Development of a Nanoscale Precipitation-Strengthened Creep ...

NORTHWESTERN UNIVERSITY

Development of a Nanoscale Precipitation-StrengthenedCreep-Resistant Aluminum Alloy Containing Trialuminide Precipitates

A DISSERTATION

SUBMITTED TO THE GRADUATE SCHOOL

IN PARTIAL FULFILLMENT OF THE REQUIREMENTS

for the degree

DOCTOR OF PHILOSOPHY

Field of Materials Science and Engineering

By

Keith Edward Knipling

EVANSTON, IL

December 2006

© Copyright by Keith Edward Knipling 2006

All Rights Reserved

i

ABSTRACT

Development of a Nanoscale Precipitation-Strengthened Creep-Resistant Aluminum

Alloy Containing Trialuminide Precipitates

Keith Edward Knipling

This research is toward developing a castable and heat-treatable precipitation-strengthened aluminum

alloy exhibiting coarsening- and creep resistance at temperatures exceeding 400°C. Criteria for selecting

alloying elements capable of producing such an alloy are established. Those systems forming Al3M tria-

luminide compounds with a cubic L12 crystal structure are favored, and based on a review of the existing

literature, these are assessed in terms of solid-solubility and diffusivity in α-Al (satisfying the need for

slow coarsening kinetics), and castability (which is discussed based on the binary phase diagrams). The

first Group 3 element, Sc, and the second Group 4 element, Zr, are shown to be most promising.

These expectations are confirmed by an initial study on the Al-Ti system, which demonstrates that

conventionally-solidified alloys are not capable of precipitation strengthening. The Al-Zr system, by con-

trast, exhibits precipitation of nanometer-scale Al3Zr (L12) producing pronounced precipitation hard-

ening when aged at 375, 400, or 425°C. The Al3Zr precipitates are coarsening resistant and have the

metastable L12 structure up to 500°C, a result of very sluggish diffusion of Zr in α-Al. Ternary additions

of Ti are also investigated, forming Al3(Zr1−xTix) (L12) precipitates with a reduced lattice parameter mis-

match with α-Al, potentially improving the coarsening resistance.

The composition of Al3(Zr1−xTix) precipitates formed at 375 or 425°C are measured directly using 3-D

atom-probe tomography. At these temperatures, the Zr:Ti atomic ratio in the precipitates is about 10 and

5, respectively, indicating that most of the available Ti fails to partition to the Al3(Zr1−xTix) phase. This

is consistent with prior studies on Al-Sc alloys, where the slower-diffusing ternary solute species make

up a small fraction of the Al3Sc-based precipitates. Despite the confirmed presence of Ti, Al3(Zr1−xTix)

precipitates exhibit no improvement in terms of coarsening resistance compared to binary Al3Zr.

Mechanical properties of the Al-Zr and Al-Zr-Ti alloys are investigated utilizing Vickers microhardness

and creep. The alloys deformed by creep at 300–400°C exhibit a dislocation climb-controlled threshold

stress, ca. 6–12 MPa. The binary Al-Zr and ternary Al-Zr-Ti alloys behave similarly under ambient- and

high temperature loading, consistent with the similar microstructures of the two alloys.

iii

ACKNOWLEDGMENTS

This research would not have been possible without the guidance of my thesis advisors, Profs.

D. C. Dunand, D. N. Seidman, and M. E. Fine. It was an honor to be your student. I thank

also Prof. M. Asta (University of California, Davis) and Dr. J. L. Murray (Alcoa) for numerous

stimulating conversations.

I am most appreciative of my colleagues in both the Dunand and Seidman research groups

who were there “in the trenches” with me. Foremost, I acknowledge Dr. D. Isheim for his in-

satiable enthusiasm for metallurgy and always making himself available for interesting discus-

sions. I especially thank R. A. Karnesky and M. van Dalen, my colleagues working on aluminum,

for putting up with me for four years. Finally, I thank my parents for their unending love and

support throughout my (lengthy) education.

This research was supported by the United States Department of Energy, Basic Sciences Di-

vision, under contract DE-FG02-02ER45997. Financial support was also provided by a Walter P.

Murphy Fellowship at Northwestern University.

v

CONTENTS

1 Criteria for Developing High-Temperature Aluminum Alloys 1

1.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

1.1.1 General strengthening requirements . . . . . . . . . . . . . . . . . . . . . . . 2

1.1.2 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

1.2 Selection Criteria . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

1.2.1 Capability of forming strengthening intermetallic phases . . . . . . . . . . . 7

1.2.2 Low solid-solubility in α-Al . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

1.2.3 Small diffusivity in α-Al . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

1.2.4 Ability to be conventionally cast . . . . . . . . . . . . . . . . . . . . . . . . . 23

1.3 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 29

1.3.1 Alloying additions with transition elements . . . . . . . . . . . . . . . . . . . 29

1.3.2 Alloying additions with lanthanide elements . . . . . . . . . . . . . . . . . . 33

1.3.3 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35

1.4 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35

2 Peritectic Solidification 37

2.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37

2.2 Nomenclature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38

2.3 Mechanisms of α Formation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

2.3.1 Peritectic reaction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40

2.3.2 Peritectic transformation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40

2.3.3 Direct solidification of supersaturated α . . . . . . . . . . . . . . . . . . . . . 41

2.4 Metastable Properitectic Phases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42

2.4.1 Compositionally metastable properitectic phases . . . . . . . . . . . . . . . 42

2.4.2 Structurally metastable properitectic phases . . . . . . . . . . . . . . . . . . 43

2.4.3 Grain refinement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44

vii

viii CONTENTS

2.5 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47

2.5.1 Mechanisms of α formation . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47

2.5.2 Possible solidification sequences as a function of cooling rate . . . . . . . . 48

2.5.3 Solute redistribution during α-Al nucleation . . . . . . . . . . . . . . . . . . 49

3 Nucleation and Precipitation Strengthening in Dilute Al-Ti and Al-Zr Alloys 53

3.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53

3.2 Experimental Procedures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55

3.2.1 Alloy compositions and preparation . . . . . . . . . . . . . . . . . . . . . . . 55

3.2.2 Aging treatments and analytical techniques . . . . . . . . . . . . . . . . . . . 56

3.3 Experimental Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57

3.3.1 Optical microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57

3.3.2 Age hardening . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59

3.3.3 Electron microscopies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61

3.3.4 Electrical conductivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

3.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

3.4.1 Diffusion kinetics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

3.4.2 Classical nucleation theory . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

3.4.3 Limitations on increasing the solute concentration . . . . . . . . . . . . . . . 69

3.4.4 Comparison with previous studies . . . . . . . . . . . . . . . . . . . . . . . . 70

3.4.5 Common industrial uses of Ti and Zr additions to Al . . . . . . . . . . . . . . 72

3.5 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

4 Atom-Probe Tomographic Studies of Precipitation in Al-0.1Zr-0.1Ti (at.%) Alloys 75

4.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75



4.3.1 Solute segregation and the precipitated microstructure . . . . . . . . . . . . 78

4.3.2 Three-dimensional atom-probe tomographic reconstructions . . . . . . . . 80

4.3.3 Quantitative chemical composition of the Al3(Zr1−xTix) precipitates . . . . 80

4.3.4 Time-of-flight mass spectra . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83

4.3.5 Correlation between measured Ti concentration proximity to precipitates . 84

CONTENTS ix

4.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 87

4.4.1 Local magnification of the Al3(Zr1−xTix) precipitates . . . . . . . . . . . . . 89

4.4.2 Compositions of the Al3(Zr1−xTix) precipitates . . . . . . . . . . . . . . . . . 92

4.4.3 Residual Ti in the α-Al solid-solution . . . . . . . . . . . . . . . . . . . . . . . 93

4.4.4 Comparison to 3-D APT studies on Al-Sc-Ti/Zr alloys . . . . . . . . . . . . . 95

4.5 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 96

5 Microstructure of Al-Zr and Al-Zr-Ti Alloys during Aging at 375, 400, or 425°C 99

5.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99



5.3.1 As-cast microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101

5.3.2 Solute segregation and the precipitated microstructure . . . . . . . . . . . . 104

5.3.3 Age hardening . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 107

5.3.4 Precipitate morphologies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 109

5.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 123

5.4.1 Comparison to Previous Studies on Al-Zr-V/Ti Alloys . . . . . . . . . . . . . 123

5.5 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 129

6 Microstructural Coarsening in Al-Zr and Al-Zr-Ti Alloys, T ≥ 450°C 131

6.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 131


6.3 Experimental Results: Isothermal Aging at 500°C . . . . . . . . . . . . . . . . . . . . 133

6.3.1 Vickers microhardness . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133

6.3.2 Electron microscopies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133

6.4 Experimental Results: Isochronal Aging, 300–600°C . . . . . . . . . . . . . . . . . . . 144

6.4.1 Vickers microhardness and electrical conductivity . . . . . . . . . . . . . . . 144

6.4.2 Transmission electron microscopy . . . . . . . . . . . . . . . . . . . . . . . . 146

6.5 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 150

6.5.1 Transformation to the equilibrium D023 structure . . . . . . . . . . . . . . . 150

6.5.2 Mechanism of microstructural coarsening . . . . . . . . . . . . . . . . . . . . 153

6.5.3 Mechanism of strengthening . . . . . . . . . . . . . . . . . . . . . . . . . . . 155

x CONTENTS

6.5.4 Effect of aging temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . 160

6.5.5 Effect of Ti additions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 161

6.6 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 163

7 Creep of Al-Zr and Al-Zr-Ti Alloys 167

7.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 167



7.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 172

7.4.1 Transformation to the equilibrium D023 structure . . . . . . . . . . . . . . . 172

7.4.2 Differences in creep behavior between Al-Zr and Al-Zr-Ti alloys . . . . . . . 173

7.4.3 Comparison to Al-Sc alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . 177

7.5 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 178

8 Final Remarks and Suggestions for Future Work 179

8.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 179

8.2 Improvements with Judicious Aging . . . . . . . . . . . . . . . . . . . . . . . . . . . 179

8.3 Alloys Containing Ta . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 180

8.4 Alloying Additions to Al-Zr Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 181

8.4.1 Al-Zr-Mg Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 182

8.4.2 Al-Zr-Sc Alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 182

References 184

A Alloy Preparation and Castability 213

A.1 Advantages of Arc-Melting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 213

A.1.1 High melt temperatures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 214

A.1.2 Inert atmosphere . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 215

A.1.3 Accelerated solidification rates . . . . . . . . . . . . . . . . . . . . . . . . . . 215

B Homogenization and Other Prior Annealing Treatments 217

B.1 Homogenization (or Solution) Anneals . . . . . . . . . . . . . . . . . . . . . . . . . . 217

B.2 Pre-Aging at 500°C . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 219

CONTENTS xi

C Length Scales of Analytical Techniques and Potential Biases 221

C.1 Length Scales of Analytical Techniques . . . . . . . . . . . . . . . . . . . . . . . . . . 221

C.2 Biases Associated with Sample Preparation . . . . . . . . . . . . . . . . . . . . . . . 222

D Study on an Al-Zr-Ti Alloy Prepared by Parameswaran et al. 225

E Potential Contamination of Fe and/or Si, and its Effects on Precipitation 229

LIST OF FIGURES

1.1 The Orowan stress, Eq.(1.1), as a function of mean precipitate radius, 〈R〉, for various vol-

ume fractions, f , of dispersed phase. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

1.2 Alloying additions to Al that form thermodynamically stable trialuminide (Al3M) inter-

metallic compounds, with the equilibrium structure indicated. . . . . . . . . . . . . . . . . 8

1.3 The (a) L12, (b) D022, and (c) D023 structures. . . . . . . . . . . . . . . . . . . . . . . . . . . . 10

1.4 Reported binary phase diagrams for dilute (<1 at.%) additions of the Group 3, 4, and 5

transition elements alloyed with Al. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16

1.5 Measured activation enthalpies (Q) of 3d transition element solutes and of 4sp non-transition

metal solutes in α-Al. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21

1.6 Calculated diffusivities at 300, 400, and 660°C (Tm of Al) of 3d transition element solutes

and of 4sp non-transition metal solutes in α-Al. . . . . . . . . . . . . . . . . . . . . . . . . . 21

1.7 Semi logarithmic plot of diffusivity in α-Al versus reciprocal temperature for the elements

which form L12 trialuminide phases with Al. . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

1.8 Comparison between eutectic and peritectic reactions. . . . . . . . . . . . . . . . . . . . . . 24

1.9 Undesirable (a) and desirable (b) characteristics of peritectic systems, as they relate to the

potential for development of castable precipitation-strengthened alloys. . . . . . . . . . . . 28

2.1 Schematic peritectic phase diagram indicating symbols used for phases, compositions, and

temperatures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

2.2 Free energy curves for the equilibrium phases at the Al-rich end of the Al-Ti system above

the peritectic temperature Tp. From Kerr et al. (1974). . . . . . . . . . . . . . . . . . . . . . . 42

2.3 Equilibrium Al-Ti phase diagram with metastable extrapolations constructed from micro-

probe analyses. From Kerr et al. (1974). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42

2.4 Effect of solidification rate and temperature of the melt on the morphology of primary Al3Zr

precipitates. As solidification rate increases, the primary phase becomes more equiaxed.

From Brodova et al. (2001). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44

2.5 Transmission electron micrographs showing the 3-D morphology of metastable cubic L12

Al3Ti primary (properitectic) precipitates. From Majumdar et al. (1993). . . . . . . . . . . . 45

xiii

xiv LIST OF FIGURES

2.6 Relation between the grain sizes and Ti concentration in Al-Ti alloys solidified at various

cooling rates. From Ohashi and Ichikawa (1973). . . . . . . . . . . . . . . . . . . . . . . . . . 46

2.7 Relation between the grain sizes and Zr concentration in Al-Zr alloys solidified at various

cooling rates. From Ohashi and Ichikawa (1973). . . . . . . . . . . . . . . . . . . . . . . . . . 46

2.8 Schematic equilibrium (full lines) and non-equilibrium (broken and dotted lines) peritectic

phase diagram. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47

2.9 Effects of cooling rate and Zr concentration on the solidification microstructures of the Al-

Zr alloys investigated by Hori et al.(1981). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48

2.10 Solidus-liquidus relationships for hypothetical dilute binary alloys. (a) k0 < 1 (b) k0 > 1. . . 49

2.11 Consequences of eutectic (k0 < 1) and peritectic (k0 > 1) solidification on the resulting

solute distribution in cast Al alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51

3.1 Macrostructure of as-cast alloys Al-0.18Ti, Al-0.22Ti, and Al-0.19Zr. . . . . . . . . . . . . . . 58

3.2 Vickers microhardness vs. aging time at 375°C for the Al-Ti and Al-Zr alloys investigated. . . 59

3.3 Vickers microhardness vs. aging time at 425°C for the Al-Ti and Al-Zr alloys investigated. . . 59

3.4 SEM micrographs of electropolished TEM foils, showing no evidence of precipitation of

Al3Ti and copious precipitation of Al3Zr (L12) after aging Al-0.22Ti and Al-0.19Zr (at.%)

alloys at 425°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 60

3.5 Centered superlattice dark-field TEM micrographs of Al-0.19Zr aged at 425°C for 400 h

displaying small, coherent, homogeneously-distributed Al3Zr (L12) precipitates within the

highly-supersaturated dendrites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61

3.6 Calculated root-mean-square diffusion distances,√

6Dt, for Ti and Zr in α-Al at 375 and

425°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 64

3.7 Equilibrium Al-rich Al-Ti binary phase diagram. . . . . . . . . . . . . . . . . . . . . . . . . . 65

3.8 Equilibrium Al-rich Al-Zr binary phase diagram. . . . . . . . . . . . . . . . . . . . . . . . . . 65

3.9 Critical cooling rate for suppressing primary Al3Ti or Al3Zr and achieving supersaturated

α-Al solid-solution. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

4.1 Centered superlattice dark-field TEM micrographs of alloy Al-0.1Zr-0.1Ti(a) aged at 375°C

for 1,600 h. A representative 50×50×200 nm3 3-D APT analysis volume is indicated for

comparative purposes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

4.2 EDS concentration profiles of Zr and Ti across three solute-rich dendrite arms in as-cast

Al-0.1Zr-0.1Ti(b). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 79

4.3 LEAP reconstruction of an Al3(Zr1−xTix) precipitate in alloy Al-0.1Zr-0.1Ti(b) aged for 400

h at 425°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81

LIST OF FIGURES xv

4.4 Proxigram displaying the distribution of Zr and Ti atoms for the precipitate reconstructed

in Figure 4.3 (Al-0.1Zr-0.1Ti(b) aged at 425°C for 400 h). The precipitate is Zr-rich, with

a small fraction of the Zr lattice sites replaced by Ti atoms. The ratio of Zr:Ti is about 5,

corresponding to Al3(Zr0.83Ti0.17). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 82

4.5 Proxigrams displaying the distribution of Zr and Ti atoms in Al3(Zr1−xTix) precipitates

formed in alloy Al-0.1Zr-0.1Ti(a) aged at 375°C for 20, 400, and 1,600 h. The ratio of Zr:Ti is

about 10, independent of aging time, corresponding to Al3(Zr0.91Ti0.09). . . . . . . . . . . . 82

4.6 Atom-probe time-of-flight mass spectrum of Al-0.1Zr-0.1Ti(b) aged for 400 h at 425°C (re-

construction shown in Figure 4.3). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 84

4.7 LEAP reconstruction of Al-0.1Zr-0.1Ti(a) aged for 1,600 h at 375°C. One large (>25 nm ra-

dius) Al3(Zr1−xTix) precipitate is detected. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 86

4.8 Proxigram showing the distribution of Zr and Ti atoms in the Al3(Zr1−xTix) precipitate dis-

played in Figure 4.7, Al-0.1Zr-0.1Ti(a) aged at 375°C for 1,600 h. The precipitate is slightly

super-stoichiometric (25–30 at.% total solute) with very little Ti (ca. 1 at.%). . . . . . . . . . 87

4.9 Isothermal sections at 375 or 425°C of a calculated ternary phase diagram for the Al-Zr-

Ti system in the Al-rich corner, showing the α-Al solid-solution in equilibrium with the

metastable Al3(Zr1−xTix) (L12) and Al3Ti (D022). . . . . . . . . . . . . . . . . . . . . . . . . . 88

4.10 A 3 nm-thick slice, cut parallel to the analysis direction, through the reconstruction in Fig-

ure 4.7. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89

5.1 Macrostructure of as-cast alloys Al-0.1Zr(b), Al-0.1Zr-0.1Ti(b), Al-0.2Zr, and Al-0.2Zr-0.2Ti. 102

5.2 SEM secondary electron images of as-solidified Al-0.2Zr. Several petal-like primary Al3Zr

precipitates (light contrast) are observed. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 103

5.3 SEM secondary electron images of as-solidified Al-0.2Zr-0.2Ti. . . . . . . . . . . . . . . . . . 103

5.4 SEM micrograph of a primary Al3(Zr1−xTix) precipitate, showing also positions of quanti-

tative EDS measurements. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104

5.5 EDS concentration profiles of Zr and Ti across three solute-rich dendrite arms in as-cast

Al-0.1Zr-0.1Ti(b) and Al-0.2Zr-0.2Ti. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 105

5.6 TEM and SEM micrographs of Al-0.1Zr-0.1Ti(a) aged at 375°C for 1,600 h. . . . . . . . . . . 106

5.7 Vickers microhardness vs. aging time for Al-0.1Zr(a) and Al-0.1Zr(b) at 375, 400, or 425°C. . 108

5.8 Vickers microhardness vs. aging time for Al-0.2Zr at 375, 400, or 425°C. . . . . . . . . . . . . 108

5.9 Vickers microhardness vs. aging time for Al-0.1Zr-0.1Ti(a) and Al-0.1Zr-0.1Ti(b) at 375, 400,

or 425°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 108

5.10 Vickers microhardness vs. aging time for Al-0.2Zr-0.2Ti at 375, 400, or 425°C. . . . . . . . . . 108

5.11 Centered superlattice dark-field TEM micrograph of an interdendritic channel between

two dendrite arms in alloy Al-0.2Zr-0.2Ti aged at 425°C for 100 h. . . . . . . . . . . . . . . . 110

xvi LIST OF FIGURES

5.12 Centered superlattice dark-field TEM micrographs of heterogeneously-nucleated spheroidal

L12 interdendritic precipitates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 110

5.13 Development of growth instabilities with increasing precipitate radius for interdendritic

Al3(Zr1−xTix) precipitates in Al-0.1Zr-0.1Ti(b) aged at 425°C for 400 h. . . . . . . . . . . . . . 111

5.14 SEM micrographs of Al-0.2Zr and TEM micrographs of Al-0.2Zr-0.2Ti, aged at 425°C for

400 h. Gradients in precipitate radius vary considerably throughout the microstructure,

even across the same interdendritic channel. . . . . . . . . . . . . . . . . . . . . . . . . . . . 113

5.15 Complementary dark-field and bright-field TEM micrographs showing loss of coherency in

the interdendritic precipitate size gradients. . . . . . . . . . . . . . . . . . . . . . . . . . . . 114

5.16 TEM micrographs of rod-shaped interdendritic Al3Zr or Al3(Zr1−xTix) precipitates, ori-

ented in 〈100〉matrix directions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 116

5.17 SEM secondary electron images of primary Al3(Zr1−xTix) precipitates in alloy Al-0.2Zr-

0.2Ti aged at 425°C for 400 h, illustrating solid-state precipitation of Al3(Zr1−xTix) (L12)

in supersaturated α-Al surrounding the primary precipitates. . . . . . . . . . . . . . . . . . 118

5.18 Centered superlattice dark-field TEM micrographs demonstrating solid-state precipitation

of spheroidal Al3(Zr1−xTix) (L12) precipitates in supersaturated α-Al solid-solution sur-

rounding a petal-like Al3(Zr1−xTix) (L12) primary precipitate in Al-0.2Zr aged at 425°C for

400 h. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 119

5.19 Centered superlattice dark-field TEM micrographs of Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b) aged

at 425°C for 1600 h. Within the dendrites, Al3(Zr1−xTix) (L12) precipitates exhibit no im-

provement in coarsening resistance as compared to Al3Zr(L12). . . . . . . . . . . . . . . . . 121

5.20 Comparison of L12-structured Al3(Zr1−xMx) (M = V and/or Ti) average precipitate radius

(〈R〉, in nm) with aging time at 425°C (M = V and/or Ti). . . . . . . . . . . . . . . . . . . . . . 125

5.21 Montage of centered superlattice dark-field TEM micrographs of Al-0.2Zr aged at 425°C for

400 h. The apparent mean radius, 〈R〉, of spheroidal Al3Zr (L12) precipitates is strongly

dependent on location in the specimen. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 127

6.1 Vickers microhardness vs. exposure at 500°C for peak-aged Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b)

alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 133

6.2 Montage of bright-field TEM micrographs of Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after

aging at 375°C for 100 h). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 134

6.3 Montage of centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(b) aged at

500°C for 100 h (after aging at 375°C for 100 h). . . . . . . . . . . . . . . . . . . . . . . . . . . 135

6.4 SEM micrograph of heterogeneously-nucleated Al3Zr (D023) precipitates on dislocations or

small-angle grain boundaries in Al-0.1Zr(b) aged at 500°C for 100 h (after aging at 375°C for

100 h). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 136

LIST OF FIGURES xvii

6.5 Schematic of the three possible orientation relationships with α-Al for the disc-like equi-

librium D023 precipitates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 137

6.6 Complementary bright-field and dark-field TEM micrographs of heterogeneously-nucleated

L12- and D023-structured Al3(Zr1−xTix) precipitates on dislocations in Al-0.1Zr-0.1Ti(b) aged

at 500°C for 100 h (after aging at 375°C for 100 h). Many of the L12-structured precipitates

exhibit attributed to antiphase boundaries (APBs) along 〈100〉 directions, indicating trans-

formation to the equilibrium D023 structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . 139

6.7 Bright-field TEM micrograph of disc-like D023 precipitates in three possible orientations

and corresponding selected area diffraction patterns (SADPs) of the equilibrium D023 phase

in Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h).The patterns are

indexed in Figure 6.9. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 140

6.8 Complementary bright-field and dark-field TEM micrographs of heterogeneously-nucleated

D023 precipitates on dislocations in Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at

375°C for 100 h). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 141

6.9 Calculated electron diffraction patterns from the equilibrium tetragonal D023 (closed cir-

cles, T subscript) and metastable cubic L12 (crosses, C subscript) Al3Zr phases in three

possible orientation relationships with α-Al (open circles). The spots not indexed (e.g. 002T,

006T) are kinematically-forbidden. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 142

6.10 Complementary dark-field and bright-field TEM micrographs of dendritic precipitates in

Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h). . . 143

6.11 Centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(b) aged at 500°C for

100 h (after aging at 375°C for 100 h). Many of the spheroidal L12-structured precipitates

exhibit APBs along 〈100〉 directions, indicating a partial transformation to the equilibrium

D023 structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 144

6.12 Vickers microhardness and electrical conductivity during isochronal aging of Al-0.1Zr(c)

and Al-0.1Zr-0.1Ti(c). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 145

6.13 Bright-field and centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(c) isochronally

aged to 450°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 146

6.14 Centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(c) isochronally aged to

525°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 147

6.15 Complementary bright-field and dark-field TEM micrographs of Al-0.1Zr-0.1Ti(c) isochronally

aged to 575°C. Most of the metastable L12 precipitates are transformed to the disc-like D023

structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 148

6.16 Complementary dark-field and bright-field TEM micrographs of dendritic precipitates in

Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c) isochronally aged to 575°C. . . . . . . . . . . . . . . . . . . 149

xviii LIST OF FIGURES

6.17 Schematic of the L12 to D023 transformation in Al3Zr precipitates. . . . . . . . . . . . . . . . 151

6.18 Schematic of four stages of microstructural coarsening occurring on the nanometer-scale

of the precipitates and the micrometer-scale of the dendrites. . . . . . . . . . . . . . . . . . 153

6.19 Centered superlattice dark-field TEM micrographs displaying the four stages of microstruc-

tural coarsening shown schematically in Figure 6.18. . . . . . . . . . . . . . . . . . . . . . . . 155

6.20 Vickers microhardness yield stress increment vs. mean precipitate radius 〈R〉 for Al-0.1Zr(c)

and Al-0.1Zr-0.1Ti(c) isochronally aged at 450, 525, and 575°C. . . . . . . . . . . . . . . . . . 159

6.21 Schematic as-solidified structure and the expected precipitation behavior for a dendritically-

solidified α solid-solution in which k0 > 1 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 161

6.22 Al-rich equilibrium binary Al-Zr phase diagram and calculated metastable L12 Al3Zr and

Al3(Zr1−xTix) solvus by Murray. Additions of Ti reduce the solid-solubility of Zr. . . . . . . . 162

7.1 Variation of strain, ε, with time for alloy Al-0.1Zr(d) crept at 300°C for four levels of stress. . 170

7.2 Double logarithmic plot of minimum creep rate at 300, 350, or 400°C vs. applied stress, for

Al-0.1Zr(d) and Al-0.1Zr-0.1Ti(d). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171

