General Considerations for Developing Creep-Resistant Al ...arc.nucapt.northwestern.edu/refbase/files/Considerations.pdfKeith Knipling General Considerations for Developing Creep-Resistant

General Considerations for Developing Creep-ResistantAl-Based Alloys

Keith Knipling

May 18, 2005

1 Introduction

Based on the behavior of nickel-based superalloys, which resist degradation of mechanical propertiesto approximately 75 % of their absolute melting temperature, it is conceivable that aluminum-basedalloys could be similarly developed which would be useful to temperatures in excess of 400 °C. Thecreep resistance of nickel-based superalloys is conferred by large volume fractions of precipitatedNi3(Al,Ti) ordered phases (γ′), which are of the L12 structure and are isomorphous with the fccNi-rich matrix (γ phase). An effective high temperature aluminum alloy should exhibit a similarstructural constitution, with suitable alloying additions to aluminum exhibiting the following quali-ties:

Capability to form strengthening intermetallic phases. As is true for γ′ in the nickel-based systems,a high-temperature aluminum alloy should contain a large volume fraction of a suitable dis-persed phase, which must be thermodynamically stable at the intended service temperature.These precipitated phases should also exhibit a similar structure to, and a low lattice misfitwith, the aluminum solid solution.

Low solid solubility in Al. A low equilibrium solid solubility at elevated temperatures is necessaryto prevent dissolution of the precipitated phases. By the lever rule, limited solubility also max-imizes the equilibrium volume fraction of the dispersed phase.

Low diffusivity in Al. Limited diffusivity of solute in Al should stifle volume diffusion-controlledcoarsening, allowing the precipitates to remain effective barriers to dislocations at elevatedtemperatures.

Ability to be conventionally cast. In the interest of affordability, the alloy should be easily producedvia conventional thermo-mechanical processing routes (i.e., it is castable).

This document is a broad ‘big picture’ review of the challenge of developing an Al-based alloy whichmeets all of the above criteria. For example, Ti is an extraordinarily slow diffuser in Al and forms a

1

Keith Knipling General Considerations for Developing Creep-Resistant Al-Based Alloys 2

D022 Al3Ti dispersed phase which may even be transformed to the high-symmetry cubic L12 struc-ture by alloying with fourth-period transition elements such as Cr, Mn, Fe, Co, Ni, Cu, and Zn [1–4].Moreover, there is no shortage of literature documenting the the thermal stability and strength athigh temperatures of Al-Ti alloys [5–23], again attributable to the very sluggish diffusion kinetics.However, the vast majority of these studies investigated alloys prepared by rapid solidification pro-cessing (RSP) [7–11] or mechanical alloying (MA) [12–23] techniques, both of which offer advantagesin obtaining homogenous fine dispersions of Al3Ti precipitates.

The reason why Al-Ti did not work in the present study is because these alloys are not amenableto conventional casting due to the terminal peritectic phase equilibria in this system. The presenceof a peritectic necessarily implies that the liquidus of the alloy progressively increases with solutecontent. Moreover, there is a tendency for solute to be scavenged in the form of primary Al3Tiphases, and preventing this properitectic precipitation involves cutting back the solute content ofthe alloy to the point that there is insufficient supersaturation to effect precipitation of Al3Ti duringpost-solidification annealing (the substantial solid solubility of Ti, Figure 6(d), does not help). Whatis worse, potent grain refinement is associated with primary precipitation of Al3Ti [24–32], which actas heterogeneous nuclei of Al during solidification (indeed, the most common use for Ti in industrialAl alloys is as a grain refiner, where the propensity for primary Al3Ti precipitation may be exploited).

These difficulties are unavoidable for any peritectic system, and it is hoped that this brief ‘big picture’analysis brings some of these points to light. We all know that the Al-Sc system offers great potentialfor developing a high-temeprature precipitation-strengthened Al-based alloy. Unfortunately, if onewishes to achieve similar performance with alternate systems, dealing with this problem of peritecticphase equilibria is practically inevitable.

2 Capability to Form Strengthening Phases

The first criterion listed above requires that any suitable system must have the capability to forma dispersed strengthening intermetallic phase. Trialuminide intermetallic compounds of the typeAl3X (X is an element of the transition metals, lanthanide, or actinide series) have many attractivecharacteristics including low density (they are nominally 75% Al on an atomic basis), high specificstrength, excellent thermal stability (they have generally very high melting points), and excellentoxidation resistance (again, mostly Al). Moreover, these Al3X intermetallics are directly analogousto the γ′ Ni3Al ordered phases in the Ni-based alloys.

It is fortunate that numerous alloying additions crystallize to form Al3X trialuminides, as showngraphically in the periodic table 1 of Figure 1. What is not so fortuitous, however, is that very few –only Sc, Er, Tm, Yb, Lu, U, and Np – form intrinsically stable cubic L12 phases.

1Groups of the periodic table will be denoted by the recommended ‘new IUPAC’ (International Union of Pure andApplied Chemistry) convention where columns are numbered with Arabic numerals from 1 to 18, corresponding to thenumber of s, p, and d orbital electrons. This system is less ambiguous than the older, but mutually confusing, schemes – i.e.,the old IUPAC and CAS (Chemical Abstract Service) designations – which label columns with Roman numerals followedby either the letter A or B. Group 3 (new IUPAC), Group III-A (old IUPAC), and Group III-B (CAS) refer to the same columnin the periodic table, and henceforth the new IUPAC designations will be used.


