Intermetallic Phase Selection in 1xxx Al Alloys

Intermetallic phase selection in 1XXX Alalloys

C.M. Allena, *, K.A.Q. O'Reilly a, B. Cantor a, P.V. Evansb

aOxford Centre for Advanced Materials and Composites, Department of Materials, University of Oxford,

Parks Road, Oxford, OX1 3PH, UKbAlcan International Limited, Banbury Laboratory, Southam Road, Banbury, Oxon, OX16 7SP, UK

Accepted 1 October 1998

CONTENTS

1. INTRODUCTION 902. BINARY Al±Fe PHASES 92

2.1. The equilibrium Al±Fe4AL13 eutectic 922.2. Metastable Al±Fe eutectic phases 93

2.2.1. Metastable Al±FeAl6 eutectic 952.2.2. Metastable Al±FeAlm eutectic 962.2.3. Metastable Al±FeAlx eutectics 972.2.4. Metastable Al±Fe2Al9 eutectic 1012.2.5. Metastable Al±FeAlp eutectic 1012.2.6. The e�ect of Si addition on the formation of Al±Fe phases 101

3. TERNARY Al±Fe±Si PHASES 1023.1. The equilibrium a-AlFeSi and b-AlFeSi phases 1023.2. Metastable a-AlFeSi and b-AlFeSi phases 105

3.2.1. Metastable cubic a-AlFeSi phase 1053.2.2. Metastable av-AlFeSi phase 1063.2.3. Metastable a0 or q1-AlFeSi phase 1063.2.4. Metastable aT-AlFeSi phase 1063.2.5. Metastable q2-AlFeSi phase 1093.2.6. Metastable b 0-AlFeSi phase 111

4. FACTORS GOVERNING PHASE SELECTION IN 1XXX ALLOYS 1114.1. Competitive growth 112

4.1.1. Transition from Fe4Al13 to FeAl6 1124.1.2. Transition from Fe4Al13 to FeAlx 1184.1.3. Transition from FeAL6 to FeAlm 1184.1.4. Transitions in Al±Fe±Si alloys 119

4.2. Competitive nucleation 1204.2.1. The transition from Fe4Al13 to FeAl6 1204.2.2. Promotion of nucleation of other phases 122

4.3. Suppression of equilibrium solidi®cation reactions 1234.4. Metastable phase diagrams and solidi®cation microstructure selection maps 124

Progress in Materials Science 43 (1998) 89±170

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* Corresponding author: Tel.: +44-1865-273774; fax: +44-1865-273764;

e-mail: [email protected].

5. FIR TREE FORMATION IN DC CASTS 1275.1. Fir tree zones 1275.2. Cooling rate 1285.3. Fir tree nucleation 1305.4. Fir tree nucleation and growth 1325.5. E�ect of solid fraction 135

6. FIR TREE PHASES IN DC CASTS 1377. TRANSFORMATION OF METASTABLE PHASES 140

7.1. The FeAl6 ÿ4 Fe4Al13 transformation 1417.1.1. Transformation mechanism and activation energy 141

7.1.1.1. Dissolution±precipitation mechanism and net activation energy 1417.1.1.2. Formation of acicular Fe4Al13 precipitates 1437.1.1.3. Continuous heating transformation 1437.1.1.4. Two step ageing 1457.1.1.5. Isothermal transformation 148

7.1.2. Transformation rate 1517.1.2.1. Microstructural scale 1517.1.2.2. E�ect of cold work 1547.1.2.3. Presence of pre-existing nuclei 154

7.2. The FeAlm ÿ4 Fe4Al13 transformation 1567.3. E�ect of Si on transformations of metastable Al±Fe phases 1577.4. Transformations involving ternary Al±Fe±Si phases 158

8. EFFECT OF IMPURITIES ON PHASE FORMATION IN Al±Fe AND Al±Fe±SiALLOYS 159

9. EFFECT OF GRAIN REFINER ADDITIONS ON Al±Fe AND Al±Fe±Si ALLOYS 1649.1. Proposed mechanisms of primary Al grain re®nement 1649.2. The roÃle of grain re®ners on secondary/ternary phase selection 166

10. SUMMARY 167REFERENCES 168

1. Introduction

Flat rolled aluminium products account for approximately 40% of the 24million tonnes annual world production of aluminium. These products arecommonly used for packaging and canning, in electrical applications (e.g.capacitor electrodes), architectural cladding, cable wrap, lithographic printing andautomotive sheet.

About 90% of ¯at rolled products are produced from the melt by the followingmanufacturing route: the melt is degassed, ®ltered and grain re®ned, then directchill (DC) cast into rectangular or cylindrical water cooled ring moulds withremovable bases. Fig. 1 shows a schematic of the DC casting process. Theremovable bases are withdrawn at a controlled rate as the metal solidi®es,resulting in the semicontinuous casting of rectangular ingots or cylindrical billets,typically 0.5±1 m in diameter and 5±10 m in length. The cast surface is oftenuneven, and the outermost00.2 m of the cast in from surface is often of a coarsergrain structure than the interior and can contain higher levels of segregates. Thecast surface is commonly scalped o� therefore, as discussed in Section 5.1, andthe remainder heat treated in the temperature range 450±6008C, in the form of apre-heat in order to e�ect microstructural homogenization prior to rolling.Homogenization reduces segregation, encourages the transformation of metastablesecondary and ternary phases into equilibrium phases, and acts to equilibrate solid

C.M. Allen et al. / Progress in Materials Science 43 (1998) 89±17090

solution levels of soluble elements, resulting in certain cases in the precipitation ofdispersoids. A series of both hot and cold rolls with intermediate annealingtreatments are then applied to produce a foil or sheet of the required ®nal gauge,typically in the range 6±150 mm (foil) or 150±3000 mm (sheet), which is thencommonly subjected to a ®nal anneal.

A range of di�erent aluminium alloys are DC cast and processed by the aboveroute. The exact compositions depend upon the ®nal application of the casting,but Cu, Zn, Mg, Mn, Si and Fe are common alloying additions. The alloys thatare the subject of this review are those designated AA1xxx by the InternationalAlloy Designation System (IADS). Commercial 1xxx series Al alloys containtypically R0.5 wt% Fe and R0.2 wt% Si, sometimes present as deliberate alloyingadditions, but also as impurities. Other common impurities are Cu, Cr, Mn, Mg,V and Zn. Al±Ti±B additions are frequently used to promote primary Al grainre®nement.

Fig. 1. Schematic of the vertical DC casting process (after Maggs).

C.M. Allen et al. / Progress in Materials Science 43 (1998) 89±170 91

The identity, size and distribution of the secondary and ternary inter-metallic phases are critical in¯uences on the material properties of the alloy [74],including strength, toughness, formability, fatigue resistance, corrosion resistanceand anodizing response [58]. Anodizing quality and etching response areespecially important in surface critical products such as lithographic printingsheet, as well as in sheet used in architectural applications. The solid solutioncontent is particularly important in controlling properties such as electricalconductivity and recrystallization characteristics. Thermodynamic consider-ations often fail to predict correctly the phase content and solid solutioncontent of the as-cast microstructure because of the non-equilibrium nature ofsolidi®cation during DC casting. The key alloy properties are controlled by solidsolution levels and secondary and ternary phase crystallography and morphology,which in turn are dependent on complex kinetic competitions for nucleation andgrowth.

In Sections 2 and 3 the wide range of both equilibrium and metastablesecondary Al±Fe and ternary Al±Fe±Si phases reported in 1xxx alloys areexamined.

2. Binary Al±Fe phases

The maximum equilibrium solid solubility of Fe in Al is very low, at 00.05wt% Fe, and Fe is usually present therefore in the form of secondary Fealuminide phases [74]. The maximum equilibrium solid solubility of Si in Al ishigher at 01.6 wt%, and low levels (00.1±0.2 wt%) of Si in the bulk are readilyaccommodated therefore by dissolution in the Al matrix and in the Fe aluminides.Consequently, the phase contents of DC cast Al±Fe and Al±Fe±Si alloys withR0.1 wt% Si are similar, although in the latter case the so called `binary' Fealuminides often contain dissolved Si. Ternary Al±Fe±Si phases, as reviewed inSection 3, only form at higher bulk concentrations of Si, typically >0.1 wt% Si inR0.2 wt% Fe containing alloys, and >0.2 wt% Si inR0.3±0.4 wt% Fe containingalloys.

2.1. The equilibrium Al±Fe4Al13 eutectic

Fig. 2 shows the equilibrium Al±Fe binary phase diagram. As shown in Fig. 2,the ®rst secondary phase to form on solidi®cation of dilute Al±Fe alloys underequilibrium conditions is given by the eutectic reaction;

Liquid ÿ4 aÿAl� Fe4Al13 �also denoted as FeAl3; or the y phase�The exact temperature and composition of the invariant point is of some debate,but Liang and Jones [65] have recently reported 655.120.18C at 1.8 wt% Fe,respectively. The eutectic temperature of 655.18C has subsequently been con®rmed


by Allen [4] using calorimetric methods. At the eutectic temperature the a-Almatrix has the maximum equilibrium solid solubility of Fe, at 00.05 wt% [74].The equilibrium secondary phase exists over a range of compositions, and is oftendenoted as having the stoichiometry of either FeAl3 or Fe4Al13. Black [15]determined from X-ray di�raction studies that Fe4Al13 has a c-face centredmonoclinic structure containing 100 atoms per unit cell [74]. Fig. 3a and b showsa typical TEM micrograph of needle shaped Fe4Al13 particles at the grainboundaries in a DC cast ingot and a typical [100] zone axis selected areadi�raction pattern (SADP) from a crystal of Fe4Al13 extracted from the Almatrix, respectively. Fe4Al13 commonly forms relatively large angular precipitatesin as-cast microstructures (Fig. 3a), which increase hardness but lead toembrittlement, reducing formability and fatigue resistance. As shown in Fig. 3b,Fe4Al13 exhibits spot streaking in certain zone axis di�raction patterns parallel tothe (00 l) reciprocal lattice vector, although it is not clear whether this is due tostacking faults or microtwinning on (001) [96, 97]. Fe4Al13 can also form pseudo10-fold twins, resulting from alternate repetition of (100) and (201) twins [56, 57].Fig. 4 shows a bright ®eld TEM micrograph of a 10-fold branched dendriticparticle present in a melt-spun Al±20 at% Fe alloy.

2.2. Metastable Al±Fe eutectic phases

As long ago as 1925 Dix [2] noted that fully eutectic microstructures could beattained in rapidly cooled alloys of Fe content well in excess of that of theequilibrium eutectic, 1.8 wt%. This indicated that large undercoolings are requiredto nucleate and/or grow the Al±Fe4Al13 eutectic under certain solidi®cation

Fig. 2. Al rich corner of the equilibrium Al±Fe binary phase diagram.


conditions. The nucleation and growth requirements of both the Al±Fe4Al13and metastable Al±Fe eutectics is discussed in further detail in Section 4.2

and 4.1, respectively. Under non-equilibrium solidi®cation conditions a range of

Fig. 3. (a) Fe4Al13 at grain boundaries in cast ingot. After Skjerpe [97]. Reproduced by kind permission

of Blackwell Science Ltd. (b). Typical [110] di�raction pattern of a faulted Fe4Al13 crystal. Faults on

{001} planes produce lines parallel to the {001} direction in reciprocal space. After Skjerpe [96].

Reproduced from Metallurgical and Materials Transactions by kind permission of TMS and ASM

international.


thermodynamically metastable Al±Fe eutectic phases that have smallerundercoolings for nucleation and growth than Al±Fe4Al13 can form in addition toAl±Fe4Al13. These are summarized in Table 1. The composition ranges of some ofthe metastable Al±Fe eutectics are shown in Fig. 5.

2.2.1. Metastable Al±FeAl6 eutecticHollingsworth et al. [42] were the ®rst to identify one of the metastable Al±Fe

eutectics displacing Al±Fe4Al13. They observed the displacement of Al±Fe4Al13 byAl±FeAl6 in continuously cast Al±2 wt% Fe. The exact solidi®cation conditions(cooling rate and solidi®cation velocity) at which this displacement occurs havesince been characterized by BaÈ ckerud [12], Adam et al. [1, 2], Hughes andJones [45], Liang and Jones [85], Gilgien et al. [35], Evans et al. [32] and Thomaset al. [105] and will be discussed further in Section 4.

Liang and Jones [65] report the eutectic temperature as 652.920.28C, with aeutectic composition of 3.0 wt% Fe. The crystal structure is c-face centredorthorhombic [96], with 28 atoms per unit cell [74]. FeAl6 is a commonconstituent of DC cast ingots and billets [112]. Fig. 6a and b show a typical TEMmicrograph of FeAl6 eutectic embedded in an Al matrix and thecorresponding [110] zone axis SADP. FeAl6 is also an important phase in Mn-containing alloys. MnAl6 and FeAl6 are isomorphous, and consequently Mn cansubstitute freely for Fe in the FeAl6 lattice, lowering its free energy. This raises thethermodynamic stability of the FeAl6 phase in Mn containing Al alloys (seeSection 8 and Table 6).

Fig. 4. TEM micrograph of 10-fold branched dendritic Fe4Al13 particle. After Kim and Cantor [56, 57].

Reproduced from Phil. Mag. A by kind permission of Taylor and Francis Ltd.


2.2.2. Metastable Al±FeAlm eutecticKosuge and co-workers [58] report that a metastable Al±FeAlm eutectic appears

at higher cooling rates (e.g. >10 K sÿ1 in wedge-shaped moulds of Al±0.6 wt%Fe) than those at which Fe4Al13 and FeAl6 form. This phase has also beenobserved in the more rapidly cooled zones of DC cast billets [96, 112]. Fig. 7a andb shows a typical TEM micrograph of a dendritic like FeAlm particle extractedfrom the Al matrix and the corresponding [110] zone axis SADP, respectively. Asshown in Fig. 7b, FeAlm exhibits incommensurate re¯ections in certain zone axisdi�raction patterns parallel to the (hh0) reciprocal lattice vector, due to stackingfaults on (110) planes [96, 98].

The eutectic temperature and composition for FeAlm in Al±Fe binaryalloys have not been determined. Allen et al. [5] have determined a eutectictemperature of 649.58C for FeAlm in Al±0.3 wt% Fe-0.1 wt% Si±0.05 wt% V,1.7 K lower than the eutectic temperature of 651.28C for Fe4Al13 measured inthe same alloy. The FeAlm phase exists over a range of compositions, with thevalue of m quoted in the range from 4.0 to 4.4. The crystal structure is bodycentred tetragonal [96], the unit cell containing in the region of 110 to 118atoms [98].

Table 1

Al±Fe phases formed in dilute Al±Fe alloys

Phase Bravais lattice Lattice parameters References

Fe4Al13 c-Centred monoclinic a=15.49 AÊ , Skjerpe [96, 97]

b=8.08 AÊ ,

c=12.48 AÊ ,

b=107.758

FeAl6 c-Centred orthorhombic a=6.49 AÊ ,

b=7.44 AÊ

c=8.79 AÊ

Hollingsworth et al. [42], BaÈ ckerud [12],

Jones [51], Adam and Hogan [1],

Hughes and Jones [45]

FeAlx (I) c-Centred orthorhombic a16 AÊ ,

b17 AÊ ,

Westengen [112], Skjerpe [96]

(x15.7±5.8)

c14.7 AÊ

FeAlx ? ? Young and Clyne [113] (x15), Evans et

al. [31] (x14.5), Wang et al. [111]

FeAlm Body centred tetragonal a=8.84 AÊ ,

b=c=31.6 AÊYoung and Clyne [113], Westengen [112],

Skjerpe [96], Skjerpe [98] (m14.0±4.4)

Fe2Al9 Monoclinic a=8.90 AÊ ,

b=6.35 AÊ ,

Simensen and Vellasamy [94], Brobak

and Brusethaug [16], Griger et al. [38]

c=6.32 AÊ ,

b=93.48

FeAlp Body centred cubic a= b= c=10.3 AÊ Ping et al. [80]


2.2.3. Metastable Al±FeAlx eutecticsWestengen [112] discovered another Al±Fe eutectic in DC cast material,

denoted Al±FeAlx, determining from EDX a stoichiometry of x15.8, withtraces (<1 wt%) of Si and Ni also being detected. Westengen was unableto determine the crystal structure of FeAlx from the irregular nature of itsdi�raction patterns, and suggested that it was heavily faulted. Skjerpe alsodetected this phase, with a similar stoichiometry of x15.7, and containing 1.9wt% Si and 0.3 wt% Ni. Fig. 8a and b show a typical TEM micrograph of anFeAlx particle and a typically irregular SADP, respectively [96, 97]. Skjerpeindexed the strongest intensity spots of his di�raction patterns to ®t a c-facecentred orthorhombic unit cell of cell parameters very similar to that of FeAl6(Table 1). HREM lattice imaging of FeAlx revealed that the di�raction patternsarise from a complex stacking sequence in real space. Given also the similarity instoichiometries (Fig. 5), Skjerpe suggested that FeAlx was a Si modi®ed version ofFeAl6. The eutectic temperature and composition of this phase have not beenreported.

Another phase also denoted FeAlx has been reported by Young and Clyne [113].Young and Clyne determined x15 from EDX data, and tentatively proposed amonoclinic crystal structure to ®t the XRD data they obtained from this phase. Ifthis structure is correct then this is not therefore the same FeAlx as observed byWestengen and Skjerpe. Young and Clyne's FeAlx displaced Fe4Al13 at coolingrates below that of FeAl6 during unidirectional solidi®cation experiments. Evanset al. [31] have similarly observed produced FeAlx at cooling rates intermediatebetween Fe4Al13 and FeAl6 during unidirectional solidi®cation, with a value of

Fig. 5. Compositions of the common binary and ternary compounds found in dilute Al±Fe±Si alloys.

After Langsrud [61].