7.3 ln(ε) vs. ln(σ − σth) for the data in Figure 7.2. . . . . . . . . . . . . . . . . . . . . . . . . . . . 172

7.4 Series of optical micrographs illustrating the macrostructure of alloy Al-0.1Zr(b). The ingot

cross-section is composed entirely of columnar grains (etched using Poultan’s reagent). . . 175

7.5 Series of optical micrographs illustrating the macrostructure of alloy Al-0.1Zr-0.1Ti(b). The

grain size is more equiaxed and refined than that of Al-0.1Zr(b), Figure 7.4 (etched using

Poultan’s reagent). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 175

7.6 Dependence of the lattice parameter of the metastable Al3(Zr1−xTix) (L12) phase on the

stoichiometric parameter x. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 176

7.7 Minimum creep rate at 300°C vs. applied stress, comparing data in Figure 7.2 to Al-Sc, Al-

Sc-Zr, and Al-Sc-Ti alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 177

7.8 Minimum creep rate at 350 or 400°C vs. applied stress, comparing data in Figure 7.2 to Al-

Sc-Ti alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 177

8.1 Vickers microhardness and electrical conductivity during isochronal aging of Al-0.1Zr, Al-

0.1Sc and Al-0.1Zr-0.1Sc (at.%). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183

8.2 Vickers microhardness and electrical conductivity during isochronal aging of Al-0.06Zr, Al-

0.06Sc and Al-0.06Zr-0.06Sc (at.%). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 183

A.1 Optical micrograph of a conventionally-solidified hyperperitectic Al-Ti alloy. Numerous

plate-like Al3Ti primary precipitates are observed. . . . . . . . . . . . . . . . . . . . . . . . . 214

LIST OF FIGURES xix

B.1 Observed Vickers microhardness of Al-0.2 at.% Zr subjected to various annealing treat-

ments prior to isothermal aging at 375°C for 100 h. . . . . . . . . . . . . . . . . . . . . . . . . 218

B.2 SEM micrographs of a conventionally-solidified Al-2.36Zn-0.26Zr (at.%) alloy following ho-

mogenization at 640°C for 650 h. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 218

C.1 Length scales of the analytical techniques used. . . . . . . . . . . . . . . . . . . . . . . . . . . 222

C.2 SEM micrograph displaying precipitate-rich dendrites around the specimen edge of an

electropolished TEM foil. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 223

C.3 Vickers microhardness value as a function of indent diagonal length for a 200 g load. . . . . 223

C.4 Bright-field and centered superlattice dark-field TEM micrographs showing preferentially

electropolishing around the precipitate-rich dendrites. . . . . . . . . . . . . . . . . . . . . . 224

D.1 Variation of spheroidal L12 precipitate radius vs. time for Al3(Zr1−xTix) and Al3(Zr1−xVx)

precipitates at 425°C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 226

D.2 Vickers microhardness vs. aging time for Al-0.19Zr-0.18Ti (at.%) at 425°C. . . . . . . . . . . 227

D.3 Centered superlattice dark-field TEM micrograph of Al3(Zr1−xTix) (L12) precipitates ob-

served after aging at 425°C for 400 h (preceded by a 500°C 1 h pre-age). The precipitates

have a mean radius 〈R〉 of 19.2±3.6 nm. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227

E.1 SEM secondary and backscattered electron images displaying Fe-rich precipitates decorat-

ing the grain boundaries in as-solidified Al-0.2Zr-0.2Ti. . . . . . . . . . . . . . . . . . . . . . 230

LIST OF TABLES

1.1 Reported lattice parameters and the corresponding mismatch with Al for cubic (L12) and

related tetragonal (D022 or D023) Al3M trialuminide intermetallic compounds. . . . . . . . . 14

1.2 Equilibrium maximum solid-solubility (Cmax) and solubility at 400°C (C400) in binary Al-M

alloys that form cubic (L12) and related tetragonal (D022 or D023) Al3M trialuminide inter-

metallic compounds. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

1.3 Measured diffusion data for selected transition metal solutes in α-Al. . . . . . . . . . . . . . 20

1.4 Measured diffusion data for selected lanthanide solutes in α-Al. . . . . . . . . . . . . . . . . 22

1.5 Invariant reactions in binary Al-M alloys which form cubic (L12) and related tetragonal

(D022 or D023) Al3M trialuminide intermetallic compounds. . . . . . . . . . . . . . . . . . . 25

3.1 Designations and verified compositions (at.%) of the Al-Ti and Al-Zr alloys investigated. . . 55

3.2 Calculation of chemical driving force, ∆Fch, effecting nucleation of Al3Ti (L12) and Al3Zr

(L12). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 68

4.1 Compositions and aging conditions of the Al-0.1Zr-0.1Ti alloys investigated by 3-D APT. . . 77

4.2 Number of precipitates intersected by 3-D APT, total number of atoms collected, number

of Ti atoms collected and resulting Ti concentration of the α-Al solid-solution. . . . . . . . 85

4.3 Calculated evaporation fields (V nm-1) for singly- (E1), doubly- (E2), and triply- (E3) charged

Al, Zr, and Ti ions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91

5.1 Compositions and aging conditions of the Al-Zr and Al-Zr-Ti alloys investigated. . . . . . . 101

5.2 Quantitative EDS analysis of the properitectic precipitate of Figure 5.4 in as-cast Al-0.2Zr-

0.2Ti. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 104

5.3 Mean precipitate radii, 〈R〉, observed in the dendrite centers from six different Al-Zr and

Al-Zr-Ti (three of each) specimens after extended aging times at 425°C. . . . . . . . . . . . . 122

5.4 Reported coarsening rates of L12 Al3Zr-based precipitates in Al-Zr, Al-Zr-Ti, and Al-Zr-Ti-V

alloys aged at 425°C . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 124

6.1 Compositions and aging conditions of the Al-Zr and Al-Zr-Ti alloys investigated. . . . . . . 132

xxi

xxii LIST OF TABLES

6.2 Mean precipitate radii, 〈R〉, of precipitates observed in the dendrite centers of Al-0.1Zr(b)

and Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h). . . . . . . . . . 143

6.3 Mean precipitate radii, 〈R〉, observed in the dendrite centers in Al-0.1Zr(c) and Al-0.1Zr-

0.1Ti(c) isochronally aged at 450, 525, and 575°C. . . . . . . . . . . . . . . . . . . . . . . . . . 150

6.4 Mean precipitate radii, 〈R〉, observed in the dendrite centers after different thermal histo-

ries. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 159

7.1 Compositions and aging conditions of the Al-Zr and Al-Zr-Ti alloys investigated by creep. . 168

7.2 Vickers microhardness of the aged Al-Zr and Al-Zr-Ti alloys before and after creep. . . . . . 169

7.3 Threshold stress, σth, of the Al-Zr and Al-Zr-Ti alloys crept at 300, 350, or 400°C. . . . . . . . 171

CHAPTER 1

Criteria for Developing High-TemperatureAluminum Alloys

Four criteria are described for the selection of alloying elements capable of produc-

ing castable, precipitation-strengthened Al alloys with high-temperature stability and

strength: these alloying elements must: (i) be capable of forming a suitable strengthen-

ing phase; (ii) exhibit low solid-solubility in α-Al; (iii) display a small solute diffusivity in

α-Al; and (iv) retain the ability for the alloy to be conventionally solidified. With regard

to criterion (i), only those systems forming Al3M trialuminide compounds with a cubic

L12 crystal structure are considered, which are chemically and structurally analogous

to Ni3Al in the Ni-based superalloys. Eight elements, clustered in the same region of

the periodic table, fulfill criterion (i): the first Group 3 transition metal (Sc), the three

Group 4 transition metals (Ti, Zr, Hf), and the four latest lanthanide elements (Er, Tm,

Yb, Lu). Based on a review of the existing literature, these elements are assessed in terms

of criteria (ii) and (iii), which satisfy the need for a dispersion in Al with slow coarsening

kinetics, and criteria (iv), which is discussed based on the binary phase diagrams.

1.1 Introduction

IMPROVED STRENGTH AT ELEVATED TEMPERATURES has been a continuing goal in aluminum

alloy development for more than three decades (e.g., [1–3] or a number of papers in [4]).

Aluminum-based alloys have several characteristics that make them especially attractive for the

development of high-temperature, high-strength alloys. As with the Ni-based superalloys, the

unit cell of Al is face-centered cubic (fcc), whose close-packed structure is more creep resistant

than more open crystalline structures. Moreover, Al alloys are naturally oxidation resistant due

to an extremely stable passivating protective oxide layer. For weight-sensitive applications, the

low density of Al alloys allow for materials with high specific strengths. Finally, Al alloys are con-

siderably more economical than existing high-temperature aerospace (e.g., Ni- and Ti-based)

alloys.

Historically, most efforts to develop high-strength, thermally-stable Al alloys have sought al-

1

2 CHAPTER 1 CRITERIA FOR DEVELOPING HIGH-TEMPERATURE ALUMINUM ALLOYS

loying elements that exhibit both limited solid-solubility and low diffusivity in α-Al. This ap-

proach was originally promoted by Adam [2] who argued, based on diffusion-controlled coars-

ening theory, that dispersed phases formed from such alloying additions would be resistant to

Ostwald ripening. Borne out of these ideas are the rapidly-solidified alloys based on the eutec-

tic Al-Fe system that, to date, represent the most promising high-temperature Al-based alloys.

These include the well-known Al-Fe-V-Si alloys developed by Skinner et al. [5–7], as well as more

complex Al-Fe based systems with ternary and often quaternary additions such as Ce, Ni, Co, Zr,

Mo, V [2,3,8–10]. These alloys, however, derive their high-temperature strength from a large vol-

ume fraction of stable precipitates that form directly from the melt during rapid solidification.

A conventional ingot metallurgy approach to alloy development offers several benefits, both

from performance and economic standpoints. Aluminum alloys compacted from powders pro-

duced, for example, by rapid solidification processing (RSP) and mechanical alloying (MA) are

prone to brittleness, in part due to residual porosity and the presence of oxides from the original

powders. Also, the powder production and subsequent compaction stages negatively impact the

commercial competitiveness of these alloys. More importantly, solid-state precipitation during

post-solidification aging offers the potential for generating a much finer dispersion of strength-

ening phases than those formed in the melt.

This review considers the general requirements — and challenges — for developing such a

castable, precipitation-strengthened, thermally-stable Al-based alloy. The current work is dis-

tinguished from other reviews, most notably by Starke and colleagues [11–13], devoted to the in-

telligent design of high-strength aluminum alloys. This work is concerned less with design, but

rather focuses on four broad criteria that a suitable alloying addition to Al must meet. Specifi-

cally, alloying additions that are (i) capable of forming a trialuminide strengthening phase, ex-

hibit (ii) low solid-solubility and (iii) low diffusivity in α-Al, and (iv) retain the ability for the alloy

to be conventionally solidified, are sought.

1.1.1 General strengthening requirements

Physical metallurgy of Ni-based and Al-based alloys

Before discussing the challenges of developing high-strength, high-temperature Al alloys, it is

valuable to consider what are certainly the most complex and successfully engineered high-

temperature alloys developed to date — the Ni-based superalloys. Modern Ni-based superalloys

SECTION 1.1 INTRODUCTION 3

are able to sustain stresses of the order of 150 MPa for thousands of hours, operating at temper-

atures ca. 0.75Tm. This extraordinary creep resistance is achieved primarily by additions of Al

to produce ordered Ni3Al precipitates (γ′), with the L12 structure, which are both isomorphous

and coherent with the fcc Ni-rich matrix (γ). The solubility in Ni of γ′-producing Al is substan-

tial , thus allowing for very large volume fractions of precipitated γ′ which, in many commercial

alloys, exceeds 70 vol.%.

Development of Al alloys by conventional solidification processing is subject to the restric-

tion that appreciable solubility (> 1 at.%) at equilibrium is limited to eight alloying elements —

Zn, Ag, Mg, Li, Ga, Ge, Cu, Si (in order of decreasing maximum solubility) — which are situated

near Al in the periodic table [14–17]. This limitation necessarily restricts the equilibrium volume

fraction of any precipitated phase (excluding those for the above elements of which only Ag,

Mg, Li, and Cu form stable intermetallic compounds with Al) to a value that is less than 1 vol.%,

which is one to two orders of magnitude smaller than typically found in Ni-based superalloys.

Precipitation strengthening mechanisms

To appreciate the ramifications of the generally very limited solid-solubility of alloying elements

in α-Al (and the concomitantly low volume fraction of the dispersed phase), it is necessary to

consider the quantitative effect of precipitate volume fraction on the predicted strengthening

increment in precipitation-strengthened alloys. The strengthening produced by the interaction

of dislocations with a dispersion of incoherent, inpenetrable particles within a matrix phase

was first described by Orowan [18], and has been reviewed thoroughly (e.g., [19–24]). The shear

stress required for a dislocation to loop around a precipitate is inversely proportional to the

edge-to-edge distance between precipitates, and the increase in yield strength ∆σor due to this

mechanism is given by [25]:

∆σor = M · 0.4 ·Gb

π√

(1− ν)·ln

(2Rb

)λ

; (1.1)

where M is the Taylor factor for the matrix, G and ν are the shear modulus and Poissons ratio of

the matrix, b is the magnitude of the Burgers vector, R is the mean planar precipitate radius (not

equal to the mean radius, 〈R〉), and λ is an effective inter-precipitate distance, which takes into

account the finite size of the precipitates. Both R and λ depend on the distribution of precipitate


sizes. For a monodispersed assembly, these parameters are given by [20, 22, 23]:

R =π

4〈R〉 (1.2)

and

λ =(√

2π

3f− 2

π

4

)〈R〉 ; (1.3)

where f is the precipitate volume fraction. Equations (1.2) and (1.3) are also good approxima-

tions for polydispersed arrays [23].

A plot of Orowan stress as a function of precipitate radius for different volume fractions, gen-

erated using Eq.(1.1), is displayed in Figure 1.1 with parameters pertinent to Al: M = 3.06 [26],

b = 0.286 nm [27], G = 25.4 GPa [27], and ν = 0.345 [26]. Because of the intrinsically low precipitate

volume fraction in Al-based alloys, it is critical that these dispersed phases be small (of the or-

der of 10 nm or less) and remain small (resist coarsening) throughout thermal exposure during

operation. This requirement does not apply to Ni-based superalloys, given their much higher

solubility for different alloying elements, thus allowing for higher volume fractions of the or-

dered strengthening phase, γ′. Also, the higher shear modulus of Ni (G = 78.9 GPa [27]) increases

the disparity in Orowan stresses between Ni-based and Al-based alloys.

The strengthening mechanisms are also somewhat different in coarse-grained Ni- and Al-

based alloys, as might be anticipated on account of the large disparity in precipitate volume

fraction, and in neither system is deformation at elevated temperature governed directly by the

Orowan mechanism of dislocation looping. During creep deformation of Ni-based superalloys,

dislocations are generally confined to the narrow γ channels between the large γ′ precipitates,

where complex dislocation networks form and inhibit further dislocation motion [28, 29]. In

Al-based precipitation-strengthened alloys, sufficient thermal energy is usually available under

creep conditions to allow glissile dislocations to circumvent coherent precipitates by climbing

out of their glide plane. The increase in length of the dislocation during the climb process re-

sults in a threshold stress (which is linearly proportional to the Orowan stress [30–32]), below

which creep deformation is not measurable. For coherent precipitates, elastic interactions due

to precipitate-matrix modulus and lattice parameter mismatches can further increase the creep

threshold stress [33]. These elastic interactions generally result in an optimum precipitate size

for creep resistance, whereas dislocation climb predicts that alloys with the smallest precipitates

SECTION 1.1 INTRODUCTION 5

Fig. 1.1: The Orowan stress, Eq.(1.1), as a function of mean precipitate radius, 〈R〉, for various volume fractions, f , of dispersedphase.

have the greatest threshold stress (since it is proportional to the Orowan stress), provided pre-

cipitates are not sheared. Nevertheless, the comparison in Figure 1.1 shows the critical need

for very fine, and thus very coarsening-resistant, precipitates in Al alloys to compensate for the

intrinsically limited volume fraction obtained during conventional casting of these alloys.

Precipitate stability

Ostwald ripening (coarsening) occurs during the latest stages of precipitation and involves the

growth of larger precipitates at the expense of smaller ones. The kinetics of this process are con-

trolled by volume diffusion, as solute atoms are transferred through the matrix from the shrink-

ing precipitates to the growing ones, if the evaporation-condensation model obtains. In their

classic work on the coarsening of a binary alloy, Lifshitz and Slyozov [34] and Wagner [35] (LSW)


showed that the average precipitate size 〈R〉 increases with time t according to:

〈R(t)〉3 − 〈R(t = 0)〉3 = kt; (1.4)

where 〈R(t)〉 is the average precipitate radius at time t, 〈R(t = 0)〉 is the average initial precipitate

radius at the onset of coarsening, and k is the rate constant given by [36]:

k ∝ Dσ(Cβ

e − Cαe

)2 . (1.5)

Here, D is the diffusivity of the rate-controlling solute, σ is the precipitate-matrix interfacial free

energy, and Cβe and Cα

e are the equilibrium solubilities (assuming a planar interface) of the solute

species in the precipitate and matrix phases, respectively.

For any creep-resistant alloy, it is essential that the dispersion of precipitates resist coarsen-

ing during prolonged exposure at elevated service temperatures. As indicated in Figure 1.1, this

requirement is especially imperative for Al-based alloys because of the generally limited solubil-

ity of most solutes in α-Al and the concomitant limited volume fractions (f . 0.01) of dispersed

phases.

1.1.2 Summary

Based on the behavior of modern Ni-based superalloys, which remain mechanically strong at

temperatures exceeding 75% of their absolute melting temperature, it is conceivable that Al-

based alloys could be analogously developed that would be useful to 425°C (0.75 Tm). The creep

resistance of Ni-based superalloys is conferred by very large volume fractions of the precipi-

tated Ni3Al ordered phase (γ′), that have the L12 structure. An effective high-temperature Al

alloy should thus exhibit a similar structural constitution, with suitable alloying additions to Al

exhibiting the following qualities:

Capability of forming strengthening intermetallic phases. As is true for γ′ in the Ni-based sys-

tems, a high-temperature Al alloy should contain a large volume fraction of a suitable dis-

persed phase, which must be thermodynamically stable and difficult to shear by disloca-

tions at the intended service temperature. These precipitated phases should also exhibit

a similar crystal structure to, and a low lattice parameter mismatch with, the α-Al solid-

SECTION 1.2 SELECTION CRITERIA 7

solution.

Low solid-solubility in α-Al. A low equilibrium solid-solubility at the intended service temper-

ature is necessary to retard volume diffusion-controlled coarsening (Eq.(1.5)) and prevent

dissolution of the precipitated phases. By the lever rule, limited solid-solubility also maxi-

mizes the equilibrium volume fraction of the dispersed phase.

Small diffusivity in α-Al. Limited diffusivity of the solutes in α-Al should also stifle volume

diffusion-controlled coarsening (Eq.(1.5)), allowing the precipitates to remain effective

barriers to dislocation motion at elevated temperatures.

Ability to be conventionally cast. This review is concerned with developing Al alloys via stan-

dard ingot metallurgy processing routes, and therefore the alloy must be amenable to con-

ventional casting.

In the following sections, potential alloying additions to Al with respect to each of the above

four criteria are systematically evaluated.

1.2 Selection Criteria for Castable Precipitation-Strengthened Alloys

1.2.1 Capability of forming strengthening intermetallic phases

The first criterion stated requires that any suitable system must have the capability to form a

fine dispersion of a secondary strengthening phase. As suggested by Fine et al. [1, 37, 38], in-

termetallic compounds formed with Al are the most promising candidates for strengthening

phases in ductile, thermally-stable dispersion-strengthened Al-based alloys. While a number

of potential Al-rich intermetallics can be used to strengthen Al [39], trialuminide compounds of

the type Al3M (where M is an element of the transition metals, lanthanide, or actinide series)

have particularly attractive characteristics that include low density (they are nominally 75% Al

on an atomic basis), high specific strength, good thermal stability (they have generally very high

melting points), and excellent oxidation resistance (again, mostly due to the high Al content).

Moreover, the trialuminides are directly analogous, in terms of chemistry, to the γ′ Ni3Al (L12)

ordered precipitates in the Ni-based alloys.

Extending the Ni-based alloy analogy further, it is desirable that these dispersed trialuminide

precipitates have the cubic L12 structure. The similarity in crystal structure between the matrix


Thermodynamically Stable Trialuminide Intermetallic Compounds

Orthorhombic (D011

), Monoclinic, or Unknown

Hexagonal (D019

, D018

, D024

) or Rhombohedral

Tetragonal (D022

or D023

)

Cubic (L12)

Bold outline indicates trialuminide is the Al-richest intermetallic phase

Fig. 1.2: Alloying additions to Al that form thermodynamically stable trialuminide (Al3M) intermetallic compounds, with the equi-librium structure indicated.

and precipitated phases allows for a coherent interface between the two phases which, in turn,

maximizes the strengthening efficacy of the dispersed phase (e.g., by allowing for elastic inter-

actions between dislocations and misfitting precipitates). Furthermore, coherency minimizes

the surface energy per unit area of the heterophase interface, conferring stability at high tem-

peratures by reducing the driving force for precipitate coarsening; i.e., the excess free energy

associated with the total interfacial area between the precipitate phase and the matrix.

A review of the published phase diagrams and crystallographic data [40–42] indicates that a

number of alloying additions crystallize to form stable Al3M trialuminides, as shown graphically

in the periodic table1of Figure 1.2. The high-symmetry cubic L12 and related tetragonal D022

1Groups of the periodic table are denoted by the recommended ‘new IUPAC’ (International Union of Pure andApplied Chemistry) convention where columns are numbered with Arabic numerals from 1–18, corresponding tothe number 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 Romannumerals followed by either the letter A or B. Group 3 (new IUPAC), Group III-A (old IUPAC), and Group III-B (CAS)


and D023 structures are prevalent among the early transition elements (Groups 3–5), with other

lower-symmetry structures obtained with a few of the transition elements of later groups (Fe,

Co, Ni, Re, Ir). Trialuminide intermetallic compounds are even more abundant among the lan-

thanide (rare earth, RE) elements; with the exception of Eu, all RE elements form thermodynam-

ically stable Al3RE compounds. This is also the case for several of the early actinide elements

(Th, U, Np, Pu). Other metastable Al3M trialuminides also exist, e.g. Al3Li, which is a potent

strengthening phase in aerospace Al-based alloys, but they are not discussed further since their

metastability does not fulfill the criterion of high-temperature stability. They may, however, find

use as ternary (or higher order) alloying additions in Al-M systems exhibiting stable Al3M trialu-

minides.

Trialuminides formed from the transition elements

Of the only seven thermodynamically stable L12 trialuminides (Figure 1.2), Al3Sc has generated

by far the most attention in the scientific literature (see Royset and Ryum [43] for a comprehen-

sive review of the role of Al3Sc in Al alloys). Besides Al3Sc, no other thermodynamically stable

L12 trialuminides exist among the transition elements.

Just below Sc in the periodic table is Y, which forms an Al3Y trialuminide with an equilib-

rium hexagonal D019 (Ni3Sn-type) structure. In rapidly-solidified hypereutectic alloys, however,

Foley et al. [44] reported the existence of a metastable cubic L12 Al3Y phase that formed during

solidification.

Also near to Sc are the Group 4 (Ti, Zr, Hf) and Group 5 (V, Nb, Ta) elements, which crystal-

lize with the body-centered tetragonal D022 (or D023 for Al3Zr) structures. In their monolithic

form, these trialuminides have received considerable attention as potential high-strength, high-

temperature structural materials, most notably Al3Ti since it is the least dense of this class (3.36

g · cm−3) [45–50]. Unfortunately, however, the low-symmetry tetragonal structure makes these

phases intrinsically brittle. The D022 and D023 structures are, however, closely related to the cu-

bic L12 structure (Figure 1.3) and much effort has concentrated on alloying these binary inter-

metallics to transform them to the higher symmetry L12 structure, in the hope that the increased

number of independent slip systems will improve toughness. For example, Al3Ti (D022) can be

transformed to the cubic L12 structure by alloying with late fourth-period transition elements

refer to the same column in the periodic table.


Fig. 1.3: The (a) L12, (b) D022, and (c) D023 structures. Adapted from Yamaguchi [45].

such as Cr, Mn, Fe, Co, Ni, Cu, or Zn [47, 51–61]. Similarly, Li, Cr, Mn, Fe, Ni, and Cu have been

added to Al3Zr to increase the stability of the cubic L12 structure [62–64], and Cu and Zn have

also been shown to stabilize the L12 structure of Al3Hf [65].

Carlsson and Meschter [66] and Xu and Freeman [67,68] have shown by ab initio calculations

that the stability of the D022/D023 structure relative to the L12 increases rapidly as the transition

metal d-electron count increases. Therefore, the likelihood of transforming the stable tetragonal

structure to a metastable cubic L12 phase is greater for the Group 4 (Ti, Zr, Hf) elements than it is

for the Group 5 (V, Nb, Ta). Indeed, while Al3Ti and Al3Zr have been successfully stabilized into

the L12 structure by alloying additions of late fourth-period transition elements, similar efforts

to produce cubic L12 Group 5 trialuminides, such as Al3Nb and Al3Ta, have proven unsuccessful

[47, 69].

Since the free energy difference between the equilibrium and metastable structures is so

small, one might expect the L12 structure of the Group 4 trialuminides (Al3Ti, Al3Zr, Al3Hf)

to be readily achievable when precipitated from solid-solution. Indeed, the decomposition se-

quence during aging of supersaturated Al-Ti [70–75], Al-Zr [14, 76–84], and Al-Hf [85–93] solid-

solutions has been reported to occur firstly by the formation of a metastable cubic L12 Al3M


phase, with prolonged exposure (hundreds of hours) at high temperatures (> 450°C) required

before this phase transforms to the equilibrium tetragonal Al3M structure. Furthermore, sin-

gle phase L12-structured Group 4 transition metal trialuminides (Al3Ti, Al3Zr, Al3Hf) have been

produced through mechanical alloying, which do not transform to their respective equilibrium

tetragonal structures until after heating at very high temperatures (485°C, 550°C, and 750°C for

Al3Ti, Al3Zr, and Al3Hf, respectively) [94].

The Group 4 cubic L12 trialuminides, while thermodynamically metastable, readily precip-

itate from supersaturated solid-solutions and are kinetically stable at temperatures well in ex-

cess of 400°C. These transition elements are also extraordinarily slow diffusers in α-Al as dis-

cussed below, and therefore show considerable promise as thermally stable secondary phases in

precipitation-strengthened, high-temperature Al-based alloys.

