Available online at www.sciencedirect.com

Scripta Materialia 60 (2009) 516–519

www.elsevier.com/locate/scriptamat

The kinetics of Fe-rich intermetallic formation in aluminiumalloys: In situ observation

Junsheng Wang,a Peter D. Lee,a,* Richard W. Hamilton,a Mei Lib and John Allisonb

aDepartment of Materials, Imperial College, South Kensington Campus, London SW7 2AZ, UKbFord Research Laboratory, Dearborn, MI 48121-2053, USA

Received 24 October 2008; revised 28 November 2008; accepted 29 November 2008

In situ synchrotron radiography of Fe-rich intermetallic formation was performed during the solidification of an Al–7.5Si–3.5Cu–0.8Fe (wt.%) alloy. Growth kinetics was quantified by segmenting the Fe-rich intermetallic phases and nucleation tempera-tures were determined by extrapolating to zero size. Fe-rich b-intermetallics nucleated between 550 and 570 �C, growing initially atan instantaneous tip velocity of 100 lm s�1, and slowing to 10 lm s�1 towards the end of growth. Final plate size was controlled bylocal Fe concentration and a-Al impingement.� 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: X-ray radiography; Image analysis; Aluminum Alloys; Intermetallics; Kinetics

During the recycling of aluminum alloys, mixedscrap is incorporated, invariably increasing the Fe con-tent to a level of 0.4–0.8wt.% [1]. This concentration issufficient to allow the formation of the highly facetedplate-like b-Al5FeSi, which can severely degrade themechanical properties (e.g., fatigue life and ductility)of any cast components [2,3] and prevent the successfulproduction of high-strength wrought products [4]. Manyexperimental studies have attempted to alter the mor-phology of b-intermetallics into a less harmful form suchas script-like a-Al8Fe2Si or a-Al15(Fe,Mn)3Si2 throughthe addition of alloying elements (e.g., Mn, Mg, Sr,Cr, Be, Ni, V and TiB2) [5–7]. However, without a great-er understanding of these phases’ growth kinetics, suchinvestigations will remain a costly trial-and-errorprocess.

In this study, the nucleation and growth of Fe-richintermetallics in an Al–7.5Si–3.5Cu–0.8Fe (wt.%) alloywas observed using synchrotron radiography. This al-lowed the nucleation temperatures and the growth ratesof the plate-like b-Al5FeSi to be quantified together withthe interaction between phases.

An infrared furnace (Fig. 1) was designed to providecontrolled melting and resolidification of Al alloy sam-ples while using high-resolution synchrotron radiogra-

1359-6462/$ - see front matter � 2008 Acta Materialia Inc. Published by Eldoi:10.1016/j.scriptamat.2008.11.048

* Corresponding author. Tel.: +44 207 594 6801; fax: +44 207 5946758; e-mail: p.d.lee@ic.ac.uk

phy to observe evolution of microstructure. A100 � 5 � 1 mm sample of Al alloy (Al–7.5Si–3.5Cu–0.8Fe) was contained in a high-purity boron nitride(BN) boat (Boron Nitride Products GmbH, Kempten,Germany) and temperature was measured with a K-typethermocouple in direct contact with the specimen. Asynchrotron (Beamline No. 16.3 at SRS Daresbury Lab-oratory, UK) provided a white, parallel beam of X-rays(�34 keV) coupled with a 4008 � 2672 pixel detector/image change producing a pixel size of 1.0 lm. The sam-ple was heated to 640 �C (fully molten) and stabilizedbefore being allowed to cool at 0.33 �C s�1. A total of900 radiographs were acquired at 2 Hz and image anal-ysis was performed using ImageJ (ImageJ, NationalInstitutes of Health, USA) as follows: (i) the back-ground image was subtracted; (ii) an edge-preserving fil-ter (Mean shift, Prof. Dr.-Ing. Kai Uwe Barthel,Internationale Medieninformatik, Berlin, Germany)was applied; (iii) a region of interest (1500 � 1000 pixels)was chosen that contained intermetallic plates whichwhere segmented using local thresholding; and (iv) thefinal morphology was quantified.

Four typical images of the solidification sequence areshown in Figure 2, illustrating the key stages of interme-tallic formation. In this alloy, the dendritic structure isdifficult to resolve as it has a similar attenuation to thatof the interdendritic liquid. However, the Fe-rich inter-metallics are clearly resolved due to their increased con-centration of highly attenuating Fe (�25 wt.% Fe)

sevier Ltd. All rights reserved.

Figure 1. Schematic of the set-up for in situ observation.