Trialuminide intermetallics

Orthorhombic (D011

), Monoclinic, or Unknown

Hexagonal (D019

, D018

, etc.)

Body-centered tetragonal (D022

or D023

)

Face-centered cubic (L12)

Fig. 1: Alloying additions to Al which form trialuminides of the type Al3X . The equilibrium structure of the intermetallic is indicated.See Tables 1 and 2 for more detailed information on the particular structures.

2.1 Trialuminides of the group 4 and 5 transition elements

As indicated in Figure 1, Sc is the only transition metal which forms a thermodynamically stableL12 phase. Just to the right of Sc, however, lie the group 4 (Ti, Zr, Hf) and group 5 (V, Nb, Ta)elements which crystallize with the closely-related D022 (or D023 for Al3Zr) structures. Moreover,these tetragonal crystal structures can be transformed to the high-symmetry cubic L12 crystal byalloying with fourth-period transition elements such as Cr, Mn, Fe, Co, Ni, Cu, and Zn [1–4].

Furthermore, in precipitation from supersaturated solid solution the formation of Al3Zr or Al3Hf inbinary alloys is typically preceded by precipitation of the metastable L12 structure. 2 The stabilityof the D022 structure relative to the L12 increases rapidly as the transition metal d-electron countincreases [38]. Therefore, the likelihood of forming L12 is higher for the group 4 (Ti, Zr, Hf) elementsthan it is for the group 5 (V, Nb, Ta).

2The existence of a similar L12 Al3Ti phase is somewhat disputed.


Table 1: Crystal structure data for Al3X trialuminides formed with the transition elements

Phase Group Pearson symbol Space group Strukturbericht Prototype Refs

Fourth period (3d) transition elements

Al3Sc 3 cP4 Pm3m L12 Cu3Au [33]

Al3Ti 4 tI8 I4/mmm D022 Al3Ti [34]

Al3V 5 tI8 I4/mmm D022 Al3Ti [34]

Al3Fe 8 mC102 C2/m · · · · · · [34]

Al3Co 9 · · · Unknown · · · · · · [35]

Al3Ni 10 oP16 Pmma D011 Fe3C [34]

Fifth period (4d) transition elements

Al3Y 3 hP8 P63/mmc D019 Ni3Sn [34]

Al3Zr 4 tI16 I4/mmm D023 Al3Zr [36]

Al3Nb 5 tI8 I4/mmm D022 Al3Ti [34]

Sixth period (5d) transition elements

Al3La 3 hP8 P63/mmc D019 Ni3Sn [35]

Al3Hf (α, < 650 °C) 4 tI8 I4/mmm D022 Al3Ti [34, 37]Al3Hf (β, > 650 °C) 4 tI16 I4/mmm D023 Al3Zr

Al3Ta 5 tI8 I4/mmm D022 Al3Ti [34]

Al3Re 7 · · · · · · · · · · · · [34]

Al3Ir 9 hP8 P63/mmc D018 Na3As [35]

2.2 Bottom line

If one wants to achieve an L12 phase the choices of alloying elements is fairly limited to distinctregions of the periodic table. Numerous intrinsically stable L12’s may be found within the lanthanideand actinide series. For transition elements, however, Al3Sc is the only one. The group 4 (Ti, Zr, Hf)and group 5 (V, Nb, Ta) form the closely-related D022 and D023 structures, which may be transformedto an L12 phase by alloying. This is more likely for the group 4 elements. In other words, stick closeto Sc in the periodic table.

3 Low Solubility

Fortunately this is generally not a problem for Al alloys; with a few exceptions [39] all elementsexhibit maximum solubilities of less than 1 at.% in Al. See, for example, the phase diagrams of thegroup 3, 4, and 5 elements alloyed with Al (Figure 6).


Table 2: Crystal structure data for Al3X trialuminides formed with elements of the lanthanide (rare earths) or actinide series

Phase Pearson symbol Space group Strukturbericht Prototype Refs

Lanthanides (rare earths)

Al3Ce hP8 P63/mmc D019 Ni3Sn [35]

Al3Pr hP8 P63/mmc D019 Ni3Sn [35]

Al3Nd hP8 P63/mmc D019 Ni3Sn [35]

Al3Pma [35]

Al3Sm hP8 P63/mmc D019 Ni3Sn [34, 35]

Al3Gd hP8 P63/mmc D019 Ni3Sn [35]

Al3Tb hR12 R3m · · · Pb3Ba [35]

Al3Dy (α, < 636 °C) hP16 P63/mmc D024 Ni3Ti [34]Al3Dy (β, > 636 °C) hR20 R3m · · · Al3Ho

Al3Ho hR20 R3m · · · Al3Ho [35]

Al3Er cP4 Pm3m L12 Cu3Au [35]

Al3Tm cP4 Pm3m L12 Cu3Au [35]

Al3Yb cP4 Pm3m L12 Cu3Au [34, 35]

Al3Lu cP4 Pm3m L12 Cu3Au [35]

Actinides

Al3Th hP8 P63/mmc D019 Ni3Sn [35]

Al3U cP4 Pm3m L12 Cu3Au [35]

Al3Np cP4 Pm3m L12 Cu3Au [35]

Al3Pu hP24 P63/mmc · · · · · · [35]

a Assumed similar behavior to Al-Nd and Al-Sm [35].