Trialuminides formed from elements of the lanthanide series

There is a monotonic decrease in the radius of the rare earth atoms across the lanthanide period,

and this variation in atomic radius has been shown to influence strongly the stability, structure,

and composition of the intermetallic compounds formed in the Al-RE systems [95–98]. With de-

creasing RE atomic radius, the structure of the corresponding trialuminide compound exhibits

progressively more cubic character [95–97]. For larger metallic radii (Z = 57–64: La, Ce, Pr, Nd,

Pm2, Sm, or Gd), the hexagonal D019 structure (Ni3Sn-type) is found. Those elements of inter-

mediate size (Z = 65–67: Tb, Dy, Ho) possess rhombohedral and hexagonal (Ba3Pb-, Ni3Ti-, and

Al3Ho-type, respectively) structures in which cubic and hexagonal stacking is mixed. For the

smallest radius RE atoms (Z = 68–71: Er, Tm, Yb, Lu) the cubic L12 (Cu3Au-type) structure is

observed.

The composition of the terminal Al-rich intermetallic phase is also strongly dependent on

the atomic radius of the RE addition. On traversing the lanthanide period from Ce (Z = 58) to Lu

(Z = 71), there is a transition in the most Al-rich intermetallic phase from Al11RE3 (also referred

to as Al4RE by some authors), being stable for early elements of the period, to Al3RE which is

stable for the late lanthanides. According to the published equilibrium phase diagrams [40–42],

Al3RE is the terminal intermetallic compound for the smaller lanthanide elements beyond Sm

(Z = 62), inclusive. The trialuminide phases of the early, larger, lanthanide series (Ce (Z = 58)

2The Al-Pm system is not well known, but is assumed to be similar to Al-Nd [40].


to Pm (Z = 61)), therefore, are not in equilibrium with α-Al and hence may be not precipitated

during aging.

The point at which this transition in composition of the Al-rich intermetallic phase occurs

has been disputed, however [98]. In the Al-Gd system, for example, Al3Gd is the equilibrium

terminal intermetallic but Savage et al. [98, 99] report the stabilization of Al4Gd/Al11Gd3 at high

cooling rates during solidification. The existence of a metastable Al4Y/Al11Y3 phase has simi-

larly been reported in the as-cast structure of rapidly-solidified Al-Y alloys [100,101], even though

Al3Y is the equilibrium terminal phase.

The transition in trialuminide structure from hexagonal/orthorhombic to cubic is also rather

sensitive to small perturbations. Figure 1.2, which displays the equilibrium structures, indicates

that thermodynamically stable cubic L12 structures are obtained for RE additions of Er (Z = 68)

and beyond. Cannon and Hall [96], however, showed that applied pressure produced a structural

transformation toward the more cubic structure in the lanthanide trialuminides, and obtained

cubic L12 structures for Al3Ho and Al3Dy through such a pressure-induced polymorphic change.

Conversely, impurities of Si may stabilize a non-equilibrium, rhombohedral (Al3Ho-type) struc-

ture for Al3Er [40, 102].

Trialuminides formed from elements of the actinide series

Consider the trialuminides of the actinide series, of which there are four equilibrium Al3M phases

(Al3Th, Al3U, Al3Np, Al3Pu) as indicated in Figure 1.2. Of particular interest are Al3U and Al3Np,

which form thermodynamically stable cubic L12 structures. None of the actinide trialuminides,

however, are in terminal equilibrium with with their respective α-Al solid-solutions [40, 41], and

they therefore cannot be precipitated from α-Al during aging. While actinide metals could be

considered as ternary elements to modify other Al3M phases, their radioactive nature prevents

practical engineering applications.

Lattice parameters of the cubic and tetragonal trialuminide phases

A small lattice parameter mismatch between the precipitate and matrix is essential for minimiz-

ing the interfacial free energy driving Ostwald ripening (Eq.(1.5)) and for maintaining a coherent

and coplanar heterophase interface between the two phases [1,37,38,84]. It is therefore useful to

compare the lattice parameters, a, among the cubic L12 trialuminides in Figure 1.2, as displayed


in Table 1.1. For the Group 4 (Ti, Zr, Hf) and Group 5 (V, Nb, Ta) elements that form metastable

cubic L12 trialuminides, data for the related equilibrium tetragonal D022 and D023 structures are

also provided. The absolute lattice parameter mismatch, δ, for the cubic L12 structures is:

δ = 100∣∣∣∣1− a

a0

∣∣∣∣ (1.6)

where a0 = 4.0496 A is lattice parameter of Al. Taking into account the mismatch along both the

a- and c- axes, the parameter δ for the tetragonal phases may be written as [37, 39, 94]:

δ =1003

[2

∣∣∣∣1− a

a0

∣∣∣∣ +∣∣∣∣1− c

n · a0

∣∣∣∣ ](1.7)

where n = 2 for D022 and n = 4 for D023.

The lattice parameters and corresponding mismatches displayed in Table 1.1 are published

values for the pure binary Al3M trialuminides, measured at room temperature, which may be

modified significantly by alloying or thermal expansion at elevated temperature.

The cubic and tetragonal phases in Table 1.1 are isostructural, respectively, and so there is

generally extensive mutual solubility between them with a concomitant shift in lattice parame-

ter, which is approximately linear with composition assuming Vegard’s law obtains. Fine et al.,

for example, showed that the lattice parameters of both the stable (D023) and metastable (L12)

Al3Zr phases could be reduced by additions of Ti, Hf, or V [84, 110, 111]. The Al-Zr-V system, in

particular, was studied extensively [84, 112–114], and the reduced lattice parameter mismatch

was observed to decrease the rate of Ostwald ripening for both metastable cubic L12 and the

equilibrium tetragonal D023 phases [37]. Similar improvements have been obtained in Al-Sc–

based alloys, where partitioning of Zr to Al3(Sc1−xZrx) [115–117] and Ti to Al3(Sc1−xTix) [118]

has been shown to offer improved stability of the precipitated Al3Sc-based trialuminides in di-

lute alloys compared to binary Al3Sc [119]. These and other Group 3 (Y), Group 4 (Ti, Zr, Hf), and

Group 5 (V, Nb, Ta) elements reduce the lattice parameter of Al3Sc (L12) [120,121], thus reducing

the mismatch with Al. As discussed later, these ternary additions to the Al-Zr [84, 112–114] and

Al-Sc [115–118] alloys are slower diffusers than either Zr or Sc, thus further improving coarsening

resistance.


Table 1.1: Reported lattice parameters and the corresponding mismatch with Al for cubic (L12) and related tetragonal (D022 orD023) Al3M trialuminide intermetallic compounds.

Phase Structure Lattice parameters Mismatch with Al Absolute mismatch, δ References(A)

Group 3 transition elements

Al3Sc L12 a = 4.103 +1.32% 1.32% [103]

Al3Y L12a a = 4.234 +4.55% 4.55% [44]


Al3Ti L12a a = 3.967 −2.04% 2.04% [94]

D022 a = 3.848 −4.98% 5.36% [39, 104]c = 8.596 +6.13%

Al3Zr L12a a = 4.08 +0.75% 0.75% [79, 94, 105]

D023 a = 4.014 −0.88% 2.89% [39, 104]c = 17.321 +6.92%

Al3Hf L12a a = 4.048 −0.04% 0.04% [94]

D022b a = 3.893 −3.87% 5.98% [106]

c = 8.925 +10.20%


Al3V L12a,c a = 3.87 −4.44% 4.44% [68]

D022 a = 3.780 −6.66% 5.35% [39, 104, 107]c = 8.321 +2.74%

Al3Nb L12a a = 4.11 +1.49% 1.49% [39]

L12a,c a = 3.92 −3.20% 3.20% [68]

D022 a = 3.844 −5.08% 5.47% [39, 104]c = 8.605 +6.25%

Al3Tad D022 a = 3.839 −5.20% 5.26% [104]c = 8.535 +5.38%

Lanthanide series (rare earths)

Al3Er L12 a = 4.215 +4.08% 4.08% [107–109]

Al3Tm L12 a = 4.203 +3.79% 3.79% [107]

Al3Yb L12 a = 4.200 +3.71% 3.71% [107]

Al3Lu L12 a = 4.187 +3.39% 3.39% [107, 109]

Actinide series

Al3U L12 a = 4.262 +5.24% 5.24% [107]

Al3Np L12 a = 4.260 +5.20% 5.20% [107]

a Metastable.b Al3Hf exists in two different crystallographic forms: a stable high temperature D023 phase and a stable low temperature D022

phase. The D022 structure is the one relevant to solid-state precipitation.c Calculated [68].d No reported metastable cubic L12 Al3Ta.


Summary

Trialuminide (Al3M-type) intermetallic compounds have many beneficial characteristics includ-

ing low density, high elastic modulus, high melting points, and are often stable with Al. They are

therefore ideal dispersed strengthening phases for high-strength thermally-stable Al-based al-

loys. The cubic L12-structured trialuminides are especially attractive since these ordered fcc

structures are commensurate with Al. While 31 elements form trialuminides when alloyed with

Al, as shown in Figure 1.2, only six elements — Sc, Er, Tm, Yb, Lu, U, and Np — form thermo-

dynamically stable cubic L12 Al3M structures. Several metastable L12 structures exist, most no-

tably among the Group 4 and Group 5 elements. The Group 4 elements (Ti, Zr, Hf) are especially

attractive since the degree of metastability of the cubic L12 trialuminide is very slight.

1.2.2 Low solid-solubility in α-Al

Figure 1.4 shows the published phase diagrams of the Group 3 (Sc, Y, La), Group 4 (Ti, Zr, Hf),

and Group 5 (V, Nb, Ta) transition elements alloyed with Al, positioned as they are situated in

the periodic table. Many of the arguments employed for considering the capacity for precip-

itation strengthening (this section) and also castability (section 1.2.4) will be made with refer-

ence to features in the equilibrium binary phase diagrams. While the particular features are

discussed later, a few patterns in Figure 1.4 are worth commenting on now. The Group 3 (Sc,

Y, La) elements exhibit terminal eutectics with Al, while the Group 4 (Ti, Zr, Hf) and Group 5

(V, Nb, Ta) are peritectics. Of the eutectic Group 3 elements, Sc is unique in that it exhibits the

highest solid-solubility and the lowest liquid solubility at the eutectic temperature. The other

eutectic-forming elements, Y and La (as well as most of the lanthanides alloyed with Al), exhibit

rather high liquid solubility of solute with very little solid-solubility, as shown in Figures 1.4(b)

and 1.4(c) for Al-Y and Al-La, respectively. These features will have a profound influence on the

conduciveness of the various systems for development into castable precipitation-strengthened

alloys, as discussed in detail later.

Requirements for precipitation strengthening

In the interest of precipitation-strengthening, a low solid-solubility is desired to maximize the

chemical driving force for nucleation and, concomitantly, the equilibrium volume fraction of


0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

650 °C

(Al)

0.078 at.%

�

��

��

��

��

��

�

(Al) + Al3Hf

L + Al3Hf

662.2 °C

��

Atomic Percent Hf

L

0.186 at.%

(Al) + Al3Hf

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

(Al) + L

�

��

��

��

��

��

��

�

(Al) + Al11

La3

640 °C

��

Atomic Percent La

L

(Al)

0.01 at.%

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

0.28 at.%

�

��

��

��

��

��

��

�

(Al) + Al3Sc

L + Al3Sc

660 °C

��

Atomic Percent Sc

L

(Al) 0.23 at.%

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

(Al) + L

�

��

��

��

��

��

�

(Al) + Al3Y

639 °C

��

Atomic Percent Y

L

(Al)

0.05 at.%

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

0.079 at.%

�

��

��

��

��

��

�(Al) + Al

3Ti

L + Al3Ti

665.4 °C

��

Atomic Percent Ti

L

(Al)

0.79 at.%

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

(Al)

0.029 at.%

�

��

��

��

��

��

��

�

(Al) + Al3Ta

L + Al3Ta

662 °C

��

Atomic Percent Ta

L

0.235 at.%

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

(Al)

0.033 at.%

�

��

��

��

��

��

�

(Al) + Al3Zr

L + Al3Zr

660.8 °C

��

Atomic Percent Zr

L

0.083 at.%

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

736 °C

688 °C670 °C

L + Al23

V4

(Al)

0.10 at.%

�

��

��

��

��

��

�

(Al) + Al21

V2

L + Al3V

662.1 °C

��

Atomic Percent V

L

0.33 at.% L + Al45

V7

L + Al21

V2

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1

300

400

500

600

700

800

900

(Al)

0.047 at.%

�

��

��

��

��

��

��

�

(Al) + Al3Nb

L + Al3Nb

661.4 °C

��

Atomic Percent Nb

L

0.066 at.%

(a) (d) (g)

(b) (e) (h)

(c) (f) (i)

Fig. 1.4: Reported binary phase diagrams for dilute (<1 at.%) additions of the Group 3, 4, and 5 transition elements alloyed with Al.References for the diagrams are: Al-Sc [42,103]; Al-Y [42]; Al-La [42]; Al-Ti [122]; Al-Zr [123]; Al-Hf [42,106]; Al-V [42]; Al-Nb [42,124];Al-Ta [42].


the precipitated phase. Moreover, according to volume diffusion-controlled coarsening theory

(Eq.(1.5)), limited solid-solubility is also favorable for retarding Ostwald ripening of the precipi-

tate dispersion, which is essential for creep resistance. As noted previously, appreciable solubil-

ity in α-Alexists only for eight elements (Zn, Ag, Mg, Li, Ga, Ge, Cu, Si) [14–17]. Solid-solubility is

therefore considered primarily from the standpoint of maximizing the potential for precipitation

strengthening.

The basic requirement for precipitation strengthening, as originally formulated in the sem-

inal work of Merica et al. [125] to explain Wilm’s serendipitous discovery of age hardening, is

decreasing solute solubility with decreasing temperature. This criterion alone is not very useful

since virtually all elements, when alloyed with Al, exhibit this behavior. A more discriminating

feature by which to compare various systems, therefore, is to consider the potential for obtaining

large volume fractions of dispersed phases which, as discussed in detail by Ryum [126], scales

with the maximum solubility, Cmax. This is exactly true only for eutectic systems as the liquid

solubility of solute in peritectic systems is often the limiting factor for determining the amount

of solute that may be quenched into solid-solution. This point is discussed in more detail when

considering castability.

A large maximum solubility, Cmax, is also essential for solutionizing the alloys in the single

phase α-Al solid-solution prior to precipitation aging, which is necessary for achieving a ho-

mogenous distribution of supersaturated solute atoms after quenching, as well as a potential

supersaturation of vacancies capable of accelerating precipitation. In addition to a large Cmax,

limited solubility at intermediate aging/service temperatures, say 400°C (C400), is favored for

driving nucleation as well as retarding Ostwald ripening of the precipitated phase.

Finally, precipitation of an Al3M trialuminide from solid-solution requires that Al3M is the

most Al-rich intermetallic compound in the system (i.e., Al3M exists in equilibrium with the

terminal α-Al solid-solution). Several of the transition metals (V, Co, La, Ir), the early lanthanides

(Ce, Pr, Nd, Pm), and the actinides (Th, U, Np, Pu) all form trialuminides that are not the terminal

intermetallic compound, and hence Al3M may not be precipitated from solid-solution in these

systems.

Table 1.2 considers the cubic L12 Al3M-forming systems in Figure 1.2 with respect to the cri-

teria favorable for precipitation strengthening. As indicated, all trialuminides except for Al3V,

Al3U, and Al3Np are in equilibrium with α-Al. Unfortunately, few of these elements exhibit ap-


Table 1.2: Equilibrium maximum solid-solubility (Cmax) and solubility at 400°C (C400) in binary Al-M alloys that form cubic (L12) and related tetragonal (D022 or D023) Al3M trialuminide intermetalliccompounds.

Element Cmax C400 α-Al-Al3M References(at.%) (at.%) equilibrium


Sc 0.23 0.01 Yes [42, 103]

Y 0.049 0.016 Yes [40, 42]


Ti 0.79 0.13 Yes [122]

Zr 0.083 0.0005 Yes [123]

Hf 0.186 0.130 Yes [42, 106]


V 0.33 < 0.15 No [42]

Nb 0.066 0.038 Yes [42]

Ta 0.235 0.05 Yes [42]


Er ≈ 0 ≈ 0 Yes [40]

Tm ≈ 0 ≈ 0 Yes [40]

Yba 0.18 < 0.1 Yes [42]

Lu ≈ 0 ≈ 0 Yes [40]

Actinide series

U 0.007 0.0037 No [40, 42]

Npb [40]

a Based on the published phase diagram in reference [42]. The original references [127,128] for this data,however, did not measure Cmax. Furthermore, reference [40] claims there is no significant solubilityof Yb in α-Al.

b No phase diagram available, although presumed to be similar to Al-U [40].

preciable solubility in α-Al. Moreover, the solid-solubility of the RE elements, several of which

form thermodynamically stable L12 trialuminides, is particularly quite low. This may be readily

explained by the substantial size difference between Al and RE atoms [95].

1.2.3 Small diffusivity in α-Al

Slow diffusion kinetics are an essential requirement for retention of strength for any alloy sub-

jected to long-term exposure at elevated temperatures. This requirement is especially true for


any Al-based system due to the intrinsically low volume fraction of dispersed phases and the

concomitant necessity for fine precipitates with exceptional coarsening resistance (Figure 1.1).

From an experimental point-of-view, it is difficult to carry out diffusion experiments in Al.

Cited difficulties include the kinetic barrier for diffusion of a radioactive Al isotope through the

passiviting oxide layer, the extremely low solid-solubility of solutes in α-Al, and a strong ten-

dency to form intermetallic compounds [129]. As a result, some of the early diffusion data for

solutes in α-Al are unreliable. The advent of readily available radioactive isotopes after 1950, and

more recent developments with ion implantation, have mitigated many of these experimental

difficulties. Nevertheless, for several elements (e.g., Ti), measurements of impurity diffusion co-

efficients are still hampered primarily by the unavailability of suitable and inexpensive radioac-

tive isotopes.

In light of these advances in experimental techniques, impurity diffusion in α-Al (especially

that of the transition elements) has received renewed interest in recent years [129–133], and

the most reliable data, with particular emphasis on the solutes forming cubic L12 Al3M inter-

metallics, are reviewed here.

Trends in diffusivity among the transition elements

It is fortunate that the transition elements are anomalously slow diffusers in α-Al, characterized

by large activation enthalpies, large pre-exponential factors, and a wide range of variation of

diffusivity values as compared with Al self-diffusion. Measured activation enthalpies for tracer

diffusion (Q) and pre-exponential factors (D0) for all of the 3d transition elements and other

selected 4d- and 5d transition elements are listed in Table 1.3. The majority of these data was

obtained from two review articles by Mehrer et al. [129, 130] and Fujikawa [131], which provide

the most authoritative theoretical interpretations for transition metal diffusion in α-Al. Data for

some of the diffusing elements were also obtained from a recent review by Du et al. [132] and

a somewhat older article by Grammatikakis [133]. Also, the comprehensive Landolt-Bornstein

review [134] was utilized.

Data for 3d transition elements and of 4sp non-transition elements (foreign atoms from the

same row of the periodic table) are depicted graphically in Figure 1.5 and 1.6. Figure 1.5 shows

the activation enthalpies for tracer diffusion versus the position in the periodic table. Figure 1.6

shows the calculated diffusivities near the melting point of Al (660°C) as well as two other tem-


Table 1.3: Measured diffusion data for selected transition metal solutes in α-Al.

Pre-exponential, D0 Activation enthalpy, Q D at 400°C Referencesm2 s−1 kJ mol−1 eV atom−1 m2 s−1 Original reference Cited in

Self-diffusion

Al 1.37× 10−5 124 1.29 3.25× 10−15 Dais et al. [135] [129]

Fourth period (3d) transition elements

Sc 5.31× 10−4 173 1.79 1.98× 10−17 Fujikawa [136] [131]

Ti 1.12× 10−1 260 2.69 7.39× 10−22 Bergner and van Chi [137] [129, 131–133]

V 1.60 303 3.14 4.85× 10−24 Bergner and van Chi [137] [129, 131–133]

Cr 10.0 282 2.92 1.29× 10−21 Rummel et al. [129] [129, 131]

Mn 8.7 × 10−3 208 2.16 6.24× 10−19 Rummel et al. [129] [129, 131]

Fe 7.7 × 10−1 221 2.29 5.41× 10−18 Rummel et al. [129] [129, 131]

Co 1.93× 10−2 168 1.74 1.76× 10−15 Rummel et al. [129] [129, 131, 132]

Ni 4.4 × 10−4 146 1.51 2.05× 10−15 Erdelyi et al. [138] [129, 131, 133, 134]

Cu 6.54× 10−5 136 1.41 1.54× 10−15 Fujikawa and Hirano [139] [129, 131, 134]

Zn 2.59× 10−5 121 1.25 1.05× 10−14 Peterson and Rothman [140] [129, 131, 133, 134]

Fifth period (4d) transition elements

Zr 7.28× 10−2 242 2.51 1.20× 10−20 Marumo et al. [141] [131, 133, 134, 142]

Mo 1.4 × 10−3 250 2.59 5.52× 10−23 van Chi and Bergner [143] [131, 133, 134]

Sixth period (5d) transition elements

La 1.40× 10−10 113 1.17 2.43× 10−19 Murarka and Agarwala [144] [145]

Hf 1.07× 10−2 241 2.50 2.11× 10−21 Minamino [131]

W 1.06× 10−3 249 2.58 5.00× 10−23 van Chi and Bergner [143] —

peratures of interest (300 and 400°C). The 3d sublevel starts to fill with Sc (Z = 21, [Ar]4s23d1)

and becomes filled at Zn (Z = 30, [Ar]4s23d10), and it is evident that valence has a strong influ-

ence on both the activation enthalpies and the corresponding diffusivities. This dependence is

most obvious in Figure 1.6, where the calculated diffusivity increases with increasing number of

d electrons from V to Co by nearly six orders of magnitude at 660°C.

It is well known that attractive or repulsive interactions among vacancies and substitutionally

dissolved solute atoms may lead to higher or lower diffusivities of solute atoms compared with


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 e

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

Ac

tiv

ati

on

en

tha

lpy

(k

J m

ol-1

)

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

Ac

tiv

ati

on

en

tha

lpy

(e

V a

tom

-1

)

Fig. 1.5: Measured activation enthalpies (Q) of 3d transi-tion element solutes and of 4sp non-transition metal so-lutes in α-Al. Data for Ga and Ge are from reference [129].

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

1 0- 2 8

1 0- 2 6

1 0- 2 4

1 0- 2 2

1 0- 2 0

1 0- 1 8

1 0- 1 6

1 0- 1 4

1 0- 1 2

1 0- 1 0

A l , 6 6 0 � C

A l , 4 0 0 � C6 6 0 � C

4 0 0 � C

3 0 0 � C

Dif

fus

ivit

y

(m2

s-1

) A l , 3 0 0 � C

Fig. 1.6: Calculated diffusivities at 300, 400, and 660°C(Tm of Al) of 3d transition element solutes and of 4sp non-transition metal solutes in α-Al. Data for Ga and Ge arefrom reference [129].

the self diffusivity of the solvent atom. This relationship between the vacancy-solute binding free

energy and the valence difference between solute and solvent atoms was originally discussed by

Lazarus [146] and further elaborated on by LeClaire [147]. While the so-called Lazarus-LeClaire

model works well for electropositive impurities in noble metals, it is known (e.g., [148, 149])

that this model is unsatisfactory for several solvents, including Al. More accurate theoretical

interpretations have been provided by Hoshino et al. [150] and most recently by Sandberg and

Holmestad [151] studying vacancy-solute interactions in α-Al based on local-density-functional

theory. The calculations indicate that while 4sp (and 5sp) solute atoms are attracted to the va-

cancy, the interaction for 3d (and 4d) transition elements is repulsive and exhibits a maximum

near the middle of the 3d and 4d rows, which is reasonably consistent with the measured activa-

tion enthalpies displayed in Figure 1.5.

Alexander and Slifkin [152] observed that the row of the diffusing species — and hence atomic

radius — also has an influence on diffusivity; the diffusivities of the Group 11 solutes, for exam-

ple, increases in the sequence Cu, Ag, Au. Nevertheless, Rummel et al. [129] noted that this

influence amounts to at most a factor of about 2, and is therefore minor compared to the strong

valency effect. Elements of the same group may be assumed to exhibit similar diffusion kinetics

in α-Al, as evidenced by the data for the homovalent Group 3 (Ti, Zr, Hf) and Group 6 (Cr, Mo,

W) solutes in Table 1.3, exhibiting a factor of about 20 difference for the diffusivity at 400°C.


Table 1.4: Measured diffusion data for selected lanthanide solutes in α-Al.

Pre-exponential, D0 Activation enthalpy, Q D at 400°C Referencesm2 s−1 kJ mol−1 eV atom−1 m2 s−1 Original Reference Cited in

Ce 1.9 × 10−10 111 1.15 4.65× 10−19 Murarka and Agarwala [144] [145, 154]

Pr 3.58× 10−11 99.4 1.03 6.94× 10−19 Murarka and Agarwala [144] [145, 154]

Nd 4.8 × 10−11 104 1.08 3.93× 10−19 Murarka and Agarwala [144] [145, 154]

Sm 3.45× 10−11 95.5 0.99 1.33× 10−18 Murarka and Agarwala [144] [145, 154]

Trends in diffusivity among the lanthanide elements

Reported diffusion data for the lanthanide or actinide elements in α-Al appears to be limited

to a study by Murarka and Agarwala [144], which reports diffusion data for the early lanthanide

elements (La, Ce, Pr, Nd, and Sm). These data are presented in Table 1.4 (Table 1.3 for La). As

indicated, the diffusivities of the lanthanides generally lie between those of the Group 3 (Sc) and

Group 4 (Ti, Zr, Hf) elements, as might be expected considering the position of the lanthanides

in the periodic table. This intermediate diffusivity of the lanthanides in α-Al is substantiated by a

decomposition study of melt spun Al-Ti, Al-Zr, and Al-Er alloys by Angers et al. [153]. The coars-

ening rates of the precipitated Al3Er (L12) phases were more rapid than those for Al3Ti (L12) or

Al3Zr (L12), most likely due to faster diffusion kinetics. The activation enthalpy for tracer diffu-

sion is very high compared to the other elements shown in Figure 1.7, and this unexpected ex-

perimental result demands that this experiment be repeated using modern analytical methods.

Even more useful would be measurements of diffusivities for the other lanthanide elements.

Summary

The transition elements are anomalously slow diffusers in α-Al, with diffusivities several orders

of magnitude smaller than that for Al self-diffusion. This anomalous behavior is ascribed to

repulsive vacancy-solute interactions which, according to experimental evidence (Figure 1.5),

reaches a maximum near the Group 5 (V, Nb, Ta) column of the periodic table. The period of the

diffusing species has only a moderate effect (compared to valence) and so homovalent solutes

of the same group in the periodic table may be assumed to obey similar diffusion kinetics in α-

Al. Diffusion of the lanthanide elements in α-Al has received comparatively little attention, but

available data for the light lanthanides indicate that diffusivities are relatively small, intermedi-


1 . 0 1 . 1 1 . 2 1 . 3 1 . 4 1 . 5 1 . 6 1 . 7 1 . 81 0

- 2 2

1 0- 2 1

1 0- 2 0

1 0- 1 9

1 0- 1 8

1 0- 1 7

1 0- 1 6

1 0- 1 5

1 0- 1 4

1 0- 1 3

1 0- 1 2

1 0- 1 1

A l

S c

S e l f - d i f f u s i o n

L a n t h a n i d e s

G r o u p 3

G r o u p 4

L a

P r

S m

��

�

��

�

��

�N d

C e

T i

H f

Z r

Dif

fus

ivit

y

(m2

s-1

)

1 0 0 0 / T ( K- 1

)

V

��

�

G r o u p 5

Fig. 1.7: Semi logarithmic plot of diffusivity in α-Al versus reciprocal temperature for the elements which form L12 trialuminidephases with Al. The data for Ce, Pr, Nd, and Sm are assumed to be representative of the L12-forming lanthanides (Er, Tm, Yb, andLu).

ate between those of the Group 3 (Sc) and Group 4 (Ti, Zr, Hf) elements.