Fig. 6. (a) FeAl6. After Westengen [112]. Reproduced by kind permission of Carl Hauser Verlag,

Munich, Germany. (b) [110] zone axis selected area di�raction pattern of FeAl6. After Westengen [112].

Reproduced by kind permission of Carl Hauser Verlag, Munich, Germany.


x14.5. The XRD trace from this phase could not be made to ®t the monoclinicstructure proposed by Young and Clyne however. Wang et al. [111] have sinceproduced a fully eutectic microstructure of Al±FeAlx in Al±3 wt% Fe±0.1 wt% Valloys directionally solidi®ed at velocities in the range 0.09±1.03 mm sÿ1, with thesame XRD trace as Evans's FeAlx, and have shown that Young and Clyne'sstructure determination was incorrect [33]. The FeAlx eutectics produced byYoung and Clyne, Evans et al. and Wang et al. are the same phase therefore, withx14.5±5.0, and in turn are di�erent to the FeAlx eutectics produced byWestengen and Skjerpe, with x15.7±5.8.

Fig. 7. FeAlm and corresponding [110] di�raction pattern. Stacking faults on {hh0} planes lead to

incommensurate re¯ections along the {hh0} direction in reciprocal space. After Skjerpe [96].

Reproduced from Metallurgical and Materials Transactions by kind permission of TMS and ASM

International.


Fig. 8. (a) FeAlx. After Skjerpe [96]. Reproduced from Metallurgical and Materials Transactions by

kind permission of TMS and ASM International. (b) Di�raction pattern from FeAlx showing

incommensurate nature of re¯ections. After Skjerpe [96]. Reproduced from Metallurgical and Materials

Transactions by kind permission of TMS and ASM International.


2.2.4. Metastable Al±Fe2Al9 eutecticFig. 9 shows a typical TEM micrograph of Fe2Al9 eutectic embedded in an Al

matrix, obtained in strip cast Al±0.5 wt% Fe±0.2 wt% Si alloy [94]. Thestoichiometry of this phase was determined by EDX. Analysis of electrondi�raction patterns indicated a monoclinic crystal structure. The solidi®cationconditions under which this phase form are unclear. Tezuka and Kamio [104]noted that in DC cast Al±0.5 wt% Fe, additions of >0.075 wt% Co promotedthe formation of the (Fe,Co)2Al9 phase. Fe2Al9 and Co2Al9 are isomorphous, andconsequently Co can substitute freely for Fe in the Fe2Al9 lattice, lowering its freeenergy. This raises the thermodynamic stability of the Fe2Al9 phase in Cocontaining Al alloys (see Section 8 and Table 6).

2.2.5. Metastable Al±FeAlp eutecticPing et al. [80] reported the formation of a metastable body centred cubic phase

FeAlp (where p14.5) in directionally chill cast Al±(0.25±0.50) wt% Fe±0.125wt% Si. This phase has yet to be observed independently.

2.2.6. The e�ect of Si addition on the formation of Al±Fe phasesAs stated in Section 2, small quantities of Si (typically <0.1 wt% Si bulk

composition inR0.2 wt% Fe containing alloys, andR0.2 wt% Si inR0.3±0.4 wt%Fe containing alloys) can be dissolved into the `binary' Fe aluminides. However,these aluminides have di�erent degrees of Si solubility (Fig. 5). Consequently, the

Fig. 9. Fe2Al9. After Simensen and Vellesamy [94]. Reproduced by kind permission of Carl Hauser

Verlag, Munich, Germany.


occurrence of FeAl6 which can only dissolve up to 0.5 wt% Si in its lattice isrestricted in Al±Fe±Si alloys. FeAl6 is replaced by Fe aluminides that canincorporate Si, such as Fe4Al13 or FeAlm [61].

3. Ternary Al±Fe±Si phases

3.1. The equilibrium a-AlFeSi and b-AlFeSi phases

Three ternary phases form under equilibrium solidi®cation conditions in diluteAl±Fe±Si alloys of su�ciently high bulk Si content, >0.1 wt% Si inR0.2 wt% Fecontaining alloys, and >0.2 wt% Si in R0.3±0.4 wt% Fe containing alloys, attemperatures below that of the liquid4 Al+Fe4Al13 eutectic reaction. Fig. 10shows the liquidus projection and associated equilibrium solidi®cation reactions inthe Al corner of the Al±Fe±Si ternary phase diagram. The three equilibriumternary phases produced by either of two ternary peritectic reactions followed by aternary eutectic reaction are;

i. Liquid+Fe4Al13 ÿ4 Al+Fe2SiAl8 (also denoted as the a phase);ii. Liquid+Fe2SiAl8 ÿ4 Al+FeSiAl5 (also denoted as the b phase); and/oriii. Liquid ÿ4 Al+Si+FeSiAl5.A range of temperatures have been measured for these three invariant points in

the ternary phase diagram: 620±6388C for the a peritectic, 611±6158C for the bperitectic and 576±5778C for the ternary eutectic [14, 74, 85]. These ranges mayre¯ect a di�culty in nucleating or growing one or more of these phases duringsolidi®cation. This point is discussed in Section 4.3. Both phases exist over a rangeof compositions, as shown in Fig. 5. The accepted stoichiometries as given aboveare those of Mondolfo [74].

Both the a and b phases can adopt a number of di�erent crystal structures.Table 2 summarizes the structural variants of these two phases. Munson [75]determined from X-ray di�raction studies that a-AlFeSi has a hexagonal crystalstructure (Table 2), in agreement with earlier single crystal studies performed byRobinson and Black [86], and this was con®rmed in the same year by Sun andMondolfo [102]. The hexagonal crystal structure of a is also known in theliterature as a 0 [86] or a2 [9, 10]. Fig. 11a and b show a typical TEM micrographof an a 0 particle in DC cast 1050 alloy and corresponding [100] zone axis SADP,respectively [112]. Dons [29] observed that a 0 survived heat treating in >99.9 wt%pure Al based DC cast commercial purity alloys, to progressively highertemperatures with increasing bulk Si content of the alloy, from 4508C in Al±0.6wt% Fe±0.15 wt% Si, to 6008C in Al±0.6 wt% Fe±0.6 wt% Si.

Phragmen [79] determined that b-AlFeSi has a monoclinic crystal structure(Table 2). b-AlFeSi is an important ternary phase in wrought aluminium alloys.Fig. 12 shows a typical SEM micrograph of b platelets on a deep etched surfaceof a DC cast alloy, showing their characteristic long thin curving morphology,which can dramatically reduce ductility.


Table 2

Common structural variants of the ternary phases Fe2SiAl8 (a) and FeSiAl5 (b)

Phase Bravail lattice Lattice parameters References

a(a1) Cubic a=12.56 AÊ (Im3)

a=12.52 AÊ (Pm3)

Cooper [23], Munson [75], Sun and

Mondolfo [102], Westengen [112], Griger

et al. [37], Turmezey et al. [109]

a 0 (a2) Hexagonal a= b=12.3 AÊ

c=26.2 AÊDons [29], Munson [75], sun and

Mondolfo [102], Griger et al. [37],

Thoresen et al. [106]

av Monoclinic a=8.90 AÊ , Dons [29]

b=6.35 AÊ ,

c=6.32 AÊ ,

b=93.48a0 (q1) c-Centred orthorhombic a=12.7 AÊ ,

b=26.2 AÊWestengen [112], Skjerpe [96], Ping et

al. [80], Ping [82]

c=12.7 AÊ

q2 Monoclinic a=12.50 AÊ , Ping et al. [80]

b=12.30 AÊ ,

c=19.70 AÊ ,

b=1118aT c-Centred monoclinic a=27.95 AÊ ,

b=30.62 AÊ ,

Dons [29], Skjerpe [96], Jensen and

Wyss [50], Turmezey et al. [109], Ping [82]

c=20.73 AÊ ,

b=97.748b Monoclinic a=6.12 AÊ Skjerpe [96]

b=6.12 AÊ ,

c=41.5 AÊ ,

b=918b 0 Monoclinic a=8.9 AÊ , Westengen [112], Skjerpe [96]

b=4.9 AÊ ,

c=41.6 AÊ ,

b=928

Fig. 10. Liquidus surface and associated equilibrium phase ®elds in the Al corner of the ternary Al±Fe±

Si phase diagram. After Skjerpe [96]. Reproduced from Metallurgical and Materials Transactions by

kind permission of TMS and ASM International.


Fig. 11. (a) Hexagonal a 0 Al±Fe±Si. After Westengen [112]. Reproduced by kind permission of Carl

Hauser Verlag, Munich, Germany. (b) [100] zone axis selected area di�raction pattern from a 0. After

Westengen [112]. Reproduced by kind permission of Carl Hauser Verlag, Munich, Germany.


3.2. Metastable a-AlFeSi and b-AlFeSi phases

Under non-equilibrium solidi®cation conditions, the liquid becomes enriched inSi, due to partitioning of the Si to the interdendritic liquid, and hence there is agreater tendency for ternary phases to form [61]. In addition, a number ofmetastable structural variants of both the a- and b-AlFeSi ternary phases arecommonly observed in commercial DC cast alloys, and are summarized in Table 2.

3.2.1. Metastable cubic a-AlFeSi phaseMunson [75] and Sun and Mondolfo [102] determined that the equilibrium

hexagonal form of a-AlFeSi was only thermodynamically stable in high purity Al±Fe±Si alloys. Additions of V, Cr, Mn, Cu, Mo and W all promote a body-centredcubic structure for the a-AlFeSi phase, also known in the literature as c [79],a1 [9, 10] or simply a [78]. Additions of Ti, Ni, Zn and Mg do not promote thecubic structure [102]. The cubic structure had also been previously observed byPhragmen [79] and Cooper [23], but had incorrectly been assumed to represent theequilibrium crystal structure for the a-AlFeSi phase. The cubic structure isisostructural with a-AlMnSi. Only 0.1 wt% Mn is required therefore to stabilizethe cubic form during solidi®cation at a cooling rate of 0.75 K minÿ1 [75].

The stabilization of the cubic form by trace elements common to commercialpurity alloys results in the cubic form being the one which is most commonlyobserved in commercial alloys. Westengen [112] observed cubic a in DC castAA1050 alloy in the more rapidly cooled outer zone of the billet. Weakh+ k+ l= odd integer spots were observed in the di�raction patterns, indicatingthat the structure may not have been body centred but primitive cubic. Dons [29]observed cubic a in both as-cast and heat-treated DC cast commercial purity

Fig. 12. SEM micrograph of b Al±Fe±Si platelets, in surface etched DC cast alloy. After Griger et

al. [37]. Reproduced by kind permission of Aluminium.


aluminium of Fe/Si ratio <1, which survived to progressively higher heattreatment temperatures with increasing bulk Si content. Fig. 13a and b shows atypical TEM micrograph of a cubic a particle extracted from a DC cast Al±0.25wt% Fe±0.13 wt% Si alloy, with a partly dendritic morphology, and acorresponding [111] zone axis SADP, respectively [97]. Skjerpe observed cubic a inthe more rapidly cooled outer 50 mm of the billet. Griger et al. [37] observedcubic a in semicontinuously cast Al±0.5 wt% Fe±0.2 wt% Si, across the entirecross-section of the billet. Turmezey et al. [109] observed that the Si content of thecubic a phase was directly proportional to the bulk Si content, suggesting thatdirect Alt Si substitution can take place in the cubic a lattice. Thoresen etal. [106] investigated Al±4 wt% (Fe, Mn)±7.5 wt% Si alloys, with varying Fe:Mnratios, in which the primary phase was a-AlFeSi, either in its cubic or hexagonalform, depending upon the bulk Mn content of the alloy [75]. The total transitionmetal content (Fe+Mn) of the cubic a phase in the Al±(Fe, Mn)±Si alloy wasless than the cubic a phase in the Al±Mn±Si alloy, indicating that vacanciesstabilise the cubic a structure when both Fe and Mn are present.

3.2.2. Metastable av-AlFeSi phaseDons [29] observed a monoclinic structural variant of a-AlFeSi in DC cast

commercial purity Al±0.2 wt% Fe±0.2 wt% Si, denoted av. Dons stated that avwas structurally related to the Fe2Al9 phase [94], the a-axis being 2.6% shorterand the c-axis 3.6% shorter than the monoclinic structure of Fe2Al9. However, theSi content of av was in the range 4.5±10.5 wt% (corresponding to the Si contentrange typically observed in a 0 and a), signi®cantly higher than the maximum Sicontent of02 wt% seen in Fe2Al9.

3.2.3. Metastable a0 or q1-AlFeSi phaseFig. 14a and b show a typical TEM micrograph of a0 particles embedded in a

DC cast AA1050 alloy and a corresponding [100] zone axis SADP,respectively [112]. Westengen observed that a0 was closely related to the cubic aform, indexing the a0 di�raction patterns as originating from a tetragonal unitcell. EDX data showed that a0 had a lower Si content than cubic a. Westengentherefore suggested that a0 was a low Si modi®cation of cubic a, as illustrated inFig. 5. Fig. 15 summarizes the subsequent EDX measurements made bySkjerpe [96] on particles in DC cast Al±0.25 wt% Fe±0.13 wt% Si, whichsupported the idea that a0 was a low Si modi®cation of cubic a. Ping et al. [80]also observed the a0 phase in DC cast Al±0.28 wt% Fe±0.13 wt% Si, but denotedit q1, which formed at a cooling rate of 010 K sÿ1. Detailed convergent beamelectron di�raction analysis by Ping and co-workers [80, 81] of a0 revealed a c-facecentred orthorhombic structure, which was also con®rmed by [96].

3.2.4. Metastable aT-AlFeSi phaseDons [29] observed a further structural variant of a-AlFeSi in DC cast

commercial purity Al±0.2 wt% Fe±0.2 wt% Si, denoted aT, whose crystal


Fig. 13. (a) Cubic a-Al±Fe±Si. After Skjerpe [97]. Reproduced from Metallurgical and Materials

Transactions by kind permission of TMS and ASM International. (b) [111] zone selected area

di�raction pattern cubic a-Al±Fe±Si. After Skjerpe [97]. Reproduced from Metallurgical and Materials



structure, although at that time unidenti®ed, appeared to be related to that ofcubic a. Fig. 16a and b shows a typical TEM micrograph of aT and cubic aparticles (marked A and B, respectively, in the ®gure) and a corresponding SADPfrom the aT particle [96]. Both Dons and Skjerpe [96] noted that certain zone axisdi�raction patterns of aT consisted of super-re¯ections imposed on the basic bcc apattern, and that other zone axis patterns from the two phases could not be

Fig. 14. (a) Tetragonal a0-Al±Fe±Si. After Westengen [112]. Reproduced by kind permission of Carl

Hauser Verlag, Munich, Germany. (b) Typical [100] di�raction pattern of a tetragonal a0 crystal. After



distinguished from each other. Given the chemical similarity between aT and a(Fig. 15), Skjerpe noted that these two phases can be confused easily. Skjerpeindexed the super-re¯ections to a previously proposed c-face-centred monoclinicunit cell [41]. Jensen and Wyss [50] proposed that aT was not a distinct phase, butwas a structural modi®cation of cubic a. They were not, however, able to identifythe speci®c structural feature giving rise to the observed super-re¯ections.Turmezey et al. [109] proposed that the super-re¯ections arise from thearrangement of the cubic phase into domains, of the order of0100 AÊ in diameter,and therefore irresolvable using conventional TEM. Ping et al. [80] disagreed withthis interpretation, arguing that the domain structure could not account for all theobserved super-re¯ections, and went on to use CBED techniques [82] to determinethat the extra re¯ections, with a rhombohedral superstructure, arose from anordering of vacancies on the Fe atom sites. As Ping noted, irradiation by a TEMbeam for 10 min removed this long range ordering, and the super-re¯ections werelost.

3.2.5. Metastable q2-AlFeSi phaseFig. 17 shows a typical TEM micrograph of q2 particles embedded in the Al

matrix from the outer zone of a directionally chill cast Al±0.28 wt% Fe±0.13 wt%

Fig. 15. EDX analyses of crystals extracted from DC cast alloy showing similarity in values for a-Al±

Fe±Si, a0-Al±Fe±Si and aT-Al±Fe±Si. After Skjerpe (1987 (II)). Reproduced from Metallurgical and

Materials Transactions by kind permission of TMS and ASM International.


Fig. 16. (a) aT-Al±Fe±Si (marked A) and cubic a-Al±Fe±Si (marked B). After Skjerpe [96]. Reproduced

from Metallurgical and Materials Transactions by kind permission of TMS and ASM International. (b)

Typical di�raction pattern of an aT-Al±Fe±Si crystal, showing super-re¯ections imposed on a cubic apattern. After Skjerpe [96]. Reproduced from Metallurgical and Materials Transactions by kind

permission of TMS and ASM International.


Si billet which had then been heat treated for 24 h at 6008C [80, 82]. This phasearose from the transformation of the a0 or q1 phase (Section 3.2.3).

3.2.6. Metastable b 0-AlFeSi phaseFig. 18 shows a typical TEM micrograph of b 0 and Si particles (single and

double headed arrows in ®gure, respectively) embedded in the Al matrix of a DCcast AA1050 [112]. The b` phase exhibited planar defects and bore a resemblanceto the b phase, the c axes of the two phases being very similar. EDX datasuggested that b 0 may be a low Si modi®cation of b.

4. Factors governing phase selection in 1xxx alloys

As noted in Section 1, due to the non-equilibrium nature of solidi®cationduring DC casting thermodynamics are usually not capable of predicting thephase and solution contents of the as cast microstructure. An understanding ofthe factors that govern phase selection in 1xxx Al alloys under conditions of non-equilibrium solidi®cation is important, since varying solidi®cation conditions canlead to variations in secondary Al±Fe and ternary Al±Fe±Si phase contents atdi�erent positions in the casting, which in turn can lead to a degradation in themechanical and surface properties of the ®nal rolled sheet or foil. The e�ect of

Fig. 17. q2-Al±Fe±Si. After Ping et al. [80]. Reproduced by kind permission of the Institute of

Materials.


phase content variations on surface properties is discussed in more detail inSection 5.1.