4 Low Diffusivity

The transition elements are anomalous diffusers in Al, characterized by high activation energies,high pre-exponential factors, and a wide range of variation of the diffusivities compared with self-diffusion in aluminum. Observed activation enthalpies (Q) and pre-exponential factors (D0) for allof the 3d transition elements and other selected 4d- and 5d transition elements are tabulated in Ta-ble 3. The majority of this data was obtained from two articles by the most prominent researchesin the field, Mehrer of Munster [40, 41] and Fujikawa of Tokyo [42], who have dealt with the prob-lem of transition metal diffusion in Al and provide the most authoritative theoretical interpretations.Data for some of the diffusing elements were also obtained from a somewhat older article by Gram-matikakis [43], which is also a review of various diffusers in Al. Also, the comprehensive compilationof Landolt-Bornstein [44] was consulted. The original source for the data, and the year in which itwas determined, is also provided in Table 3.


Table 3: Reported diffusion data for transition metal solutes in Al

Pre-exponential, D0 Activation enthalpy, Q D at 400 °C Referencesm2/s kJ/mol eV/atom m2/s Original Reference Cited in

Self-diffusion

Al 1.88× 10−5 126 1.31 3.12× 10−15 Least-squares fit to data in [44]


Sc 5.31× 10−4 173 1.79 1.98× 10−17 Fujikawa (1997) [45] [40, 42, 43]

Ti 1.12× 10−1 260 2.69 7.39× 10−22 Bergner (1977) [46] [40, 42, 43]

V 1.60 303 3.14 4.85× 10−24 Bergner (1977) [46] [40, 42, 43]

Cr 10.0 282 2.92 1.29× 10−21 Rummel (1995) [40] [40, 42]

Mn 8.7 × 10−3 208 2.16 6.24× 10−19 Rummel (1995) [40] [40, 42]

Fe 7.7 × 10−1 221 2.29 5.41× 10−18 Rummel (1995) [40] [40, 42]

Co 1.93× 10−2 168 1.74 1.76× 10−15 Rummel (1995) [40] [40, 42]

Ni 4.4 × 10−4 146 1.51 2.05× 10−15 Erdelyi (1978) [40, 42–44]

Cu 6.54× 10−5 136 1.41 1.54× 10−15 Fujikawa (1989) [40, 42, 44]

Zn 2.59× 10−5 121 1.25 1.05× 10−14 Peterson (1970) [40, 42–44]


Zr 7.28× 10−2 242 2.51 1.20× 10−20 Fujikawa (1973) [42–44]

Mo 1.4 × 10−3 250 2.59 5.52× 10−23 Bergner (1983) [42–44]


Hf 1.07× 10−2 241 2.50 2.11× 10−21 Minamino [42]

4.1 Diffusion data

Data for 3d transition elements and of 4sp non-transition elements (foreign atoms from the same rowof the periodic table) are depicted graphically in Figures 2 and 3. Figure 2 shows the activation en-thalpies versus the position of foreign element in the periodic table (i.e. valence). Figure 3 showsthe calculated diffusivities near the melting point of aluminum (660 °C) as well as two other temper-atures of interest (300 and 400 °C). It is evident that valence is an important factor in determiningdiffusivities and activation enthalpies. This trend is most obvious in Figure 3, where the calculateddiffusivity increases with increasing number of d electrons from V to Co by nearly six orders of mag-nitude at 660 °C. Indeed, the notion of a valence dependence on the diffusivity is not a new one, andwas first proposed by Lazarus in 1954 [47] and later modified by LeClaire [48]. While the so-called


1 . 0 0

1 . 2 5

1 . 5 0

1 . 7 5

2 . 0 0

2 . 2 5

2 . 5 0

2 . 7 5

3 . 0 0

3 . 2 5

S c T i V C r M n F e C o N i C u Z n G a G e1 0 0

1 2 5

1 5 0

1 7 5

2 0 0

2 2 5

2 5 0

2 7 5

3 0 0

Activ

ation

energ

y, Q

(kJ m

ol-1 )

Q f o r A l s e l f - d i f f u s i o n

Activation energy, Q (eV atom-1)

Fig. 2: Empirical activation enthalpies (Q) of 3d transitionelement solutes and of 4sp non-transition metal solutes inAl.

S c T i V C r M n F e C o N i C u Z n G a G e1 0 - 2 81 0 - 2 71 0 - 2 61 0 - 2 51 0 - 2 41 0 - 2 31 0 - 2 21 0 - 2 11 0 - 2 01 0 - 1 91 0 - 1 81 0 - 1 71 0 - 1 61 0 - 1 51 0 - 1 41 0 - 1 31 0 - 1 21 0 - 1 1

A l , 6 6 0 �C

A l , 4 0 0 �C

6 6 0 � C

4 0 0 � C

3 0 0 � C

Diffu

sivity

, D (m

2 /s)

A l , 3 0 0 �C

Fig. 3: Calculated diffusivities at 300, 400, and 660 °C (Tm ofAl) of 3d transition element solutes and of 4sp non-transitionmetal solutes in Al.