The diffusion behavior of elements capable of forming precipitated cubic L12 trialuminides

— Al3M formed with elements of the Group 3, Group 4, Group 5, and lanthanide series and

which are in equilibrium with α-Al in Table 1.2 — is presented in the Arrhenius diagram of Fig-

ure 1.7. The diffusion data for the early lanthanides (Ce, Pr, Nd, and Sm) is assumed to be repre-

sentative of the late lanthanides (Er, Tm, Yb, Lu), which form Al3M (L12) trialuminides.

1.2.4 Ability to be conventionally cast

This review is concerned with developing precipitation-strengthened Al-based alloys produced

via conventional ingot metallurgy routes, which requires that any suitable system be amenable


k0 = C

m a x / C

L p > 1k

0 = C

m a x / C

e < 1

L + βL + α

L

α

Tm o f p u r e s o l v e n t

Cm a x

Ce C

L p C

m a x

Te

mp

era

ture

C o m p o s i t i o n

T0

Te

α + β

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

L + α

L + β

Tp

T0


( a ) ( b )

α + β

α

L

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

Fig. 1.8: Comparison between eutectic (a) and peritectic (b) reactions. In the context of the present discussion, “α” refers to α-Aland “β” refers to Al3M.

to conventional casting. Castability is considered only in general terms, primarily with respect

to features in the reported binary phase diagrams that the various elements form with Al.

The transition elements, as well as those of the lanthanide and actinide series, form rather

complex binary systems when alloyed with Al, in which one or more intermetallic phases oc-

cur [15, 17]. In these systems, eutectic phase equilibria generally exists between the liquid, the

α-Al terminal solid-solution, and the Al-rich intermetallic phase. The Al-Cu, Al-Ni, and Al-Sc

systems are familiar examples. However, α-Al solid-solutions are formed via the peritectic reac-

tion between the liquid and the Al-rich intermetallic phase with the Group 4 (Ti, Zr, Hf), Group 5

(V, Nb, Ta) and Group 6 (Cr, Mo, W) transition elements. Details of the reactions are provided in

Table 1.5 for the systems considered; i.e., those that form cubic (L12) or related tetragonal (D022

or D023) Al3M trialuminide phases.

Peritectic alloys

To appreciate the complications associated with the existence of a peritectic reaction in the

phase diagram, consider Figure 1.8 which compares the essential differences between eutectic

and peritectic systems. These distinctions are enumerated below:

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

phases in equilibrium). The eutectic reaction involves decomposition of a single phase


Table 1.5: Invariant reactions in binary Al-M alloys which form cubic (L12) and related tetragonal (D022 or D023) Al3M trialuminideintermetallic compounds.

Reaction Type Reaction Temp. Liquid solubility Solid-solubility k0 References(°C) Ce or CLp (at.%) Cmax (at.%)


Sc Eutectic 660 0.28 0.23 0.82 [42, 103]

Y Eutectic 639 ≈ 3 0.049 ≈ 0.02 [40, 42]


Ti Peritectic 665.4 0.079 0.79 10 [122]

Zr Peritectic 660.8 0.033 0.083 2.5 [123]

Hf Peritectic 662.2 0.078 0.186 2.4 [42, 106]


V Peritectic 662.1 0.10 0.33 3.3 [42]

Nb Peritectic 661.4 0.047 0.066 1.4 [42]

Ta Peritectic 662 0.029 0.235 8.1 [42]


Er Eutectic 655 ≈ 1 ≈ 0 ≈ 0 [40]

Tm Eutectic 645 1.74 ≈ 0 ≈ 0 [40]

Yb Eutectic 625 3.98 0.18a 0.045a [40, 42]

Lu Eutectic ≈ 650 ≈ 2 ≈ 0 ≈ 0 [40]

Actinide series

U Eutectic 646 1.7 0.007 0.004 [40, 42]

Npb [40]

a Based on the published phase diagram in reference [42]. The original references [127,128] for this data, however, did not measureCmax. Furthermore, reference [40] claims there is no significant solubility of Yb in α-Al.

b No phase diagram available, although presumed to be similar to Al-U [40].

liquid into two different solid phases (L → α + β), while the peritectic reaction is the for-

mation of a single solid phase 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 solute-poor

α solid-solution, whereas for a peritectic the first solid to form can be the solute-rich β

phase.

3. Reaction temperature. The reaction temperature is less than the melting point of pure sol-

vent in eutectic alloys; for peritectic systems the opposite is true.


4. Liquid-solid partition coefficient. The equilibrium partition coefficient, k0, for solidification

of the α solid-solution is less than unity in a eutectic system, while it is greater than unity

in a peritectic system. This parameter dictates the solute distribution in cast alloys and

therefore influences precipitation of dispersed phases during post-solidification aging.

The key difference — and limitation with regard to castability — of the Al-based peritec-

tic systems compared to the eutectics in Table 1.5 relates to the solidification sequence. In an

alloy of peritectic composition (i.e., with composition exceeding CLp, Figure 1.8(b)) solidified

under equilibrium conditions, the first solid to form is the solute-rich primary (or properitec-

tic) Al3M phase. Consequently there is a strong tendency to lose a significant amount of solute

to this primary phase when such peritectic alloys are conventionally cast. Moreover, as shown

in Figure 1.8, there is a necessary liquidus elevation with solute content in peritectic systems.

Thus, in order to melt (completely) a peritectic alloy (i.e., form a single-phase liquid) it is nec-

essary to heat (and hold) the melt above this elevated temperature to dissolve completely the

primary Al3M phase. This increase in melting temperature is significant in the Al-based peritec-

tic systems (e.g., Al-Ti, Al-Zr, Al-Nb) considering the extraordinarily high melting points of the

trialuminide phases (1380°C, 1580°C, and 1680°C for Al3Ti, Al3Zr, and Al3Nb, respectively).

For developing a creep-resistant alloy, peritectic systems present an additional challenge due

to the potent grain refinement associated with primary Al3M precipitation in cast alloys. It is well

known [124, 155–168] that minor additions of the Group 4 (Ti, Zr, Hf), Group 5 (V, Nb, Ta), and

Group 6 (Cr, Mo, W) elements may be used to refine the grain structure in cast Al alloys. Though

the actual mechanism is disputed [156, 157, 159, 166, 168, 169], the marked grain refinement is

generally attributed to the presence of primary Al3M precipitates, which act as heterogeneous

nuclei during solidification of the melt. To avoid rapid diffusional creep associated with a refined

grain structure it is necessary to suppress nucleation of the properitectic Al3M phase which, for

a given cooling rate, is achieved by reducing the solute content of the alloy [170–172]. Therefore,

the already-limited solid-solubility of most alloying elements in α-Al is further reduced by the

necessity to avoid properitectic precipitation in the peritectic systems.

What characteristics of the equilibrium phase diagram are conducive to properitectic sup-

pression? In alloys produced by RSP, the resistance of peritectic alloys to primary Al3M formation

has been shown [173, 174] to be related to the undercooling necessary to reach the metastable

liquidus of the α-Al solid-solution; i.e., the temperature difference, at a composition of inter-


est, between the β liquidus and extrapolated α-Al liquidus. While RSP is outside the scope of

the present discussion, it is useful to invoke similar arguments in evaluating peritectic alloys in

terms of conventional castability.

A minimum undercooling for quenching into the metastable α-Al phase field is favored in

systems that have a shallow β liquidus and a steep α-Al liquidus. In the six Al-based peritectic

systems of interest, the peritectic reaction temperatures are only a few degrees above the melt-

ing point of pure Al (Figure 1.4, Table 1.5), and so the slopes of the α-Al liquidus curves in these

systems are shallow. Therefore, the slope of the β liquidus is the primary influence determin-

ing the ease of properitectic suppression. A second approach to minimizing this undercooling

(without changing the slope of the β liquidus) is achieved by a general increase in the solubility

of β in the liquid — i.e., a shift to the right of the liquidus curve.

Castability in peritectic alloys is therefore dictated primarily by the liquid solubility of so-

lute, whereas the propensity for precipitation strengthening, as discussed previously, is deter-

mined by the solid-solubility; high liquid solubilities being favored for castability and low solid-

solubilities required for precipitation. These criteria are conveniently expressed in terms of the

equilibrium solid-liquid partition coefficient, k0, which is the ratio of solute composition of the

solid and liquid phases in local equilibrium during solidification [175–177]3. Making the usual

assumption that both the liquidus and solidus boundaries are straight lines, k0 is then constant

at all temperatures and may be expressed in terms of the solid and liquid compositions at the

peritectic temperature: k0 = Cmax/CLp. All peritectic alloys exhibit partition coefficients greater

than unity (Figure 1.8(b)), and so maximum liquid solubility and minimum solid-solubility is

obtained in systems with partition coefficients approaching unity (i.e., minimum disparity be-

tween liquid and solid-solubilities).

Figure 1.9 shows undesirable and desirable features of hypothetical peritectic systems based

on the preceding arguments. In Figure 1.9(a) it is apparent that large undercoolings are nec-

essary to obviate the primary β phase when casting peritectic alloys of this type. Moreover, on

account of the large disparity in liquid and solid-solubilities (k0 � 1), it is difficult to achieve

a significant supersaturation of solute for post-solidification aging. This reduced driving force

3The partition coefficient k0 is usually defined where the compositions are in wt.%. Compositions are reportedin at.% since, assuming equal molar volumes between Al3M and Al (reasonable considering the data in Table 1.1),the supersaturation (in at.%) is a direct measure of the equilibrium volume fraction of precipitated phase. Further-more, the difference between k0 defined by compositions in wt.% or at.% is negligible for the dilute concentrationsconsidered presently.


CL p

Cm a x

CL p

Cm a x

~Tp~ T

m

S t e e p β l i q u i d u s

k0> > 1

L + β

L + α

L

α

Te

mp

era

ture


Tm

α + β

S t e e p α s o l v u s

S h a l l o w β l i q u i d u s

~k0~ 1

L + α

L + β

Tp> > T

m

Tm


( a ) ( b )

α + βα

L

S h a l l o w α s o l v u s

Fig. 1.9: Undesirable (a) and desirable (b) characteristics of peritectic systems, as they relate to the potential for development ofcastable precipitation-strengthened alloys.

for precipitation is exacerbated by a relatively steep α-Al solvus, i.e., the solid-solubility does not

decrease appreciably with temperature.

The hypothetical phase diagram in Figure 1.9(b) exhibits features favorable for conventional

solidification processing. By virtue of the large liquid solubility of the β phase, a considerable

amount of solute may be added before precipitating primary β during solidification. Further-

more, there is a relatively low solubility of solute at the peritectic reaction temperature (k0 ≈ 1),

which diminishes rapidly with decreasing temperature, providing a maximum supersaturation

of solute for precipitation strengthening. It is even conceivable, on account of the large liquid

solubility coupled with the very low solid-solubility, that nonperitectic compositions (i.e., with

composition less than CLp [171]) may provide appreciable precipitation strengthening in a sys-

tem resembling Figure 1.9(b). This is most desirable since the possibility of nucleating properi-

tectic β, which is responsible for most of the complications of casting peritectic alloys (loss of

solute during solidification and grain refinement), is eliminated.

Finally, the reaction temperature Tp in the system of Figure 1.9(b) is much greater than that

in Figure 1.9(a). Not only does this tend to maximize the slope of the α-Al liquidus (desirable for

minimizing the undercooling necessary to suppress properitectic β) but, more importantly, such

a feature is favorable in any high-temperature alloy since a high reaction temperature extends

the temperature range where the reinforced two-phase (α+β) solid is thermodynamically stable.

SECTION 1.3 DISCUSSION 29

This could, in principle, extend the service temperature of the alloy beyond the melting point of

pure Al.

Eutectic alloys

Casting eutectic alloys (of hypoeutectic composition) is considerably less complicated than for

peritectic alloys since the first solid to form is the solvent-rich α solid-solution (Figure 1.8(a)).

Nevertheless, it is worthwhile commenting on a few features in the eutectic phase diagram as

they relate to the potential of a given system for developing high-strength high-temperature al-

loys.

As in any alloy undergoing dendritic solidification, microsegregation of solute is minimized

for solid-liquid partition coefficients, k0, approaching unity. In eutectic alloys, k0 = Cmax/Ce is

less than unity since the α-Al liquidus has a greater solubility than the α-Al solidus, Figure 1.8(a).

The disparity in solute solubility is what drives microsegregation, and so systems in which k0 is

near unity are especially amenable to casting, since the degree of microsegregation is small.

For high-temperature applications, a eutectic system poses a potential limitation since, by

definition, the reaction temperature is less than the melting point of pure Al (Figure 1.8(a)). In an

extreme case, the eutectic reaction temperature could conceivably limit the service temperature

of the alloy since its melting point is reduced. Deep Al-M binary eutectic temperatures exist for

elements such as Ga and Zn, but not for those forming stable or metastable L12 trialuminides

(Figure 1.4, Table 1.5).

1.3 Discussion

1.3.1 Alloying additions with transition elements

As indicated in Table 1.1, transition elements that form cubic L12 trialuminides are limited to

the early elements near Sc in the periodic table (Ti, Zr, Hf, V, Nb, Y), with Al3Sc as the only

thermodynamically stable L12 structure. While a metastable cubic L12 Al3Y has been reported

[44], this investigation was on highly supersaturated melt spun alloys and precipitation of Al3Y

(L12) was observed to commence during solidification. The likelihood of precipitating a similar

phase in a conventionally-solidified alloy during post-solidification aging seems unlikely.

The Group 4 (Ti, Zr, Hf) and Group 5 (V, Nb, Ta) elements form metastable cubic L12 trialu-


minides, with the degree of metastability increasing from the former to the latter group (section

1.2.1). Therefore, the Group 5 trialuminides such as Al3Nb and Al3Ta are poor candidates for

cubic L12 modification. Indeed, as indicated in Table 1.1, the existence of Al3Ta (L12) is un-

known to the authors. This is unfortunate, since both the Al-Nb and Al-Ta peritectic systems

(Figures 1.4(h) and 1.4(i), respectively) exhibit favorable characteristics (small k0 and limited

solid-solubility, C400) and, considering they are also Group 5 elements, should be very slow dif-

fusers in α-Al (Figure 1.7). Furthermore, unlike the Group 4 (Ti, Zr, Hf) elements, reports of

precipitation during aging of Al-Nb and Al-Ta alloys are unknown to the authors. Among the

transition elements, therefore, Sc and the neighboring Group 4 (Ti, Zr, Hf) elements seem to of-

fer the most potential for producing a precipitated dispersion of L12 precipitates formed during

solid-state aging.

Stable L12 Sc trialuminide precipitates

Aside from being the only system forming a thermodynamically stable cubic L12 trialuminide,

Al-Sc exhibits other unique features that make it attractive for developing conventionally-cast

precipitation-strengthened alloys. As displayed in Figure 1.4(a) and Table 1.5, the eutectic Al-Sc

system features an exceptionally high reaction temperature (660°C, within 1°C of the melting

point of pure Al), with a relatively large solid-solubility (0.23 at.% Sc) and comparatively low

liquid solubility (0.28 at.% Sc) at the reaction temperature. This similarity in solid and liquid

solubilities leads to an equilibrium partition coefficient, k0, very near unity (k0 = 0.82, Table 1.5),

which minimizes segregation during solidification and makes this system especially amenable

to conventional casting. The extremely high eutectic temperature does not limit the service tem-

perature of Al-Sc alloys for high-temperature applications. In fact, the liquid and solid eutectic

solubilities are so close and the reaction temperature so high that the first published phase di-

agrams for Al-Sc indicated a terminal peritectic reaction like the neighboring Al-Ti, Al-Zr, and

Al-Hf systems [43].

The maximum solid-solubility of Sc in α-Al is relatively large, exceeded for the elements in

Table 1.2 only by Ti and V (and similar to Ta), and it diminishes significantly with decreasing tem-

perature (C400 = 0.01 at.% Sc, Figure 1.4(a) and Table 1.2), thus maximizing the volume fraction of

Al3Sc achievable during post-solidification aging. Of the systems in Table 1.2, only Al-Ti and Al-

Ta exhibit a comparably high decrease in solid-solubility with temperature. This fortuitous con-


fluence of desirable characteristics — a stable L12 trialuminide, a solvus boundary conducive to

precipitation strengthening, eutectic phase equilibria favoring conventional solidification, and a

very high eutectic temperature — makes the Al-Sc system particularly well-suited for developing

castable, precipitation-strengthened alloys. Additions of Sc, forming coherent nanoscale Al3Sc

(L12) precipitates, provides the highest increment of strengthening (at room temperature) per

atomic percent of any alloying element added to Al [178]. In addition to the marked precipita-

tion hardening response, these precipitates are also very effective at inhibiting recrystallization

and maintaining a fine uniform microstructure in wrought Al alloys [43].

The potent strengthening from precipitated Al3Sc (L12), the high eutectic temperature, and

the relatively small diffusivity of Sc in α-Al (Figure 1.7) indicate that Al-Sc alloys possess signifi-

cant potential for developing conventionally-solidified creep-resistant Al-based alloys. Indeed,

recent studies by Dunand, Seidman, and colleagues have shown conventionally-solidified Al-

Sc alloys exhibiting remarkably high coarsening and creep resistance at 300°C [179–181], which

may be improved with ternary additions of Mg [182–184], Zr [115–117], and Ti [118]. The latter

two elements, together with many other candidates with diffusivitivies below that of Sc [120],

segregate to the Al3Sc phase without changing its L12 structure. These Sc substitutions also re-

duce the relatively high cost of Sc additions.

Metastable L12 trialuminides of the Group 4 transition elements

The Group 4 transition elements (Ti, Zr, Hf) are extremely slow diffusers in α-Al (Figure 1.7),

and therefore offer potentially significant improvements in thermal stability compared to Al-Sc

alloys. While the L12 trialuminides formed from the Group 4 elements are metastable, these

are slow to transform to their respective equilibrium tetragonal structures (section 1.2.1) and

therefore seem promising as thermally stable dispersed phases in high-temperature Al-based

alloys. The Group 4 elements, however, form peritectics with Al which, as discussed in section

1.2.4, introduces significant complications to produce alloys by conventional ingot metallurgy

routes.

The Al-Ti system (Figure 1.4(d)) seems especially attractive due to the very sluggish diffusion

kinetics of Ti in α-Al (Figure 1.7). Moreover, of all the systems in Table 1.2, Ti has the highest

maximum solid-solubility in α-Al (0.79 at.% Ti), suggesting that relatively high volume fractions

of precipitated Al3Ti may be obtained. The Al-Ti system has, however, a comparatively limited


liquid solubility (k0 = 10, Table 1.5). Hence, for alloy compositions approaching the maximum

solubility there will be a strong tendency for primary Al3Ti precipitation during solidification

and so appreciable Ti concentrations are unlikely to remain in solid-solution after conventional

solidification. Moreover, the solid-solubility of Ti is still rather large at appropriate aging tem-

peratures (C400 = 0.13 at.% Ti), limiting both the chemical driving force for nucleation and the

volume fraction of the dispersed phase (if formed). As reviewed in Chapter 3, virtually all precip-

itation studies for the Al-Ti system have investigated alloys prepared by nonequilibrium means:

RSP and MA, techniques that circumvent the difficulties encountered during conventional so-

lidification of Al-Ti alloys.

In the Al-Hf system (Figure 1.4(f)) there is a relatively small disparity between the liquid and

solid-solubilities at the peritectic temperature (k0 = 2.4, Table 1.5), implying that alloys with com-

positions corresponding to the maximum solubility (0.186 at.% Hf, Table 1.2) may be obtained

with conventional casting techniques. This solid-solubility, however, diminishes only slightly

with decreasing temperature (C400 = 0.130 at.% Hf, Figure 1.4(f) and Table 1.2), which limits the

potential for precipitation strengthening. Indeed, in all of the precipitation studies on Al-Hf al-

loys cited previously [85–93], the alloys were generally highly supersaturated and produced by

relatively rapid (nonequilibrium) chill casting methods. Ryum [85], for example, investigated

the decomposition of supersaturated 0.27 at.% Hf solid-solutions produced by chill casting. Al-

though the cooling rate was not reported, it is reasonable to believe it was about 3 × 103°C s-1,

which was the solidification rate reported by Hori et al. [86, 87, 91, 92] on investigations of chill

cast alloys containing up to 0.79 at.% Hf. It was reported [86–88] that the solid-solubility of Hf in

α-Al could be extended to approximately 0.5 at.% Hf at this cooling rate, and such high values of

supersaturations, well in excess of the maximum solid-solubility, were required to effect contin-

uous precipitation of Al3Hf (L12) [88], with a pronounced precipitation hardening response in

alloys aged at 350 to 450°C [86, 87]. Other studies investigated even more supersaturated alloys

(up to 1.0 at.% Hf) obtained by rapid solidification at a rate of approximately 107°C s-1 [88,90,93].

Unlike Ti and Hf, Zr exhibits negligible solubility in α-Al at temperatures of interest for post-

solidification aging (C400 < 0.001 at.% Zr, Figure 1.4(e) and Table 1.2). Consequently, even very

dilute alloying additions may produce an appreciable precipitation hardening response as shown

by Hori et al. [185, 186], who reported precipitation hardening in alloys containing 0.07 at.% Zr

aged at 350 to 450°C. Ichikawa and Ohashi [187] similarly observed a significant age hardening


response when dilute chill cast alloys, containing as little as 0.12 at.% Zr, were aged at 300 to

500°C. Furthermore, Nes [79] has reported precipitation of Al3Zr (L12) in very dilute (hypoperi-

tectic, 0.05 at.% Zr) alloys.

Moreover, the role of dilute additions of Zr to commercial wrought alloys as recrystallization

inhibitors is well known [78,188–191], where fine (20–30 nm) dispersions of coherent Al3Zr (L12)

precipitates are used to pin grain and subgrain boundaries during annealing. The Zr concen-

tration required to inhibit recrystallization in conventionally cast commercial alloys is very low,

typically 0.03–0.06 at.% Zr [191].

1.3.2 Alloying additions with lanthanide elements

There would seem to exist significant potential for the use of late rare earth elements (Er, Tm,

Yb, Lu) in Al, as they form thermodynamically stable L12 trialuminides (Figure 1.2), are likely

slower diffusers than Sc (Figure 1.7), and, unlike the Group 4 elements, exhibit the preferred

eutectic phase equilibrium (Table 1.5), which is favorable for conventional ingot metallurgy

processing. Indeed, a number of Al-RE systems — Al-La [100, 192–194], Al-Ce [195, 196], Al-

Nd [195,197,198], Al-Gd [98,99,198,199], Al-Sm [99], and Al-Er [99,153,192,198,200–202], in par-

ticular — have received considerable attention for development into dispersion-strengthened

alloys for elevated temperature applications. All of these studies, however, investigated highly-

supersaturated (sometimes hypereutectic) alloys prepared by RSP. The attraction to the RE ele-

ments stems from their generally large solubility in the liquid state, very limited solid-solubility,

and small solid state diffusivity in α-Al. The large liquid solubility is conducive to solid-solubility

extension by RSP, while the very limited solid-solubility and small diffusivity retards volume

diffusion-controlled coarsening of the dispersed intermetallic phases during thermal exposure.

As in the Al-Fe–based alloys described previously, the strengthening dispersions in these al-

loys form during solidification, either as lamellar eutectic constituents, as a primary phase in

hypereutectic alloys, or are precipitated in the solid state on cooling from solidification. The

microstructures in these rapidly-solidified alloys, as monitored by microhardness, are generally

stable beyond 300°C [98,192,193,195–197,199,201,202], as might be anticipated due to the small

diffusivity and solid-solubility of the RE elements in α-Al.

Since the strengthening dispersions form on solidification, precipitation hardening during

post solidification aging is generally not observed in these alloys. This lack of hardening is a


consequence of the very limited equilibrium solubility of the RE elements in α-Al (Table 1.2),

which assures almost complete precipitation of solute on cooling. Even in highly supersatu-

rated Al-La (0.8 at.% La [192, 193]) and Al-Er (0.7 at.% Er [192, 202]) alloys produced by RSP, the

amount of La and Er retained in solid-solution was negligible as indicated by lattice parameter

measurements using x-ray diffraction (XRD) of the as-solidified α-Al solid-solution.

Solid-state precipitation of Al3Er (L12) has, however, been reported by Angers et al. [153] in

extremely supersaturated (hypereutectic) rapidly-solidified Al-Er alloys containing 6.25 at.% Er.

Furthermore, a weak indication of additional precipitation hardening (beyond the strengthening

provided by the as-solidifed dispersions) was shown in other studies on Al-Er alloys [201], as well

as Al-Gd alloys [98, 199], but the origin of this hardness increase was not reported.

More support for Al-Er as a promising system capable of precipitation strengthening may

be provided by Nie et al. [203–205], who indicated that coherent nanoscale Al3Er (L12) disper-

sions can be potent recrystallization inhibitors in Al alloys, comparable in effect to the more

commonly-used Al3Zr and Al3Sc phases in commercial alloys. Although their experimental

techniques were not explicit about this point, micrographs indicate that coherent Al3Er precip-

itates were not precipitated during aging but rather formed during cooling after solidification,

similar to what Foley et al. [44] reported for metastable Al3Y (L12) precipitation described pre-

viously.

While the Al-RE systems show considerable promise for development by RSP, their conducive-

ness to conventional solidification is limited. Indeed, the very properties that make the Al-RE

systems attractive for RSP, namely a very high liquid solubility and low solid-solubility (in other

words k0 � 1), limits their potential when these alloys are produced by conventional casting

techniques. Because of the very limited solid-solubility, a supersaturation of solute is impossible

to achieve under moderate cooling rates since there will be a strong driving force for precipita-

tion during cooling, as shown in studies on Al-Er [203, 204] and Al-Y [44]. A partition coefficient

deviating far from unity also predicts that these alloys will be prone to significant solute seg-

regation during solidification, as substantiated by Ruder and Eliezer [193, 202] comparing the

microstructures of Al-La and Al-Er solidified conventionally and with RSP. Moreover, due to the

very limited maximum solid-solubility of the RE elements in α-Al (Table 1.2), post-solidification

homogenization is not possible.

SECTION 1.4 CHAPTER SUMMARY 35

1.3.3 Summary

It is unfortunate that the slowest diffusers in α-Al — i.e., the Groups 4–6 transition elements —

are also the only elements in the periodic table forming terminal peritectics with Al (all other

alloying additions exhibit eutectic or monotectic phase equilibria) [15]. The existence of a peri-

tectic reaction in the phase diagram reduces significantly the conduciveness of these systems to

conventional casting. It is equally unfortunate that all of the late lanthanide elements, with the

possible exception of Yb (see notes in Table. 1.2 and 1.5), exhibit negligible solid-solubility in α-

Al. This severely limits the potential for precipitation strengthening since virtually all available

solute is precipitated out of solid-solution during post-solidification cooling.