J. Wang et al. / Scripta Materialia 60 (2009) 516–519 517

compared to the primary a-Al phase (�0.05 wt.% Fe)(Fig. 2). It should be noted that the intermetallics, whichappear needle-like in the two-dimensional radiographs,are actually plate-like as shown by serial sectioning [8].In order for a phase to be visible, a through-thicknessattenuation variation of >1% is required. The thin(<10 lm) b-plates are therefore only observable whenaligned edge-on to the beam. As the temperature de-creases from 570 �C (Fig. 2a) to 540 �C (Fig. 2b),plate-like Fe-rich intermetallics (dark regions) were ob-served to nucleate. The growth of the four intermetallicshighlighted in Figure 2b was quantified, illustrating awide range of nucleation temperatures (Fig. 3). This isalso the temperature range over which the Al–Si eutecticphase will nucleate and grow [5]. Unfortunately, thisphase has an X-ray attenuation nearly identical to theprimary phase, and hence cannot be quantified fromthe radiographs.

Starting at temperatures of about 540 �C, porosityforms (the white phase in Fig. 2c), growing in the spacebetween dendrite arms as the liquid is drawn out tofeed volumetric shrinkage and by the evolution ofhydrogen [9–11]. The porosity level is relatively highfor two reasons: (i) the sample was not degassed, and(ii) the conditions are similar to a hot spot in a cast-

Figure 2. Digital radiographs showing the Fe-rich intermetallics (b), porosity(a) ts = 90 s, 570 �C; (b) ts = 180 s, 540 �C; (c) ts = 180.5 s, 539.8 �C; (d) ts =

ing—i.e. solidification is occurring inwards from bothedges, rather than unidirectionally, with no externalsource of molten metal feed. As the interdendriticliquid is displaced, the liquid around the b-plates is alsodisplaced, partially exposing them. The associatedchange of attenuation contrast from b-plate/liquid tob-plate/void causes better resolution of the b-plates,producing an artificial jump in growth rate, as will bediscussed later. The pores reach a size of up to500 lm within a single frame, hence growing at a veloc-ity in excess of 1 mm s�1. These spurts of pore growthwere observed to occur over a range of temperaturesfrom 540 �C until solidification ended at �510 �C(Fig. 2d).

One may conclude from the radiographs in Figure 2that intermetallic growth is initially very rapid, growingmostly at the edges of the plates where attachment ismost favorable. Quantitatively, this is illustrated bythe steep growth curves of maximum length (Fig. 3).These b-phase plates nucleate at a solid fraction (�0.4)sufficiently great that the primary dendrite grains are al-ready well developed. The nucleation of these plates isalso difficult, resulting in the interdendritic liquid beingsupersaturated in Fe. This large undercooling is similarto that found for the formation of the similar intermetal-lic phase in the a-Al/Al3Fe eutectic reaction in binaryAl–Fe [12]. Once nucleated, the plates grow with a tre-mendous burst of speed, with an initial instantaneouspeak velocity observed to be in excess of 100 lm s�1

(in a single frame, it grows �50 lm, or v = 50 lm/0.5 s). The b-Al5FeSi phase grows as faceted platesdue to its large entropy of fusion (�4.97 J mol�1 K�1)[13], with the largest intermetallic in Figure 2 reachinga length of 170 lm in less than 5 s after initial observa-tion, as shown quantitatively in Figure 3a–A. This inter-metallic continued to grow to 370 lm over a further 20s, as shown in the inset in Figure 3a–B. This first stage,‘‘lateral growth”, ceases at this point when the rapidlyexpanding plate-like b -phase impinges on the surround-ing primary phase dendrites/grains. The average growth

(P), and Al–Al2Cu eutectic (Cu) at different solidification temperatures:270 s, 510 �C.

Figure 3. (a) Quantified evolution in length (Lmax) of individual plate-like Fe-rich intermetallics. Inset radiographs illustrate growth stages ofthe largest intermetallic plate at: (A) 564 �C, (B) 558 �C, (C) 531 �Cand (D) 502 �C. (b) The evolution in cross-sectional area of individualFe-rich intermetallic plates.

518 J. Wang et al. / Scripta Materialia 60 (2009) 516–519

velocity of the four quantified intermetallics during thelateral growth stage is 34 ± 20 lm s�1.