Lazarus-LeClaire model works well for electropositive impurities in noble metals, it is well known(see, e.g., [49,50]) that this model is unsatisfactory for Al, Mg, Pb, or the transition metals as solvents.As noted by Rummel [40], the row of the diffusing species - and hence atomic size - has minor influ-ence as compared to the valence effect. Therefore, elements of the same group obey similar diffusionkinetics in Al, as seen in the data for Ti, Zr, and Hf in Table 3.

4.2 Bottom line

Slow diffusion kinetics are an essential requirement for the retention of strength at elevated tempera-tures. Based on the preceding discussion, it is evident why the group 4 and 5 transition elements areattractive for potential alloying elements for use in high temperature Al alloys. These elements formtrialuminide compounds which are closely related to L12 and are also considerably slower diffusersthan Sc in Al. Group 4 → slower than Sc. Group 5 → much slower than Sc.

5 Castability

Up to this point the group 4 (Ti, Zr, Hf) and group 5(V, Nb, Ta) sound great.

1. They form Al3X trialuminide phases which are closely related to the desired cubic L12 struc-ture.

2. The are much slower diffusers than Sc.

The final requirement – and perhaps the one which complicates things the most – is that these alloysmust be castable.


Table 4: Invariant reactions involving in binary aluminum alloys alloyed with the transition elements which form Al3X

Group Reaction Type Reaction Temp. Liq. solubility Sol. solubility Partition coeff. (Al)-Al3X Refs(°C) (at.%) (at.%) equilibrium


Sc 3 Eutectic 660 0.28 0.23 0.82 Yes [33, 34]

Ti 4 Peritectic 665.4 0.079 0.79 10.0 Yes [51]

V 5 Peritectic 662.1 0.10 0.33 3.3 No [34]

Fe 8 Eutectic 655 0.91 0.026 0.03 Yes [34, 35, 39]

Co 9 Eutectic 657 0.46 < 0.009 < 0.02 No [34, 35, 39]

Ni 10 Eutectic 639.9 2.7 0.023 0.009 Yes [34, 35]


Y 3 Eutectic 639 2.47 < 0.05 < 0.02 Yes [34, 39]

Zr 4 Peritectic 660.8 0.033 0.083 2.52 Yes [36]

Nb 5 Peritectic 661.4 0.047 0.066 1.40 Yes [34]


La 3 Eutectic 640 ≈ 2.5 0.01 ≈ 0.004 No [34, 35]

Hf 4 Peritectic 662.2 0.078 0.186 2.38 Yes [34]

Ta 5 Peritectic 662 0.029 0.235 8.10 Yes [34]

Re a 7 Unknown ≈ 0.26 [34]

Ir 9 Eutectic 650 � 0.1 � 0.1 No [34]

a Because of the large difference in melting points (Al: 640.5 °C, Re: 3186 °C), it is difficult to bring alloys to equilibrium and consequentlythe Al-Re phase diagram is not well known.

5.1 Phase equilibria: generalizations

The transition elements, as well as those of the lanthanide (rare earth) and actinide series, form rathercomplex binary systems when alloyed with Al, in which one or more intermetallic phases occur.In these systems, a eutectic reaction generally occurs involving the liquid, the Al terminal solidsolution, and the Al-rich intermetallic phase. However, Al solid solution is formed via the peritecticreaction between the liquid and the Al-rich intermetallic phase with the group 4 (Ti, Zr, Hf), group5 (V, Nb, Ta) and group 6 (Cr, Mo, W) transition elements. The reactions are summarized in Table 4(for the transition elements) and Table 5 (for the actinides and lanthanides) for the systems we haveconsidered thus far.

5.2 Eutectics versus peritectics

There are other distinctions between eutectic and peritectic systems, summarized below:

1. The reaction itself. Both eutectic and peritectic reactions represent invariant points (three phases


Table 5: Invariant reactions in binary aluminum alloys alloyed with elements of the lanthanide (rare earth) or actinide series which formAl3X trialuminide compounds

Reaction Type Reaction Temp. Liq. solubility Sol. solubility Partition coeff. (Al)-Al3X Refs(°C) (at.%) (at.%) equilibrium

Lanthanides (rare earths)

Ce Eutectic 640 ≈ 4 ≈ 0.01 ≈ 0.0025 No [35]

Pr Eutectic 640 ≈ 5 0.009 ≈ 0.002 No [34, 35]

Nd Eutectic 640 ≈ 8 0.007 ≈ 0.0009 No [34, 35]

Pm a Eutectic ≈ 640 ≈ 8 ≈ 0.007 ≈ 0.0009 No [35]

Sm Eutectic 633 3.07 0.15 0.049 Yes [34, 35]

Gd Eutectic 650 ≈ 5 � 0.1 � 0.02 Yes [35]

Tb Eutectic 644 1.8b ≈ 0.005 ≈ 0.003 Yes [35]

Dy Eutectic 636 ≈ 2.5 < 0.1 < 0.04 Yes [35]

Ho Eutectic 650 1.8 < 0.1 < 0.06 Yes [35]

Er Eutectic 655 1 < 0.1 < 0.10 Yes [35]

Tm Eutectic 645 1.74 < 0.1 < 0.06 Yes [35]

Yb Eutectic 625 3.98 0.18 0.045 Yes [34]