The Group 3 element, Sc, exhibits a unique eutectic phase equilibrium with Al, which makes

it particularly conducive to developing conventionally-solidified precipitation strengthened al-

loys. Moreover, Al3Sc is a thermodynamically stable L12 trialuminide, which exists in two-phase

equilibrium with Al to the high eutectic temperature of 660°C. Unfortunately, the very high cost

of Sc limits its application. Moreover, compared to the other transition elements, Sc is only a

moderately slow diffuser in α-Al (Figure 1.7).

Of the slower-diffusing peritectic-forming elements that form L12 trialuminides (Groups 4

and 5), the Group 4 (Ti, Zr, Hf) are strongly favored since the L12 structure in these systems is

only slightly metastable. Of the Group 4 elements, only Zr exhibits negligible solubility at suit-

able aging temperatures (Figure 1.4), both maximizing the volume fraction of the precipitated

Al3Zr phase as well as improving its resistance to coarsening.

1.4 Chapter Summary

This review has summarized basic criteria required for a conventionally-solidified precipitation-

strengthened Al-based alloy for high-temperature applications:

(i) Solid-state precipitation upon aging of coherent trialuminide Al3M with the L12 crystal

structure, to achieve high strengthening and low coarsening;

(ii) Shallow α-Al solvus curve, to maximize the volume fraction of precipitated Al3M (and thus

increase strengthening); and the concomitant low solid-solubility at the aging temperature

to minimize coarsening;


(iii) Small diffusivity of M in α-Al, to minimize Al3M coarsening and the associated loss of

strength;

(iv) Solid-liquid partition coefficient (k0) near unity, to minimize segregation and accommo-

date conventional solidification. For peritectic systems, a shallow Al3M liquidus boundary

is desirable for minimizing the casting temperature and suppressing Al3M primary precip-

itation during solidification.

Criterion (i) narrows the potential candidates to only eight elements, all situated near Sc in

the periodic table: the first Group 3 element (Sc), the three Group 4 elements (Ti, Zr, Hf) and the

four heaviest rare-earth elements (Er, Tm, Yb, Lu). Among these elements, Sc and Zr stand out

for the following reasons:

• The Al-Sc system exhibits a unique combination of a shallow solvus curve conducive to

precipitation strengthening, an eutectic phase equilibrium favoring conventional solidi-

fication, and thermodynamically stable Al3Sc with the L12 structure. Sc is, however, the

fastest diffuser in α-Al and the priciest of the above eight elements.

• The Al-Zr system is characterized by one of the smallest diffusion rates and lattice param-

eter mismatch between Al3M and Al, as well as price among the eight candidates. The

L12 structure of Al3Zr, however, is metastable. Furthermore, the Al-Zr system is peritectic

which limits the concentration of Zr retained in solid-solution after conventional solidifi-

cation, ultimately limiting the strength attainable by precipitation strengthening. Unlike

the other Group 4 systems (Al-Ti and Al-Hf), the solid-solubility of Zr in α-Al is negligi-

ble and so precipitation strengthening can be obtained in conventionally cast alloys. The

steep α-Al solvus, however, limits the possibility of post-solidification homogenization.

CHAPTER 2

Peritectic Solidification

The peritectic reaction is the formation of the peritectic phase (α) by reaction of the pri-

mary phase (β) directly with the liquid (L) at the peritectic temperature (Tp). Despite the

existence of a peritectic reaction in an equilibrium phase diagram the possibility of the

reaction taking place is dependent on the average alloy composition and the cooling

rate from the melt, with increasing cooling rates and decreasing solute concentration

generally limiting the extent of the reaction. Still faster rates of cooling prevent nucle-

ation of the properitectic β phase, and supersaturated α may be formed directly from

the melt. Terminal peritectic equilibrium introduces several complications with regard

to castability, which were introduced in Chapter 1. This chapter is about understanding

these complications, and more importantly, how to avoid them.

2.1 Introduction

IT WAS SHOWN IN CHAPTER 1 that the slowest diffusers in α-Al occupy the Groups 4–6 in the

periodic table. It is unfortunate that these same elements — Ti, Zr, Hf, V, Nb, Ta, Cr, Mo,

W — are also the only elements exhibiting peritectic phase equilibria with the terminal α-Al

solid-solution. As discussed (Section 1.2.4), this peritectic equilibrium limits the feasibility for

conventional solidification of these alloys because of the associated:

1. Liquidus elevation. The liquidus of a peritectic alloy rises continuously with increasing so-

lute concentration, Figure 1.8(b).

2. Solute loss from primary precipitation of Al3M. In a peritectic alloy (of peritectic compo-

sition) the first solid form is the solute-rich properitectic β phase (Al3M for the systems

discussed). Under equilibrium conditions, the solute concentration of the liquid phase

follows the liquidus with the last solid to form having a composition necessarily less than

the maximum solid-solubility.

37

38 CHAPTER 2 PERITECTIC SOLIDIFICATION

3. Grain refinement associated properitectic precipitation. When Al3M forms primarily in the

melt there is a concomitant refinement of the as-cast grain size of the alloy. This is dele-

terious for creep resistance since a refined grain structure can promote rapid diffusional

creep.

While the increase in melting temperature is unavoidable, the extent of properitectic Al3M

precipitation (and hence the amount of solute lost to this phase, as well as the degree of grain

refinement) is highly dependent on the concentration of solute in the alloy and the cooling rate

at which the alloy is solidified. It is necessary, therefore, to understand the kinetics are of the

peritectic reaction and what non-equilibrium solidification sequences are possible. The seminal

paper on this topic, which was on dilute Al-Ti alloys, is that of Kerr et al. [170]. The first line of this

paper aptly reads: “The peritectic transformation appears to have puzzled students and authors

alike.”

2.2 Nomenclature

Establishing clear definitions will be beneficial before discussing in detail the possible solidifi-

cation sequences in peritectic systems. Peritectic solidification has called the attention of many

researchers, most notably Kerr et al. [170,171] and St. John et al. [172,206–209], and consequently

ambiguities in nomenclature have arisen [171]. We will use the nomenclature introduced by St.

John with terminology suggested by Kerr and Kurz [171].

A peritectic reaction as part of a schematic binary phase diagram is shown in Figure 2.1. The

peritectic reaction involves decomposition of β, and hence this phase is termed the properitectic

phase; similarly, the product phase, α, is called the peritectic phase. The term peritectic composi-

tion refers to the composition Cαp of the peritectic phase α at temperature Tp, and corresponds

to the maximum solid-solubility of solute in the α phase. This is an invariant point and is analo-

gous to the eutectic composition in eutectic systems, since only at the composition Cαp can the

high temperature β and L phases be replaced completely by the lower temperature α phase on

cooling. The liquid composition CLp is termed the peritectic liquid.1

At temperatures above Tp all alloys to the right of the peritectic liquid CLp will initially form

some β phase directly from the liquid, whereas alloys with less than CLp will not. Since the pres-1As discussed by Kerr and Kurz [171], CLp has also been termed the ‘peritectic point’ by others [170, 175] or ‘peri-

tectic limit’ by St. John [172, 206, 208]. Moreover, ‘peritectic point’ has also been used to refer to Cαp by St. John [206].

SECTION 2.3 MECHANISMS OF α FORMATION 39

A l

H y p e r p e r i t e c t i cH y p o p e r i t e c t i cN o n p e r i t e c t i c

Tm

Tp

β

α + β

α + L

L + β

Te

mp

era

ture

C o m p o s i t i o n ( % M )

L

α

CL p

Cα p

Cβ p

Fig. 2.1: Schematic peritectic phase diagram indicating symbols used for phases, compositions, and temperatures. During theperitectic reaction, primary phase (β) of composition Cβp reacts with liquid (L) of composition CLp to form peritectic α phase withcomposition Cαp.

ence of β (the properitectic phase) is required for the reaction, these alloys are thus designated

as peritectic (C0 > CLp) and nonperitectic (C0 < CLp). Peritectic alloys are further distinguished

as hypoperitectic (C0 < Cαp) and hyperperitectic (C0 > Cαp), as indicated in Figure 2.1.

2.3 Mechanisms of α Formation: The Peritectic Reaction, Peritectic

Transformation, and α Nucleation Directly from the Melt

Solidification of dilute Al alloys involves formation of α-Al solid-solution from the melt (L). For

a peritectic alloy there are three possible mechanisms of forming α-Al: (i) the peritectic reaction;

(ii) the peritectic transformation; and (iii) direct nucleation of supersaturated α-Al.


2.3.1 Peritectic reaction

Consider a peritectic alloy (i.e. of initial composition between CLp and Cβp) cooled slowly from

the melt. The equilibrium freezing sequence begins with the nucleation of the primary phase β

below the liquidus, with subsequent growth of this phase as the L + β mixture is cooled. With-

drawal of solute from the melt will cause the liquid composition to follow the β liquidus with de-

creasing temperature until the interdendritic liquid has composition CLp at the peritectic tem-

perature Tp. At this point, liquid of composition CLp and primary β of composition Cβp react

together to give α:

L + β → α.

The Gibbs Phase Law requires that, at Tp, the three phases remain in local equilibrium until the

reaction is complete, resulting in final equilibrium quantities of α, β, and L as dictated by the

bulk alloy composition C0 and the phase diagram. For an alloy precisely of peritectic composi-

tion, Cαp, all of the α phase is produced by decomposition of properitectic β via the peritectic

reaction, with no β remaining thereafter. For hypoperitectic alloys (i.e. those with initial com-

position between CLp and Cαp), there is enough liquid remaining at Tp to consume all of the

primary β which has formed during cooling between the β liquidus and Tp. Hyperperitectic al-

loys (between Cαp and Cβp) have insufficient liquid to react with all of the properitectic β so that

some primary phase remains in the final solidified alloy [209].

2.3.2 Peritectic transformation

In reality, the peritectic reaction can continue only until each primary β particle is completely

encapsulated with α, at which point the three phase contact required for equilibrium is no longer

possible. However, the α phase may continue to form at the expense of β by diffusion through

the peritectic envelope enclosing each primary particle. This mechanism is termed the peritectic

transformation.The rate of α formation by this mechanism is diffusion-controlled, and St. John

and Hogan [206] showed that the peritectic transformation is negligible in the Al-Ti system, with

the majority of the peritectic α solid-solution formed directly from the melt. Similar behavior is

expected for the other Al-based peritectic systems. Since the peritectic transformation does not

occur to any significant extent, properitectic β formed in the liquid is not consumed and hence

the solute contained in this phase cannot be recovered back into α.

SECTION 2.3 MECHANISMS OF α FORMATION 41

2.3.3 Direct solidification of supersaturated α

Consider the schematic free energy curves in Figure 2.2. Under equilibrium conditions a hy-

perperitectic liquid of composition C0 will lower its energy by the nucleation of properitectic β

with the attendant change in the liquid composition to that of the β liquidus (this is the process

of primary β precipitation described in Section 2.3.1). The nucleation of the primary β phase,

however, requires a severe fluctuation in composition (25 at.% solute for a trialuminide primary

Al3M phase compared to C0 . 0.1 at.% for the dilute Al alloys studied in this thesis). As shown in

Figure 2.2, the free energy of the system can be lowered metastably by forming supersaturated

α, which may be favored kinetically under accelerated cooling rates.

In Figure 2.2, the common tangent to the α and liquid free energy curves touches them at

higher solute concentrations than the equilibrium liquid composition, and these compositions

will correspond to extrapolations of the equilibrium solidus and liquidus curves since both the

equilibrium phase boundaries and the metastable extensions are defined by common tangents

to the same curves [170]. This point was demonstrated convincingly by Kerr et al. [170, 210] in

their seminal study on Al-Ti alloys. Using microprobe and concurrent thermal analyses, these

researchers showed that with relatively high cooling rates a supersaturated α solid-solution nu-

cleated above Tp (Figure 2.3), with the measured liquidus and solidus compositions following

their respective extrapolated values.

Considering the equilibrium β liquidus, which increases rapidly with solute concentration

(Figure 2.1), the amount of undercooling required to reach the extrapolated α-liquidus, and

hence the cooling rate required to suppress properitectic formation, increases with solute con-

centration. There is consequently a critical cooling rate above which β phase nucleation is com-

pletely suppressed in favor of the α solid-solution, with suppression of the properitectic β phase

becoming more difficult with increasing solute concentration. Nucleation of the peritectic phase

α therefore occurs independently of the peritectic reaction, and crystallization of α can occur

above the peritectic temperature Tp.2

2This elevation of the reaction temperature is exactly opposite to that observed for eutectic systems, where thetemperature of thermal arrest is depressed as the cooling rate of the melt is increased.


Fig. 2.2: Free energy curves for the equilibrium phases atthe Al-rich end of the Al-Ti system above the peritectic tem-perature Tp. The relative driving forces for the equilibrium(nucleation of properitectic β) and nonequilibrium (nucle-ation of peritectic α) possibilities are illustrated for compo-sition C0. From Kerr et al. [170].

Fig. 2.3: Equilibrium Al-Ti phase diagram with metastableextrapolations constructed from microprobe analyses.From Kerr et al. [170].

2.4 Intermediate Cooling Rates: Metastable Properitectic Phases

As discussed, for any peritectic alloy the amount of properitectic precipitated in the melt de-

creases with increasing solidification rate and/or decreasing solute content; for sufficiently rapid

cooling, the formation of the properitectic Al3M phase may be suppressed completely, result-

ing in the solidification of the peritectic α phase directly from the melt. At intermediate cool-

ing rates, however, nucleation of the peritectic α phase may be preceded by precipitation of a

metastable properitectic phase, β′. As discussed below, these metastable properitectic phases

(with the L12 structure) are the most desirable form of primary Al3M for grain refinement of

commercial Al alloys, and consequently understanding the morphology, structure, and forma-

tion mechanism of these metastable phases as a function of cooling rate has received consider-

able attention in the literature (e.g., Figures 2.9).

2.4.1 Compositionally metastable properitectic phases

As discussed above, the direct nucleation of supersaturated α-Al is favored at fast cooling rates

because of the significant composition fluctuation required to precipitate the equilibrium tri-

SECTION 2.4 METASTABLE PROPERITECTIC PHASES 43

aluminide properitectic β (Figure 2.2). For the same reason, metastable sub-stoichiometric

(Al-rich) properitectic phases may be favored under intermediate cooling rates since the com-

position fluctuation required to form these AlxM phases is less than that of the stoichiometric

trialuminide (Al3M) properitectic. This phenomenon has been studied most extensively in Al-

Ti alloys, where the occurrence of Al3Ti [71, 160, 211, 212], Al17Ti4, Al15Ti4 [210], Al9Ti [213],

and AlxTi [163] metastable properitectic phases have been reported. Recent convincing evi-

dence [214, 215] suggests that the Ti concentration in these metastable primary phases is ap-

proximately 20 at.%, corresponding to Al4Ti, which is also in reasonable agreement with that

reported originally by Kerr et al. [170, 210].

2.4.2 Structurally metastable properitectic phases

In addition to a possible sub-stoichiometric composition, intermediate cooling rates can also

lead to structurally metastable properitectic phases. For the Al-based peritectic systems consid-

ered (i.e., transition elements of the Groups 4–6 alloyed with Al), the primary β phase formed

under slow cooling of the melt is of equilibrium tetragonal (D022 or D023) crystal structure, ex-

hibiting a coarse plate- or needle-like shape in cross-section, typically 10–100 µm in length. At

smaller solute concentrations and higher freezing rates these properitectic phases become finer

and more equiaxed [80,210,211,216], eventually exhibiting a morphology which is generally de-

scribed as ‘cuboidal’ or ‘petal-like’ (Figure 2.4). Attending this change in morphology is a change

in crystal structure from the equilibrium tetragonal to a metastable cubic L12 structure, which

was first identified by Ohashi and Ichikawa investigating Al-Zr and Al-Ti alloys [160, 217].

The internal structure of these metastable L12 properitectic phases has been disputed. Sev-

eral investigations on metastable L12 Al3Ti [71,160,219], Al3Zr [160,216,220], and Al3Hf [86,88]

described the primary precipitates as two-phase aggregates, consisting of fine discontinuous

lamellae of Al3M (L12) in their respective solid-solutions. Majumdar et al. [72, 214], however,

showed that the petal-like Al3Ti primary precipitates were homogenous single phase. These

differences in morphology have been attributed to breakdown of a stable growth front of Al3M

during solidification [215].

Coarser metastable L12 cuboidal primary precipitates exhibit prominent extensions along

〈111〉 body diagonals, Figure 2.5, which has been reported for primary Al3Ti (L12) [71, 72, 214],

Al3Zr (L12) [216,220,221], and Al3Hf (L12) [88,90]. An explanation for this preferred 〈111〉 growth


Fig. 2.4: Effect of solidification rate and temperature of the melt on the morphology of primary Al3Zr precipitates formed in ahyperperitectic Al-Zr alloy (0.45 at.% Zr). As solidification rate increases, the primary phase becomes more equiaxed. (a) Solidifica-tion rate of 2 × 102°C s-1, cooled from 1100°C. (b) Solidification rate of 2 × 104°C s-1, cooled from 1100°C. (c) Solidification rate of2 × 104°C s-1, cooled from 1300°C. The efficacy of the grain refinement from the cuboidal particles is readily apparent. In additionto solidification rate, overheating the melt above the liquidus also affects the morphology of the primary phase, which is related tothe degree of primary Al3Zr dissolution in the melt. From Brodova et al. [218].

is that {111} planes of any L12 Al3M phase have Al and M atoms in the ratio 1:1, while the {200}

planes have only Al atoms. Thus, for these properitectic L12 phases to grow in 〈100〉 directions,

two kinds of atomic planes must be placed one upon the other. Growth in the 〈111〉 direc-

tion, however, simply consists of sequential deposition of the same {111} planes consisting of

equiatomic fractions of Al and M atoms. It is argued [216, 220, 221] that the latter may be favor-

able kinetically. As these primary phases coarsen this preferred 〈111〉 growth commonly involves

the development of secondary branching and an extension into dendritic growth along the cube

diagonals. Therefore, the morphology of these cuboidal particles is often described as dendritic,

with four main branches generating from a solute-rich center [222, 223].

2.4.3 Grain refinement

As discussed previously (Section 1.2.4), primary Al3M precipitates are efficient heterogeneous

nucleants of α-Al during solidification, and additions of the Groups 4–6 elements (particularly

SECTION 2.4 METASTABLE PROPERITECTIC PHASES 45

Fig. 2.5: Transmission electron micrographs showing the 3-D morphology of metastable cubic L12 Al3Ti primary (properitectic)precipitates. (a) [111]β′ , (b) [111]β′ , and (c) [001]β′ orientations. From Majumdar et al. [214].

Ti) are routinely used as grain refiners in cast Al alloys. The grain size decreases as solute con-

tent increases, with significant grain refinement occurring when the solubility limit (for a given

cooling rate) is exceeded and the Al3M phase precipitates primarily. This grain refinement phe-

nomenon is especially well documented in the Group 4 systems (Al-Ti [160, 210, 219, 224], Al-

Zr [80, 160, 216, 220, 222], and Al-Hf [86–88, 223]), and is shown in Figures 2.6 and 2.7 for Al-Ti

and Al-Zr alloys, respectively, solidified under various cooling rates.

The most effective form of the properitectic Al3M phase for grain refinement is the metastable

L12 structure [80, 87, 216, 220, 223] (see Figure 2.4(c), in which each grain contains a petal-like

Al3Zr primary precipitate), whose cubic structure is commensurate with- and exhibits a small

lattice parameter mismatch to fcc α-Al. The lattice parameters are 4.041 A for Al3Ti [72,160,214,

225], 4.073 A for Al3Zr [217], and 4.051 A for Al3Hf [223]. The metastable cubic primary precip-

itates are oriented with the cube-on-cube relationship with their respective surrounding α-Al

solid-solutions, (100)β′‖(100)α and [001]β′‖[001]α [72, 160, 213, 214, 217, 222, 226], indicating that

α-Al grows epitaxially from the Al3M primary phase.3

3This is demonstrated also in Figure 5.18.


Fig. 2.6: Relation between the grain sizes and Ti concentration inAl-Ti alloys solidified at various cooling rates. From Ohashi andIchikawa [160].

Fig. 2.7: Relation between the grain sizes and Zr concentra-tion in Al-Zr alloys solidified at various cooling rates. FromOhashi and Ichikawa [160].

Grain refinement in eutectic systems

The mechanism of grain refinement just described is not unique to peritectic alloys. Al-Sc alloys,

which are eutectic, are also known for their potent grain refinement in cast alloys [43, 227–231].

As with the peritectic alloys, grain refinement relies on the primary precipitation of the Al3Sc

phase in the melt. Considering the schematic eutectic phase diagram in Figure 1.8(a), however,

primary Al3Sc precipitation may occur only when the alloy composition is hypereutectic. In

a peritectic alloy, Figure 1.8(b), primary precipitation of Al3M occurs generally at much lower

solute concentrations. This, along with the much higher costs of Sc, explains why peritectic

systems are favored industrially as grain refining inoculants to Al alloys. However, unlike in the

peritectic alloys, in the Al-Sc system the L12 intermetallic compound is an equilibrium phase

and grain refinement will consequently occur at any cooling rate.


Tα’

Tp ’

A l

Tm

Tp

Te

mp

era

ture

C o m p o s i t i o n ( % M )

C0

Tβ ’

Tβ

Fig. 2.8: Schematic equilibrium (full lines) and non-equilibrium (broken and dotted lines) peritectic phase diagram. Adapted fromNes et al. [232].

2.5 Chapter Summary

2.5.1 Mechanisms of α formation

During the equilibrium solidification of a peritectic alloy, properitectic β (Al3M) is the first solid

phase to form. At the peritectic temperature Tp this phase and the liquid undergo a peritectic

reaction (L + β → α) to complete the solidification process. This ’textbook’ definition of the

peritectic reaction is rarely realized in practice as it is stifled by the formation of an envelope

of α solid-solution encapsulating the primary β dendrites. The three-phase contact required

for equilibrium is no longer maintained, and further transformation must occur by diffusion

through this peritectic α layer. This mechanism, denoted as the peritectic transformation, does

not occur to any significant extent in the Al-Ti, Al-Zr, or Al-Hf systems and so virtually all of the

α solid-solution is formed by solidification directly from the melt (with or without prior properi-


Fig. 2.9: Effects of cooling rate and Zr concentration on the solidification microstructures of the Al-Zr alloys investigated by Hori etal. [216, 220]. Open circles indicate α solid-solution (no observed primary phase); half-filled circles indicate angular primary phase(composed of fine dendritic metastable L12 Al3Zr) + α solid-solution; filled circles indicate needle-like primary phase (equilibriumD022 Al3Zr) + α solid-solution. From Hori et al. [216, 220].

tectic β precipitation).

2.5.2 Possible solidification sequences as a function of cooling rate

Consider a hyperperitectic alloy of initial composition C0, as shown in Figure 2.8. Under slow

cooling from the melt the equilibrium primary β phase will be nucleated at temperatures below

Tβ , resulting in coarse plate- or needle-like primary precipitates of equilibrium Al3M surrounded

by a (solute-depleted) α-Al solid-solution. Increased cooling rates can suppress the formation of

the equilibrium phase and at temperatures below Tβ′ the melt will also be supersaturated with

respect to the metastable Al3M L12 phase. Although this phase has a higher free energy than

the equilibrium β, its formation may be favored due to a lower surface energy with the melt.

By a further increase in cooling rate the formation of metastable β′ may also be suppressed,

causing the melt eventually, at temperatures below Tα′ , to be supersaturated with respect to α-Al


S o l i d u s

CS C

L C

L C

S

So l i d u s

Te

mp

era

ture


L i q u i d u s

T

k0 < 1 k

0 > 1

L iq u

i du s

T


( a ) ( b )

Fig. 2.10: Solidus-liquidus relationships for hypothetical dilute binary alloys where the solutes either lower (k0 < 1) or raise k0 > 1the melting point of the alloy relative to the pure solvent.

solid-solution. This latter scenario results in direct precipitation of α-Al, which is most desirable

for solid-state precipitation during post-solidification aging, since no solute is lost to primary

phases and the solid-solution is metastably supersaturated.

In real alloys, the situation is as depicted in Figure 2.9, which shows the interrelation of cool-

ing rate and solute content on the observed properitectic morphology and microstructure of

Al-Zr alloys [216,220]4. To a specific cooling rate corresponds a critical solute concentration below

which primary precipitation of properitectic (whether equilibrium β or metastable β′) will not

occur.

2.5.3 Solute redistribution during α-Al nucleation

Solidification is a first-order phase transformation, characterized by a discontinuous change in

equilibrium atom fractions of solute in the liquid and solid phases; the composition of the solid

necessarily differs from that of the liquid in equilibrium with it, as conveyed quantitatively by

the solidus and liquidus phase boundaries [233, 234]. When direct nucleation of α-Al occurs

in peritectic alloys, the solidification follows the dotted lines of the metastable extrapolated α

liquidus and solidus extensions, Figure 2.8.

The difference in composition at the growing solid-liquid interface, assuming that local equi-

4Similar studies have been performed on Al-Ti alloys and other Al-Zr alloys, as discussed in Figure 3.9


librium exists5, can be described by the equilibrium distribution (or partition) coefficient [175–

177]:

k0 =CS

CL

which is simply the ratio of the concentrations of the solid and the liquid specified by the solidus

and liquidus, respectively (Figure 2.10). As discussed in Section 1.2.4, k0 may be expressed in

terms of the solid and liquid compositions at the eutectic or peritectic reaction temperature,

Table 1.5.

It is of interest to compare the solute redistribution that occurs during solidification of α-Al

solid-solution in eutectic (k0 < 1) and peritectic (k0 > 1) systems. On paper, the two situations

are no different. The starting condition is exactly the same, with the initial solid being of compo-

sition k0C0 in both cases. During solidification, solvent (k0 < 1) or solute (k0 > 1) is rejected. As

long as the partition coefficient k0 is constant, some liquid will remain until an invariant tem-

perature is reached. For eutectic systems this the eutectic point, where the remainder of the

liquid then solidifies at eutectic composition (Figures 2.11(a) and (b)); for peritectic systems this

is the melting point of pure Al, resulting in precipitate-free interdendritic regions during post-

solidification aging (Figures 2.11(c) and (d)). As discussed in subsequent chapters, this segre-

gation of solutes influences the distribution, morphology, and composition of the precipitates

formed during post-solidification aging. It is also important to recognize that the nonuniform

distributions of precipitates displayed in Figure 2.11(c) and 2.11(d) are not unprecedented; sim-

ilar dendritic precipitate distributions have been observed during the decomposition of super-

saturated Al-Ti [70, 72, 236], Al-Zr [76, 77, 80], and Al-Hf [86–88, 90, 223] solid-solutions — all

terminal peritectic systems with α-Al.

5This is generally a valid assumption for solid/liquid interface [175–177].