The intermetallic is now constrained by the primarydendrites and can no longer grow in the most favorableorientation, with the rate of change in maximum lengthremaining constant almost until the end of solidification(see Figure 3a). In this second stage of growth, the b-plates grow by thickening, expanding in the less-favor-able growth direction normal to the faces of the plate(compare Figure 3a-B with Figure 3a-C). This growthmechanism is similar to the ledge-wise growth mecha-nism observed by Laird and Aaronson [14]. To illustratethis growth quantitatively, the change in cross-sectionalarea for each intermetallic is plotted in Figure 3b.Although these growth curves are not completelysmooth, this is due to the changes being slow in compar-ison to the resolution (1 lm pixel�1). The actual changein cross-sectional area will be more parabolic, withgrowth limited by attachment kinetics and diffusion be-tween the two phases (a-Al and b-Al5FeSi) [15]. Thesource of Fe for growth is provided by the continuedformation of the primary phase and also the formationof the Al–Si eutectic which has been shown to occurover these temperatures [5,16], and has an even lowersolubility than the primary phase (Fe solubility in theAl-phase is less than 0.05 wt.% at this temperature whilethe Si-phase has nil solubility [17]). This second stage ofintermetallic growth, ‘‘attachment and diffusion limited”

growth, continues until the Al–Al2Cu eutectic phaseforms at approximately 510 �C. During the secondstage, porosity was also observed visible as the whitephase in the high-magnification inset, Figure 3a-C. Not-ing again that the porosity was observed to form aroundintermetallics (perhaps nucleating on them), the poresgrow to a size of hundreds of microns in a single 0.5 sframe. In the same frame, the b-plates appear to thicken(see step growth in Fig. 3b), but this growth is probablyboth an artefact of the change in X-ray attenuation con-trast and partially due to thickening. Because the inter-metallics are now constrained along the preferredgrowth direction of the plate edges, this growth isthrough thickening via ledge growth and hence is at leastan order of magnitude slower than pore growth (takingseveral seconds, Fig. 3b, as compared to a single frame).

The third and final stage, ‘‘step growth”, occurs whensolidification is completed by the formation of the Al–Al2Cu eutectic, as shown in the inset in Figure 3a-D.This eutectic was observed to nucleate at �510 �C, againin liquid regions close to Fe-rich intermetallics. The evo-lution of this phase has only a minor effect on the sizeand aspect ratio of the b-plates (see Fig. 3a and b), sig-naling the end of solidification and the growth of the b-intermetallics from the liquid.

In summary, the morphology of the plate-like b-Al5-

FeSi phase was observed to evolve in three distinctgrowth stages:� Stage I—lateral growth: once nucleated, the b-

Al5FeSi plates grow rapidly from the Fe-supersatu-rated interdendritic liquid via the following eutecticreaction:

Liquid ! a�AlþAl5FeSi � 570 ! 555 �C

� Stage II—attachment and diffusion-limited growth:prevented by impingement from growing on thefavorable sites, growth continues via ledge growthon the plate faces, limited by this attachment kineticsand the rejection/diffusion of solute:

Liquid ! ða�Alþ SiÞ þAl5FeSiðþPoresÞ� 565 ! 510 �C

� Stage III—final step growth: formation of the Al–Sieutectic and Al2Cu phases providing two final spurtsof growth when Fe is rejected by each of the followingreactions:

Liquid! ða�Alþ SiÞ þAl2CuþAl5FeSi

� 510 ! 500 �C

These observed transformation temperatures in theAl–7.5Si–3.5Cu–0.8Fe quaternary alloy system agreequalitatively well with CALPHAD calculations (Ther-mo-Calc TCCR, Thermo-Calc Software, Stockholm,Sweden) using the database in Ref. [18], as shown by com-paring them with calculated ones in Figure 4. The nucle-ation temperature of each of the four intermetallics can

500

510

520

530

540

550

560

570

580

590

600

610

Tem

pera

ture

, C

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0Mole Fraction of Solid

1

2

3

4

5

599.96 C

581.30565.44

504.96

Scheil Model

Equilibrium Solidification

º

Cº

Cº

Cº

Figure 4. Solidification path of the Al–7.5Si–3.5Cu–0.8Fe (wt.%)quaternary alloy calculated from the CALPHAD model (Thermo-CalcTCCR) using Scheil and equilibrium equations. (1) Liquid; (2)liquid + primary Al; (3) liquid + b-Al5FeSi + primary Al; (4)liquid + b-Al5FeSi + primary Al + Al–Si eutectic; and (4) liquid + b-Al5FeSi + primary Al + Al–Si eutectic + Al–Cu eutectic.

J. Wang et al. / Scripta Materialia 60 (2009) 516–519 519

be estimated by extrapolating their size to zero, giving val-ues of 558, 564, 565 and 566 �C (Fig. 3a), or an averagenucleation temperature of 563 ± 4 �C. Note that theexperimentally measured first intermetallic formation oc-curs at a lower temperature (�566 �C) than predicted(�581 �C) because of the kinetic effects (e.g., undercool-ing required for nucleation) and the restriction ofresolution.