Lu Eutectic ≈ 650 ≈ 2 < 0.1 < 0.05 Yes [35]

Actinides

Th Eutectic 630 2.9 0.8 0.28 No [35]

U Eutectic 646 1.7 0.007 0.004 No [34, 35]

Np c [35]

Pu Eutectic 650 1.6 0.0057 0.004 No [34, 35]

a Assumed similar behavior to Al-Nd and Al-Sm [35].b The eutectic has also been estimated to occur at ≈ 3.5 at.% Tb [35].c No phase diagram available, although presumed to be similar to Al-U [35].

in equilibrium). The eutectic reaction involves decomposition of a single phase liquid into twodifferent solid phases (L → α+β), while the peritectic reaction is the formation of a single solidphase by the reaction of a different solid phase with the liquid (L + β → α).

2. Solidification sequence. In a dilute eutectic system the first solid to form is the α solid solution,whereas for a peritectic alloy the first solid to form is the solute-rich β phase (Figure 4).

3. Reaction temperature. As indicated in Figure 4, the reaction temperature is less than the meltingpoint of pure solvent in eutectic alloys; for peritectic systems the opposite is true.

4. Liquid-solid partition coefficient. The equilibrium partition coefficient, k0, for solidification ofthe α solid solution is fundamentally different for eutectic and peritectic systems (Figure 5);k0 is less than unity in a eutectic system and is greater than unity in a peritectic system. Thisparameter dictates the solute distribution in cast alloys and therefore influences precipitationof dispersed phases during post-solidification heat treatment.


5.3 Bottom line

Reviewing the data in Tables 4 and 5, it is clear that, for the majority of the systems which formtrialuminides, the invariant reactions are eutectics. The exception to this rule, unfortunately, involvesthe group 4 (Ti, Zr, Hf) and group 5 (V, Nb, Ta) elements 3 – the very set of alloying additions which,up to this point, seemed to offer the most potential to compete with Al-Sc.

6 Summary and Impressions

While an initial glance at the periodic table (Figure 1) indicates numerous systems which form tria-luminides, few of these offer potential for developing castable precipitation-strengthened Al-basedalloys. Considering the structure of the trialuminides alone, only a handful of transition elements –Sc and the group 4 and 5 elements – seem feasible. Of these, Al-Sc is the only eutectic system.

6.1 Alternatives to Sc

The group 4 (Ti, Zr, Hf) and group 5 (V, Nb, Ta) transition elements offer distinct advantages onseveral counts:

1. They’re cheaper. Compared to Sc, everything is cheap.

2. They may not be L12, but they’re close. The group 4 and group 5 transition elements which isvery closely related to the desirable cubic L12 structure.

3. Much slower diffusers. Consider the data in Figure 3 and Table 3.

However, there are several factors which limit the potential of these systems

1. Castability. These are peritectics. This is not a complete show stopper, but it certainly compli-cates things. Peritectics imply the following:

• Necessary liquidus elevation• Tendency to lose solute in the liquid phase during solidification• Propensity for grain refinement, which ultimately limits the amount of solute which can

be used in the alloy

2. Solid solubility. This is interrelated with castability, since the liquid phase equilibria dictateshow much solute can be quenched into solid solution. Consider the phase diagrams in Figure6, which are for the group 3, 4, and 5 transition elements alloyed with Al. Using Al-Sc as abenchmark (Figure 6(a)), most of the other systems of interest exhibit substantially higher solidsolubilities (see Figures 6(d)–6(i)). Of the group 4 and 5 elements, Al-Zr (Figure 6(e)) and Al-Ta(Figure 6(i)) seem to offer the most potential.

3The group 6 (Cr, Mo, W) elements are also peritectics, but these do not form trialuminides.


6.2 The very unique Al-Sc system

It’s been said that Sc has the highest strengthening effect of any alloying element currently added toaluminum alloys (see, e.g. [57]). Perhaps now we know why. Consider the following:

1. Forms a stable cubic L12. Sc is the only transition element whose Al3X trialuminide phase is astable L12.

2. Al-Sc isn’t just a eutectic – it’s a great eutectic. Of all the eutectic systems in Tables 4 and 5,Al-Sc has the highest eutectic temperature as well as a partition coefficient, k0, very near unity.The eutectic temperature is important for a high-temperature alloy since, by definition, this islower than the melting point of Al. In an extreme case, this could limit the service temperatureof the alloy since its melting point is reduced. A partition coefficient near unity minimizessegregation, making the alloys especially amenable to casting.

3. It’s light. Probably not a huge deal considering the generally very dilute alloys we are inter-ested in, but of all the elements considered in Tables 1 and 2, Sc is the lightest (2.99 g · cm−3).The next lightest elements are Ti and Y, each about 4.5 g · cm−3. The other elements with stableL12 trialuminides – Er, Tm, Yb, Lu, U, and Np – are substantially heavier (9.07, 9.32, 6.97, 9.84,19.05, and 20.45 g · cm−3, respectively).