Fig. 2.11: Consequences of eutectic (k0 < 1) and peritectic (k0 > 1) solidification on the resulting solute distribution in cast Alalloys. Panels (a) and (b) are optical micrographs from ternary Al-Mg-Si eutectic alloys cooled at ca. 5°C s-1, from [235]. Panel (c)is an SEM micrograph showing relief contrast between precipitate-rich and precipitate-free regions in an electropolished Al-Zr-Tialloy aged at 425°C. Panel (d) is a dark-field TEM micrograph, showing the dendritic distribution of nanonscale Al3(Zr, Ti) (L12)precipitates in an Al-Zr-Ti alloy aged at 375°C.

CHAPTER 3

Nucleation and Precipitation Strengthening inDilute Al-Ti and Al-Zr Alloys

Two conventionally-solidified Al-Ti alloys (0.18 and 0.22 at.%Ti) exhibit no hardening

after aging up to 3,200 h at 375 or 425°C. This is due to lack of precipitation of Al3Ti,

as confirmed by electron microscopy and electrical conductivity measurements. This

behavior is in contrast to that of an Al-0.19Zr (at.%) alloy, which displays strong age

hardening at both temperatures due to precipitation of Al3Zr (L12). These observations

are interpreted within the context of the equilibrium phase diagrams. The disparity in

solid and liquid solubilities of Ti in α-Al is much greater than that of Zr in α-Al; the rel-

atively small liquid-solubility of Ti in α-Al limits the amount of solute retained in solid-

solution during solidification, while the comparatively high solid-solubility reduces the

supersaturation effecting precipitation during post-solidification aging. The lattice pa-

rameter mismatch of Al3Ti (L12) with α-Al is also larger than that of Al3Zr (L12), further

hindering nucleation of Al3Ti. Classical nucleation theory indicates that the minimum

solute supersaturation required to overcome the elastic strain energy accompanying ho-

mogeneous nucleation cannot be obtained during conventional solidification of Al-Ti

alloys (unlike for Al-Zr alloys), thus explaining the absence of Al3Ti precipitation and

presence of Al3Zr precipitation.

3.1 Introduction

AS DISCUSSED IN CHAPTER 1, the Group 4 transition metals (Ti, Zr, Hf) constitute a group

of alloying additions to Al that show particular promise for developing creep-resistant,

thermally-stable Al-based alloys. In each of these systems, an ordered Al3M (where M = Ti,

Zr, or Hf) trialuminide may be precipitated from supersaturated solid-solution during post-

soldification aging. While the equilibrium structure of these Al3M trialuminides is tetragonal

(D022 for Al3Ti and Al3Hf1 and D023 for Al3Zr), decomposition of supersaturated Al-M solid-

solutions occurs initially by the formation of nanometer-scale metastable cubic L12 Al3M pre-1Al3Hf exists in two different crystallographic structures: a stable high temperature D023 phase and a stable low

temperature D022 phase [85]. The D022 structure is the one relevant to solid-state precipitation, while D023 is therelevant structure when considering primary precipitation from the liquid.

53

54 CHAPTER 3 NUCLEATION AND PRECIPITATION STRENGTHENING IN DILUTE AL-TI AND AL-ZR ALLOYS

cipitates exhibiting small lattice parameter mismatches with α-Al, which transform to their re-

spective equilibrium structures after prolonged exposure (ca. 102–103 h) to elevated tempera-

tures (ca. 450°C). Moreover, these transition elements are anomalously slow diffusers in α-Al,

with very limited equilibrium solid-solubilities, enabling precipitated Al3M to be resistant to

Ostwald ripening in accordance with volume diffusion-controlled coarsening theory.

For high-temperature applications, Ti seems especially promising among the Group 4 el-

ements since it has the smallest atomic weight and is also the slowest diffuser in α-Al (Sec-

tion 1.2.3). Indeed, Al alloys reinforced by Al3Ti dispersions have garnered considerable in-

terest as potential lightweight high-temperature structural materials, and the thermal stabil-

ity and strength at high temperatures of alloys precipitation-strengthened with Al3Ti is well-

documented [73, 237–258]. These studies, however, investigated alloys prepared by rapid solid-

ification processing (RSP) [73, 237–243] or mechanical alloying (MA) [237, 238, 244–258] tech-

niques, whose nonequilibrium processing routes circumvent the difficulties encountered dur-

ing conventional solidification of Al-Ti alloys.

The problems with casting these alloys arise from the fact that dilute additions of Ti or Zr (as

well as Hf) to Al exhibit peritectic phase equilibria with the terminal α-Al solid-solution. The

first solid to form under equilibrium conditions is the properitectic, or primary, Al3M (M = Ti,

Zr, Hf) ordered phase (for alloys of peritectic composition, i.e., those enriched beyond the min-

imum liquid-solubility of solute). These solute-rich primary phases readily nucleate and grow

into coarse (ca. 100 µm) precipitates during conventional casting, leaving the remaining melt,

and ultimately the solidified α-Al solid-solution, substantially depleted in solute. This limits the

potential for precipitation strengthening since the supersaturation — and therefore the chem-

ical driving force for solid-state nucleation, as well as the equilibrium volume fraction of the

dispersed phase (if formed) — is significantly reduced.

The equilibrium structure of the primary precipitates is tetragonal (D022 for Al3Ti and D023

for Al3Zr and Al3Hf1), but under certain conditions primary Al3M may precipitate also with the

metastable cubic L12 structure [160]. These metastable primary precipitates are isostructural to,

and exhibit a small lattice parameter mismatch with, the α-Al solid-solution, and therefore act as

efficient heterogeneous nucleants during solidification of α-Al. The resulting grain refinement

in cast alloys is a well known and frequently exploited phenomenon industrially for Ti additions

to Al [155–157,160], but has also been observed in Al-Zr [80,160,216,220,222] (and Al-Hf [86–88,

SECTION 3.2 EXPERIMENTAL PROCEDURES 55

Table 3.1: Designations and verified compositions (at.%) of the Al-Ti and Al-Zr alloys investigated.

Designation Bodycote ATI Wah Chang

Al-0.18Ti 0.175 0.178

Al-0.22Ti 0.237 0.211

Al-0.19Zr — 0.186

223]) alloys whenever primary precipitation of Al3M (L12) occurs.

Both consequences of properitectic precipitation — solute depletion from primary Al3M

precipitation or potent grain refinement from metastable L12 Al3M precipitation — are unde-

sirable for high-strength high-temperature applications since the reduced solute concentration

retained in solid-solution leads to smaller volume fractions of precipitated strengthening phases

formed during aging, and the possibility of grain refinement is also deleterious for creep resis-

tance since a refined grain structure promotes rapid diffusional creep.

This study compares the feasibility of developing conventionally-cast and aged, precipitation-

strengthened, dilute Al-based alloys alloyed with the Group 4 elements Ti or Zr. Most studies

related to conventionally-solidified Al-Ti alloys have focused on the thermal stability of the rel-

atively coarse primary Al3Ti phase formed during solidification [212, 259–262]. As noted in ref-

erences [263,264], the scientific literature pertaining to decomposition or age hardening of Al-Ti

solid-solutions produced by conventional ingot metallurgy processing is rather limited. The pri-

mary objective, therefore, is to report and discuss the precipitation behavior in Al-Ti alloys, with

the data for Al-Zr provided primarily as a comparison. In Chapter 5 the precipitation behav-

ior and stability of Al3Zr precipitated from conventionally-solidified Al-Zr alloys is discussed in

detail.

3.2 Experimental Procedures

3.2.1 Alloy compositions and preparation

Two Al-Ti alloys (Al-0.18Ti and Al-0.22Ti) and one Al-Zr alloy (Al-0.19Zr) were investigated (com-

positions indicated in Table 3.1). Small (ca. 7 g) buttons of these alloys were prepared by arc-

melting in a gettered purified argon atmosphere using a non-consumable tungsten electrode


and a water-cooled copper cathode as the crucible. The charges were melted a minimum of

ten times, and inverted between melts to ensure alloy homogeneity. The moderately enhanced

cooling rate associated with the chilled crucible was estimated to be of the order of 10–100°C s-1,

comparable to continuous casting.2 For dilute Al-Ti and Al-Zr alloys, as well as other peritec-

tic systems, to a specific cooling rate corresponds a critical solute concentration below which

primary precipitation of Al3M (M = Ti or Zr) does not occur [170–172]. The compositions in Ta-

ble 3.1 were expected, based on solidification studies (Figure 3.9, discussed below), to be near

the threshold for avoiding primary precipitation of Al3M under the moderate cooling rates in-

vestigated.

Both Al-Ti alloys were prepared by melting nominally3 99.99 at.% Al (Atlantic Equipment

Engineers, Bergenfield, NJ) with pure Al3Ti master alloy. A first advantage of this approach is

that the melting point of Al3Ti (1350°C) is lower than that of pure Ti (1688°C). Additionally, this

ordered structure is a line compound whose composition and constitution is entirely homoge-

neous. This is not true of commercial master alloys (containing typically 5–10 wt.% Ti) that ex-

hibit coarse (ca. 100 µm) primary Al3Ti precipitates. For the dilute alloys in this investigation and

the small mass of the button ingots, such inhomogeneous master alloys introduce unacceptable

uncertainties in chemical composition. The monolithic Al3Ti master alloy was prepared by arc-

melting amounts of pure constituents (99.99% Al, Atlantic Equipment Engineers; 99.99+% Ti,

Alfa Aesar, Ward Hill, MA) corresponding to the stoichiometry of the Al3Ti phase; chemical ho-

mogeneity and structural uniformity of the button ingot was verified with X-ray diffraction.

The Al-0.19Zr alloy was prepared from a dilute master alloy containing 1.9 wt.% Zr, which was

dilution cast from a commercial 10 wt.% Zr master alloy (KB Alloys, Reading, PA). The chemical

compositions of both Al-Ti alloys were verified independently by Bodycote Materials Testing

(Skokie, IL) and ATI Wah Chang (Albany, OR), while the composition of Al-0.19Zr was checked

by ATI Wah Chang only.

3.2.2 Aging treatments and analytical techniques

The as-cast alloys were isothermally aged at 375 or 425°C, within the range of temperatures

reported to exhibit a strong age hardening response for Al-Ti alloys produced by RSP [73, 74,

2The solidification rate of the arc-melter is discussed briefly in Appendix A.3See Appendix E.

SECTION 3.3 EXPERIMENTAL RESULTS 57

241, 265, 266] or MA [246] and Al-Zr alloys produced by RSP [14, 264, 267–270] or chill-casting

[187, 190, 271]. Precipitation of Al3Ti or Al3Zr during aging was monitored by a number of tech-

niques, described below.

First, Vickers microhardness measurements were performed at room temperature on met-

allographically polished sections using a load of 200 g and a dwell time of 5 s. Second, pre-

cipitation of Al3Ti or Al3Zr was assessed directly by transmission electron microscopy (TEM).

TEM foils were prepared by mechanical grinding sections of aged specimens to a thickness of

ca. 100 µm. Discs of 3 mm diameter were punched from these sections and thinned to perfo-

ration by twin-jet electropolishing at 20 V d.c. (Struers TenuPol-5) using a 10 vol.% solution of

perchloric acid in methanol at -40°C. Electropolished TEM foils were also examined using a LEO

1525 high-resolution field-emission gun scanning electron microscope (SEM) operated at 3 kV

with a short working distance (3 mm) using a secondary electron in-lens detector. As discussed

below, the alloys are highly segregated on the micrometer scale, which is not easily discerned in

the TEM. The advantage of SEM as a complement to TEM is the much larger field-of-view from

both the larger range in magnification and also since observation of specimens in the SEM are

not subject to the limitation that the foil be electron transparent. Furthermore, the resolution of

the SEM, ca. 1 nm, is capable of resolving nanometer-scale Al3Ti or Al3Zr precipitates.

Finally, for the Al-Ti alloys, which did not exhibit detectable age hardening or evidence for

precipitation of Al3Ti by electron microscopy, electrical conductivity measurements were also

performed to monitor the decomposition of the supersaturated solid-solutions during extended

aging at 425°C. The conductivity measurements were performed using a SIGMATEST 2.069 (Fo-

erster Instruments, Pittsburgh, PA) eddy current apparatus at room temperature. Five measure-

ments were recorded, each corresponding to a different frequency (60, 120, 240, 480, 960 kHz),

on each specimen. For consistency, a single specimen of each alloy (aged for various times) was

used for conductivity measurements.

3.3 Experimental Results

3.3.1 Optical microscopy

Figure 3.1 shows the as-cast macrostructure of the alloys studied. The solidification macrostruc-

ture is typical of cast alloys, with coarse columnar grains, originating at the bottom surface of the


Fig. 3.1: Macrostructure of as-cast alloys Al-0.18Ti, Al-0.22Ti, and Al-0.19Zr (etched using Poultan’s reagent), showing various de-grees of grain refinement.

ingot (which was in contact with the chilled copper crucible of the arc-melter), growing upward

toward a zone of equiaxed grains at the center of the ingot. The relative sizes of the columnar and

equiaxed zones is strongly dependent on the solute content of the alloy and, more precisely, the

extent of properitectic Al3M (M = Ti or Zr) precipitation, since these primary phases are potent

grain refiners, as discussed above.

Alloy Al-0.18Ti is the most dilute of those studied (Table 3.1), and the extent of the columnar

zone, which comprises nearly half of the ingot cross-section in this alloy, is also greatest. Within

the equiaxed zone, the grains are relatively coarse, with grain sizes ranging from 0.5–1.0 mm

across. The effect of a small increase in Ti concentration in these conventionally-solidified al-

loys is demonstrated for Al-0.22Ti, which exhibits a dramatic refinement in grain size with a fine

columnar zone that adjoins an even finer equiaxed region of small (50–100 µm) grains in the

upper half of the ingot. Moderate grain refinement is exhibited also for Al-0.19Zr, which is inter-

mediate between that of alloy Al-0.18Ti and alloy Al-0.22Ti. Nearly the entire ingot cross-section

of Al-0.19Zr is equiaxed, with grain sizes ranging from 0.15–0.30 mm in size.

The pronounced grain refinement observed in Figure 3.1 for Al-0.22Ti is due to copious

precipitation of ‘petal-like’ primary Al3Ti, which were observed in metallographically polished


0 . 1 1 1 0 1 0 0 1 0 0 0

1 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

2 0

3 0

4 0

5 0

6 0

7 0

8 0

A l - 0 . 2 2 T i

A l - 0 . 1 8 T i

��A l - 0 . 1 9 Z r

1 d a y 1 2 4 8 1 6 w e e k s

Vic

ke

rs

mic

ro

ha

rd

ne

ss

(M

Pa

)

�� A s - c a s t V

ick

ers

mic

ro

ha

rd

ne

ss

(H

V2

00)

Fig. 3.2: Vickers microhardness vs. aging time at 375°C forthe Al-Ti and Al-Zr alloys investigated.

0 . 1 1 1 0 1 0 0 1 0 0 0

1 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

2 0

3 0

4 0

5 0

6 0

7 0

8 0

A l - 0 . 2 2 T i

A l - 0 . 1 8 T i

��

A l - 0 . 1 9 Z r

1 d a y 1 2 4 8 1 6 w e e k s

Vic

ke

rs

mic

ro

ha

rd

ne

ss

(M

Pa

)

�� A s - c a s t

Vic

ke

rs

mic

ro

ha

rd

ne

ss

(H

V2

00)

Fig. 3.3: Vickers microhardness vs. aging time at 425°C forthe Al-Ti and Al-Zr alloys investigated.

specimens by SEM. This ‘petal-like’ morphology is characteristic of the metastable L12 Al3M

phase [160], whose cubic structure is commensurate with fcc α-Al and acts as an effective het-

erogeneous nucleant of α-Al during solidification. The presence of these primary precipitates

indicates also that, for the conventional casting conditions used in this study, exceeding ca. 0.2

at.% solute (Ti or Zr) results in primary precipitation of Al3M and hence does not increase the

amount of solute retained in solid-solution.

3.3.2 Age hardening

Figures 3.2 and 3.3 display the observed microhardness of the Al-Ti and Al-Zr alloys after isother-

mal aging at 375 and 425°C, respectively. Each data point represents a minimum of 20 measure-

ments, with the standard deviation of these measurements indicated by the error bars. At both

temperatures, the Al-Ti alloys exhibit a negligible age hardening response after extended aging

times up to 3,200 h (19 weeks). This lack of strengthening contrasts with the significant precipita-

tion hardening response of the Zr-containing alloy, which commences at times as short as 0.1 h

(6 min) at 425°C, and achieves peak strength within 25 h at both 375 and 425°C. As described

below, the pronounced strengthening is due to precipitation of small (< 10 nm) coherent Al3Zr

(L12) precipitates; no evidence for precipitation of a similar Al3Ti phase is observed. The as-

cast hardness values of Al-0.22Ti and Al-0.19Zr are both 242 MPa, approximately 20 MPa greater

than that of Al-0.18Ti. This moderate strength increase is probably attributable to solid-solution


Fig. 3.4: SEM micrographs of electropolished TEM foils. (a) As-cast Al-0.22Ti, showing preferential etching around the solute-enriched dendrites. (b) Al-0.22Ti aged at 425°C for 1,600 h, showing no evidence of precipitation of Al3Ti within the solute-enricheddendrites. (c) Al-0.19Zr aged at 425°C for 400 h, showing similar dendritic distributions of solute. (d) Magnified view of Al3Zrprecipitate-rich dendritic regions and precipitate-free interdendritic channels.

strengthening from dissolved Ti or Zr in the as-cast alloys [14, 272]. While there is a refinement

in the as-cast grain size with increasing solute concentration (Figure 3.1), Hall-Petch strength-

ening is not applicable because the scale of the microhardness indent (ca. 125 µm)4 relative to

the grain sizes in Figure 3.1.


Fig. 3.5: Centered superlattice dark-field TEM micrographs of Al-0.19Zr aged at 425°C for 400 h. Images obtained under dynamicaltwo-beam conditions, the specimen oriented near the [110] zone axis and imaged with the (110) L12 superlattice reflection asindicated in the inset diffraction pattern. (a) Precipitate-rich dendritic regions separated by a precipitate-free interdendritic channelsimilar to those in Figures 3.4(c) and 3.4(d). (b) Small, coherent, homogeneously-distributed Al3Zr (L12) precipitates within thehighly-supersaturated dendrite. The precipitates in this region have a mean radius of 〈R〉 = 6.7 ± 1.7 nm.

3.3.3 Electron microscopies

The origin of the precipitation hardening response indicated in Figures 3.2 and 3.3 was investi-

gated directly for alloys Al-0.22Ti and Al-0.19Zr using SEM, Figure 3.4. The dendritic distribution

of solute species formed during solidification in both alloys is revealed by preferential etching

of the foil surface during electropolishing. Ryum [85] observed similar resistance to electropol-

ishing in the dendritic cells of as-cast specimens of Al-Hf alloys. Figure 3.4(a) shows the as-cast

structure for Al-0.22Ti, where the relief contrast around the dendrite arms is due to variations

in solute concentration in solid-solution between the dendritic and interdendritic regions of the

alloy. The center of the dendrites are enriched in solute [70, 170, 210, 273], and are apparently

more resistant to electropolishing. Figure 3.4(b) shows the same Al-0.22Ti alloy after extended

aging (1,600 h) at 425°C. The dendritic cells are delineated as before, but no evidence for precipi-

tation of Al3Ti within the supersaturated cells is observed, consistent also with the negligible age

hardening of the Al-Ti alloys demonstrated in Figures 3.2 and 3.3. Analyses of the same speci-

men by TEM, both in real space (strain contrast) and reciprocal space (selected area electron

diffraction), also provided no evidence for Al3Ti precipitates, or any other precipitated phase.

4See also Figure C.3 in Appendix C.


Figures 3.4(c) and 3.4(d) show similar SEM micrographs for Al-0.19Zr aged 400 h at 425°C.

As with the Ti-containing alloys, dendrites are clearly delineated from the electropolishing pro-

cess. Unlike the Al-Ti alloys, however, the solute-rich dendrites decompose during aging to form

regions with a high number density of homogeneously-distributed, nanometer-scale Al3Zr pre-

cipitates. Figure 3.5 shows superlattice dark-field TEM micrographs of the same specimen, ex-

hibiting again precipitate-rich and precipitate-free regions associated with the initial dendritic

Zr distribution. Within the highly-supersaturated dendrites, Figure 3.5(b), the Al3Zr precipitates

have the metastable cubic L12 structure, are coherent with the α-Al solid-solution, and are small

(〈R〉 = 6.7 ± 1.7 nm). It is those precipitates that are responsible for the marked precipitation

hardening response of Al-0.19Zr presented in Figures 3.2 and 3.3.

3.3.4 Electrical conductivity

The electrical conductivity of a disordered binary substitutional alloy is a parabolic function of

the absolute solute concentration [274], and is therefore a sensitive means of monitoring the de-

composition of a supersaturated solid-solution during aging. Electrical conductivity measure-

ments were performed on both Al-Ti alloys (Al-0.18Ti and Al-0.22Ti) during extended aging (up

to 1,600 h) at 425°C. No detectable change in conductivity was observed, indicating that solute

atoms remained dissolved in supersaturated solid-solution and did not precipitate as Al3Ti, cor-

roborating the lack of precipitation observed by hardness and electron microscopies. As shown

in Chapter 5, the decomposition of more dilute concentrations than those studied presently are

easily detected by electrical conductivity. Moreover, Royset and Ryum [275] recently monitored

in detail the decomposition kinetics of dilute Al-0.12Sc (at.%) alloys using a similar eddy current

apparatus.

3.4 Discussion

The hardness data in Figures 3.2 and 3.3 demonstrate that a conventionally-solidified Al-0.19Zr

alloy exhibits pronounced precipitation hardening at 375 and 425°C, unlike Al-Ti alloys with sim-

ilar solute concentrations. The strong age hardening response of the Zr-containing alloy is due to

precipitation of small (< 10 nm) coherent Al3Zr (L12) precipitates, shown by SEM (Figures 3.4(c)

and 3.4(d)) and TEM (Figure 3.5). No evidence for precipitation of a similar Al3Ti phase is ob-


served in the Al-Ti alloys by SEM or TEM. Moreover, there is no appreciable change in electrical

conductivity after aging Al-0.18Ti or Al-0.22Ti at 425°C for 1,600 h, which indicates that the Ti

concentration in α-Al solid-solution is unchanged, confirming the lack of precipitation of Al3Ti.

3.4.1 Diffusion kinetics

The measured tracer diffusivity of Ti in α-Al is smaller than that of Zr in α-Al, and it is conceivable

that sluggish diffusion might account for the lack of Al3Ti precipitation. The diffusivities, D, of

Ti and Zr in α-Al are given by an Arrhenius relationship:

D = D0 exp (−Q/RgT ) , (3.1)

where Rg is the ideal gas constant, T is the absolute temperature, Q is the activation enthalpy for

solute diffusion, and D0 is the pre-exponential factor. With values Q = 260 and 242 kJ mol-1 and

D0 = 1.12 × 10−1 and 7.28 × 10−2 m2 s-1 for Ti and Zr, respectively (Table 1.3), Figure 3.6 shows

the calculated root-mean-squared (RMS) diffusion distances,√

6Dt (where t is the aging time),

for Ti and Zr in α-Al at 375 and 425°C; the diffusivity of Ti atoms at 425°C is comparable to that

of Zr atoms at 375°C.

The hardness curve of Figure 3.2 demonstrates that the Al-Zr binary alloy reaches near-peak

hardness by 6.25 h of aging at 375°C, which corresponds to an RMS diffusion distance of 18 nm

for Zr atoms. For the slower-diffusing Ti atoms, a similar RMS diffusion distance can be achieved

after 115 h at 375°C or 4 h at 425°C, indicating that Ti atoms diffuse fast enough to support the

nucleation and growth of Al3Ti precipitates. Atom-probe tomography studies on Al-Sc-Ti alloys

aged at 300°C [118] and Al-Zr-Ti alloys aged at 375 and 425°C (Chapter 4) demonstrate that Ti

segregates to the precipitated Al3(Sc1−xTix) and Al3(Zr1−xTix) precipitates, further evidence that

the limited mobility of Ti is not the limiting factor explaining the unobserved precipitation of

Al3Ti.

3.4.2 Classical nucleation theory

The inability to nucleate Al3Ti is a result of an insufficient supersaturation of Ti in solid-solution

following conventional solidification. Figures 3.7 and 3.8 display dilute equilibrium binary phase

diagrams for the Al-Ti and Al-Zr systems based on experimental data. From a casting stand-


1 1 0 1 0 0 1 0 0 0

1 0

1 0 0

1 0 0 0

6 . 2 5 2 5 1 0 0 4 0 0 1 6 0 0 3 2 0 0

��

��

��

��

��

�

Dif

fus

ion

dis

tan

ce

(n

m)

A g i n g t i m e ( h )

C a l c u l a t e d r o o t - m e a n - s q u a r e d

d i f f u s i o n d i s t a n c e , ( 6 D �t ) 1 / 2

Fig. 3.6: Calculated root-mean-squared diffusion distances,√

6Dt, for Ti and Zr in α-Al at 375 and 425°C to t = 3,200 h, comparableto the aging conditions in the hardness curves of Figures 3.2 and 3.3.

point, a suitable alloying element with Al should produce a low liquidus temperature or, equiv-

alently, exhibit high solubility in the liquid. This liquidus criterion facilitates melting into the

single-phase liquid and also reduces the tendency for primary phase precipitation during so-

lidification. For precipitation strengthening, limited solid-solubility is required to maximize the

chemical driving force for nucleation as well as the volume fraction of the precipitated phase.

From Figures 3.7 and 3.8, the disparity in the liquid and solid-solubilities of Ti in α-Al is much

greater than that of Zr.

The equilibrium phase boundaries in Figures 3.7 and 3.8 correspond to the tetragonal-structured

trialuminides: D022 for Al3Ti and D023 for Al3Zr. The solubility limits of Ti and Zr in α-Al, in equi-

librium with their respective metastable L12 structure Al3M trialuminides, have been evaluated

in recent ab initio calculations by Liu and Asta [276, 277]:

CAl3Tiα = 2.376 · exp (−4397.9/T )

CAl3Zrα = 10.523 · exp (−8285.9/T )

(3.2)

where CAl3Tiα and CAl3Zr

α are the solubilities (expressed in atomic fraction) of Ti and Zr in equilib-


0 0 . 2 0 . 4 0 . 6 0 . 8 1 . 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

M e t a s t a b l e L 1 2 S o l v u s

0 . 0 7 9 a t . %

( A l ) + A l

3T i

L + A l3T 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 . %

C o n c e n t r a t i o n s o f T i s t u d i e d

Fig. 3.7: Equilibrium Al-rich Al-Ti binary phase dia-gram (adapted from Murray [122]). Metastable L12 Al3Tisolvus calculated by Liu and Asta [276, 277].