Features of non-equilibrium solidi®cation include: competitive growth of twophases; competitive nucleation of two phases; suppression of solidi®cationreactions involving solid state di�usion of slowly di�using species over largedistances; and non-equilibrium solute partitioning. These points are considered inturn below in Section 4.1±4.4, respectively. The role of alloy chemistry, speci®callyof impurities and grain re®ning additions, is considered in Sections 8 and 9.2,respectively.

4.1. Competitive growth

4.1.1. Transition from Fe4Al13 to FeAl6One phase, say A, can be kinetically displaced by another, B, if the growth

temperature (Tg) of A is depressed to below that of B, i.e. if

Tg;A < Tg;B

assuming that both phases can nucleate under the given solidi®cation conditions.Assume that this displacement occurs either at some critical growth velocity Ucrit

or at some critical cooling rate (dT/dt)crit. For a given temperature gradient across

Fig. 18. b 0 Al±Fe±Si (single headed arrow) and Si (double headed arrow) particles. After



the solid/liquid interface, G, the cooling rate dT/dt and the growth velocity U arerelated under one dimensional steady-state growth conditions by:

dT

dt� dT

dx� dxdt� G �U �1�

so that the critical cooling rate and solidi®cation velocity for the displacement ofphase A by phase B are related by:

dT

dtcrit� G �Ucrit �2�

Which of the two parameters out of (dT/dt)crit and Ucrit is the more fundamentalin determining whether the transition from phase A to phase B will occur willdepend upon the growth kinetics of the two phases.

BaÈ ckerud [12] studied the variation of eutectic growth temperatures Tg withcooling rate for Al±Fe4Al13 (Tg,EU1) and Al±FeAl6 (Tg,EU2) in Al±0.5±4.0 wt% Fealloys. Fig. 19 shows his experimental variation of Tg with the square root ofcooling rate for both the eutectic phases. He observed that, with increasingcooling rate during solidi®cation, both Tg,EU1 and Tg,EU2 decreased, but bydi�erent amounts. Below the critical cooling rate of (dT/dt)crit=3.3 K sÿ1,Tg,EU1>Tg,EU2, and above the critical cooling rate of (dT/dt)crit=3.3 K sÿ1,Tg,EU1<Tg,EU2. BaÈ ckerud did not report the corresponding growth velocity atwhich the transition from Al±Fe4Al13 eutectic to Al±FeAl6 occurred, nor indicatewhether cooling rate or solidi®cation velocity was the more fundamental indetermining whether the transition occurred.

Adam and Hogan [1, 2] in unidirectional solidi®cation experiments on Al±2.0±4.0 wt% Fe alloys examined the variation in preferred eutectic phase and themorphology of that phase with changing growth velocity U and temperature

Fig. 19. Variation of eutectic growth temperature for Fe4Al13 and FeAl6 with the square root of

cooling rate. After BaÈ ckerud [12].


gradient G at the solidi®cation front. Fig. 20 shows their experimental variation ofeutectic microstructure observed with changing U and G. At high G/U, thepreferred eutectic is Al±Fe4Al13, which grows with a lamellar morphology. Thelong platelet axis of the lamellae parallel to the growth direction was [010](Fe4Al13). The aluminium matrix grew with [100] (Al) parallel to the growthdirection, with a 2158 divergence. No consistent orientation relationship betweencrystallographic planes of the Fe4Al13 and the Al was observed, however. As G/Udecreases, the eutectic becomes increasingly branched and the spacing between theFe4Al13 particles decreases. For constant G/U, the value of G determines thedegree of branching. Conversely for constant G/U, changing G does not alter thespacing between the Fe4Al13 particles. A further decrease of G/U is accompaniedby a change in the preferred eutectic, from a faceted Fe4Al13 morphology to a rodlike FeAl6 morphology. The cooling rate at which this occurs in an Al-2 wt% Fealloy of 2.5 K sÿ1 is close to BaÈ ckerud's earlier measurement of 3.3 K sÿ1.However, as shown in Fig. 20 the transition occurs over only a small range ofsolidi®cation velocities and is relatively independent of G. The value of Ucrit

Fig. 20. Growth rate and temperature gradient conditions imposed on unidirectionally solidi®ed Al±2

wt% Fe alloys and the resultant eutectic reactions. After Adam and Hogan [1]. Reproduced by kind

permission of the Institute of Materials Engineering Australasia.


determined by Adam and Hogan was 0100 mm sÿ1. With increasing solidi®cationvelocity therefore the undercooling required by Al±Fe4Al13 for growth becomesgreater than that required by Al±FeAl6, and Al±Fe4Al13 can be kineticallydisplaced by the metastable Al±FeAl6 eutectic. As indicated above, this transitioncan only take place if nucleation conditions permit. There is no unique coolingrate for the transition however, and from Eq. (2) it can be seen that the value of(dT/dt)crit measured will depend upon the value of G in each case.

The FeAl6 rods grow parallel to the [001]FeAl6 axis, and Adam andHogan [1] determined that they have a unique orientation relationship with the Almatrix:

��100�FeA16==��111�A1

�130�FeA16==��11�1�A1

�310�FeA16==�002�A1

This interface is semicoherent. Hughes and Jones [45] observed an increasedtendency of the FeAl6 rods to cross-link with decreasing U.

Hughes and Jones [45] observed a displacement of Fe4Al13 by FeAl6 at agrowth velocity of 0.10 mm sÿ1 in unidirectionally solidi®ed Al±Fe of Fe contentin the range 2.2±6.1 wt%. The critical velocity was independent of Fe alloyingcontent. This velocity equated to a cooling rate of 2 K sÿ1, in reasonableagreement with earlier measurements by Adam and Hogan and BaÈ ckerud detailedabove.

Later work by Liang and Jones [65] arrived at a similar critical velocity valueof 0.11 mm sÿ1. Fig. 21 shows their experimental variation of Tg with thesquare root of solidi®cation velocity for both the Al±Fe4Al13 (Eu1) andAl±FeAl6 (Eu2) eutectics. The data is plotted similarly to that of BaÈ ckerud, exceptthat the x ordinate is solidi®cation velocity instead of cooling rate, the formerdetermining the solidi®cation conditions under which one eutectic will displaceanother as shown by Adam and Hogan above. For both eutectics theexperimental variation of growth temperatures, Tg, with solidi®cation velocity, U,®t the form;

Tg � TEU ÿ BU1=2 �3�where TEU is the extrapolated equilibrium solidi®cation temperature (i.e. Tg atU=0) and B is a constant that varies from eutectic to eutectic. A considerationof the parameters that determine B is beyond the scope of this review, but it is acomplex function of the eutectic interphase spacing corresponding to growth atminimum undercooling or maximum velocity, the liquidus slopes of the eutecticphases near the eutectic composition, the volume fractions of the eutectic phases,the solute concentration gradient between the eutectic phases, the interfacialenergies between the eutectic phases, and between each of those phases and themelt, the heats of fusion of the eutectic phases, the solute di�usivity in the melt,and the contact angles between each of the eutectic phases and the melt. The


extrapolated equilibrium solidi®cation temperatures (i.e. TEU) of Al±Fe4Al13 andAl±FeAl6 are 655.120.18C and 652.920.28C respectively [35, 65]. Theextrapolated TEU of Al±Fe4Al13 agrees well with that from the equilibrium phasediagram. Allen [4] has recently determined the eutectic temperatures of Al±Fe4Al13and Al±FeAl6 calorimetrically, by remelting unidirectionally solidi®ed samples at2 K minÿ1, to be 655.120.38C and 653.320.58C, respectively, again in excellentagreement with Liang and Jones's reported values.

Evans et al. [32] observed the variation of phase content in Al±Fe based alloysat solidi®cation velocities <2 mm sÿ1. Fig. 22 shows the qualitative variation inphase content with solidi®cation velocity in Al±0.5±3.0 wt% Fe and the e�ectof Si and grain re®ner addition to Al±0.3 wt% Fe. In hypoeutectic alloys(Al±0.5±1.5 wt% Fe in ®gure) they observed the coexistence of Al±Fe4Al13 andAl±FeAl6 over the solidi®cation velocity range 00.2±0.6 mm sÿ1, independent ofFe alloying content. Fe4Al13 appeared to grow at solidi®cation velocities >0.1 mmsÿ1, unlike in hypereutectic compositions (Al±3.0 wt% Fe in ®gure). Continuingthis work, Thomas et al. [105] have recently demonstrated that this observed

Fig. 21. Experimental determinations of variation of growth temperatures with solidi®cation velocity

for Al±Fe4Al13 eutectic (Eu1) and Al±FeAl6 eutectic (Eu2). After Liang and Jones [65]. Reproduced by

kind permission of Carl Hauser Verlag, Munich, Germany.


mixture of Fe4Al13 and FeAl6 phases over the solidi®cation velocity range 00.17±0.5 mm sÿ1 results from partial transformation of the FeAl6 to Fe4Al13 in thesolid state during cooling following solidi®cation (see Section 7.1.2.3), and doesnot result directly from competitive growth driving solidi®cation of the twophases. In other words, at the growth front FeAl6 displaces Fe4Al13 at growthvelocities Ucrit>00.17 mm sÿ1, similar to the values reported for hypereutecticalloys (00.10±0.11 mm sÿ1). Behind the growth front however the as-solidi®edFeAl6 does not survive and transforms to Fe4Al13, the proportion transformingdecreasing with increasing growth velocity and hence decreasing time available fortransformation, until the growth velocity exceeds 00.6 mm sÿ1 and all the FeAl6survives.

The requirement for increased undercooling for growth of the Al±Fe4Al13eutectic phase at high solidi®cation velocities compared with that of Al±FeAl6 hasbeen proposed to be due to di�erences in the eutectic morphology and growthmechanism between the two eutectics. From the Hunt and Jackson classi®cationof eutectics [47] the morphology of the eutectic adopted during growth dependsupon the entropies of melting of both of the eutectic components. Al±FeAl6 isclassi®ed as a non-faceted/non-faceted eutectic, i.e. both Al and FeAl6 have lowentropies of melting. The eutectic grows accordingly with a ®brous rod-likemorphology. Al±Fe4Al13, however, is classi®ed as a faceted/non-faceted eutectic,i.e. the Fe4Al13 phase of the eutectic has a high entropy of melting and thereforegrows with a faceted, plate-like morphology.

Adam and Hogan [1, 2] determined experimentally the following relationshipsbetween the interphase spacing, l (in mm) and the growth velocity, U (in mm sÿ1),for the Al±Fe4Al13 and Al±FeAl6 eutectics in hypereutectic alloys;

l2:6U � 318 mm3:6 sÿ1 for Al±Fe4Al13 �4�

Fig. 22. Qualitative diagram of phase selection with changing growth velocity and composition. After

Evans et al. [32]. Filled space, Fe4Al13; hollow space, FeAl6; shaded space, FeAlm. Reproduced by kind

permission of PV Evans.


and

l2U � 10:2 mm3 sÿ1

for Al±FeAl6 �5�

Tiller [107] proposed that for the relationship l nU=constant, the value of nwas related to the rate determining step for eutectic growth. For n=2 (i.e. as forAl±FeAl6) di�usion of the solute to or from the growing interface is the ratedetermining step. For n>2 (i.e. as for Al±Fe4Al13) atomic attachment to thegrowing eutectic is the rate determining step. In other words, with increasingsolidi®cation velocity, and hence reduced time for growth, a greater degree ofundercooling is required to overcome the barrier to atomic attachment to (andhence growth of) the faceted Fe4Al13 morphology than to the ®brous FeAl6morphology [37].

4.1.2. Transition from Fe4Al13 to FeAlxYoung and Clyne [113] used unidirectional solidi®cation, with thermal gradients

of 3±8 K mmÿ1 on hypoeutectic alloys of unspeci®ed purity, to study thevariation of phase content with solidi®cation velocity. Fig. 23 summarizes their®ndings, converting their results into cooling rate ranges using Eq. (1) (Section4.1.1). They observed the displacement of Fe4Al13 by the FeAlx phase (Section2.2.3) at 0.7±1 K sÿ1, suggesting a displacement at 00.15±0.2 mm sÿ1. Theyclaimed that this was an intermediate stage in the transition of Fe4Al13 to FeAl6,and that FeAlx had previously been mistaken for FeAl6 due to their similarappearance. Young and Clyne proposed that Fe4Al13 was detected at cooling ratesup to 00.9 K sÿ1, FeAlx in the range 0.5±6 K sÿ1, FeAl6 at cooling rates >3 Ksÿ1, and FeAlm at >10 K sÿ1 (from the ®ndings of Kosuge and co-workers(Section 2.2.2)). However, the appropriateness of assigning ®xed cooling ratesto these transitions has already been questioned in Section 4.1.1 above.Additionally, there are no other reports in the literature of the appearance ofFeAlx at cooling rates intermediate between those of Fe4Al13 and FeAl6.Shillington [92] has however produced FeAlx by unidirectional solidi®cation of afully molten specimen at all growth velocities investigated R2 mm sÿ1.The uncertainty in the solidi®cation parameters necessary to reproducibly growFeAlx indicates that the occurrence of this phase is sensitive to the nucleationconditions present preceeding growth. This is discussed in further detail in Section4.2.2.

4.1.3. Transition from FeAl6 to FeAlmThere is general agreement in the literature that under more extreme

solidi®cation conditions FeAl6 becomes kinetically displaced by FeAlm. However,there does not appear to be a unique critical cooling rate for this transition,which was the same for the case for the displacement of Fe4Al13 by FeAl6.Miki et al. [72] observed FeAlm at cooling rates >10 K sÿ1 in Al±0.6 wt% Fecast in wedge moulds, and Kosuge and Mizukami [59] at >20 K sÿ1, in the samealloy unidirectionally solidi®ed. Asami et al. [59] were unable to produce FeAlm


at cooling rates below 35 K sÿ1 in Al±0.5 wt% Fe. Few controlled growthrate experiments have concentrated on growing FeAlm. However, Evans et al. [32]have recently observed the appearance of FeAlm at >1.33±2 mm sÿ1 inunidirectionally solidi®ed Al±0.3 wt% Fe, but only in the presence of 0.1 wt% Siand Ti:B grain re®ner addition (as previously shown in Fig. 22). Lockyer [66] hasalso observed the appearance of FeAlm in twin roll cast Al±0.55 wt% Fe±0.15wt% Si.

4.1.4. Transitions in Al±Fe±Si alloysIn Al±Fe±Si alloys, the possibilities exist that Al±Fe phases can be displaced by

ternary Al±Fe±Si phases with increasing solidi®cation velocity or cooling rate, orthat the values of the critical solidi®cation velocities or cooling rates for thetransitions (Section 4.1.1±4.1.3) between Al±Fe phases can be altered.

In unidirectionally solidi®ed Al±0.29 wt% Fe±0.17 wt% Si Brobak andBrusethaug [16] noted the displacement of Fe4Al13 by a-AlFeSi over the range01±2 mm sÿ1, equivalent to 05.2±5.5 K sÿ1. In semicontinuously cast Al±0.5

Fig. 23. E�ect of cooling rate on formation of Al±Fe eutectics in hypoeutectic Al±Fe alloys. After

Young and Clyne [113]. Reproduced by kind permission of Elsevier Science Ltd.


wt% Fe±0.2 wt% Si, Griger et al. [37] noted the displacement of Fe4Al13 by a-AlFeSi at cooling rates of >2.3 K sÿ1. Griger et al. interpreted this displacementto be due to the di�erences in eutectic morphology between Fe4Al13 and a-AlFeSi,an argument similar to that which had been used to explain the displacement ofFe4Al13 by FeAl6 (see Section 4.1.1). For the non-faceted a-AlFeSi phase di�usionof solute is the rate determining step for growth. For faceted Fe4Al13 atomicattachment is the rate determining step. With increasing solidi®cation velocity,and hence reduced time for growth, a greater degree of undercooling is requiredto overcome the barrier to atomic attachment to (and hence growth of) thefaceted Fe4Al13 morphology than to the non-faceted a morphology, and adisplaces Fe4Al13. Griger et al. [37] also used a similar argument to propose thatthe faceted b-AlFeSi phase would be displaced during growth at high solidi®cationvelocities by the non-faceted a-AlFeSi.

As noted in Section 4.1.3, Evans et al. [32] have observed the appearance ofFeAlm at >1.33±2 mm sÿ1 in unidirectionally solidi®ed Al±0.3 wt% Fe±0.1 wt%Si, but only in the presence of Ti:B grain re®ner addition. In Al±0.3 wt% Fe±0.1wt% Si with or without grain re®ner addition, Fe4Al13 and FeAl6 are seen tocoexist over the solidi®cation velocity range 00.17±1.33 mm sÿ1, although thisapparent coexistence may arise from the solid state transformationFeAl64Fe4Al13 during cooling, as was proposed to occur by Thomas et al. [105]in hypoeutectic Al±Fe alloys.

4.2. Competitive nucleation

4.2.1. The transition from Fe4Al13 to FeAl6As was seen in Section 4.1, most experiments to date have concentrated upon

establishing the critical cooling rates or critical solidi®cation velocities for thetransition of one eutectic phase to another. In the case of the transition from Al±Fe4Al13 eutectic to Al±FeAl6 eutectic it has been shown in Section 4.1.1 thatsolidi®cation velocity and not cooling rate is the more fundamental parameter. Itwas assumed therefore that as competitive growth accounts for the displacementof Al±Fe4Al13 by Al±FeAl6 that both phases can nucleate relatively easily at lowundercoolings.