The two main drawbacks to Al-Sc are the following

1. Expensive. The exorbitant cost of Sc precludes extensive experimentation as well as widespreadapplications

2. Only moderately slow diffusion. Consider the data in Figure 3 and Table 3.


L + βL + α

L

α

T m o f p u r e s o l v e n t

C m a x C m a x

Temp

eratur

e

C o m p o s i t i o n

T 0

T e

α+ β

E u t e c t i c p o i n t

L + α

L + β

T p

T 0

C o m p o s i t i o n( a ) ( b )

α+ β

α

L

P e r i t e c t i c p o i n t

Fig. 4: Comparison between eutectic (a) and peritectic (b) reactions. While the form of the solvus curves is the same in (a) and (b), thetemperature of vanishing solubility, T0, is higher in the case of a peritectic system because of the elevated temperature of the reactionisotherm.

S o l i d u s

C S C L C L C S

S o l i d u s

Temp

eratur

e

C o m p o s i t i o n

L i q u i d u sT

k 0 < 1 k 0 > 1L i q u

i d u s

T

C o m p o s i t i o n( a ) ( b )

Fig. 5: Solidus-liquidus relationships for hypothetical dilute binary alloys. (a) k0 < 1 typical of a eutectic system (b) k0 > 1 typical of aperitectic system.


0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

0 . 2 8 a t . %

�

��

( A l ) + A l 3 S c

L + A l 3 S c6 6 0 °C

��

��

��

A t o m i c P e r c e n t S c

L

( A l ) 0 . 2 3 a t . %

(a) Al-Sc. Al3Sc: cP4, Pm3m, L12.Adapted from [33, 34].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

( A l ) + L

�

β��

( A l ) + α A l 3 Y

6 3 9 °C

��

��

��

A t o m i c P e r c e n t Y

L

( A l )0 . 0 5 a t . %

(b) Al-Y. Al3Y: hP8, P63/mmc,DO19. Adapted from [34].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

( A l ) + L

�

��

( A l ) + A l 1 1 L a 3

6 4 0 °C

��

��

��

A t o m i c P e r c e n t L a

L

( A l )0 . 0 1 a t . %

(c) Al-La. Al3La: hP8, P63/mmc,DO19. Adapted from [34].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

0 . 0 7 9 a t . %

�

��

( A l ) + A l 3 T i

L + A l 3 T i

6 6 5 . 4 °C

��

��

��

A t o m i c P e r c e n t T i

L

( A l )0 . 7 9 a t . %

(d) Al-Ti. Al3Ti: tI8, I4/mmm,D022. Adapted from [51].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

( A l )

0 . 0 3 3 a t . %

�

��

( A l ) + A l 3 Z r

L + A l 3 Z r

6 6 0 . 8 °C

��

��

��

A t o m i c P e r c e n t Z r

L

0 . 0 8 3 a t . %

(e) Al-Zr. Al3Zr: tI16, I4/mmm,D023. Adapted from [36].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

6 5 0 °C

( A l )

0 . 0 7 8 a t . %

�

β��

( A l ) + α A l 3 H f

L + β A l 3 H f

6 6 2 . 2 °C

��

��

��

A t o m i c P e r c e n t H f

L

0 . 1 8 6 a t . %( A l ) + β A l 3 H f

(f) Al-Hf. α Al3Hf : tI18,I4/mmm, D022. β Al3Hf :tI16, I4/mmm, D023. Adaptedfrom [34, 37].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

7 3 6 °C6 8 8 °C6 7 0 °C

L + A l 2 3 V 4

( A l )

0 . 1 0 a t . %

�

��

( A l ) + A l 2 1 V 2

L + A l 3 V

6 6 2 . 1 °C

��

��

��

A t o m i c P e r c e n t V

L

0 . 3 3 a t . % L + A l 4 5 V 7L + A l 2 1 V 2

(g) Al-V. Al3V: tI8, I4/mmm,D022. Adapted from [34].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

( A l )

0 . 0 4 7 a t . %

�

��

( A l ) + A l 3 N b

L + A l 3 N b

6 6 1 . 4 °C

��

��

��

A t o m i c P e r c e n t N b

L

0 . 0 6 6 a t . %

(h) Al-Nb. Al3Nb: tI8, I4/mmm,D022. Adapted from [29, 34].

0 0 . 1 0 . 2 0 . 3 0 . 4 0 . 5 0 . 6 0 . 7 0 . 8 0 . 9 13 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

( A l )

0 . 0 2 9 a t . %

�

� � ��

( A l ) + A l 3 T a

L + A l 3 T a

6 6 2 °C

��

��

��

A t o m i c P e r c e n t T a

L

0 . 2 3 5 a t . %

(i) Al-Ta. Al3Ta: tI8, I4/mmm,D022. Adapted from [34].

Fig. 6: Reported binary phase diagrams for dilute (<1 at.%) additions of the Group 3, 4, and 5 transition elements alloyed with Al.


References

[1] K. S. Kumar, “Ternary Intermetallics in Aluminum Refractory-Metal X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn),” Int Mater Rev,35(6), 293–327 (1990).

[2] C. Amador, et al., “Theoretical and Experimental-Study of Relaxations in Al3Ti and Al3Zr Ordered Phases,” Phys Rev Lett, 74(24),4955–4958 (1995).

[3] Y. Nakayama and H. Mabuchi, “Formation of Ternary L12 Compounds in Al3Ti-Base Alloys,” Intermetallics, 1(1), 41–48 (1993).