0 0 . 2 0 . 4 0 . 6 0 . 8 1 . 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

9 0 0

Met a

st a

bl e

L1

2 S

ol v

us

6 6 0 . 8 ° C

( A l )

0 . 0 3 3 a t . %

( A l ) + A l3Z r

L + A l3Z r

��

��

��

�

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

L

0 . 0 8 3 a t . %

C o n c e n t r a t i o n o f Z r s t u d i e d

0 . 1 4 7 a t . %

Fig. 3.8: Equilibrium Al-rich Al-Zr binary phase dia-gram (adapted from Murray [123]). Metastable L12 Al3Zrsolvus calculated by Liu and Asta [276, 277].

rium with their respective metastable L12 precipitates. These calculated solvus boundaries are

indicated by the dotted lines in Figures 3.7 and 3.8.

According to classical nucleation theory [278–282], the steady-state nucleation current (nu-

cleation rate per unit volume) may be written:

J ∝ exp(−∆F ∗

RgT

)exp

(− Q

RgT

), (3.3)

where ∆F ∗ is the net reversible work required for the formation of a critical nucleus, given by:

∆F ∗ =16πσ3

3(∆Fch −∆Fel)2. (3.4)

The steady-state nucleation current, Eq.(3.3), is thus a function of the chemical driving force,

∆Fch; the reduction of this driving force by the elastic strain energy, ∆Fel; and the activation

enthalpy for diffusion, Q. As discussed below, these three parameters hinder the nucleation of

Al3Ti compared to Al3Zr.


Chemical driving force, ∆Fch, for nucleation of Al3M

The metastable L12 solvus boundaries in Figures 3.7 and 3.8 demonstrate that the solid-solubility

of Ti in α-Al is significantly greater than that of Zr, which dictates the chemical driving force for

homogenous nucleation of metastable L12 Al3M (M = Ti or Zr). Assuming ideal solution behav-

ior (valid for the dilute concentrations discussed), this driving force is proportional to the natural

logarithm of the supersaturation ratio, C0/Cα, where C0 is the concentration (atomic fraction)

of M dissolved in solid-solution and Cα is the solid-solubility of M in α-Al given by Eqs.(3.2).

Following Doherty [283], the chemical driving force per unit volume driving precipitation of the

ordered Al3M phase may be estimated as:

∆Fch = − RgT

VAl3M

(CAl3M − Cα

1− Cα

)ln

C0

Cα, (3.5)

where CAl3M = 0.25 is the atomic fraction of solute in the ordered trialuminide, and VAl3M =

Na (aAl3M)3 /4 (Na is Avogadro’s number) is the molar volume of the precipitated phase.

Figures 3.7 and 3.8 indicate that, for a given solute concentration (C0), the supersaturation

driving precipitation of Al3Zr is significantly greater than that for Al3Ti, which partially explains

the unobserved precipitation of Al3Ti because of the reduced supersaturation ratio C0/Cα en-

tering Eq.(3.5). The solute concentration C0 is, however, not constant throughout the alloys,

because of the segregation visible in the SEM micrographs of Figure 3.4. For both Al-Ti and Al-

Zr, the dendrites are supersaturated with respect to the overall bulk alloy composition, C0, and

so in both cases the local supersaturation of solute can be greater than the bulk compositions

indicated in Table 3.1 and Figures 3.7 and 3.8.

The degree of this segregation was not measured directly (e.g., by electron microprobe anal-

ysis, EPMA), but may be estimated from the equilibrium solid-liquid partition coefficient, k0 =

CS/CL, which quantifies the difference in composition of the solid and liquid phases in local

equilibrium during solidification [175–177]. Making the usual assumption that both the liq-

uidus and solidus boundaries are straight lines, k0 may be expressed in terms of the solid and

liquid compositions at the peritectic temperature. From Figures 3.7 and 3.8, k0 = 10 for Ti and

k0 = 2.5 for Zr. Therefore, the first solid to form is predicted to be enriched by a factor of 10

and 2.5 for Ti and Zr, respectively. This is in reasonable agreement with results of Kerr and col-

leagues [170,210], who measured the concentration of Ti across the dendritic cells in Al-Ti alloys


containing up to 0.28 at.% Ti using EPMA, and observed an enrichment in Ti of 5–6 times in the

dendrite centers. Setiukov and Fridlyander [273] similarly observed that the Ti concentration in

the central zones of the dendritic cells exceeded its average value by 6–8 times in alloys contain-

ing 0.02–0.17 at.% Ti. All present Al-Ti and Al-Zr alloys are therefore locally supersaturated well

beyond their respective equilibrium solubilities of solute, and so an insufficient supersatura-

tion of solute, ∆Fch, alone cannot account for the absence of Al3Ti precipitates. As indicated in

Eq.(3.4), the influence of elastic strain energy, ∆Fel, must also be considered since this reduces

the net driving force for nucleation.

Reduction of the driving force by the elastic strain energy, ∆Fel

Solid-state reactions involve strain energy since precipitating a second phase, which is fully or

partially coherent with the matrix, requires the straining of the lattices in both the matrix and

precipitated phases. This strain energy, ∆Fel, diminishes the net driving force, Eq.(3.4), and

hence the rate of nucleation, Eq.(3.3). Considering only dilatational strains and assuming elastic

isotropy, the strain energy per unit volume for a coherent inclusion may be written [279, 280]:

∆Fel = 2µ

(1 + ν

1− ν

)δ2, (3.6)

where µ = 25.4 GPa [27] and ν = 0.345 [26] are the shear modulus and Poisson’s ratio of Al at

room temperature.5 The average lattice parameter mismatch at room temperature, δ, of both

the tetragonal and metastable cubic (L12) phases are significantly greater for Al3Ti (δ = 2.04%

and 5.36% for L12 and D022, respectively) than for Al3Zr (δ = 0.75% and 2.89% for L12 and D023,

respectively), Table 1.1. Thus, the elastic strain energy impeding nucleation of Al3Ti (L12) is over

7 times greater than that of Al3Zr (L12).

Critical solute concentration, C0, to overcome ∆Fel

Compared to Al3Zr (L12), nucleation of Al3Ti (L12) is hindered by: (i) a reduced chemical driving

force (∆Fch) for a given C0; and (ii) a greater elastic strain energy opposing nucleation (∆Fel). To

5Eq.(3.6) makes the simplifying assumption that the elastic constants of the Al3M precipitate and α-Al matrix arecomparable; Barnett et al. [284] developed a similar expression that considers separately the elastic constants of bothphases, which has been used to estimate ∆Fel in nucleation studies on the similar Al3Sc system [285,286]. Neverthe-less, the error in considering only the properties the α-Al matrix is minor and Eq.(3.6) is of sufficient accuracy for ourpurposes here.


Table 3.2: Calculation of chemical driving force, ∆Fch, effecting nucleation of Al3Ti (L12) and Al3Zr(L12).

VAl3M Cα at 425°C (Eq.(3.2)) CAl3M ∆Fch (Eq.(3.5))(m3mol−1) (atomic fraction) (atomic fraction) (MJ m-3)

Al3Ti (L12) 9.40×10−6 0.00436 0.25 152.3 ln (229.4 · C0)

Al3Zr (L12) 1.03×10−5 0.00007 0.25 141.9 ln (13594.0 · C0)

explain why nucleation of Al3Zr (L12) is observed and that of Al3Ti (L12) is not, consider the

relative magnitudes of ∆Fch and ∆Fel entering Eq.(3.4) in both systems. The chemical driving

force, ∆Fch (Eq.(3.5)), is a strong function of C0 and is calculated using the parameters in Ta-

ble 3.2. Equating ∆Fch and ∆Fel gives a critical solute concentration necessary to overcome the

elastic strain energy barrier associated with Al3Ti (∆Fel = 45 MJ m-3) and Al3Zr (∆Fel = 6 MJ m-3)

nucleation. For Al-Ti alloys, the critical value of C0 is 0.580 at.% Ti and for Al-Zr it is 0.008 at.%

Zr.

Table 3.1 indicates that the bulk compositions of both of the Al-Ti alloys studied are less than

the critical threshold of 0.580 at.% Ti whereas Al-0.19Zr is supersaturated well beyond 0.008 at.%

Zr, thus explaining the observed precipitation of Al3Zr (L12) and the lack thereof for Al3Ti. This

explanation is actually not strictly complete because, as discussed, the alloys are segregated and

hence locally supersaturated beyond the bulk compositions in Table 3.1. Nevertheless, it is the

combination of a small chemical driving force, ∆Fch, coupled with a large elastic strain energy

barrier, ∆Fel, in the Al-Ti system, as compared to the Al-Zr system, that explains their disparate

precipitation behavior.

Activation enthalpy for solute diffusion, Q

Equation(3.3) indicates that the nucleation current is strongly dependent on temperature. The

solute supersaturation (and, correspondingly, ∆Fch) increases with decreasing temperature

(Eqs.(3.2)) while the solute diffusivity D diminishes rapidly (Eq.(3.1)). These competing influ-

ences result in an intermediate temperature for optimum nucleation kinetics [282]. As discussed

in Section 3.4.1, the activation enthalpy for solute diffusion, Q, for Ti in α-Al (260 kJ mol-1) is

greater than that of Zr (242 kJ mol-1), which tends to increase the temperature for maximum nu-

cleation current for Al3Ti. As discussed, ∆Fch driving precipitation of Al3Ti is significantly less


0 0 . 2 5 0 . 5 0 0 . 7 5 11 0

0

1 01

1 02

1 03

1 04

1 05

T i

Z r

( a )

( d )

( b )

� �

��

��

��

T i o r Z r c o n c e n t r a t i o n ( a t . % )

( a )

( c )

Fig. 3.9: Critical cooling rate for suppressing primary Al3Ti (solid lines) or Al3Zr (dashed lines) and achieving supersaturated α-Alsolid-solution, as reported by: (a) Ohashi and Ichikawa [160]; (b) Kerr et al. [170]; (c) Hori et al. [71,219]; and (d) Hori et al. [216,220].

than that for Al3Zr for a given C0 (Figures 3.7 and 3.8). This reduced driving force is exacerbated

by the higher Q for Ti diffusion since at the temperatures where Ti is mobile, the supersaturation

is not large enough to effect nucleation of Al3Ti.

3.4.3 Limitations on increasing the solute concentration

The discussion in Section 3.4.2 indicates that an increased concentration of Ti (C0) beyond those

studied presently is required for precipitation of Al3Ti. This, however, is not possible under

conventional casting conditions. As demonstrated in Figure 3.1, the marked grain refinement

observed with increasing Ti concentration is due to primary precipitation of Al3Ti in the melt,

and increasing the alloy composition does not increase the amount of Ti retained in α-Al solid-

solution. The interrelation between solidification rate, solute concentration, and solidified mi-

crostructure of the Al-Ti and Al-Zr binary systems has been the subject of a number of stud-

ies [71, 160, 170, 216, 219, 220], which are summarized in Figure 3.9. These curves show the mea-


sured solidification rate, as a function of solute concentration, necessary to suppress primary

precipitation of Al3Ti or Al3Zr, thereby obtaining supersaturated α-Al solid-solution. Figure 3.9

indicates that in order to achieve a concentration of 0.580 at.% Ti supersaturated in α-Al solid-

solution, required to just offset the elastic strain energy associated with nucleation of Al3Ti(L12),

accelerated cooling rates of 8×103 °C s-1 (based on the data of Hori et al. [71,219]) are required. To

achieve a greater driving force for nucleation higher cooling rates are required still. Nucleation of

Al3Zr (L12) requires a small supersaturation to overcome the elastic strain energy barrier, which

is easily achieved for even the slowest cooling rates in Figure 3.9.

3.4.4 Comparison with previous studies

Decomposition behavior of Al-Ti solid-solutions

In contrast to the closely-related Al-Zr and Al-Hf systems, the precipitation behavior of Al3Ti in

Al-Ti alloys has been debated in the scientific literature. As with Al-Zr and Al-Hf alloys, the de-

composition sequence of supersaturated Al-Ti solid-solutions has been reported to occur firstly

by the precipitation of a metastable Al3Ti (L12) phase, which then transforms to the equilibrium

tetragonal (D022) structure after prolonged aging (Chapter 1). Asboll and Ryum [236], however,

observed only precipitation of the equilibrium D022 phase in chill cast Al-0.34Ti (at.%) alloys

aged between 400–550°C. They also noted that the decomposition kinetics of Al3Ti were slug-

gish compared to that of Al alloyed with other elements of the same subgroup, Zr or Hf. Pandey

and Suryanarayana [93] similarly compared the decomposition behavior of rapidly-solidified Al

alloyed with the Group 4 elements (Ti, Zr, Hf) and noted that, while a metastable L12 Al3M or-

dered phase is observed in Al-Zr and Al-Hf alloys, this is not the case for Al-Ti alloys. This result

is in accord with other studies [243,262,269] reporting only precipitation of the equilibrium D022

phase in Al-Ti alloys aged in the broad temperature range of 200–600°C. Park and Kim [269], how-

ever, reported that the equilibrium D022 structure was preceded by another metastable tetrag-

onal Al3Ti phase in Al-1.13Ti (at.%) alloys aged at 400°C. Asboll and Ryum [236] also observed

tetragonal Al3Ti precipitates that did not have the D022 structure, in addition to those of the

equilibrium D022 phase.

Other studies report no precipitation during aging of supersaturated Al-Ti solid-solutions.

Ohashi et al. [287] found no evidence for precipitation of Al3Ti, using microhardness and elec-

trical resistivity measurements, in Al-Ti alloys containing up to 0.45 at.% Ti produced by RSP


during 100 h of aging in the range of 400–640°C. St. John et al. [261] similarly observed no solid-

state precipitation of Al3Ti when aging an Al-1.1Ti (at.%) alloy for 24 h at 435°C; their alloy did,

however, contain coarse primary Al3Ti precipitates, indicating the amount of Ti initially in solid-

solution was reduced.

Nucleation of coherent Al3Ti (L12) precipitates has been reported only for highly supersatu-

rated alloys produced by nonequilibrium means (RSP). Ohashi and Ichikawa [70] observed Al3Ti

(L12) precipitates in alloys containing 0.3–1.4 at.% Ti aged at 400–500°C. Muddle and cowork-

ers [72–75] similarly showed precipitation of coherent Al3Ti (L12) in melt-spun alloys containing

2.9–3.5 at.% Ti aged at 300–500°C. Similar supersaturations cannot be achieved by conventional

casting, Figures 3.1 and 3.9, since Ti precipitates out in the liquid as properitectic Al3Ti at mod-

erate cooling rates.

Age hardening in the Al-Ti and Al-Zr systems

As with the decomposition behavior from solid-solution, the age hardening response accom-

panying Al3Ti precipitation is similarly disputed. Dobatkin et al. [268] observed that, among

several transition metals alloyed with Al, Al-Ti alloys exhibit one of the weakest hardening re-

sponses in isochronal (2 h) aging experiments. This lack of hardening, however, was not due to

the absence of precipitation since decomposition of the supersaturated solid-solution, as mon-

itored by electrical resistivity, was commenced upon aging at ca. 300°C. In similar isochronal

aging experiments, Park and Kim [269] also confirmed precipitation of Al3Ti (D022) by X-ray

diffraction and TEM in an Al-1.13Ti (at.%) aged at 400°C, but observed no accompanying in-

crease in microhardness. Precipitation of Al3Ti with no apparent increase in strength was shown

more conclusively by Asboll and Ryum [236], who aged a chill-cast Al-0.34Ti (at.%) alloy at tem-

peratures between 400–550°C. No increase in hardness was observed after up to 24 days, but

precipitation of irregularly-shaped lath Al3Ti (D022) precipitates was confirmed by TEM to oc-

cur within the highly supersaturated dendritic cells of the alloy. The precipitates were relatively

coarse (200–500 nm in length) and exhibited several, non-specific, orientation relationships with

the surrounding α-Al solid-solution; other studies [72,260] have also reported an ill-defined ori-

entation relationship between Al3Ti (D022) and α-Al. As discussed, both the metastable L12

and equilibrium D022 structures of Al3Ti exhibit a large lattice parameter mismatch with α-Al.

This precludes the formation of a coherent Al3Ti strengthening phase, and may explain the


reported negligible age hardening even when precipitation of Al3Ti is confirmed. The ques-

tionable age hardening behavior in Al-Ti alloys contrasts significantly with the otherwise similar

Al-Zr system, in which a strong age hardening effect, even in conventionally cast alloys, is well-

documented [78, 185–187, 267].

3.4.5 Common industrial uses of Ti and Zr additions to Al

Perhaps the most illustrative comparison of Al-Ti and Al-Zr alloys is found by considering their

primary industrial roles. As discussed previously, minor additions of Ti are used to refine the

grain structure in commercial aluminum alloys, where primary Al3Ti precipitates act as hetero-

geneous nuclei during solidification of the melt [155–157, 160]. A solid-state counterpart to this

grain refinement effect is realized when dilute additions of Zr are added to commercial wrought

alloys as recrystallization inhibitors [78, 168, 188–191]. The precipitation of fine (20–30 nm) dis-

persions of coherent Al3Zr (L12) precipitates, formed during post-solidification aging, are effec-

tive barriers to grain boundary migration. Titanium is thus added for the purpose of forming

primary precipitates in the melt whereas Zr additions form coherent dispersoids precipitated

from supersaturated solid-solution. This suggests that: (i) Zr, compared to Ti, is easier to retain

in solid-solution during conventional solidification; and (ii) Al3Zr compared to Al3Ti is more

readily precipitated during post-solidification aging. These considerations indicate that the Al-

Zr system is a better candidate than the Al-Ti system for developing a castable precipitation-

strengthened alloy.

3.5 Chapter Summary

This investigation has compared the decomposition of supersaturated, conventionally-cast Al-

0.2Ti and Al-0.2Zr (at.%) alloys, upon aging at 375 or 425°C. The Zr-containing alloy exhibits a

strong age hardening response at both temperatures, due to copious precipitation of coherent,

nanometer-scale Al3Zr (L12) precipitates within the supersaturated dendritic cells of the alloy.

Precipitation of a similar Al3Ti phase is unobserved during extended aging times (up to 3,200 h

at 425°C). Compared to Al3Zr, nucleation of Al3Ti (L12) is hindered by: (i) a reduced chemical

driving force (∆Fch) driving nucleation (for a given alloy composition, C0); (ii) a larger elastic

strain energy opposing nucleation (∆Fel); and (iii) a greater activation enthalpy for solute diffu-


sion (Q). The minimum supersaturation to effect precipitation of Al3Ti cannot be achieved by

conventional solidification of binary Al-Ti alloys, since increasing the Ti concentration causes

precipitation of primary Al3Ti, and thus does not contribute to a supersaturation of solute in

solid-solution. The Al-Zr system is more promising system for developing a castable alloy capa-

ble of precipitation hardening.

CHAPTER 4

Atom-Probe Tomographic Studies ofPrecipitation in Al-0.1Zr-0.1Ti (at.%) Alloys

Atom-probe tomography was utilized to measure directly the chemical compositions of

Al3(Zr1−xTix) precipitates with metastable L12 structure formed in Al-0.1Zr-0.1Ti (at.%)

alloys upon aging at 375 or 425°C. At these aging temperatures, the Zr:Ti atomic ratio

in the precipitates is about 10 and 5, respectively, indicating that Ti remains mostly in

solid-solution rather than partitioning to the Al3(Zr1−xTix) precipitates. This is inter-

preted as being due to the very small diffusivity of Ti in α-Al, consistent with prior stud-

ies on Al-Sc-Zr and Al-Sc-Ti alloys where the slower-diffusing Zr and Ti atoms make up

a small fraction of the Al3(Sc1−x, Zr/Tix) precipitates. Unlike those alloys, however, the

present Al-Zr-Ti alloys exhibit no interfacial segregation of Ti at the matrix/precipitate

heterophase interface, a result that may be affected by a significant disparity in evap-

oration fields of the α-Al matrix and Al3(Zr1−xTix) precipitates, and/or a lack of local

equilibrium at the interface. The analyses are complicated by an initially heterogeneous

distribution of solute atoms after casting, resulting in an inhomogeneous dendritic dis-

tribution of Al3(Zr1−xTix) precipitates after aging.

4.1 Introduction

IT WAS SHOWN IN CHAPTERS 1 AND 3 that the Al-Zr system shows particular promise for devel-

oping creep-resistant, thermally-stable Al-based alloys. It is well established that decom-

position of supersaturated Al-Zr solid-solutions occurs initially by the formation of nanometer-

scale Al3Zr precipitates with a metastable cubic L12 structure, which transform to the equilib-

rium D023 phase after prolonged aging at elevated temperatures (> 450°C). The high stability of

the L12 metastable phase is attributed to slow diffusion kinetics and a small lattice parameter

mismatch with α-Al (Chapter 1).

Fine and coworkers showed that this lattice parameter mismatch is reduced by additions of

Ti, V, or Hf [84,110,111], resulting in improved coarsening resistance of L12 Al3(Zr1−xVx) [84,112–

114] and Al3(Zr1−xTix) [288] precipitates, which they ascribed to a reduced matrix/precipitate

75

76 CHAPTER 4 ATOM-PROBE TOMOGRAPHIC STUDIES OF PRECIPITATION IN AL-0.1ZR-0.1TI (AT.%) ALLOYS

interfacial free energy. These findings motivated subsequent studies on Al-Zr-V [270,289], Al-Zr-

Ti [264, 271, 290–296], and Al-Zr-Ti-V [289, 297, 298] alloys, in which the precipitated ternary and

quaternary trialuminides were reported to offer improved coarsening resistance and thermal

stability compared to binary Al3Zr precipitates [84].1 This improved stability was attributed to

a reduced lattice parameter mismatch, as suggested by Fine, but in none of these studies (with

the exception of reference [271]) was the composition of the precipitates (and hence the lattice

parameter mismatch) measured directly.

Three-dimensional atom-probe tomography (3-D APT) [299–303] is based on the field-

evaporation of surface atoms, as ions in different charge states, that are identified chemically

by time-of-flight mass spectrometry and whose positions are deduced from the coordinates of

ion impacts on a 2-D position-sensitive detector. Small, nanometer-scale concentration fluctu-

ations are examined on the sub-nanometer-scale in real space, without data deconvolution, and

without comparison to standards. In this respect, 3-D APT is uniquely suited to examine directly

the effects of microalloying in nanoscale precipitation processes [304, 305].

In this chapter, the compositions of nanometer-scale Al3(Zr1−xTix) precipitates formed upon

aging at 375 or 425°C for various aging times are measured directly by 3-D APT. The effect of the

precipitate compositions on the microstructures and mechanical properties of these alloys is

discussed in Chapters 5 and 7.


Two ternary Al-0.1Zr-0.1Ti (all compositions are given in at.%) alloys were investigated, desig-

nated Al-0.1Zr-0.1Ti(a) and Al-0.1Zr-0.1Ti(b). The alloys were prepared by melting nominally2

99.99 at.% Al (Atlantic Equipment Engineers, Bergenfield, NJ) with pure Al3Ti master alloy and

a dilute Al-0.57Zr (at.%) master alloy employing non-consumable electrode arc-melting in a

gettered purified argon atmosphere. The monolithic Al3Ti master alloy was prepared by arc-

melting amounts of the pure constituents — 99.99% Al (Atlantic Equipment Engineers, Bergen-

field, NJ); 99.99+% Ti (Alfa Aesar, Ward Hill, MA — to obtain the stoichiometric Al3Ti phase;

the chemical homogeneity and structural uniformity of the Al3Ti button ingot was verified with

X-ray diffraction. The Al-0.57Zr master alloy was dilution cast from a commercial 10 wt.% Zr

1These studies are discussed in detail in Chapter 5.2See Appendix E.


Table 4.1: Compositions and aging conditions of the Al-0.1Zr-0.1Ti alloys investigated by 3-D APT.

Nominal comp. (at.%) Verified comp. (at.%) Aging conditionsZr Ti Zr Ti Temperature (°C) Time (h)

Al-0.1Zr-0.1Ti(a) 0.1 0.1 — — 375 20, 100, 400, 1600

Al-0.1Zr-0.1Ti(b) 0.1 0.1 0.089 0.099 425 400

master alloy provided by KB Alloys (Reading, PA). The chemical composition of alloy Al-0.1Zr-

0.1Ti(b) was confirmed by bulk chemical analysis performed by direct current plasma emission

spectroscopy (ATI Wah Chang, Albany, OR). As-cast specimens were isothermally aged at 375 or

425°C, within the range of temperatures shown to produce a strong precipitation strengthening

response for Al-Zr alloys (Chapter 5). The alloy designations, compositions, and aging condi-

tions investigated by 3-D APT are summarized in Table 4.1.

Needle-like 3-D APT specimens were prepared by a two-step electropolishing procedure.

Specimen blanks, excised from aged specimens and mechanically ground to approximately

0.2×0.2×10 mm3, were initially shaped into long needle-like tips with a slight taper angle us-

ing a solution of 10 vol.% perchloric acid in acetic acid at ca. 10 V d.c. at room temperature. Final

electropolishing involved formation of a neck near the tip apex, controlled by carefully limit-

ing the amount of chemical solution in contact with the area of the neck using a loop-polishing

apparatus. For this final tip preparation procedure, a solution of 2 vol.% perchloric acid in bu-

toxyethanol at room temperature with an applied potential of 3–8 V d.c. was employed. The

resulting specimen is a sharply-pointed tip with an end radius of curvature less than 50 nm.

A new class of atom probe tomographs, the local electrode atom-probe (LEAP®) [306–308],

was used for the majority of 3-D APT analyses. Pulsed field-evaporation was conducted under

ultrahigh vacuum (UHV) conditions (<10−10 Torr gauge pressure) at a specimen temperature

of 30 K utilizing a pulse fraction (ratio of the pulse voltage to the steady-state d.c. imaging volt-

age) of 15–20% and a pulse frequency of 200 kHz. A fixed flight path of 80 mm was used for

all analyses. Some 3-D APT analyses were also performed using an energy-compensated optical

position-sensitive atom-probe [309–311] (referred to in the text as conventional 3-D APT), with a

pulse fraction of 20%, 1.5 kHz pulse frequency, and a specimen temperature of 30 K in a vacuum

of <10−9 Torr gauge pressure.

Post-analysis data visualization and evaluation were performed with IVAS (Imago Scientific


Fig. 4.1: Centered superlattice dark-field TEM micrographs of alloy Al-0.1Zr-0.1Ti(a) aged at 375°C for 1,600 h. (a) Low mag-nification image showing dendritic distribution of Al3(Zr1−xTix) precipitates. (b) Magnified view of small (3–5 nm radius), co-herent, homogeneously-distributed L12 Al3(Zr1−xTix) precipitates within the highly-supersaturated dendrites. A representative50×50×200 nm3 3-D APT analysis volume is indicated for comparative purposes.

Instruments, Madison, WI) and custom software developed at Northwestern University called

ADAM or currently APEX [312]. Quantitative precipitate compositions and solute composition

profiles were obtained using the proximity histogram method [313, 314], measuring solute con-

centration profiles in the α-Al matrix and Al3(Zr1−xTix) precipitate phases with respect to a 5

at.% Zr isoconcentration surface delineating the two phases.