BaÈ ckerud [12] had originally advanced a nucleation based argument however toexplain the observed transition of Al±Fe4Al13 eutectic by Al±FeAl6 eutectic onincreasing cooling rate. One phase, say A, can also be kinetically displaced byanother, say B, if the nucleation temperatures are such that

Tn;A < Tn;B

assuming that both phases can grow under the given solidi®cation conditions. Inhis study of the variation of preferred eutectic phase with cooling rate in Al±0.5±4.0 wt% Fe alloys, BaÈ ckerud proposed that the Al±Fe4Al13 eutectic was not thepreferred eutectic phase at high cooling rates (as previously shown in Fig. 19). Heproposed that it required a high undercooling to nucleate, irrespective of whether


it was the primary phase or secondary phase, and interpreted this as evidence thatFe4Al13 was poorly nucleated by Al, suggesting that Fe4Al13 nucleated onimpurity particles in the liquid instead. This would explain Adam and Hogan'sfailure to observe an orientation relationship between Fe4Al13 and Al (Section4.1.1). Adam and Hogan's results (Fig. 20) do however demonstrate that growthvelocity, and not cooling rate, chie¯y determines which eutectic phase will form,suggesting that the role of competitive nucleation in the kinetic displacement ofFe4Al13 by FeAl6 is minor.

Unlike Adam and Hogan [1], Donnelly and Rudee [28] and Ping et al.(1988) have observed orientation relationships between Fe4Al13 and the Almatrix, albeit when Fe4Al13 had precipitated from solid solution. This may suggestthat Al nucleates Fe4Al13. Donnelly and Rudee observed the orientationrelationship:

�001�Fe4Al13==�100�Al

�100�Fe4l13==�001�Al�210��

with the long platelet axis of the lamellae parallel to [010] (Fe4Al13) (as Adam andHogan). They claimed this orientation relationship was that of the lowestdisregistery between the two crystal lattices. Ping observed three di�erentorientation relationships:

�001�Fe4Al13==�001�Al

�200�Fe4Al13==�200�Al

�020�Fe4Al13==�020�Al

�100�Fe4Al13==�100�Al

�020�Fe4Al13==�020�Al

�001�Fe4Al13==�002�Al

�100�Fe4Al13==�10 �1�Al

�20�Fe4Al13==�020�Al

�001�Fe4Al13==�202�Al

arguing that there was good atomic matching between the two lattices at the(200)Fe4Al13 // (200)Al interface. These three orientation relationships are not inagreement with Donnelly and Rudee's observations however. Lendvai et al. [62, 63]and Shoji and Fujikura [93] similarly observed the precipitation of acicularFe4Al13 from the matrix on heat treatment of DC cast samples, which coarsen toa more irregular morphology with time. This also suggests that there may be adegree of coherency initially between the lattices of the Fe4Al13 phase and the Almatrix, which may promote nucleation of the Fe4Al13 by Al. Recent work byAllen et al. [4] on the secondary heterogeneous nucleation behaviour of micronsized liquid eutectic puddles entrained in a solid Al matrix supports BaÈ ckerud'shypothesis however. In the absence of impurities, undercoolings of 010±15 K atcooling rates as low as 0.03 K sÿ1 are measured prior to the onset of Al±Fe4Al13eutectic solidi®cation. If nucleation of Al±Fe4Al13 eutectic from the liquid by thesurrounding solid Al matrix were facile, no such undercoolings would beobserved.

Conversely, BaÈ ckerud proposed that the Al±FeAl6 eutectic needed no detectableundercooling to nucleate, indicating that Al is an e�cient nucleant for FeAl6. Thiscorrelates with the observation by Adam and Hogan [1] (Section 4.1.1) that thereis a consistent orientation relationship between FeAl6 and Al, corresponding to a


semicoherent interface. The lattice mis®t strain between the FeAl6 and Al latticesis indeed smaller than that between Fe4Al13 and Al [58]. Lattice mis®t strain maynot be the only determinant of nucleating e�ciency however [114], andthermodynamic factors may also play a role in the relative ease of nucleation ofthe two eutectics, eg. their latent heats of fusion or Al/intermetallic surface energyvalues [8, 100], which will in¯uence the degree of adsorption between the eutecticand the Al matrix [18±20, 57, 58].

A nucleation based argument can still be invoked however to explain the kineticdisplacement of Al±Fe4Al13 by Al±FeAl6 with increasing cooling rate, even if Al isa good nucleant for Fe4Al13 at low cooling rates. Fig. 24 illustrates schematicallythe isothermal nucleation diagrams (or INDs) for the Al±Fe4Al13 and Al±FeAl6eutectics [68]. Assume that nucleation of Al±Fe4Al13 is easier at highertemperature and lower undercooling than for Al±FeAl6, contrary to BaÈ ckerud'sproposal, and that growth of Al±Fe4Al13 is harder at higher temperature andlower undercooling than for Al±FeAl6, as proposed in Section 4.1.1. At lowcooling rates Al±Fe4Al13 will be the preferred eutectic phase. At higher coolingrates however the converse is true, and Al±FeAl6 displaces Al±Fe4Al13.

4.2.2. Promotion of nucleation of other phasesNucleation of phases other than Fe4Al13 and Al±FeAl6 has been little

considered, and consequently few details of the nucleants for, and associatedkinetics of, most common DC cast phases are available. Impurity particles presentin the alloy or grain re®ning additions made during casting may providenucleation sites for phases. Ho and Cantor [40] have demonstrated that bulk levelsof P as low as 2 ppm can a�ect the nucleation of eutectic Al±Si in hypoeutectic

Fig. 24. Curves for the start of nucleation of Fe4Al13 and FeAl6. At dT/dt=(dT/dt)1, Fe4Al13 nucleates

®rst at TN1. At dT/dt=(dT/dt)2, FeAl6 nucleates ®rst at TN2. The critical cooling path, at dT/dt=

(dT/dt)crit passes through point P, the intersection of the two TTT curves (after Maggs et al. [68]).


Al±Si. Recently, Hsu et al. (1998) have also demonstrated that 0.4 wt% Ca in6xxx series Al alloys can promote the formation of a-AlFeSi via the formation ofCaAl2Si2 during solidi®cation. Kosuge [58] and Griger et al. [37] have alsoproposed that Al±Ti±B grain re®ner addition increases the number density ofnucleation sites for FeAlm in 1xxx series alloys, whereas Tezuka and Kamio [104]noted that Al±Ti±B addition to Al±0.3±0.5 wt% Fe±0.1- 0.15 wt% Si promotedboth FeAlm and a-AlFeSi. Maggs et al. [68] has proposed that there may be asmall lattice mismatch d between the hexagonal lattice of TiB2 and phases withorthogonal crystal axes such as cubic a and FeAlm. Evans et al. [32] have similarlyobserved that the appearance of FeAlm in unidirectionally solidi®ed specimensrequires the presence of both Si and Al±Ti±B grain re®ner addition (as shownpreviously in Fig. 22). Recent work by Allen et al. [6] has shown that thenucleation of FeAlm can be promoted by both V impurity and Al±Ti±B grainre®ner addition at cooling rates as low as 0.03 K sÿ1. Shillington [92] has notedthat the solidi®cation conditions necessary to produce the FeAlx phase inunidirectionally solidi®ed specimens changes depending on whether thatsolidi®cation proceeds from a solid seed or directly from the melt, indicating thatthe formation of the FeAlx phase is sensitive to the nucleation conditions presentpreceeding growth. What nucleants are favourable for FeAlx formation remainsunclear however.

The roles of impurities and grain re®ner additions on the relative ease ofnucleation of phases are discussed in further details in Sections 8 and 9.2,respectively.

4.3. Suppression of equilibrium solidi®cation reactions

Si rich phases such as a-AlFeSi, b-AlFeSi and even eutectic Si are commonlyobserved in DC cast microstructures. The equilibrium peritectic and ternaryeutectic reactions that form these phases (Section 3.1) involve di�usion over largedistances. Suppose that an a-AlFeSi phase is undergoing the peritectic reactionliquid+ a4 b+Al. The reaction would commence at the liquid/a interface,forming b and Al at the interface. The b and Al would then act as a di�usionbarrier between the a and the liquid, slowing the reaction rate. The reaction ratewould be slowed still further as the reaction proceeds as the b+Al layer thickens.Under non-equilibrium solidi®cation conditions (e.g. DC casting) there isinsu�cient time for the reaction to proceed to completion and b phases with acores would result.

Cored phases are not commonly observed in DC cast microstructures however.The a and b phases must therefore form by alternative mechanisms. Langsrud [61]suggested that during DC casting a-AlFeSi and b-AlFeSi form metastably bydirect precipitation from the solute enriched interdendritic liquid viapseudoeutectic reactions, which occur at undercoolings below their respectiveequilibrium peritectic temperatures. This may explain the spread of measuredtemperatures for the corresponding invariant reactions [14, 85].


4.4. Metastable phase diagrams and solidi®cation microstructure selection maps

If under equilibrium conditions a solute preferentially partitions to the liquidphase (e.g. Fe or Si in Al), during non-equilibrium solidi®cation the liquid canbecome more enriched in solute than would be predicted from the equilibriumphase diagram by the lever rule [61]. As solidi®cation time decreases, even if localequilibrium is maintained at the solid/liquid interface, there is insu�cient time forsolute di�usion in the solid (known as back di�usion) to maintain thecomposition of the solid the same everywhere. The mean solid composition thenbecomes solute depleted with respect to the equilibrium solidus composition, andthe liquid conversely becomes solute enriched. Fig. 25 illustrates schematically theevolution of the concentrations in the liquid of Fe, Cl,Fe, and Si, Cl,Si, withfraction solid for two di�erent solidi®cation models. Where full di�usion occurs inthe solid the Lever rule applies (Fig. 25). Where no di�usion occurs in the solidthe Scheil model applies (Fig. 25). If solute enrichment of the liquid occurs, phasesricher in Si than would be predicted from the equilibrium phase diagram for thegiven bulk composition can form, such as a-AlFeSi, b-AlFeSi and eutectic Si(Section 3).

Fig. 25. Evolution of Fe content in liquid (top) and Si content in liquid (bottom) with fraction solid, in

cases of full solid di�usion (Lever rule) and no solid state di�usion (Scheil model). After Langsrud [61].


Based upon this, Langsrud [61] has constructed a series of psuedo-metastablephase diagrams for given non zero cooling rates. Fig. 26 illustrates these diagramsin the form of projected liquidus surfaces (as shown by solid lines) in the Alcorner of the Al±Fe±Si ternary. Phase ®eld boundaries are shifted to lower Si andhigher Fe with increasing solidi®cation rate, to account for the observation of Sirich phases in the commercial DC cast microstructure (Section 4.3) and theincrease in the undercooling required by the Al±Fe4Al13 eutectic reactionrespectively (Section 4.1.1 and 4.2.1). Proposed phase ®elds for metastable Al±Feeutectics (Section 2.2) are inserted at high cooling rates. On these metastablephase diagrams Langsrud then imposed calculated solidi®cation paths, allowing

Fig. 26. Metastable projected liquidus diagrams for Al±Fe±Si (after Langsrud [61]). Dashed lines

indicate the position of the equilibrium liquidus. The topmost diagram is marked with solidi®cation

paths with full solid state back di�usion (Lever rule) and no solid-state back di�usion (Scheil model).


for varying degrees of back di�usion, from cases with full back di�usion (lever

rule) to cases with no back di�usion (Scheil model). The degree of back di�usion

to be considered is determined by solidi®cation rate (which in turn also determines

the positions of the metastable phase ®elds, Fig. 5). These diagrams illustrate that

secondary/ternary phase content is particularly sensitive to the solidi®cation rate

for dilute Al±Fe±Si alloys.

Another method, which concentrates more on the variation of microstructure

with solidi®cation velocity rather than solidi®cation sequence with cooling rate,

was proposed by Hughes and Jones [45]. Fig. 27 plots the variation of

microstructural type with bulk Fe content and solidi®cation velocity from their

results of unidirectional solidi®cation experiments on Al±Fe alloys, of Fe contents

in the range 2.2±6.1 wt%. This type of plot is known as a solidi®cation

microstructure selection map (SMSM). Gilgien et al. [35], combining their data

with previous work done by Gremaud et al. [36] then developed this map to

higher velocities using laser surface remelting techniques. Fig. 28 shows their

resultant SMSM. Such SMSMs can be used as a guide to predict DC cast

microstructures, casting velocities being towards the lower ends of those mapped

(01±2 mm sÿ1). The maps shown in Figs. 27 and 28 assume that the temperature

gradient at the solidi®cation front G has a negligible e�ect on phase selection.

This is at least true in the case of the transition of Fe4Al13 to FeAl6, as

demonstrated by Adam and Hogan in Section 4.1.1.

Fig. 27. Microstructure selection map from experimental results for growth regimes of Al±Fe4Al13eutectic (Eul) and Al±FeAl6 eutectic (Eu2). After Hughes and Jones [45]. Reproduced from the Journal

of Materials Science by kind permission of Kluwer Academic Publishers.


5. Fir tree formation in DC casts

5.1. Fir tree zones

As was seen in Sections 2 and 3, a wide range of both equilibrium andmetastable Al±Fe and Al±Fe-Si phases are observed in model and commercial1xxx alloys. The non-equilibrium solidi®cation conditions that prevail duringconventional casting mean that equilibrium thermodynamic considerations areusually not capable of predicting cast phase content.

As was seen in Section 4, constant solidi®cation velocity unidirectionalsolidi®cation studies can characterize phase contents with di�erent solidi®cationconditions. However, it is di�cult to apply these results directly to DC casting.During unidirectional solidi®cation the substrate remains the same for di�erentsolidi®cation conditions and thus determines which phases nucleate during theinitial stages of solidi®cation. During DC casting, both solidi®cation velocity andcooling rate vary both with position across the solidifying ingot cross-section andwith solid volume fraction (i.e. with solidi®cation time). Thus the immediatelypreceeding solidi®cation history is continually determining which nucleants areavailable for subsequent solidi®cation. This continuous nature of DC castingtherefore leads to discrepancies between DC cast phase contents and thoseproduced using controlled solidi®cation rate techniques for a given alloycomposition. This is explored further in Section 6.

As mentioned in Section 1, the varying solidi®cation conditions during DCcasting are important as they can produce corresponding changes in phasecontent, grain size and solid solution content. On sectioning and chemical etching

Fig. 28. Solidi®cation microstructure selection map for Al±Fe alloys. After Gilgien et al. [35].

Reproduced by kind permission of Elsevier Science Ltd.


a DC cast ingot or billet, changes in phase content are revealed as transitionalzones between areas with di�erent responses to a given etchant. These zones arereferred to as Altenpohl zones [7] or `®r-tree' zones, due to their irregular tree-likeoutline. Following scalping and rolling of the DC cast ingot or billet, these ®r treezones appear on the surface as bands after surface treatment (commonly etching).This banding is highly undesirable in any surface critical products, such asarchitectural cladding or lithographic printing sheet. Positional variations insecondary phase content can also lead to a loss in formability during subsequentdownstream processing, as di�erent phases will break up di�erently on rollingdepending on their relative hardnesses.

Fig. 29 illustrates a horizontal section through a DC cast billet of commercialpurity Al, which has been caustic macro-etched to reveal the ®r-tree zone. Theposition and width of the ®r tree zone and the phases it contains depends uponboth the alloy composition and casting parameters. The variation in phase contentcan also result in reduced strength, toughness and formability. Consequently thebillet or ingot has to be scalped and/or heat-treated to homogenise itsmicrostructure prior to rolling (see Section 7).

5.2. Cooling rate

Kosuge and co-workers [58, 59] measured dendrite arm spacing (DAS) acrossthe width of a DC cast Al±0.6 wt% Fe billet to estimate the variation of thecooling rate with position. Fig. 30 illustrates their results, showing a sharpdecrease in cooling rate over the ®rst few mm of the billet thickness to a minimumvalue 010 mm in from the billet surface. The cooling rate then rises again to asecond maximum 020 mm in from the surface, then decreases slowly toward thebillet centre.

The variation in measurements of cooling rate such as shown in Fig. 30 is notsurprising given the cooling methods employed during DC casting.Westengen [112], Kosuge [58], Brusethaug et al. [17], Lakner et al. [60],Langsrud [61] and Cigdem and Bennett [22] have all considered the thermalevolution of the billet or ingot during DC casting. Fig. 31 shows a schematic ofthe cross-section through a typical DC cast billet during solidi®cation. As moltenmetal is poured into the mould initial contact of the liquid metal with the water-cooled mould wall produces a chill layer <1 mm thick. The extent of chillinglocally will depend upon the intimacy of the quenched melt/mould wallcontact [58]. Solidi®cation shrinkage produces an àir-gap' which acts as a barrierto conductive heat transport, reducing the average cooling rate to a minimumvalue, over a shell zone 05±10 mm thick (as illustrated). Coarse DAS zones arecorrespondingly observed in this region of the billet [58]. The enthalpy released onsolidi®cation can act to raise the temperature locally (i.e. recalescence), and in theextreme case remelting of the Al dendrite tips can occur in this shell zone [58].Rapid cooling is re-initiated when secondary cooling of the billet takes e�ect bythe impingement of a curtain of pressurised water jets onto the solidi®ed surface


Fig.29.Etched

cross

sectionthroughDC

cast

ingotshowinguneven

response

toetchingboughtaboutbyanuneven

distributionin

secondary

phases,

knownasthe`®rtree'structure.ReproducedbykindpermissionofP.V.Evans.


of the billet. The maximum cooling rate re-attained is dependent principally uponthe casting speed, i.e. the rate of withdrawal of the solidifying billet. The coolingrate then decreases gradually once more towards the billet centre, as the heattransport path length increases [112].

5.3. Fir tree nucleation

Kosuge and co-workers [58, 59] argued that competitive nucleation betweenAl±Fe eutectic phases could be used to explain how the ®r-tree zone originates

Fig. 30. Schematic variation of cooling rate with position in DC cast ingot. After Kosuge [58].