[4] M. Takeda, T. Kikuchi, and S. Makihara, “Stabilizing Effects of Third Elements on an L12-Al3Ti Compound,” J Mater Sci Lett, 18(8),631–634 (1999).

[5] H. Jones and W. M. Rainforth, “The Coarsening of Dispersed Al3Ti in Aluminum-Based Matrices,” Metall Mater Trans A, 34(2),419–421 (2003).

[6] W. E. Frazier, “PM Al Alloys: Hot Prospects for Aerospace Applications,” Adv Mater Process, 134(5), 42–46 (1988).

[7] J. F. Nie, S. Sridhara, and B. C. Muddle, “Characterization of Dispersed Intermetallic Phases in Rapidly Quenched Al-Ti-Ce Alloys,”Metall Trans A, 23(12), 3193–3205 (1992).

[8] J. F. Nie, A. Majumdar, and B. C. Muddle, “Development of High-Temperature Dispersion Strengthening in Rapidly QuenchedAl-Ti-X Alloys,” Mater Sci Eng A, 179, 619–624 (1994).

[9] W. E. Frazier and M. J. Koczak, “Coarsening Behavior of Rapidly Solidified Aluminum-Titanium Alloys,” in High Strength PowderMetallurgy Aluminum Alloys II, edited by G. J. Hildeman and M. J. Koczak (TMS-AIME, 1986), pp. 353–366.

[10] J. M. Wu, S. L. Zheng, and Z. Z. Li, “Thermal Stability and its Effects on the Mechanical Properties of Rapidly Solidified Al-Ti Alloys,”Mater Sci Eng A, 289(1-2), 246–254 (2000).

[11] Y. Wang, et al., “Microstructural Evolution and Microhardness of a Melt-Spun Al-5Ti-1B Alloy During Annealing,” Mater Sci Eng A,366(1), 17–24 (2004).

[12] R. D. Schelleng, “Mechanical Property Control of Mechanically Alloyed Aluminum,” JOM, 41(1), 32–35 (1989).

[13] G. X. Liang, Z. C. Li, and E. Wang, “Thermal Stability and Mechanical Properties of Mechanically Alloyed Al-10Ti Alloy,” J Mater Sci,31(4), 901–904 (1996).

[14] W. E. Frazier and M. J. Koczak, “Mechanical and Thermal-Stability of Powder-Metallurgy Aluminum Titanium-Alloys,” ScriptaMetall, 21(2), 129–134 (1987).

[15] G. S. Murty, M. J. Koczak, and W. E. Frazier, “High Temperature Deformation of Rapid Solidification Processed / MechanicallyAlloyed Al-Ti Alloys,” Scripta Metall, 21(2), 141–146 (1987).

[16] J. A. Hawk, K. R. Lawless, and H. G. F. Wilsdorf, “Tensile Strength of MA Aluminum Alloys with Titanium Additions,” Scripta Metall,23(1), 119–124 (1989).

[17] P. K. Mirchandani, R. C. Benn, and K. A. Heck, “Structure-Property Relationships in Elevated Temperature High Stiffness Mechan-ically Alloyed Al-Ti Based Alloys,” in Lightweight Alloys for Aerospace Applications, edited by E. W. Lee, E. H. Chia, and N. J. Kim(TMS-AIME, 1989), pp. 33–58.

[18] P. W. Sonawane, et al., “Mechanically Alloyed Al-Ti Alloy for Superplastic Forming Characteristics,” Mater Sci Forum, 88-90, 647–654(1992).

[19] K. M. Lee and I. H. Moon, “High Temperature Deformation Behavior of Mechanically Alloyed Al-Ti Alloys,” Scripta Metall Mater,26(1), 123–126 (1992).

[20] T. D. Bayha, R. J. Kilmer, and F. E. Wawner, “The Fracture Characteristics of Al-9Ti/SiCp Metal Matrix Composites,” Metall Trans A,23(6), 1653–1662 (1992).

[21] N. Q. Wu, et al., “Microstructure and Sliding Wear Behavior of PM Alloy Al-10Ti after Thermal Exposure,” Wear, 203, 155–161 (1997).

[22] S. H. Wang and P. W. Kao, “The Strengthening Effect of Al3Ti in High Temperature Deformation of Al-Al3Ti Composites,” Acta Mater,46(8), 2675–2682 (1998).


[23] I. C. Barlow, H. Jones, and W. M. Rainforth, “Evolution of Microstructure and Hardening, and the Role of Al3Ti Coarsening, DuringExtended Thermal Treatment in Mechanically Alloyed Al-Ti-O Based Materials,” Acta Mater, 49(7), 1209–1224 (2001).

[24] F. A. Crossley and L. F. Mondolfo, “Mechanism of Grain Refinement in Aluminum,” Transactions of the AIME, 191, 1143–1148 (1951).

[25] J. Marcantonio and L. F. Mondolfo, “The Nucleation of Aluminium by Several Intermetallic Compounds,” J Inst Metals, 98, 23–27(1970).

[26] J. Marcantonio and L. J. Mondolfo, “Grain Refinement in Aluminum Alloyed with Titanium and Boron,” Metall Trans, 2(2), 465–471(1971).

[27] T. Ohashi and R. Ichikawa, “Grain Refinement in Aluminum-Zirconium and Aluminum-Titanium Alloys by Metastable Phases,” ZMetallkd, 64(7), 517–521 (1973).

