa, Z. ~~eru~~~. 64. 517 (1973). 4. H. W. Kerr, J. Cisse and G. F. Boiling. .&a Met. 22. 677 (19741. 5. N. Ryuni,Rem ,Met 17. 269 (1969). 6. E. NW, Acru Mer. d. 499 (1972]_