Fig. 31. Cross sectional schematic of direct chill casting arrangement, from edge to centre of billet

being cast.


and how its thickness depends on DC casting speed. Fig. 32 illustratesschematically the variation of cooling rate E with position in the ingot for twodi�erent casting speeds, 100 mm minÿ1 and 60 mm minÿ1. Kosuge and co-workers assumed that at any position in the billet the Al±Fe phase with thehighest nucleation temperature nucleates ®rst and then grows. Assuming thatthere exists a critical cooling rate (dT/dt)crit or E * such that the undercooling fornucleation of one phase, say FeAlm, becomes less than that required to nucleate

Fig. 32. Variation of cooling rate with position in ingot for cases of high casting rate (top) and low

casting rate (bottom). At high casting rates the renucleation of FeAlm towards the centre of the cast

becomes favourable. After Brusethaug et al. [17].

Fig. 33. With increasing cooling rate deeper undercoolings below the equilibrium eutectic temperature

can be attained producing metastable eutectics (left). The nucleation temperatures and preferred

eutectic phase thus vary as shown (right). After Brusethaug et al. [17].


another phase, say FeAl6, then whenever the cooling rate exceeds E * FeAlm willdisplace FeAl6.

The existence of a critical cooling rate above which the preferred phase to benucleated changes is as proposed above in Section 4.2.1 both by BaÈ ckerud(Fig. 19) and Maggs et al. (Fig. 24). Brusethaug et al. [17] also proposed that thenucleation temperature for a given Al±Fe phase varied with cooling rate, asillustrated in Fig. 33. Consequently, in the zones where the cooling rates arehighest then FeAlm can displace phases that require higher undercoolings tonucleate, such as FeAl6 (as illustrated in Fig. 32) and Fe4Al13. A higher number of®r tree zones of greater width are anticipated with increasing casting speed. AsFig. 32 shows, only at high casting speeds is the maximum cooling rate reattainedin the zone of the billet cooled by the water jets above the critical cooling rate E *

for the transition from FeAl6 to FeAlm. At high casting speeds a ®r tree is formedin that zone, its width increasing with increasing casting speed.

As Brusethaug et al. [17] noted, this simple nucleation based model for ®r treeformation does not account for the irregular shape of the ®r tree zone, for thecommonly observed intimate coexistence of di�erent Al±Fe phases, for ®r treezones at low casting speeds, for equilibrium phases in rapidly cooled parts of thebillet or for metastable phases in more slowly cooled regions [16, 37, 80] (seeSection 6). Nor can it account for the e�ect of impurities and grain-re®ningadditions on the position of the ®r tree zone. These e�ects are discussed further inSections 8 and 9.2 below.

The nucleation model fails because it assumes that whichever secondary/ternaryphase nucleates ®rst will subsequently grow. As has been seen already in Section4.1, whether a given phase can grow or not at a particular temperature howeverdepends not on local cooling rate but upon the local imposed solidi®cationvelocity.

5.4. Fir tree nucleation and growth

Brusethaug et al. [17] argued that both the nucleation and growth of thecompeting Al±Fe phases need to be considered in order to explain the irregularshape of the ®r tree boundary. Fig. 34 illustrates the case for the competitionbetween FeAlm and FeAl6, where if the cooling rate exceeds E * then FeAlm isnucleated in preference to FeAl6 (i.e. right hand side of Fig. 34 as bottom ofFig. 32). Fig. 34 (left hand side) shows the sequence of nucleation and growthevents as solidi®cation proceeds. In the outermost chill zone the cooling rate isabove E * and FeAlm ®rst nucleates and then grows. As the cooling rate decreasesaway from the outermost chill zone of the billet, the cooling rate falls below E *

and FeAl6 nucleates ahead of the growing FeAlm, along the isothermcorresponding to the FeAl6 nucleation temperature, but only where suitablenucleation sites exist. The FeAl6 nuclei can only grow ahead of the FeAlm,however, if their growth temperature is also greater than that of FeAlm. In


otherwords, given the conditions;

Tn�FeAl6� > Tg�FeAlm�; andTg�FeAl6� > Tg�FeAlm�; andsuitable nucleation sites for FeAl6 exist along the T � Tn �FeAl6� isotherm

then FeAl6 can nucleate ahead of and overgrow the advancing FeAlm, producinga tree-like shaped transition from one phase to another.

Brusethaug et al. [17] also considered the variation of solidi®cation velocitythrough the solidifying DC billet, the variation of the relative values of eutecticgrowth temperatures of the competing phases with solidi®cation velocity, and theavailability of nucleation sites for new phases and hence the importance ofimmediate previous local solidi®cation history. Just as the cooling rate at a givenposition in a DC cast billet or ingot is determined by the magnitude of the localmaximum heat extraction at that position, so to the growth velocity is determinedby the direction of the local maximum heat extraction. Fig. 35 illustratesschematically the variation of solidi®cation velocity with position in the solidifyingDC billet. The solidi®cation velocity at a given position is the projection of theimposed casting velocity, UDC, onto the direction of maximum heat extraction atthat position, i.e. perpendicular to the liquid/solid interface at that position.Fig. 36 shows the liquid/solid sump shapes (top), the corresponding velocitypro®les (centre) and corresponding distribution of ®r tree zones (bottom) forrectangular ingots (left hand side of ®gure) and cylindrical billets (right hand side)considered by Brusethaug.

Fig. 37 illustrates schematically the cooling rate (top) and solidi®cation velocity(centre) pro®les that give rise to a solidi®cation sequence (bottom) wherebyìslands' of FeAl6 in FeAlm are produced. Assume that the critical cooling rateabove which FeAlm nucleation is preferred to FeAl6 is (dT/dt)crit and the critical

Fig. 34. Cooling rate variation with position in ingot (right) and corresponding nucleation and growth

sequence that occurs during DC casting producing the ®r-tree structure. After Brusethaug et al. [17].


solidi®cation velocity above which FeAlm grows in preference to FeAl6 is Ucrit. Asdescribed above, FeAl6 nucleates and grows ahead of the advancing FeAlm whenthe cooling rate and solidi®cation velocity fall below these critical values. Atsu�ciently fast casting speeds the maximum in cooling rate reattained is above(dT/dt)crit and the nucleation of FeAlm on the newly growing FeAl6 becomespreferred once more. If the sump pro®le is such that the solidi®cation velocity alsorises above Ucrit in this region, then FeAlm can overgrow FeAl6, leading to theproduction of ìslands' of FeAl6 in FeAlm dominated regions. This central rapidgrowth zone in the DC casting explains the general observed intimate coexistenceof Al±Fe phases in DC cast billets. Conversely, slow growth in rapidly cooledzones can lead to the production of equilibrium phases (e.g. Fe4Al13) even at highcooling rates. This agrees with reports of the equilibrium Fe4Al13 phase beingproduced at cooling rates of up to 10 K sÿ1 [16, 96].

As demonstrated above, the distribution of phases in the as-cast billet isdetermined by the values of the critical solidi®cation velocity (Ucrit) and critical

Fig. 35. Cross section through solidifying DC billet, illustrating the variation in solidi®cation direction

and corresponding solidi®cation velocity pro®le (after Maggs et al. [68]).


cooling rate ((dT/dt)crit) at which the growth and nucleation, respectively, of agiven phase will dominate over that of another phase. These critical parameters,as well as being functions of heat extraction rate, are also in¯uenced by the localmelt and solid chemistry, and hence bulk composition and the presence ofimpurities and of grain re®ning additions, which are especially thought to alterlocal nucleation conditions (ref. Section 9.2).

5.5. E�ect of solid fraction

Although the above models consider both nucleation and growth, the assumedvariations of cooling rate with position, based on DAS measurements, are notnecessarily representative of the cooling rate experienced locally during the wholesolidi®cation time. DAS values only give the average cooling rate experiencedduring the solidi®cation of the primary Al dendrites, i.e. between zero and 085±95% fraction solid. The DAS value does not provide any information on thecooling rate during the formation of the secondary and ternary phases during the®nal stages of solidi®cation. More complex numerical modelling of the

Fig. 36. Schematic sump pro®les (top), corresponding velocity pro®les (centre) and corresponding ®r-

tree patterns (bottom) in rectangular ingots (left) and cylindrical billets (right). After Brusethaug et

al. [17].


Fig. 37. Determination of phase content by both nucleation (controlled by cooling rate) and growth

(controlled by solidi®cation velocity) (after Brusethaug et al. [17]).


solidi®cation process is necessary, the details of which are beyond the scope ofthis review. In essence however, by modelling heat and mass ¯ows during castingthe shapes and vertical separations of the isotherms or fraction solid isoplanesduring casting can be determined. The time taken at some position in the cast tosolidify from one fraction solid, say fs,1, to a second fraction solid, say fs,2 (>fs,1)is then the vertical distance between the fs,1 and fs,2 isoplanes at that positiondivided by the casting speed, i.e. the speed at which the billet or ingot is beingwithdrawn vertically. The average cooling rate over that time interval is then thedi�erence between the two isotherms (T1ÿT2) corresponding to the fs,1 and fs,2isoplanes divided by the solidi®cation time. This variation of cooling rate not onlywith position but with solidi®cation time further complicates the analysis of phasecontent formation during DC casting.

6. Fir tree phases in DC casts

A number of workers have characterised the majority phases that form òutside'the ®r-tree zone, i.e. towards the cast surface, and ìnside', i.e. towards the centreof the billet or ingot, and the critical values of cooling rate and/or solidi®cationvelocity over which ®r tree formation occurs.

Simensen and Vellasamy [94], Westengen [112], Kosuge [58], Skjerpe [96],Brusethaug et al. [17] and Maggs et al. [68] examined sections of DC cast orlaboratory scale DC simulated cast billets. Using techniques such as TEM, EDXand di�raction pattern analysis they determined which phases were present insideand outside the ®r tree zone. The local cooling rates at which given phases wereformed were determined by techniques such as measurement of the secondarydendrite arm spacing (DAS) (e.g. [96]), or by insertion of thermocouples into thebillet or ingot during casting (e.g. [17]). However, as Adam and Hogan [112]illustrated for the case of the displacement of Fe4Al13 by FeAl6, the transitionsfrom one to phase to another do not necessarily occur at one unique cooling rate,and can also be dependent upon solidi®cation velocity. Consequently,determination of the variation of phase content with cooling rate based solelyupon the analysis of DC cast samples is potentially unreliable. The localsolidi®cation velocities however could not be easily determined in these DC castsamples. In addition, there can be large variations of both cooling rate andsolidi®cation velocity over even short distances (of the order of cm).

As was seen in Section 4, accurate determination of the variation of phasecontent under more controlled conditions of cooling rate and solidi®cationvelocity is possible using unidirectionally solidi®ed Bridgman grownsamples [1, 16, 32, 45, 80, 105, 113]. Fig. 38 illustrates a schematic of a verticalBridgman growth furnace, from which a molten specimen is drawn at a controlledrate, imposing unidirectional solidi®cation at a single solidi®cation velocity. Understeady-state conditions the cooling rate experienced can be determined if thetemperature gradient across the solid-liquid interface is known, from Eq. (1).


Steady-state solidi®cation velocities and cooling rates of up to 2 mm sÿ1 and 40 Ksÿ1 can be achieved using Bridgman growth, using typical temperature gradientsof 10±20 K mmÿ1.

The cooling rate ranges over which the Al±Fe eutectics form, as determinedfrom both Bridgman grown and DC cast samples, are as detailed in Table 3 (see

Fig. 38. Schematic of Bridgman growth furnace apparatus.

Table 3

Cooling rate ranges of formation of the common Al±Fe eutectics in hypoeutectic Al-Fe alloys

Phase Cooling rate range, K sÿ1 References

Fe4Al13 0.1±3 Young and Clyne [113], Kosuge [58], Ping et

al. [80], Skjerpe [96], Brusethaug et al. [17], Evans

et al. [32], Maggs et al. [68]

FeAlx 0.4±5 Young and Clyne [113], Westengen [112],

Skjerpe [96]

Fe2Al9 1±6 Simensen and Vellasamy [94], Brobak and

Brusethaug [16], Griger et al. [38]

FeAl6 2±11 Adam and Hogan [1], Hughes and Jones [45],

Young and Clyne [113], Westengen [112],

Kosuge [58], Ping et al. [80], Brusethaug et al. [17],

Maggs et al. [68]

FeAlm >11 Young and Clyne [113], Westengen [112],

Kosuge [58], Ping et al [80], Brusethaug et al. [17]


Table 4

Solidi®cation velocities of phase change transitions

Transition Alloy Solidi®cation velocity

range of transition

References

Fe4Al134FeAl6 Hypoeutectic Al±Fe,

hypereutectic Al±Fe

00.1±0.2 mm sÿ1 Hughes and Jones [45],

Liang and Jones [65],

Evans et al. [32], Thomas

et al. [105]

Fe4Al134FeAlx Hypoeutectic Al±Fe 00.15±0.2 mm sÿ1 Young and Clyne [113]

Fe4Al134FeAlm Hypoeutectic Al±Fe+0.1

wt% Si+Al±Ti±B

>1.33 mm sÿ1 Evans et al [32]

Fe4Al134? Hypoeutectic Al±Fe+0.1

wt% Si

00.2±0.5 mm sÿ1 Brobak and Brusethaug [16]

Fe4Al134a-AlFeSi Hypoeutectic Al±Fe+0.2

wt% Si

01±2 mm sÿ1 Brobak and Brusethaug [16]

Table 5

Secondary and ternary phases detected in DC billet sections, DC simulations and directionally solidi®ed

Bridgman growth specimens

Researcher and

method

Alloy Phases òutside' ®r-tree

(and cooling rate)

Phases ìnside' ®r-tree (and

cooling rate)

Westengen [112],

DC billet sections

Al±0.25±Fe±0.13 Si

(cp)+gran re®ner

a0, m, a (10 K sÿ1) 6, x, 3 b, Si (5 K sÿ1)

Skjerpe [96], DC

billet sections

Al±0.2 Fe±0.1 Si (cp) a, m, aT, av (6±8 K sÿ1) 3, x, m, a, a0 (1 K sÿ1)

Brusethaug et

al. [17], DC billet

sections

Al±0.29 Fe±0.16 Si (cp)

Al±0.35 Fe±0.07 Si (cp)

m, 3, a, 6 (5±7 K sÿ1)

6, 3, m (5±7 K sÿ1)

3, a (1±2 K sÿ1)

3, 6 (1±2 K sÿ1)

Maggs et al. [68],

DC simulation

Al±0.2 Fe±0.1 Si (hp)

Al±0.4 Fe±0.1 Si (hp)

Al±0.4 Fe±0.2 Si (hp)

a, x (3±16 K sÿ1)

6, 3 (3±16 K sÿ1)

a, 6 (3±16 K sÿ1)

3, x (0.7±3 K sÿ1)

3 (0.7±3 K sÿ1)

3, 6 (0.7±3 K sÿ1)

Ping et al. [80],

Bridgman growth

Al±0.4 Fe±0.1 Si (cp)

Al±0.2 Fe±0.1 Si (cp)

Al±0.2 Fe±0.1 Si (hp)

3, p, 6, m, a (10 K sÿ1)

a, a0 (10 K sÿ1)

a, 3, p (10 K sÿ1)

3 (1 K sÿ1)

3 (1 K sÿ1)

3 (1 K sÿ1)

Brobak and

Brusethaug [17],

Al±0.29 Fe±0.16 Si (cp) a (5±6 K sÿ1) 3 (0.25±3 K sÿ1)

Bridgman growth Al±0.35 Fe±0.07 Si (cp) unidenti®ed? (3±6 K sÿ1) 3 (0.25±2 K sÿ1)

Note: cp=commercial purity base alloy, hp=high purity, 3=Fe4Al13, 6=Fe4Al13, 6=FeAl6,

m=FeAlm, x=FeAlx, p=FeAlp. Dominant secondary/ternary phases are in bold type.


also Section 4.1). The results of solidi®cation velocity measurements made onBridgman samples by Hughes and Jones [45], Young and Clyne [113], Brobak andBrusethaug [16], Liang and Jones [65], Evans et al. [32] and Thomas et al. [105]are summarized in Table 4 (see also Section 4.1).

However, as seen in Section 5, given the complex variation of cooling rate,solidi®cation velocity, immediate previous solidi®cation history and nucleationconditions with position in the solidifying DC cast billet, it is not surprising todiscover that discrepancies exist between the unidirectional solidi®cation data andphase content analyses from DC cast samples. These discrepancies are highlightedin Table 5, which is a comparison of the phases observed both inside and outsidethe DC ®r tree structure for the given 1xxx bulk alloy compositions, to thosepredicted by Bridgman growth experiments. The cooling rate and solidi®cationvelocity ranges quoted in Tables 3 and 4 are at best therefore only indicative ofthe general trends in change of phase content with casting parameters. The valuescan vary with bulk Fe and especially Si contents of the alloy, the presence of traceimpurities and/or grain re®ning additions, and the casting processemployed [17, 38]. The transitions from one phase to another, observed in DCcastings, are not as sharply de®ned as results from Bridgman growth experimentswould suggest, although this may partly be due to transformations that take placeduring subsequent cooling of the billet or ingot after solidi®cation [105], and thefact that the solid on which nucleation takes place is continually changing duringDC casting (see Section 5).

There is general agreement however in the literature that the boundary of the ®rtree zone observed on caustic etching corresponds to the appearance of the Al±FeAlm eutectic at05±10 K sÿ1 in the more rapidly solidi®ed zones of the ingot orbillet [17, 96, 112], although Maggs et al. [68] note that other phase transitions thatdo not give rise to a visible ®r tree zone also exist. The electrochemical potentialof the Al±FeAlm is di�erent to that of other phases that dominate in themore slowly cooled zones (e.g. Al±FeAl6, Al±Fe4Al13), and these di�erences inelectrochemical potential lead to an uneven etching response as noted previously.However, at similar cooling rates a and a0 dominate over the metastable Al±Feeutectics at low Fe/Si ratios (e.g. Al±0.2 Fe±0.1 Si). Although these phases do notproduce a marked ®r tree [68] they cannot be readily removed using equilibrationby heat treatment [96, 97, 112]. This highlights the importance of controlling the ascast phase content, as subsequent downstream processing is not always capable ofremoving undesirable features of the as cast microstructure.