[28] D. H. St. John and L. M. Hogan, “Al3Ti and Grain Refinement of Aluminum,” Journal of the Australasian Institute of Metals, 22(3-4),160–166 (1977).

[29] T. W. Clyne and M. H. Robert, “Stability of Intermetallic Aluminides in Liquid Aluminium and Implications for Grain Refinement,”Met Technol, 7(MAY), 177–185 (1980).

[30] D. G. McCartney, “Grain Refining of Aluminum and Its Alloys Using Inoculants,” Int Mater Rev, 34(5), 247–260 (1989).

[31] K. T. Kashyap and T. Chandrashekar, “Effects and Mechanisms of Grain Refinement in Aluminium Alloys,” B Mater Sci, 24(4), 345–353 (2001).

[32] B. S. Murty, S. A. Kori, and M. Chakraborty, “Grain Refinement of Aluminium and Its Alloys by Heterogeneous Nucleation andAlloying,” Int Mater Rev, 47(1), 3–29 (2002).

[33] J. L. Murray, “The Al-Sc (Aluminum-Scandium) System,” J Phase Equil, 19(4), 380–384 (1998).

[34] H. Okamoto, Phase Diagrams of Dilute Binary Alloys (ASM International, 2002).

[35] Binary Alloy Phase Diagrams, volume 1 (ASM International), 2nd edition.

[36] J. L. Murray, A. Peruzzi, and J. P. Abriata, “The Al-Zr (Aluminum-Zirconium) System,” J Phase Equil, 13(3), 277–291 (1992).

[37] J. L. Murray, A. J. McAlister, and D. J. Kahan, “The Al-Hf (Aluminum-Hafnium) System,” J Phase Equil, 19(4), 376–379 (1998).

[38] K. S. Kumar, “Advanced Intermetallics,” in Physical Metallurgy and Processing of Intermetallic Compounds, edited by N. S. Stoloff andV. K. Sikka (Chapman and Hall, 1996), pp. 392–440.

[39] J. E. Hatch, Aluminum: Properties and Physical Metallurgy (American Society for Metals, Metals Park, OH, 1984).

[40] G. Rummel, et al., “Diffusion of Implanted 3d-Transition Elements in Aluminum. Part I: Temperature Dependence,” Z Metallkd, 86(2),122–130 (1995).

[41] G. Rummel, et al., “Diffusion of Implanted 3d-Transition Elements in Aluminum. Part II: Pressure Dependence,” Z Metallkd, 86(2),131–140 (1995).

[42] S. I. Fujikawa, “Solid State Diffusion in Light Metals,” J Japan Inst Light Metals, 46(4), 202–215 (1996).

[43] J. Grammatikakis, K. Eftaxias, and V. Hadjicontis, “Interconnection of the Diffusion-Coefficients of Various Elements in Aluminum,”J Phys Chem Solids, 49(10), 1275–1277 (1988).

[44] O. Madelung, editor, Landolt-Bornstein: Numerical Data and Functional Relationships in Science and Technology, volume 26 of Diffusion inSolid Metals and Alloys (Springer-Verlag, Berlin, 1990).

[45] S. I. Fujikawa, “Impurity Diffusion of Scandium in Aluminum,” Def Diff Forum, 143-147, 115–120 (1997).

[46] D. Bergner and N. van Chi, “Untersuchungen zur Diffusion von 3d-Metallen in Al,” Wiss Z P’adagogischen Hochschule, 15(3) (1977).

[47] D. Lazarus, “Effect of Screening on Solute Diffusion in Metals,” Phys Rev, 93(5), 973–976 (1954).

[48] A. D. LeClaire, “On the Theory of Impurity Diffusion in Metals,” Phil Mag, 7, 141–167 (1962).


[49] P. Shewmon, Diffusion in Solids (The Minerals, Metals, and Materials Society, Warrendale, PA, 1989), second edition.

[50] J. Philibert, Atomic Movements - Diffusion and Mass Transport in Solids, Monographies de physique (Les Editions de Physique, Les Ulis,France, 1991).

[51] J. L. Murray, “Personal communication regarding the equilibrium Al-Ti phase diagram.” (2005).

[52] W. J. Boettinger, “Structure of Directionally Solidified Two-Phase Sn-Cd Peritectic Alloys,” Metall Trans, 5(9), 2023–2031 (1974).

[53] J. Cisse, H. W. Kerr, and G. F. Bolling, “Nucleation and Solidification of Al-Ti Alloys,” Metall Trans, 5(3), 633–641 (1974).

[54] H. W. Kerr, J. Cisse, and G. F. Bolling, “Equilibrium and Nonequilibrium Peritectic Transformations,” Acta Metall, 22(6), 677–686(1974).

[55] I. Maxwell and A. Hellawell, “Constitution and Solidification of Peritectic Alloys in System Al-Ti,” Acta Metall, 23(8), 895–899 (1975).

[56] D. H. St. John and L. M. Hogan, “Thermal Analysis of Peritectic Alloys,” J Mater Sci, 17(8), 2413–2418 (1982).

[57] C. B. Fuller, D. N. Seidman, and D. C. Dunand, “Creep Properties of Coarse-Grained Al(Sc) Alloys at 300 °C,” Scripta Mater, 40(6),691–696 (1999).

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