Unlike the quasi-peritectic growth theory proposedby Sha et al. [4], no intermediate script-like a-AlFeSiphase was observed in this study, a phase which theyhypothesized later transformed via a peritectic reactioninto b-Al5FeSi, even though the cooling rates were com-parable (0.33 as compared to 0.17 �C s�1). However, theintermetallic growth velocities observed here are in goodagreement with prior observations by Liang and Jones[12] for the similar phase formation of Al3Fe in binaryAl–Fe alloys (0.005–0.1 mm s�1 as compared to 0.01–0.4 mm s�1).

The kinetics of Fe-rich intermetallics is successfullyinvestigated by in situ observation of microstructure for-mation in an Al–7.5Si–3.5Cu–0.8Fe (wt.%) alloy usingsynchrotron X-ray radiography. To explain the quanti-tative data, a three-stage mechanism of b-Al5FeSi for-mation is proposed. In the first stage (lateral growth)Fe-rich, b-Al5FeSi, intermetallics nucleate at 550–570 �C and form faceted, plate-like shapes within 20 s,growing rapidly until constrained by the surrounding

primary dendrites. A second stage of attachment anddiffusion-controlled growth follows where the platesthicken by ledge-wise growth as solute diffuses fromthe solidifying primary and Al–Si eutectic phases. Inthe final stage of growth, first pores and then a-Al/Al2Cu eutectics form, rejecting Fe, promoting a finalstepwise thickening of the intermetallic plates, slightlyreducing their aspect ratio.

The author(s) acknowledge the support fromDorothy Hodgkin Postgraduate Awards, and fundingfrom Ford Motor Company. The authors also acknowl-edge many useful discussions both with colleagues atImperial College London and useful synchrotron radi-ography data from SRS Daresbury Laboratory BeamNo. 16.3 (Beamtime Award No. 50264).

[1] T.A. Burns, Foseco Foundryman’s Handbook, PergamonPress, Oxford, 1986.

[2] J.Z. Yi, Y.X. Gao, P.D. Lee, T.C. Lindley, Mater. Sci.Eng. A 386 (2004) 396.

[3] J. Allison, M. Li, C. Wolverton, X. Su, JOM (2006) 28.[4] G. Sha, K.A.Q. O’Reilly, B. Cantor, J.M. Titchmarsh,

R.G. Hamerton, Acta Mater. 51 (2003) 1883.[5] L. Lu, A.K. Dahle, Metall. Mater. Trans. A 36 (2007)

819.[6] H.F. Samuel, Metall. Mater. Trans. A 29 (1998) 2871.[7] C.M. Allen, K.A.Q. O’Reilly, B. Cantor, Acta Mater. 49

(2001) 1549.[8] C.M. Dinnis, J.A. Taylor, A.K. Dahle, Scripta Mater. 53

(2005) 955.[9] P.D. Lee, J.D. Hunt, Conference on modeling of casting,

Welding and advanced solidification processes VII, Lon-don, 1995, Minerals, Metals and Materials Society/AIME, Warrendale, PA, 1995, p. 585.

[10] P.D. Lee, J.D. Hunt, Scripta Mater. 36 (1997) 399.[11] P.D. Lee, J.D. Hunt, Acta Mater. 49 (2001) 1383.[12] D. Liang, P. Korgul, H. Jones, Acta Mater. 44 (1996)

2999.[13] Z.Z.-K. Liu, Metall. Mater. Trans. A Phys. Metall. Mater.

Sci. 30 (1999) 1081.[14] C. Laird, H.I. Aaronson, Acta Mater. 17 (1969) 505.[15] H.I. Aaronson, T. Furuhara, M.G. Hall, J.P. Hirth, J.F.

Nie, G.R. Purdy, J.W.T. Reynolds, Acta Mater. 54 (2006)1227.

[16] A.L. Narayanan, F.H. Samuel, J.E. Gruzleski, Metall,Mater. Trans. A 25 (1994) 1761.

[17] C.M. Allen, K.A.Q. O’Reilly, B. Cantor, P.V. Evans,Prog. Mater. Sci. 43 (1998) 89.

[18] Y. Du, J.C. Schuster, Z.-K. Liu, R. Hu, P. Nash, W. Sun,W. Zhang, J. Wang, L. Zhang, C. Tang, Z. Zhu, S. Liu,Y. Ouyang, W. Zhang, N. Krendelsberger, Intermetallics16 (2008) 554.