7. Transformation of metastable phases

Fortunately, the ®r tree structure in DC cast billets can frequently be reducedor eliminated by homogenization by heat treatment. Consequently it is desirableto have a full understanding of the transformation kinetics for the each of the


metastable Al±Fe and Al±Fe±Si phases so that homogenization treatments may beoptimized accordingly.

Using a range of techniques including TEM, MoÈ ssbauer spectroscopy, XRDand electrical resistivity measurements Hughes and Jones [46], Kosuge [58],Lendvai et al. [62, 63], MurgaÂ s et al. [76], Griger et al. [37, 38], Shoji andFujikura [93], Kim and Cantor [56] and Shillington [92] have observed themicrostructural equilibration processes that take place during heat treatment ofAl±Fe and Al±Fe±Si alloys in the temperature range 0300±6508C. They haveconcentrated on the precipitation of excess solute from solid solution and thetransformation of metastable phases: FeAl64Fe4Al13; FeAlm4Fe4Al13; a-AlFeSi 4 Fe4Al13; and a-AlFeSi 4 b-AlFeSi.

7.1. The FeAl6 ÿ4 Fe4Al13 transformation

7.1.1. Transformation mechanism and activation energy

7.1.1.1. Dissolution±precipitation mechanism and net activation energy. Kosuge andco-workers [58] isothermally homogenised Al±0.6 wt% Fe containing only Al andFeAl6 for 1 h in the range 300±6408C. Fig. 39 shows schematically their results forthe variation of the normalised intensities of X-ray peaks from FeAl6 and Fe4Al13with homogenization temperature after 1 h of isothermal holding. With increasinghomogenization temperature, the proportion of FeAl6 decreases from 5008C, andcompletely converts to Fe4Al13 by 6408C. Kosuge did not detect Fe4Al13 until05608C, as shown in Fig. 39. He suggested that in the range 500±5608C FeAl6disappears without any formation of Fe4Al13. However, as MurgaÂ s et al. [76]noted, precipitation of Fe from solid solution in some form must occur at thesame rate as FeAl6 dissolves, otherwise the Al solid solution would rapidly super-saturate in Fe. Kosuge's XRD technique would not however detect the Fe4Al13 if

Fig. 39. Schematic of variation of normalised X-ray intensity of FeAl6 and Fe4Al13 with

homogenization temperature in Al±0.6 wt% Fe alloy. After Kosuge [58].


it were present as a dispersion of ®ne precipitates. At no temperatures were anyphases other than Fe4Al13 or FeAl6 detected, indicating that there are no inter-mediate Al±Fe species involved in the transformation.

Kosuge observed that the FeAl6 particles ruptured and dissolved into the matrixwhile the Fe4Al13 nucleated and grew remotely from the dissolving FeAl6. Fe4Al13rarely nucleated at FeAl6/Al interfaces. The preference of Fe4Al13 to nucleateremotely from the dissolving FeAl6 may arise from the anticipated high level ofFe in solid solution resulting from the rapid solidi®cation experienced duringsolidi®cation [92].

Kosuge ®tted the decrease of FeAl6 with time using the Arrhenius relation;

I � I0 exp�ÿkt� �6�

where I=X-ray peak height of FeAl6 at time t, I0=X-ray peak height of FeAl6at time t=0, and k is some rate constant. The rate constant in turn was assumedto ®t the Arrhenius relation;

k � k0 exp�ÿQnet=RT� �7�

where k0 is some constant, Qnet is the net activation energy for transformation, Tis temperature and R is the molar gas constant. An Arrhenius relation was ®ttedto the normalized XRD intensity data for various homogenisation times andtemperatures, as shown in Fig. 40. By plotting the time taken for the X-ray peakheights of FeAl6 to decrease by half (i.e. I/I0=0.5) against the reciprocal oftemperature, 1/T, Kosuge and co-workers obtained a net activation energy oftransformation Qnet of0290240 kJ molÿ1.

From their microstructural observations, Kosuge and co-workers proposed thattransformation preceeded via di�usion of Fe from the dissolving FeAl6 to theprecipitating Fe4Al13, with bulk di�usion of Fe in Al between the two speciesbeing the rate determining step of the reaction. Kosuge did not state with which

Fig. 40. Schematic of variation of normalized X-ray intensity of FeAl6 and Fe4Al13 with

homogenization time (hours) in Al±0.6 wt% Fe alloy for the temperatures shown. After Kosuge [58].


determination in the literature of the activation energy for bulk Fe di�usion in

aluminium, Qdi�, he compared his value of Qnet, but his value of 290240 kJ

molÿ1 does correlate reasonably well with the value of Qdi� determined by

Hood [43] of 260 kJ molÿ1 from 59Fe isotope di�usion measurements in Al in the

temperature range 520±6608C.

7.1.1.2. Formation of acicular Fe4Al13 precipitates. Lendvai et al. [63] also followed

the transformation of FeAl6 to Fe4Al13, on homogenization of DC cast Al±0.5

wt% Fe in the temperature range 400±6208C using electrical resistivity measure-

ments and TEM. The resistivity decreased down to a saturation value with

increasing homogenization time (the time to achieve that saturation decreasing

with increasing homogenization temperature) indicating that the as cast solid sol-

ution was supersaturated and began to precipitate Fe from solution in the form of

the equilibrium Fe4Al13 phase. By ®tting an Arrhenius relation to the variation in

ageing time necessary for the resistivity to attain its saturation value with tempera-

ture, Lendvai determined a value of Qpptn for the precipitation of Fe from solid

solution of 0135220 kJ molÿ1. As Lendvai noted, this was signi®cantly lower

than some determinations of Qdi�, but did agree with the determinations by SoÈ ren-

sen and Trumpy [99] and SzaboÂ et al. [108]. Lendvai therefore concluded that the

kinetics of the Fe4Al13 precipitation are rate controlled by di�usion of Fe through

the matrix to the precipitating Fe4Al13, rather than by the attachment kinetics of

the Fe to the Fe4Al13 lattice.

TEM of specimens aged for 50 h at 400±4508C showed 01 mm long acicular

precipitates orientated with respect to the Al matrix, surrounded by mis®t

dislocations indicating that the precipitates were initially coherent with the matrix.

No morphological change in the FeAl6 was apparent. Griger et al. (1989) noted

the same precipitation e�ect on homogenising DC cast Al-0.5 wt% Fe±0.02±0.9

wt% Si for 14 h at 6058C. Fig. 41 shows a SEM micrograph of the deep etched

surface of the DC cast ingot, showing the acicular precipitates. On further ageing

some of these precipitates grew up to 10 mm long and 1±2 mm wide, while the

others grew into larger less acicular particles, perhaps brought about by a

twinning mechanism. The acicular precipitates were observed both near to and

remote from the dissolving FeAl6, with their [020] axes parallel to the [100]

direction of the Al matrix. Shoji and Fujikura [93] also homogenised DC cast

1200 alloy for 30 min at 5008C, and observed the precipitation of 00.2 mm long

needles.

7.1.1.3. Continuous heating transformation. Further evidence for the existence of

these acicular precipitates has been provided by calorimetric studies. Lendvai et

al. [62] calorimetrically analysed the melting of DC cast Al±0.5 wt% Fe alloy,

known to contain only Al and FeAl6, both in the as-cast condition and following

equilibration by melting and solidifying at 0.3 K minÿ1 in the calorimeter. On

melting the equilibrated alloy two single endotherms were detected, corresponding

to the eutectic melting of Fe4Al13 at 654.720.38C, and the subsequent melting of


the Al matrix, as shown in Fig. 42. On melting the as cast alloy, four endothermswere observed, as shown in Fig. 43: a single peak at 6528C corresponding to theeutectic melting of a remnant of untransformed FeAl6; a close double peak corre-sponding to the melting of two distinct morphologies of Fe4Al13 present [101] overthe temperature range 654±6568C; and, the melting endotherm of the Al matrix.The two morphologies of Fe4Al13 were proposed to form during heating in thecalorimeter prior to melting, in accord with Lendvai's other observations [63].Allen et al. [4] have recently corroborated these results, analysing the melting ofAl±0.5 Fe Bridgman grown samples containing only FeAl6 and Al in the as growncondition. Fig. 44 shows typical melting endotherms from an as-grown sample(top), an equilibrated sample (centre) and a sample isothermally held for 20 min at06488C (bottom). The melting endotherm from the as grown sample containsfour peaks corresponding to the four peaks seen in the endotherm from Lendvai'sas-cast sample. The melting endotherm from the equilibrated sample contains twopeaks, again corresponding to eutectic melting of Fe4Al13 followed by melting ofthe Al matrix, similar to the two peaks seen in the endotherm from Lendvai's as-cast sample. This indicates that after equilibration only one morphology ofFe4Al13 was present. The peak corresponding to the melting of the acicular pre-cipitates in the endotherm from the isothermally held sample is smaller than thatin the endotherm from the as grown sample, indicating that isothermal holdinggradually removes the acicular precipitates.

Indeed, after a few hours at 600±6208C Lendvai observed that the acicularprecipitates were absent and the morphology of the FeAl6 had changed

Fig. 41. SEM micrograph of surface etched DC cast alloy showing precipitation of Fe4Al13 needles

after 14 h isothermal hold at 6058C. After Griger et al. [37]. Reproduced by kind permission of

Aluminium.


considerably. Lendvai proposed that at these higher temperatures the FeAl6 hadbegun to dissolve and that Fe4Al13 (that had nucleated at lower temperature) grewto such a size that it was incoherent with the Al matrix (as observed by Griger).

7.1.1.4. Two step ageing. Lendvai [63] therefore investigated two step ageing treat-ments, whereby specimens were ®rst aged for 50 h at 4508C to nucleate Fe4Al13precipitates, then held at temperatures in the range 500±6208C. Fig. 45 shows thevariation of the ratio of specimen resistance measured at 273 K (in icewater) tothe resistance measured at 78 K (in liquid nitrogen) during this second step of age-ing, denoted the electrical resistance ratio (ERR). The ERR ®rst decreased, thenincreased again, indicating a resaturation of the Fe content in solution in the Almatrix. This reversion occurred due to the observed partial redissolution of theacicular precipitates. This is anticipated as the maximum solid solubility of Fe inAl increases with increasing temperature up to the equilibrium eutectic tempera-ture of 6558C. Lendvai therefore proposed that the acicular Fe4Al13 acts as an

Fig. 42. Melting thermograms (at 0.31 K minÿ1) from Al±Fe alloys solidi®ed under equilibrium

conditions compared to the melting thermogram from a DC cast Al±Fe alloy (sample 1) previously

equilibrated by melting and solidifying at 0.31 K minÿ1 also. After Lendvai et al. [63]. Reproduced by

kind permission of Aluminium.


Fig. 43. Melting thermograms (at 0.31 K minÿ1) from two as DC cast A1±Fe alloys. After Lendvai et

al. [63]. Reproduced by kind permission of Aluminium.

Fig. 44. Melting endotherms from Bridgman grown Al±0.5 wt% Fe grown at 1 mm sÿ1: as grown

(top); after melting and resolidi®cation (centre) and; after isothermal holding 20 min at approx. 6458C.


intermediary step that ®rst precipitates from solid solution and then partiallyredissolves. Fitting an Arrhenius relation to the second step ageing data yields avalue for Qnet1200240 kJ molÿ1, 01.5 times higher than the Qpptn valuemeasured for the earlier precipitation stage. Lendvai proposed that this highervalue re¯ected an additional chemical binding energy that the Fe atoms had toovercome prior to dissolution from the Fe aluminide lattices (either the dissolvingFeAl6 or the reverting Fe4Al13). This chemical binding term was not involvedduring precipitation of Fe from solid solution in the earlier lower temperaturestage of the transformation, where a lower value of activation energy had beendetermined. Lendvai concluded that in his samples dissolution of the FeAl6 wasthe rate determining step during the transformation of FeAl6 to Fe4Al13. However,Hughes and Jones [46] observed that FeAl6 readily dissolves during transform-ation, which suggests that Lendvai's interpretation is incorrect. Of note also isthat Lendvai's determination of Qnet is much lower than that determined by

Fig. 45. Changes in the electrical resistance ratio of DC cast Al±0.5 wt% Fe during ageing following

®rst-step ageing of 50 h at 4508C. After Lendvai et al. [63]. Reproduced by kind permission of

Aluminium.


Kosuge, and subsequently also by Shillington (see below), suggesting that Lend-vai's results were in error.

7.1.1.5. Isothermal transformation. MurgaÂ s et al. [76] also followed the transform-ation of FeAl6 to Fe4Al13, on homogenization of DC cast Al±0.58 wt% Fe in thetemperature range 580±6508C using MoÈ ssbauer spectroscopy. MoÈ ssbauer spectrawere deconvoluted to obtain plots of area fraction of spectrum components (eitherFe in FeAl6, Fe in Fe4Al13 or Fe in Al solid solution) versus time. No change ofFe level in solid solution with time was detected. This indicates that the dissol-ution rate of the FeAl6 and Fe4Al13 needles was matched by the precipitation rateof Fe4Al13 elsewhere. However, whereas other workers [13, 48] have measuredsupersaturated Fe levels as low as 0.9 wt% Fe in rapidly solidi®ed Al±Fe alloys,Forder et. al. [34] were not able to detect Fe levels nearer to equilibrium (00.05wt%) in Al±0.5 wt% Fe Bridgman grown at 1 mm sÿ1, suggesting that MoÈ ssbauerspectroscopy is not suited to studying low levels of Fe in solid solution.

MurgaÂ s assumed that the area of the spectrum component attributable to Fe inFeAl6 (A) compared with that at time t=0 (A0) was ®tted by Avrami kinetics:

A � A0 expÿ�kt�n �8�(i.e. Arrhenius kinetics are a special case where n=1, see Eq. (6)), where k is arate constant and n is the Avrami exponent. Fig. 46 shows the decrease of areafraction (A/A0) at di�erent homogenization temperatures in the range 580±6358Cwith isothermal holding time. At temperatures up to 6358C the area fraction A/A0

decreases exponentially with time. Fig. 47 shows the decrease of area fractionA/A0 with isothermal holding time during homogenization at 6508C. At 6508C the

Fig. 46. Area fraction of FeAl6 as compared to that found in the spectrum of an as-cast sample versus

annealing time at 580±6358C. After Murgas et al. [76]. Reproduced from Metallurgical and Materials



variation of A/A0 is more sigmoidal in shape. MurgaÂ s argued that this change in

reaction kinetics indicated that there were relatively few Fe4Al13 nuclei at 6508C,given the only 5 K undercooling below its equilibrium eutectic temperature [12].

This illustrates the importance of having suitable nuclei to catalyse the

transformation, which will a�ect the value of the rate constant k. The Avrami

exponent for 580±6358C was calculated as 0.71, which yielded an activation

energy for transformation Qnet of 0320245 kJ molÿ1, slightly greater than

Kosuge's determination. MurgaÂ s concluded that Fe di�usion through the Al

matrix was not the rate controlling step for the FeAl64Fe4Al13 transformation,

and that the reaction had some degree of interface control as well. As stated

above, the passage of Fe from the dissolving FeAl6 into solid solution is unlikely

to be the rate determining step as FeAl6 readily dissolves [46]. Equally unlikely is

that the direct passage of Fe from FeAl6 to Fe4Al13 is the rate determining step,

as direct transformation of FeAl6 to Fe4Al13 is rarely observed [58]. The only

remaining possibility for the rate determining step is Fe di�usion to Fe4Al13 and

the passage of Fe from solid solution into the Fe4Al13 lattice. This is considered in

more detail by Shillington [92] below.

Shillington [92] followed the FeAl64Fe4Al13 transformation in Bridgman

grown specimens grown at 1mmsÿ1, in which XRD detected only Al and FeAl6.

By analysing previously prepared known composition mixtures of Fe4Al13 and

FeAl6 extracted from the Al matrix Shillington demonstrated that levels <5% of

a given phase could not be readily detected using XRD, i.e. <0.05% of a phase

present in the bulk. XRD would therefore be incapable of detecting low levels of

any pre-existing Fe4Al13, which may act to catalyse the transformation (MurgaÂ s).

Fig. 47. Area fraction of FeAl6 as compared with that found in the spectrum of an as-cast sample

versus annealing time at 6508C. After Murgas et al. [76]. Reproduced from Metallurgical and Materials



The Bridgman grown specimens were isothermally aged for 1±128 h in thetemperature range 550±6258C. In the specimens, the FeAl6 was locatedpredominantly at the Al cell boundaries. SEM of extracted particles revealedpinching o� and spherodisation of the transforming FeAl6 after 4 h at 6008C.Using SEM in-situ, Fe4Al13 was observed to precipitate predominantly on the Alcell boundaries forming large faceted particles after 32 h at 6008C. Shillingtonfound no evidence for nucleation of Fe4Al13 directly on FeAl6 when examiningextracted particles. Fitting Arrhenius kinetics to his experimental data, Shillingtondetermined a value for the activation energy for the transformation ofQnet=280230 kJ molÿ1, in close agreement with that previously obtained byKosuge et al. of 290240 kJ molÿ1. Given the observed ease with which FeAl6spherodized and dissolved, and lack of evidence for direct nucleation of Fe4Al13on FeAl6, Shillington, as Kosuge had done, suggested that the transformationproceeded by a precipitation-dissolution mechanism, the rate determining stepbeing di�usion of Fe through the Al matrix to the precipitating Fe4Al13.

Shillington also investigated the precipitation of Fe from solid solution bygrowing pre-existing Fe4Al13 particles in the temperature range 550±6008C.Shillington followed the change in Fe solid solution levels using thermoelectricpower (TEP) measurements. He determined the activation energy for the growthof Fe4Al13 on pre-existing nuclei to be 207220 kJ molÿ1, 070 kJ molÿ1 lowerthan his determination for the activation energy of the transformation of FeAl6 toFe4Al13, Qnet. This di�erence in activation energies had also been observed byLendvai, although both of his determinations were lower than those ofShillington. Shillington argued that if Fe atom attachment kinetics were the ratedetermining steps for both the precipitation of Fe4Al13 from solid solution and thetransformation of FeAl6 to Fe4Al13, then the activation energies for the twoprocesses should be the same. MurgaÂ s's assertion that his determination of Qnet

was higher than that of Kosuge's due to an additional activation energyassociated with atomic attachment kinetics at the Al/Fe4Al13 is therefore incorrect,and it seems likely that MurgaÂ s's determination was infact the same as bothKosuge's and Shillington's, within the errors of his experimental set-up.Shillington therefore argued that Fe di�usion through the Al matrix was therate determining step in the precipitation of Fe4Al13 from solid solution,i.e. that Qdi�1207 kJ molÿ1. This is lower than Hood's earlier determination of260 kJ molÿ1, but Shillington argued that due to the high sensitivity of TEPmeasurements to low levels of Fe in solid solution that his method had smallererrors.

Shillington explained the 070 kJ molÿ1 di�erence between Qdi� and Qnet asfollows. In di�usion rate controlled transformations the ¯ux of solute atomsbetween the dissolving metastable phase and the precipitating stable phase (J)depends not only on the di�usion coe�cient of the solute in the matrix D, butalso the solute concentration gradient between the two phases (DC/Dx). Aconcentration di�erence exists because the Fe solid solution concentrationsbetween Al in equilibrium with Fe4Al13 and Al in equilibrium with FeAl6 are notthe same. Shillington showed that these concentrations are temperature dependent,


and can be ®tted by an Arrhenius relation of the form:

DC � DC0 exp�ÿQc=RT� �9�where Qc is the activation energy associated with the change of concentrationdi�erence with temperature, which Shillington determined to be 170 kJ molÿ1.Consequently, the net ¯ux of solute atoms can be written as:

J�T� ' DDC=Dx � �D0DC0=Dx� exp�ÿQdiff �Qc�=RT� �10�Consequently Qnet is equal to the sum of the activation energies for solutedi�usion through the matrix Qdi� and the activation energy associated with theconcentration di�erence between the dissolving and precipitating species Qc, i.e.Qnet=Qdi�+Qc.

7.1.2. Transformation rateAs shown in Eqs. (6) and (7) above, the rate of transformation depends not

only on the activation energy Q and temperature T, but also on the value of therate constant k0.

7.1.2.1. Microstructural scale. Tonejc [108] studied the decomposition behaviour ofsplat quenched Al±(3.6±10) at% Fe. Isothermal annealing produced a microstruc-ture of Al and FeAl6. Further annealing was then carried out at temperatures inthe range 300±5508C. FeAl6 did not begin to transform until after 76 h at 3008Cand was still present after 670 h. FeAl6 began to transform within 30 min at 4008Cand had all transformed within 16 h. FeAl6 had completely transformed within15 min at 5008C.

Hughes and Jones [46] unidirectionally solidi®ed Al±3.0 wt% Fe at 1.24 mmsÿ1 to produce specimens containing Al and FeAl6 rods parallel to the directionof growth. These specimens were homogenized isothermally for up to 1000 h inthe range 500±6408C. Fig. 48a and b shows typical SEM micrographs of deepetched sections transverse to the growth direction after: (a) 0 h and; (b) 15 h ofhomogenization at 5008C. As Fig. 48a and b shows, with increasinghomogenization time, coarsening of the FeAl6 rods occurs, commencing in thecentres of the eutectic cells and proceeding out to the boundaries. Hughes andJones also observed that the FeAl6 rods developed well de®ned facets duringhomogenization. Fig. 49a and b shows typical SEM micrographs of deep etchedsections parallel to the growth direction after: (a) 1 h and; (b) 280 h at 5008C. AsFig. 49a and b shows, the FeAl6 rods simultaneously pinch o� and spherodize(Shillington). Complete disappearance of the FeAl6 does not occur in 1000 h at5008C, but takes place in 500±750 h at 6008C. Concurrent with the morphologicalchanges to FeAl6, Fe4Al13 grows from ®rst the eutectic grain boundaries, and thenthe eutectic cell boundaries. Evidence for growth of the Fe4Al13 from within theeutectic cell centres was not seen until the later stages of transformation,suggesting that in Hughes and Jones's samples Fe4Al13 nucleates preferentially atthe eutectic grain and cell boundaries. Fe4Al13 was not seen to nucleate directly on


Fig. 48. Scanning electron micrographs in sections transverse to Bridgman growth direction showing

coarsening of FeAl6 on isothermal holding at 5008C after: (a) 0 h and (b) 15 h. After Hughes and

Jones [46]. Reproduced from the Journal of Materials Science by kind permission of Kluwer Academic

Publishers.


Fig. 49. Scanning electron micrographs parallel to Bridgman growth direction show necking down,

pinching o� and spherodisation of FeAl6 rods on isothermal holding at 5008C for (a) 1 h and (b) 280 h.

After Hughes and Jones [46]. Reproduced from the Journal of Materials Science by kind permission of

Kluwer Academic Publishers.


FeAl6, consistent with Kosuge's observations, and a sheath of Al was present

between the two Al±Fe eutectics throughout the transformation. The nucleation of

Fe4Al13 is faster than the disappearance of FeAl6, occurring within 13±14 h at

5008C and 4±5 h at 6008C.Of note is that the disappearance of FeAl6 in Hughes and Jones's samples is

some three orders of magnitude slower than in Tonejc's samples, and the

nucleation of Fe4Al13 some two orders slower. The increased rate of Fe4Al13precipitation in Tonejc's samples suggests that supersaturation of Fe in solid

solution stimulates nucleation of Fe4Al13. The faster transformation rate of FeAl6to Fe4Al13 observed by Tonejc may also have arisen from the ®ner microstructural

scale in the splat quenched samples than in the unidirectionally solidi®ed samples

of Hughes and Jones. This reduces the mean Fe di�usion distance and increases

the area through which FeAl6 dissolves, hence reducing the total time required for

transformation. The transformation rate in Kosuge's samples was also faster than

in Hughes and Jones's, and this can again be explained by the ®ner microstructure

in the DC cast samples compared with the unidirectionally solidi®ed samples. The

transformation rate in Shillington's unidirectionally solidi®ed samples was also

faster than that of Hughes and Jones. Shillington's samples were hypoeutectic

however, unlike the hypereutectic samples of Hughes and Jones. In both the

hypoeutectic Bridgman grown specimens of Shillington, and the DC cast

specimens of Kosuge, the aluminide phases are likely to be ®ner and more

separated in the as-solidi®ed condition than in a hypereutectic Bridgman grown

specimen where the aluminides are primary phases. This re®nement of

microstructure may explain why the transformations as measured by Shillington

and Kosuge proceeded more rapidly than that measured by Hughes and Jones.

7.1.2.2. E�ect of cold work. Kosuge and co-workers also noted that cold working

broke up the FeAl6 phase and accelerated the transformation. This again increases

the surface area through which FeAl6 can dissolve. Shoji and Fujikura [93] noted

that the introduction of dislocations by cold working did not promote the nuclea-

tion of Fe4Al13, as recovery and recrystallization removed the dislocations again

prior to any signi®cant precipitation of Fe4Al13 at higher temperatures.

7.1.2.3. Presence of pre-existing nuclei. The data of Shillington [92] predicts a half-

life of transformation of02.2 h at 6358C and 01.6 h at 6508C in Bridgman grown

Al±0.5 Fe alloys grown at 1 mm sÿ1. Allen et al. [4] have since observed 090%

transformation however during continuous heating at 2 K minÿ1 between 6428Cand the equilibrium eutectic melting temperature of 6558C, i.e. within 6.5 min, in

identical specimens. This may indicate a rapid acceleration in transformation kin-

etics at temperatures close to the equilibrium eutectic melting temperature. How-

ever, MurgaÂ s et al. [76] has previously suggested that isothermal transformation

would be slower at high temperatures due to the di�culty in nucleating Fe4Al13 at

low undercoolings. However, Fe4Al13 nuclei could have already been present in

Allen's specimens, as the specimens were heated through a range of undercoolings


prior to melting. MurgaÂ s noted that the presence of Fe4Al13 nuclei appeared to

catalyse the transformation reaction.

Thomas et al. [105] have observed the transformation during growth in

Bridgman grown Al±0.5±1.5 wt% Fe samples grown at the lower speeds of 0.25±

0.50 mm sÿ1. FeAl6 is observed to grow from the liquid at its growth front, then

partially transforms in the solid state in under 12 s to Fe4Al13 035±50 K below the

eutectic front. Fig. 50 shows two typical TEM micrographs of microtomed

sections taken transverse to the growth direction from close to the eutectic growth

front (left-hand side) and far behind the front (right-hand side). Fig. 51 shows

three typical SEM micrographs of sections taken transverse to the growth

direction from close to the eutectic front (left), 4 mm behind the front (centre) and

far behind the front (right). As both Figs. 50 and 51 show, no precipitation of

Fe4Al13 remote from the FeAl6 or changes in FeAl6 morphology on

transformation were observed. Thomas proposed that Fe4Al13 nucleated on Al

dendrites ahead of the FeAl6 growth front, contrary to the BaÈ ckerud's and Allen's

assertion that Al is not a suitable nucleant for Fe4Al13. These nuclei undercooled

by some 35±50 K before growing. The transformation was proposed to proceed

via di�usion in advance of a Fe4Al13/FeAl6 interface, there being no Fe di�usion

through the matrix thus minimizing the distance of solute transport and

accounting for the rapidity of the observed transformation. The formation of

these nuclei was proposed to occur over the solidi®cation velocity range 00.17±

0.5 mm sÿ1. In similar specimens grown at 1 mm sÿ1 neither Thomas nor

Shillington observed the formation of any Fe4Al13 in the as grown condition. On

subsequent isothermal holding, full transformation took 8 h at 6258C, suggestingthat the Fe4Al13 nuclei do not form at a growth velocity of 1 mm sÿ1, and that in

the absence of these nuclei the transformation is signi®cantly slower.

Discrepancies in transformation rate measurements between seemingly identical

specimens highlights particularly that the transformation mechanisms, the

corresponding rate determining steps, what determines which mechanism

Fig. 50. TEM micrographs of microtomed sections showing intermetallics examined in regions: near the

eutectic front (all particles are FeAl6) (left-hand image) and far behind the eutectic front (50% of

particles are FeAl6) (right-hand image). After Thomas et al. [105]. Reproduced by kind permission of

R. Thomas.


dominates, and the rate constant for that mechanism depend not only onsolidi®cation rate and hence microstructural scale and Fe solid solution levels, butalso the thermal history of the specimen and the presence or otherwise ofnucleants that catalyse the transformation.

7.2. The FeAlm ÿ4 Fe4Al13 transformation

Kosuge and co-workers [58] homogenised Al±0.8 wt% Fe±0.1 wt% Mgcontaining only Al and FeAlm for 1 h at a constant temperature in the range 300±6408C. Fig. 52 shows schematically their results for the variation of thenormalized intensities of X-ray peaks from FeAlm and FeAl3 with homogenizationtemperature after 1 h of isothermal holding. With increasing homogenizationtemperature, the proportion of FeAlm after 1 h as determined by XRD began todecrease from 5008C, and had completely converted to Fe4Al13 by 6408C. NoFe4Al13 was detected until 05608C, suggesting that in the range 500±5608C FeAlmdisappeared without any detectable formation of Fe4Al13, although XRD wouldnot be capable of detecting the Fe4Al13 if it were present as a dispersion of ®neprecipitates. An Arrhenius relation was ®tted to the normalized XRD intensitydata for various homogenization times and temperatures, as shown in Fig. 53.This ®tting yielded an activation energy for the FeAlm4Fe4Al13 transformationequal to that for the FeAl64Fe4Al13 transformation, indicating that Fe di�usionthrough the matrix was again the rate determining step. Kosuge proposed that thetransformation mechanism was the same as that that he had proposed forFeAl64Fe4Al13, namely that Fe di�uses from the dissolving FeAlm particles tothe remotely precipitating Fe4Al13 particles. Little evidence for the nucleation ofFe4Al13 on FeAlm was seen.

Fig. 51. SEM micrographs in sections transverse to the Bridgman growth direction showing similar

intermetallic morphologies in regions: (a) near to the eutectic front (all particles are FeAl6); (b)04 mm

behind the eutectic front (80% of particles are FeAl6) and; (c) far behind the eutectic front (50% of

particles are FeAl6). After Thomas et al. [105]. Reproduced by kind permission of R. Thomas.


Griger et al. [37, 38] also followed the transformation of FeAlm to Fe4Al13 onhomogenization of DC cast 1xxx alloys in the temperature range 450±6208C usingXRD, for homogenization times in the range 0.17±24 h. As was the case for theFeAl64Fe4Al13 transformation, the FeAlm4Fe4Al13 transformation proceeded bya dissolution-reprecipitation mechanism also, although somewhat faster thanFeAl64Fe4Al13. Transformation of FeAlm commenced after only 2 h at 400±4508C, and was completed within 2±3 h at 5008C.

7.3. E�ect of Si on transformations of metastable Al±Fe phases

Kosuge and co-workers [58] homogenized Al±0.5 wt% Fe±(0.08±0.11) wt% Sialloys, each containing Al, Fe4Al13, FeAl6 and FeAlm. The transformations ofboth metastable Al±Fe phases occurred more rapidly than in Si-free alloys, which

Fig. 52. Schematic of variation of normalized X-ray intensity of FeAlm and Fe4Al13 with

homogenization temperature in A1±0.8 wt% Fe alloy. After Kosuge [58].

Fig. 53. Schematic of variation of normalized X-ray intensity of FeAlm and Fe4Al13 with

homogenization time (hours) in Al±0.8 wt% Fe alloy for the temperatures shown. After Kosuge [58].


Kosuge attributed to the presence of Fe4Al13 as-cast, which then acted as pre-existing nuclei for the transformations. In the Al±0.5 wt% Fe±0.08 wt% Si alloy,complete transformation of FeAlm and FeAl6 took place in 5 and 8 h at 6008C,respectively, decreasing to the order of minutes at temperatures >6308C. Kosugedetermined that this acceleration was due to a change in the rate constant term ko(see Eq. (6)) as opposed to any change in the activation energy for transformationQ. The acceleration of the Fe4Al13 ÿ4 FeAl6 transformation in the presence of Sihas also been widely interpreted to be due to the decrease in the thermodynamicstability of the FeAl6 lattice on addition of Si, the solubility of Si in FeAl6 beingvery low [38, 58, 61, 112], as well as due to the introduction of pre-existing Fe4Al13nuclei as cast (see above). Ping et al. [80] observed the transformation of FeAl6,FeAlm and FeAlp to Fe4Al13 within 10 min at 6008C in DC cast Al±0.51 wt% Fe±0.13 wt% Si.

Shillington's [92] consideration of the e�ect of solute concentration gradient onthe transformation rate (Eq. (8), end of Section 7.1.1.2) can also be used toexplain the acceleration of transformations in the presence of Si. Di�erences in Sicontent between Si-tolerant and Si-intolerant phases such as Fe4Al13 and FeAl6would act to increase the solute concentration (and hence di�usive ¯ux) betweenthe two species, and hence accelerate the transformation rate between FeAl6 toFe4Al13.

7.4. Transformations involving ternary Al±Fe±Si phases

Westengen [112] noted that the formation of as cast phases such as a0 whichproduce a ®r-tree structure cannot always be removed by homogenizing. It isequally important therefore to understand the kinetics and mechanisms oftransformations involving Al±Fe±Si phases.

Griger et al. [37] examined the e�ect of 0.17±24 h of homogenising at 450±6308C on DC cast Al±0.5 wt% Fe±0.2 wt% Si and Al±0.5 wt% Fe±1.0 wt% Si,containing cubic a-AlFeSi and b-AlFeSi, respectively. Cubic a-AlFeSi took 1 weekto transform to Fe4Al13 at 5758C, and even at 6058C (015±25 K below themeasured a peritectic invariant point) cubic a-AlFeSi was still detected after 24 h.Complete transformation took place in only 10 min at 6308C, however, indicatinga steep temperature dependence of reaction rate. During the transformation thecubic a-AlFeSi ®rst coarsened and then dissolved, while Fe4Al13 precipitated fromsolid solution as acicular precipitates which then coarsened and changed shape.This mechanism appears to be similar to that for the transformation of themetastable Al±Fe phases. b-AlFeSi did not transform. Some particles dissolved,others coarsened, and secondary b, 02±5 mm in length, precipitated from solidsolution after 6 min at 6208C [109]. Griger suggested that b-AlFeSi did nottransform as it was an equilibrium phase for Al±0.5 wt% Fe±1.0 wt% Si. BothGriger and Turmezey et al. [109] noted that the (00 l) type electron di�ractionpattern re¯ections for b increased with intensity on increasing homogenisation,while the other re¯ections diminished. This indicates an ordering process occurs


on (001) type planes. Eutectic Si present in both of Griger's alloys as castdissolved back into solid solution on heat treatment.

Ping et al. [80] studied the e�ect of homogenization treatments of 10 min to24 h at 6008C on DC cast Al±(0.26±0.51) wt% Fe±0.13 wt% Si. On ageing for24 h, the q1 variant of a-AlFeSi transformed to the q2 variant.

8. E�ect of impurities on phase formation in Al±Fe and Al±Fe±Si alloys

Previous workers have attributed discrepancies between DC cast and Bridgmangrown data (ref. Section 6, Table 5) to be due to the e�ect of impurity elementson phase selection. Ping et al. [80] attributed the e�ect of impurities to beresponsible for the scarcity of the metastable Al±Fe eutectics in his specimens atall observed cooling rates (01±10 K sÿ1). Skjerpe's [96] observation of FeAlmforming at cooling rates down to 1 K sÿ1 was also proposed to be due toimpurities. Brobak and Brusethaug [16] were unable to reproduce by Bridgmangrowth the observed dominance of FeAl6 and FeAlm in DC cast billets ofidentical Fe and Si content [17]. Evans et al. [32] have since demonstratedhowever that grain re®ner addition appears to be necessary for FeAlm formationduring Bridgman growth (Section 9.2). This may explain the widely observedoccurrence of FeAlm in commercial DC cast ingots and billets which containgrain re®ner. In his experiments on Bridgman growth, Shillington [92] alsonoted that phase content varied for a given solidi®cation velocity depending onwhether or not a solid seed or a fully molten specimen was used at the startof solidi®cation. This suggests that phases already present in the solid seedin¯uence the nucleation of phases in the overlying liquid at the start of Bridgmangrowth.

Bridgman growth experiments are chie¯y concerned with the variation of phasecontent with changing growth conditions. The demonstrated sensitivity of phaseselection to alloy impurity content and grain re®ning additions suggests thatnucleation can be as important as growth in in¯uencing phase selection.

Due to the low bulk alloying content of 1xxx series alloys, impurity elementscan have a large e�ect on phase selection. Allen et al. [5] rapidly solidi®ed a rangeof model and commercial Al±0.3±Fe±0.1 Si alloys to produce microstructures ofsubmicron secondary phases entrained in an Al matrix. The secondary phaseswere then melted using slow heating rate calorimetry, resulting in a dispersion ofeutectic liquid droplets entrained in a solid Al matrix. Fig. 54 shows typicalcalorimeter endotherms obtained at a heating rate of 2 K minÿ1 during themelting of the secondary phases from a variety of model alloys with deliberateadditions of impurities. The similarity in endotherms indicates that impurities donot signi®cantly a�ect melting behaviour. The microstructure obtained once allthe secondary phases have been melted is e�ectively a re®ned version of thepartially molten structure that exists during the ®nal stages of conventionalcasting, i.e. of separated liquid puddles between primary Al dendrite arms. Fig. 55


shows a typical SEM micrograph of the microstructure of the Al±0.3±Fe±0.1 Sibase alloy quenched from a temperature just above the point at which all thesecondary phases were molten. The ®ne scale eutectic structures visible in Fig. 55indicate the location of the eutectic liquid puddles prior to the quench. The ®nescale eutectic structure forms during the quench. Due to the divided nature of theeutectic liquid a large number of nucleants are required to catalyse nucleation ofsolidi®cation of the liquid on cooling.

The dispersion technique has been used successfully before to study nucleationunder impurity free conditions, by segregating impurities into an insigni®cantfraction of the total liquid. Allen et al. however made deliberate impurityadditions to high purity (>99.995 wt% Al base) alloys to study the e�ect ofimpurities on the nucleation of solidi®cation. Fig. 56a shows typical calorimeterexotherms obtained on resolidi®cation at 2 K minÿ1 after the melting of thesecondary phases from a variety of Al±0.3±Fe±0.1 Si alloys with varying basepurity, Al±Ti±B grain re®ner, V and P levels (as indicated in ®gure). Whereasimpurities do not signi®cantly a�ect melting behaviour (Fig. 54), Fig. 56a showsthat they do a�ect solidi®cation behaviour. The exotherm of solidi®cation onset inthe range 0652±6548C labelled `1' in Fig. 56a corresponds to the resolidi®cationof the equilibrium phase in Al±0.3±Fe±0.1 Si, Fe4Al13, which is present in allalloys. Of particular note is that in the grain re®ned high V commercial alloy

Fig. 54. partial melting endotherms from high purity model Al 1xxx melt spun ribbons with various

additions.


FeAlm also forms, labelled `2' in Fig. 56a. Fig. 56b shows typical calorimeterexotherms obtained on resolidi®cation at 2Kminÿ1 after the melting of thesecondary phases from a variety of high purity Al±0.3±Fe±0.1 Si model alloyswith varying additions of V and Al based grain re®ners (as indicated on Figure).The exotherm labelled `1' in Fig. 56b corresponds to the resolidi®cation ofFe4Al13, which is present in all alloys (as in Fig. 56a). Additions of V impurity aslow as 10±100 ppm (typical of commercial alloys) combined with a 1:1000addition of Al±Ti±B grain re®ner promote the formation of FeAlm, labelled `2' inFig. 56b, at bulk cooling rates as low as 2 K minÿ1. FeAlm in turn has beenproposed to be one of the phases responsible for the formation of the ®r treezone, as previously discussed in Section 6.

Maggs et al. [68] proposed that impurities can alter both the nucleation andgrowth kinetics of a given phase. Impurities or grain re®ner additions (Section 9.2)may promote nucleation by providing nucleation sites for a given phase.Impurities may segregate to the growth front produce solutal undercooling andhence depress growth velocity. Conversely, impurities may promote twinningwhich in turn promote the growth of a faceted eutectic (e.g. Fe4Al13), byproviding a higher density of re-entrant edges and corners to which atoms canattach more easily.

Impurity elements can also a�ect solid solution contents [106], recrystallizationtemperatures, mechanical properties, surface electrochemistry and electricalresistivity [24, 30]. The transformation rates of metastable phases on heattreatment can also be a�ected, as was noted in Section 7.3.

Fig. 55. Secondary electron micrograph of Fe and Si intracellular phases entrained in an Al matrix,

formed when a melt-spun ribbon is partially remelted and resolidi®ed.


Interpreting the e�ect of given impurities on phase selection is not astraightforward matter, due to variations in base alloy purity, bulk composition,impurity addition levels and solidi®cation parameters used by various workers.Low level impurity additions (<1000 ppm) to high purity base alloys (>99.995wt% pure Al base) are the easiest to interpret, as the e�ect of the impurityaddition is then less likely to be masked by that of one or more impurities presentin the base alloy. These results are summarized in Table 6.

Fig. 56. (a) Resolidi®cation exotherms from Al±0.3 wt% Fe±0.1 wt% Si melt-spun ribbons of di�erent

purities. Peak `1'=Fe4Al13, peak `2'=FeAlm. (b) Resolidi®cation exotherms from high purity Al±0.3

wt% Fe±0.1 wt% Si melt-spun ribbons with di�erent mixed additions of V and grain re®ner. Peak

`1'=Fe4Al13, peak `2'=FeAlm.


Impurity elements that are not strongly partitioned to the secondary/ternary phases, e.g. Mn, are not anticipated to have an e�ect at low leveladditions (<500 ppm). Higher levels of Mn (e.g. 1±2 wt%), such as those presentin 3xxx series alloys do dramatically a�ect phase contents, however, by alteringthe relative free energies and hence thermodynamic stabilities of the phases.Conversely, impurities that are strongly partitioned to the interdendritic liquidduring the ®nal stages of solidi®cation can a�ect solidi®cation behaviour andphase selection even at very low bulk levels, as demonstrated by the e�ect of 2ppm of P in Al-Si alloys [40], and can even form compounds in their own right(e.g. AlP). Hsu et al. [44] have also demonstrated that 0.4 wt% Ca in 6xxx seriesAl alloys can promote the formation of a-AlFeSi via the formation of CaAl2Si2during solidi®cation.

Fig. 56(b).


9. E�ect of grain re®ner additions on Al±Fe and Al±Fe±Si alloys

A full review of primary grain re®nement of Al alloys is outside the scope ofthis review, and has previously been given by McCartney [67]. A brief summary isgiven however as various workers [6, 32, 37, 58, 68, 104] have presented evidencethat grain re®ners can in¯uence secondary/ternary phase selection.

9.1. Proposed mechanisms of primary Al grain re®nement

Determining the mechanism by which grain re®ning additions work is not astraightforward matter. The rapidity of subsequent growth after nucleation due tothe high atomic mobility in the melt, the submicron size of the nucleants, theobservation that only 1±2% of grain re®ner particles actually nucleate Al [39], andthe non-isothermal casting conditions all act to hinder determination of thenucleation mechanism [89].

Grain re®ner master alloys, e.g. Al±Ti±B, contain three phases: Al, TiAl3 andTiB2. Two chief mechanisms for the action of the TiAl3 and TiB2 particles thatare released on addition of a grain re®ner master alloy to an Al melt areproposed [67]:

i. The `peritectic' theory: Crossley and Mondolfo [26] proposed that TiAl3particles produce solid Al via the peritectic reaction;

L� TiAl3 4 a-Al

Table 6

E�ect of trace elements/impurity additions to high purity alloys

Phase Phase promoted

([) or inhibited (x)

Elements References

Fe4Al13 [ Mg, Cu, V, Ti, Ca Tezuka and Kamio [104], Maggs et

al. [68], Wang et al. [111]

FeAl6 [ Mn BaÈ ckerud [12], Kosuge [58], Maggs et

al. [68]

FeAl6 x Si Griger et al. [38], Langsrud [61]

FeAlx [ V Wang et al. [111]

FeAlm [ Si, Cu, V, TiB2* Kosuge [58], Griger et al. [38], Tezuka

and Kamio [104], Maggs et al. [68],

Evans et al. [32], Allen et al. [5]

FeAlm x Mg Tezuka and Kamio [104]

Fe2Al9 [ Co Tezuka and Kamio [104]

a [ Mn, V, Mo, W, Cr, Cu, TiB2* Munson [75], Barlock and Mondolfo [14],

Skinner et al. [95], Wang et al. [110],

Maggs et al. [68]

Note: See section 9.2.


This solid Al then acts as a highly e�cient nucleant for further Al solidi®cation.However, once released from the master alloy the TiAl3 particles are in a bulk

hypoperitectic composition, and thus start to dissolve.

ii. The `carbide-boride' theory: Cibula [21] proposed that the insoluble TiB2

particles (or TiC particles from Al±Ti±C master alloys, or Al±Ti with C impurity)directly nucleate Al, having a low lattice mismatch with Al.

Neither theory explains the importance of excess solute Ti above that required

to form TiB2 in Al±Ti±B additions, nor the poisoning e�ects of impurities such asZr or V in Al±Ti±B. McCartney [67] noted that the excess Ti may take part in aTiAl3±TiB2 association, slowing TiAl3 dissolution by locally preserving TiAl3either in shells or cavities of TiB2 particles, or adsorbing onto TiB2 particles. Nñssand Berg [77] and Jones [53±55] advanced a theory, referred to by Jones as

hypernucleation, to account for the de®ciencies of the two models above. Jonesproposed that excess solute Ti atoms in the Al melt di�use down a chemicalpotential gradient to the interface between the TiB2 and the liquid, the free energy

of a Ti atom at the interface being lower than a Ti atom in solution in the melt,forming a layer of Ti in Al solid solution on the surface of the TiB2 particle. This

layer then acts as a highly e�cient nucleant for Al, with little mismatch betweenthe Al(Ti) solution phase and solid Al. Conversely, the incorporation of dissimilar

sized elements such as Zr into this adsorbed solution layer can reduce itsnucleation potency for pure Al, by increasing the lattice mismatch between thelayer and solid Al.

Marcantonio and Mondolfo [70], BaÈ ckerud [12] and Cornish [25] proposed

instead that solid Al was nucleated by a TiAl3 layer adsorbed on the TiB2

particles. Schumacher and Greer [87±91] have observed this nucleation process,albeit in the devitri®cation of rapidly solidi®ed amorphous Al±Ni±Y based alloys.

Fig. 57a±d shows: (a) bright ®eld TEM micrograph in [110] orientation of a TiB2

particle; (b) dark ®eld using (001) TiB2 spot; (c) dark ®eld using (111) Al spot

showing Al crystallites nucleated on the TiB2 particle; and (d) dark ®eld usingstreak passing through the (000) spot corresponding to the thin TiAl3 layer asvisible in (d). Based on this evidence, Schumacher and Greer therefore proposed

that the nucleation of crystalline Al from the amorphous matrix occurs on TiB2

particles coated with a thin TiAl3 layer only a few monolayers thick. The TiAl3layer has a close lattice match to the Al and is also coherent to the TiB2. In theabsence of excess Ti, a pronounced decrease in nucleating e�ciency of crystalline

Al from the amorphous alloys is seen, as the TiB2 particles then lose their TiAl3coating. It is proposed therefore that during conventional solidi®cation ofcommercial Al alloys the TiB2 particles become coated by TiAl3 in the melt, but

only in the presence of excess solute Ti. Nñss and Berg [77], Schumacher andGreer [87±91], Mohanty and Gruzleski [73], and Rao et al. [83] proposed that

certain elements can dissolve in this TiAl3 layer, altering its thermodynamic andstructural characteristics, or forming complex compounds, and hence a�ectinggrain re®ning e�ciency. This would then explain the poisoning of Al±Ti±B re®ner

alloys by elements such as V and Zr.


9.2. The roÃle of grain re®ners on secondary/ternary phase selection

Grain re®ning additions are proposed to a�ect secondary/ternary phaseselection in three ways. Firstly, TiB2 or TiC particles that do not nucleate Al maybe partitioned into the interdendritic spaces (being insoluble in Al), where theymay a�ect the solidi®cation of the secondary/ternary phases [37, 58]. The localchemistry of these interdendritic spaces (i.e. solute element and impurity levels)may prove to be as important for secondary phase selection as it is in determiningthe e�ect of grain re®ning particles on primary Al grain re®nement (Section 9.1).Secondly, primary grain re®nement may result in a greater number density of

Fig. 57. TEM micrographs of a boride particle and corresponding h110iTiB2 zone axis di�raction

pattern: (a) bright ®eld; (b) dark ®eld image using (001)TiB2 spot; (c) dark ®eld image using (111)A1 spot

showing aluminium crystals nucleating on the boride particle and (d) dark ®eld image using the streak

passing through (000) corresponding to the thin TiAl3 layer between the A1 crystallites and the boride

particle, as imaged. After Schumacher and Greer [91]. Reproduced from Light Metals 1996 by kind

permission of TMS and ASM International.


interdendritic liquid spaces towards the ®nal stages of solidi®cation. Withincreasing division of the liquid volume, nucleation and hence impurities play amore important role in in¯uencing secondary phase selection. This is the basis forthe entrained droplet technique [3]. Thirdly, primary grain re®nement may changethe shape of the interdendritic liquid channels (e.g. from long channels betweencolumnar dendrites to more convoluted shapes between equiaxed grains), forcingthe secondary phases that form in these channels to change their growthmorphology. This may in¯uence which is the preferred secondary phase under agiven set of solidi®cation conditions [71].

Kosuge [58] and Griger et al. [37] proposed that grain re®ners increased thenumber density of nucleation sites for phases such as FeAlm to explain theobserved promotion of FeAlm in grain re®ned DC castings. Tezuka andKamio [104] also observed the promotion of FeAlm and a-AlFeSi inunidirectionally solidi®ed Al±0.3±0.5 wt% Fe±0±0.4 wt% Si with grain re®neraddition. Maggs et al. [68] similarly proposed that there may be a small latticemismatch d between the hexagonal lattice of TiB2 and metastable phases withorthogonal crystal axes such as a-AlFeSi and FeAlm in 1xxx alloys, explaining theobserved promotion of these phases in DC simulator casts. Evans et al. [32] havenoted that Al±Ti±B addition to Bridgman grown Al±0.3 wt% Fe±0.1 wt% Sipromotes the formation of FeAlm at velocities >1.33 mm sÿ1. FeAlm is nothowever promoted in the absence of Si [71]. Allen et al. [6], using the entraineddroplet method, have demonstrated that Al±Ti±B addition to Al±0.3 wt% Fe±0.1wt% Si containing <100 ppm V impurity also promotes FeAlm, with or withoutan excess of Ti being present. FeAlm promotion also occurs on the addition ofAl±B or Al±Ti±C grain re®ners to Al±0.3 wt% Fe±0.1 wt% Si containing 0100ppm V. This suggests that, although it is the TiB2 present in Al±Ti±B added tocommercial alloys that is involved in the promotion of FeAlm, this promotion isnot grain re®ner speci®c. Also, as this promotion does not require there to be anexcess level of Ti this suggests that an adsorped layer of TiAl3 on the TiB2

particles in the commercial alloys is not responsible for nucleation of the FeAlm.As yet however no further details of the mechanism of FeAlm promotion by TiB2

addition have been elucidated.

10. Summary

A wide range of Al±Fe secondary and Al±Fe±Si ternary phases, bothequilibrium and metastable, are reported in the literature to form during thesolidi®cation of DC cast 1xxx Al alloys. Variations in casting parameters, namelylocal solidi®cation velocity and cooling rate, and alloy composition, can a�ectphase selection during the non-equilibrium solidi®cation experienced in DCcasting, by a�ecting the kinetics of nucleation and growth of each of the phases.

Variations in secondary and ternary phase content with position in the cast isparticularly undesirable in surface critical applications, where an uneven response


to surface treatment results in a defect known as a ®r tree zone. The occurence ofthe metastable FeAlm phase is particularly associated with this phenomenon.Recent work has shown that the formation of this phase during DC casting isassociated with one or more of: (i) the Si content of the alloy; (ii) the V impuritylevel and (iii) the presence of grain re®ning additions.

Unidirectional solidi®cation, as employed in the Bridgman growth technique,provides an experimental means of studying the variation of phase selection withchanges in solidi®cation velocity. Dispersion of the solidifying liquid volume, asemployed in the entrained droplet technique provides an experimental means ofstudying the roÃ le of nucleation in phase selection with variations in cooling rateand impurity content. Bridgman growth and the entrained droplet technique aretherefore complementary in the study of phase selection processes in 1xxx seriesAl alloys. The further study of the nucleation aspects of phase selection in thesealloys is especially important, given the reported a�ects of both impurity elementsand grain re®ner additions on phase selection.

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Intermetallic Phase Selection in 1xxx Al Alloys