4.3.1 Solute segregation and the precipitated microstructure

Figure 4.1 shows the precipitated microstructure of the alloys, obtained after aging alloy Al-

0.1Zr-0.1Ti(a) for 1,600 h at 375°C. The nanometer-scale Al3(Zr1−xTix) (L12) precipitates are non-

uniformly distributed throughout the alloy, reflecting the microsegregation of Zr and Ti solute

atoms in the as-cast alloys. Both solute species, when alloyed with Al, raise the melting point

of the alloy relative to that of pure Al. Consequently, the liquidus and solidus boundaries of

the α-Al solid-solution have positive slopes and k0, the equilibrium partition coefficient for so-

lidification, is greater than unity for both systems. The first solid to form during solidification

is therefore richer in Zr and Ti compared to the bulk alloy composition, resulting in solute-


0 5 1 0 1 5 2 0 2 5 3 0 3 5

0

0 . 1

0 . 2

0 . 3

0

0 . 1

0 . 2

0 . 3

0 . 4

C0

C0

Z r

T i

Co

mp

os

itio

n

(at.

%)

D i s t a n c e ( µm )

Fig. 4.2: Concentration profiles of Zr and Ti, measured by EDS, across three solute-rich dendrite arms in as-cast Al-0.1Zr-0.1Ti(b).Both Zr and Ti segregate to the dendrite centers (k0>1).

rich dendritic cells surrounded by solute-depleted interdendritic channels. Upon aging, only

the enriched dendritic cells contain enough solute in solid-solution to cause precipitation of

Al3(Zr1−xTix).

The initial distribution of solute was measured in as-cast Al-0.1Zr-0.1Ti(b) by quantitative

energy dispersive X-ray spectroscopy (EDS) using a JEOL JSM-7000F scanning electron micro-

scope (SEM) equipped with a ThermoNORAN EDS detector. The composition profile spanning

three dendrites across 35 µm, Figure 4.2, demonstrates that the dendrite centers are enriched

nearly twofold in Zr and threefold in Ti, compared to their respective bulk compositions (Ta-

ble 4.1). The interdendritic regions are similarly virtually solute-free.

The initial inhomogeneous supersaturation of solute atoms after solidification influences

not only the spatial distribution of precipitates, but also the precipitate size and morphology, as

well as the nucleation mechanism. Consider the precipitates labeled A and B in Figure 4.1(b).

Within the center of the dendrites (A), where the Zr and Ti supersaturations are greatest, the

precipitates are small (3–5 nm radius), coherent, and are homogeneously distributed. Other

precipitates outside the dendrites (B) are larger (ca. 25 nm radius) and are inhomogeneously


distributed, indicating that they are most likely heterogeneously nucleated. The nucleation and

growth of these precipitates during aging at 375, 400, or 425°C is described in Chapter 5.

A typical 50×50×200 nm3 analysis volume for 3-D APT is displayed in Figure 4.1(b). As

compared to the relatively coarse (micrometer-scale) dendritic distribution of solute and pre-

cipitates, the 3-D volume analyzed by 3-D APT is exceedingly small and therefore the locally-

measured solute composition of the alloy, as well as the likelihood of intersecting precipitates,

varies among 3-D APT analyses. Indeed, it is demonstrated below that the solute concentration

measured during an analysis is correlated to the probability of intersecting a precipitate. This

also means the local alloy composition from which the Al3(Zr1−xTix) precipitates formed is a

priori unknown.

4.3.2 Three-dimensional atom-probe tomographic reconstructions

Figure 4.3 presents a 3-D reconstruction of an Al3(Zr1−xTix) precipitate intersected during a

LEAP tomographic analysis. The reconstructed volume is 39×39×29 nm3, containing 7.69× 105

identified atoms represented by dots. Figure 4.3(a) displays the Al atoms, with only 25% shown

for the sake of clarity. The paucity of atoms in the precipitate interior and the shell of atomic

enrichment at the matrix/precipitate interface are reconstruction artifacts, characteristic of a

strong local magnification effect arising from the disparate evaporation fields of the α-Al matrix

and Al3(Zr1−xTix) precipitates. This artifact, and its possible influence on the measured precip-

itate compositions, is discussed below. Figure 4.3(b) shows only the solute (Zr and Ti) atoms,

indicating that most of the solute atoms partition to the Al3(Zr1−xTix) phase. While the precip-

itates are Zr-rich (i.e., primarily Al3Zr as discussed below), Figure 4.3(c) shows unambiguously

that Ti atoms partition to the precipitate, forming a ternary Al3(Zr1−xTix) intermetallic.

4.3.3 Quantitative chemical composition of the Al3(Zr1−xTix) precipitates

The reconstructions in Figure 4.3 indicate that both Zr and Ti partition to the Al3(Zr1−xTix) pre-

cipitates. This information is conveyed more quantitatively in Figure 4.4, which exhibits a prox-

imity histogram, or proxigram for short [313, 314], displaying quantitatively the composition of

the Al3(Zr1−xTix) precipitate depicted in Figure 4.3. Similar to a linear composition profile, the

proxigram is a plot of concentration as a function of distance, indicating the partitioning be-

havior of solute atoms between the α-Al matrix and the Al3(Zr1−xTix) precipitates. Rather than


Fig. 4.3: LEAP reconstruction of an Al3(Zr1−xTix) precipitate in alloy Al-0.1Zr-0.1Ti(b) aged for 400 h at 425°C. (a) Reconstructionshowing only Al atoms, with only 25% displayed for the sake of clarity. (b) Reconstruction showing Zr and Ti atoms, indicating thatthe solute atoms partition to the Al3(Zr1−xTix) precipitate. (c) Reconstruction showing only Ti atoms.

using a single linear direction, however, the proxigram considers each point of 3-D space with

respect to a reference surface (a 5 at.% Zr isoconcentration surface), which is physically more

representative of the actual concentration profiles. Hence, a proxigram displays in 2-D informa-

tion that is characteristic of 3-D.

Figure 4.4 indicates that, compared to the nominal value of 25 at.% solute expected for a tri-

aluminide (Al3M-type) ordered phase, the precipitate in Figure 4.3 is sub-stoichiometric, con-

taining approximately 20 at.% total solute. Similar sub-stoichiometric compositions have been

observed in metastable L12 primary (formed in the melt) trialuminides in Al-Ti alloys, with pre-

cipitate compositions (determined by EDS) containing approximately 20 at.% Ti, corresponding

to Al4Ti [170, 210, 214, 215].3

This departure from stoichiometry may be attributable to unpositioned solute ions (Zr and

Ti) field-evaporated from the precipitate phase. Sha and Cerezo [315] showed that it is necessary

to include both correctly-positioned as well as unpositioned ions in order to obtain stoichiomet-

ric compositions of Al3Zr precipitates formed in a commercial 7050 Al alloy. The precipitates

require a much higher electric field for field-evaporation than the α-Al solid-solution (discussed

below), which generates more multiple ionization events during field-evaporation resulting in

higher proportions of these ions not being positioned. Assuming the fraction of unpositioned

ions is independent of solute ion species, the ratio of solute atoms measured in the Al3(Zr1−xTix)

precipitates herein should reflect the correct precipitate composition. Despite the overall sub-

3See also Section 2.4.1.


- 3 - 2 - 1 0 1 2 3 4 5 6 7

0

5

1 0

1 5

2 0

2 5

3 0

3 5

��

� ��

��

�

Zr,

Ti

co

mp

os

itio

n

(at.

%)

D i s t a n c e f r o m t h e i n t e r f a c e ( n m )

α� ��

Fig. 4.4: Proxigram displaying the distribution of Zr andTi atoms relative to the position of a 5 at.% Zr isoconcen-tration surface for the precipitate reconstructed in Fig-ure 4.3 (Al-0.1Zr-0.1Ti(b) aged at 425°C for 400 h). Theprecipitate is Zr-rich, with a small fraction of the Zr lat-tice sites replaced by Ti atoms. The ratio of Zr:Ti is about5, corresponding to Al3(Zr0.83Ti0.17).

- 3 - 2 - 1 0 1 2 3 4 5

0

5

1 0

1 5

2 0

2 5

3 0

3 5

��

��

�

Zr,

Ti

co

mp

os

itio

n

(at.

%)


��!��!��!��

��α��

Fig. 4.5: Proxigrams displaying the distribution of Zr andTi atoms in Al3(Zr1−xTix) precipitates formed in alloyAl-0.1Zr-0.1Ti(a) aged at 375°C for the times indicated.The 20 h aging time results are based on four precipitatesand a total of 4.96 × 105 atoms, the results for 400 h arebased on 10 precipitates and a total of 9.09 × 105 atoms,and the results for 1,600 h are based on four precipitatesand a total of 3.07 × 106 atoms. The ratio of Zr:Ti isabout 10, independent of aging time, corresponding toAl3(Zr0.91Ti0.09).

stoichiometric precipitate composition, Figure 4.4 demonstrates that Ti is a minor constituent

in the Al3(Zr1−xTix) precipitate, with an average Zr:Ti ratio of 5 (corresponding to x = 0.17). This

is significantly greater than the overall Zr:Ti ratio of unity in the alloy (Table 4.1).

Figure 4.5 shows proxigrams for precipitates formed in Al-0.1Zr-0.1Ti(a) aged for 20, 400, or

1,600 h at 375°C. Compared to Figure 4.4, these concentration profiles are closer to stoichiome-

try (25 at.% total solute) and, as with Figure 4.4 (425°C aging temperature), Ti is clearly a minor

constituent in the Al3(Zr1−xTix) precipitates at 375°C. After 20 h at 375°C, the concentration of

Ti in the precipitate is 3–4 at.%, comparable to that displayed in Figure 4.4 for Al-0.1Zr-0.1Ti(b)

aged 400 h at 425°CFigures 4.4 and 4.5 indicate that the composition of the Al3(Zr1−xTix) precip-

itates varies little with aging temperature (425°C, Figure 4.4), or aging time (20–1,600 h at 375°C,

Figure 4.5). At 375°C, the compositions of the Al3(Zr1−xTix) precipitates displayed in Figure 4.5

each have a Zr:Ti ratio of ca. 10, corresponding to x = 0.09.


4.3.4 Time-of-flight mass spectra

Figure 4.6(a) shows the time-of-flight mass spectrum for the entire analysis reconstructed in Fig-

ure 4.3. The prominent Al peaks correspond to singly- (Al1+, 27 amu) and doubly-charged (Al2+,

13.5 amu) ions produced during pulsed field-evaporation; Al has only one isotope. The peaks for

the Zr and Ti solute atoms, while less pronounced (a result which is expected, considering the

compositions in Table 4.1), are nevertheless clearly identifiable. Zirconium is field-evaporated

in both the doubly- (Zr2+, 45–48 amu) and triply-charged (Zr3+, 30–32 amu) states, whereas Ti

atoms are observed only in the doubly-charged (Ti2+, 23–25 amu) state. For all species, the

stable isotopes are clearly distinguished from one another, demonstrating the excellent mass

resolution of the 3-D APT technique.

Compare now the mass spectrum in Figure 4.6(a) to the one displayed in Figure 4.6(b), which

represents an analyzed volume from the α-Al solid-solution below the precipitate displayed in

Figure 4.3. The Zr peaks, which are distinct in Figure 4.6(a), are no longer visible in Figure 4.6(b),

indicating that all detectable Zr atoms partition to the Al3(Zr1−xTix) phase. The Ti peaks, by con-

trast, are still quite prominent in Figure 4.6(b). By counting the number of atoms corresponding

to these peaks, the concentration of Ti in α-Al solid-solution is determined to be 0.109 ± 0.006

at.%, which is somewhat greater than the overall composition of alloy Al-0.1Zr-0.1Ti(b) (Ta-

ble 4.1). Considering the microsegregation of solute demonstrated in Figure 4.2, this result is

not surprising since the as-cast alloys are highly segregated, with some regions significantly en-

riched in solute and other regions depleted. The compositions in Table 4.1 are bulk chemical

compositions whereas those measured by 3-D APT are locally sampled. Furthermore, as shown

in the next section, there is a correlation between measured solute concentration in the α-Al

solid-solution and proximity to precipitates, as anticipated. The mass spectrum in Figure 4.6(c)

represents the ions field-evaporated from the core of the Al3(Zr1−xTix) precipitate displayed in

Figure 4.3; as anticipated, the Zr and Ti peaks from this solute-rich phase are distinct.

There is also a shift in charge states of the field-evaporated ions between the two phases.

The spectrum of field-evaporated ions from the α-Al solid-solution matrix, Figure 4.6(b), shows

that 95% of the Al ions are field-evaporated in the singly-charged state (Al1+), whereas in the

Al3(Zr1−xTix) precipitates, Figure 4.6(c), Al is almost entirely doubly-charged (Al2+). These

higher charge states are indicative of the larger evaporation fields required to field-evaporate

the ions from Al3(Zr1−xTix), leading to a pronounced local magnification effect.


1 5 2 0 2 5 3 0 4 5 5 0

1 01

1 02

1 03

1 04

1 01

1 02

1 03

1 04

1 01

1 02

1 03

1 04

1 05

2 7A l

1 +

2 7A l

2 +

Nu

mb

er o

f e

ve

nts

M a s s - t o - c h a r g e - s t a t e r a t i o ( a m u )

5 0T i

2 +

4 9T i

2 +

4 6T i

2 +

4 7T i

2 +

4 8T i

2 +

9 1Z r

3 +

9 2Z r

3 +

9 4Z r

3 +

9 6Z r

3 +

9 0Z r

3 +

9 4Z r

2 +

9 2Z r

2 +

9 1Z r

2 +

9 0Z r

2 +

( c ) A l3( Z r

1 - x, T i

x) p r e c i p i t a t e

2 7A l

1 +

2 7A l

2 +

4 8T i

2 +

5 0T i

2 +

( b ) α - A l s o l i d s o l u t i o n

( a ) E n t i r e a n a l y s i s

5 0T i

2 +

4 9T i

2 +

4 6T i

2 +

4 7T i

2 +

4 8T i

2 +

9 1Z r

3 +

9 2Z r

3 +

9 4Z r

3 +

9 6Z r

3 +

9 0Z r

3 +

2 7A l

1 +

2 7A l

2 +

9 1Z r

2 +

9 2Z r

2 +

9 0Z r

2 + 9 4Z r

2 +

9 6Z r

2 +

Fig. 4.6: Atom-probe time-of-flight mass spectrum of Al-0.1Zr-0.1Ti(b) aged for 400 h at 425°C (reconstruction shown in Figure 4.3).(a) Mass spectrum for the entire analysis, representing 7.67× 105 identified atoms, of which 5.21× 103 atoms (0.679± 0.009 at.%)are identified as Zr and 2.16 × 103 atoms (0.282 ± 0.006 at.%) are Ti. (b) Mass spectrum indicating the composition of the matrix(α-Al solid-solution). This mass spectrum represents 2.09 × 105 identified atoms, of which 224 atoms (0.107 ± 0.007 at.%) are Ti;there is no evidence of Zr (<0.010 at.%), indicating that all detectable Zr atoms have partitioned to the Al3(Zr1−xTix) precipitate.(c) Mass spectrum indicating the core composition of the Al3(Zr1−xTix) precipitate. This spectrum represents 8.88× 103 atoms, ofwhich 1,508 atoms (17.00 ± 0.40 at.%) are Zr and 335 atoms (3.77 ± 0.20 at.%) are Ti. This corresponds to the plateau compositionin Figure 4.4.

4.3.5 Correlation between measured Ti concentration proximity to precipitates

Considering the initial dendritic distribution of solute species after solidification (Figure 4.2),

and the resulting non-uniform distribution of Al3(Zr1−xTix) precipitates after aging (Figure 4.1),

it is anticipated that 3-D APT analyses would be frequently precipitate-free. Indeed, at least

half of all performed analyses fail to intersect precipitates but rather sample interdendritic

precipitate-free volumes of α-Al solid-solution. These analyses, though precipitate-free, are nev-

ertheless valuable since they provide a direct, atom-by-atom chemical verification of the local


Table 4.2: Number of precipitates intersected by 3-D APT, total number of atoms collected, number of Ti atoms collected, andresulting Ti concentration of the α-Al solid-solution (expressed in terms of the bulk value, C0 = 0.099 at.% Ti).

Alloy Number of precipitates Total number of atoms (×106) Ti atoms C/C0

Al-0.1Zr-0.1Ti(a), 375°C, 100 h 0 0.287 — 0.00± 0.07

Al-0.1Zr-0.1Ti(a), 375°C, 100 h 0 0.814 252 0.31± 0.04

Al-0.1Zr-0.1Ti(a), 375°C, 1,600 h 0 1.272 416 0.33± 0.03

Al-0.1Zr-0.1Ti(a), 375°C, 1,600 h 1 0.455 176 0.39± 0.05

Al-0.1Zr-0.1Ti(a), 375°C, 1,600 h 0 1.300 576 0.44± 0.03

Al-0.1Zr-0.1Ti(b), 425°C, 400 h 0 2.254 1, 369 0.61± 0.04

Al-0.1Zr-0.1Ti(a), 375°C, 20 h 0 0.275 169 0.61± 0.08

Al-0.1Zr-0.1Ti(a), 375°C, 1,600 h 6 3.700 2, 539 0.69± 0.03

Al-0.1Zr-0.1Ti(a), 375°C, 100 h 1 0.632 606 0.96± 0.07

Al-0.1Zr-0.1Ti(a), 375°C, 1,600 h 1 0.547 551 1.01± 0.08

Al-0.1Zr-0.1Ti(a), 375°C, 400 h 1 0.161 164 1.02± 0.12

Al-0.1Zr-0.1Ti(b), 425°C, 400 h 1 0.575 627 1.09± 0.07

Al-0.1Zr-0.1Ti(a), 375°C, 20 h 4 0.496 657 1.32± 0.08

Al-0.1Zr-0.1Ti(a), 375°C, 1,600 h 4 3.068 4, 115 1.34± 0.06

Al-0.1Zr-0.1Ti(a), 375°C, 1,600 h 10 0.909 1, 272 1.40± 0.08

Al-0.1Zr-0.1Ti(b), 425°C, 400 h 1 0.986 2, 120 2.15± 0.10

solute concentration in the particular volume sampled. As shown in Figure 4.1, the spatial dis-

tribution of Al3(Zr1−xTix) precipitates is correlated directly to the initial dendritic distribution

of solute atoms formed during solidification. Figure 4.6 indicates that, during aging, nearly all of

the Zr atoms partition to the precipitates, whereas a substantial fraction of the Ti is observed in

the α-Al solid-solution around the precipitates (in a typical volume analyzed by 3-D APT, ca. 104–

106 nm3). Since Zr and Ti partition in the same manner (Figure 4.2) there should be a correlation

between measured Ti concentration in the α-Al solid-solution and proximity to the precipitate-

rich dendrites in the microstructure. Table 4.2 compares the results of several LEAP and conven-

tional APT analyses for various aging conditions of the Al-0.1Zr-0.1Ti alloys, arranged in order

of increasing measured Ti concentration in the α-Al solid-solution. For all but three analyses

that intersected precipitates, Table 4.2, the Ti concentration measured locally in the α-Al solid-

solution is enriched relative to the overall value of 0.099 at.% Ti (i.e., C/C0 > 1). Similarly, all

analyses that do not detect precipitates have smaller Ti concentrations than the nominal value

(C/C0 < 1), with the α-Al solid-solutions of some analyses containing undetectable amounts


Fig. 4.7: LEAP reconstruction of Al-0.1Zr-0.1Ti(a) aged for 1,600 h at 375°C. The dimensions of the rectangular parallelepipedare 28×27×30 nm3, containing 4.42×105 identified atoms. One large (>25 nm radius) Al3(Zr1−xTix) precipitate is detected. (a)Reconstruction showing only Al atoms, with only 25% of the atoms displayed for clarity. (b) Reconstruction showing Zr and Ti atoms.(c) Reconstruction showing only Ti atoms.

of Ti (C/C0 → 0). There appears to be a critical value, between C/C0 = 0.61 and 0.69 as indi-

cated by the horizontal rule in Table 4.2, which separates the precipitate-free analyses from the

precipitate-containing ones.

There is one outlier at C/C0 = 0.39, Table 4.2, which represents an analysis that intersected

a precipitate, but whose surrounding α-Al solid-solution is depleted in Ti compared to its over-

all concentration in the alloy. A portion of the corresponding 3-D reconstruction is shown in

Figure 4.7, in which a large (R > 25 nm) Al3(Zr1−xTix) precipitate is intersected at the end of

the analysis. The reconstructed precipitate is comparable in radius to precipitate B in the TEM

micrograph in Figure 4.1(b)); the large radius of this precipitate suggests that the analyzed vol-

ume is interdendritic, consistent with the depleted measured Ti concentration in the α-Al solid-

solution (Table 4.2).

This hypothesis is supported by the composition, proxigram in Figure 4.8, of the precipi-

tate reconstructed in Figure 4.7. Not only is the matrix depleted in Ti (C/C0 = 0.39, Table 4.2),

but the measured Ti concentration in the Al3(Zr1−xTix) precipitate is also small (ca. 1 at.% Ti)

compared to the values in the proxigrams displayed in Figures 4.4 and 4.5. That there is little Ti

in both the α-Al solid-solution and precipitate phases suggests that this analysis intersected an

interdendritic precipitate (e.g., B in Figure 4.1(b)).


- 3 - 2 - 1 0 1 2 3 4 5 6 7 8 9 1 0 1 1 1 2 1 3 1 4 1 5

0

5

1 0

1 5

2 0

2 5

3 0

3 5

α� ��

��

� ��

��

�

Zr,

Ti

co

mp

os

itio

n

(at.

%)


��

Fig. 4.8: Proxigram showing the distribution of Zr and Ti atoms in the Al3(Zr1−xTix) precipitate displayed in Figure 4.7, Al-0.1Zr-0.1Ti(a) aged at 375°C for 1,600 h. The precipitate is slightly super-stoichiometric (25–30 at.% total solute) with very little Ti (ca. 1at.%).

4.4 Discussion

Because of the significant microsegregation of solute species in these alloys (Figure 4.1), each

localized region sampled by a given 3-D APT analysis can be thought of as probing a unique and

a priori unknown alloy composition. Figure 4.9 shows calculated [316] isothermal sections of

the Al-Zr-Ti phase diagram at 375°C (solid lines) or 425°C (dashed lines), which indicates three-

phase equilibria among α-Al, Al3(Zr1−xTix) (L12), and Al3Ti (D022) in the Al-rich corner. Su-

perimposed are the data from the EDS linescan in Figure 4.2, showing measured ternary alloy

compositions across solute-rich dendritic and solute-depleted interdendritic regions in the al-

loy. While the local alloy composition varies significantly, Figure 4.9 illustrates that at both tem-

peratures (375 or 425°C), most of the alloy is in the α-Al + Al3(Zr1−xTix) (L12) two-phase field

with a few points (from the interdendritic regions) lying in the single-phase α-Al solid-solution.

No points are situated in the α-Al + Al3(Zr1−xTix) (L12) + Al3Ti (D022) three-phase triangle.

This is consistent with the observed precipitated microstructure (Figure 4.1), which shows only

Al3(Zr1−xTix) (L12) precipitates concentrated in the dendrite centers and precipitate-free α-Al


0 0 . 0 5 0 . 1 0 0 . 1 5 0 . 2 0 0 . 2 5 0 . 3 0

0

0 . 0 5

0 . 1 0

0 . 1 5

0 . 2 0

B u lk

r at i o

of s

o lu t

e s

α��

α��"��&��&��

�

Zr

(at.

%)

T i ( a t . % )α��

α��"��&��&��

� ��(��(�

�%�� '� �! #�$� �

�� '� �! #�$� ��

Fig. 4.9: Isothermal sections at 375 or 425°C of a calculated ternary phase diagram for the Al-Zr-Ti system in the Al-rich corner,showing the α-Al solid-solution in equilibrium with the metastable Al3(Zr1−xTix) (L12) and Al3Ti (D022). The locally-measuredalloy compositions (by EDS, Figure 4.2) are indicated by the circles. Phase diagram calculated by Murray [316].

solid-solution in the interdendritic channels.

The bulk ratio of solute atoms in the alloy is nearly unity (Table 4.1), yet the data in Figures 4.2

and 4.9 show a bias in the measured Zr:Ti ratio of approximately 2:3. This apparent enrichment

of Ti in the locally sampled region of the alloy (35 µm linescan, Figure 4.2) could represent a

macrosegregation of solutes. It is also possible that this apparent enrichment of Ti is an artifact

of the EDS technique. The data presented in Figure 4.2 were not calibrated with respect to known

standards, and are therefore strictly semi-quantitative. Nevertheless, the mean alloy composi-

tion measured by EDS is comparable to the bulk alloy composition (Figure 4.9), indicating that

the EDS results are fairly accurate.

The variation in alloy composition, shown in Figures 4.2 and 4.9, makes the discussion of the

measured compositions of the Al3(Zr1−xTix) precipitates difficult since the local alloy composi-

tion from which these precipitates form is a priori unknown. The situation is complicated fur-

ther by a significant disparity in evaporation fields between the α-Al and Al3(Zr1−xTix) phases,

resulting in distortions in atomic density in the reconstructed volume, as well as potential bi-

ases in the measured precipitate compositions. Nevertheless, consistent with the proxigrams

presented in Figures 4.4, 4.5, and 4.8 (as well as others measured), two conclusions are drawn:

(i) the precipitated phase is predominantly Zr-rich with a small amount of Ti partitioning to the


Fig. 4.10: A 3 nm-thick slice, cut parallel to the analysis direction, through the reconstruction in Figure 4.7. This volume contains6.12×103 identified atoms, with all atoms displayed.

Al3(Zr1−xTix) precipitates; and (ii) there is no evidence of interfacial segregation of solute at the

α-Al/Al3(Zr1−xTix) heterophase interface.

4.4.1 Local magnification of the Al3(Zr1−xTix) precipitates

Before discussing the measured precipitate compositions, the disparate evaporation fields of the

Al3(Zr1−xTix) and α-Al phases, and the potential effects on measured solute concentrations in

these phases and at their interface, must be discussed. Figures 4.3 and 4.7 exhibit noticeable

aberrations in the reconstructed volumes, including: (i) reduced atomic density of the recon-

structed precipitates; (ii) a shell of enriched atomic density at the matrix/precipitate interface;

and (iii) compression of the precipitate dimension in the analysis direction. These artifacts are

attributable to a local magnification effect, which can occur during field-evaporation of two-

phase systems when the phases have disparate evaporation fields, leading to variations in lo-

cal curvature of the specimen surface, and ultimately aberrations in the trajectories of field-

evaporated ions [301, 303, 317]. These aberrations are not accounted for in the reconstruction

procedure, leading to distortions in the reconstructed volumes.

The apparent local atomic density in the 3-D APT reconstruction provides an indication


of the severity of local focusing or defocusing effects resulting from ion trajectory aberra-

tions [318, 319]. The differences in the reconstructed atomic density of the α-Al solid-solution

and Al3(Zr1−xTix) precipitates are clear in Figure 4.10, which shows all identified atoms in a thin

(3 nm) slice parallel to the analysis direction through the reconstruction in Figure 4.7. The {111}-

type atomic planes in the α-Al matrix are well resolved and with the correct interplanar spacing,

indicative of an accurately reconstructed volume using reconstruction parameters appropriate

for the α-Al matrix. The same is not true for the reconstructed volume inside the precipitate.

The spheroidal precipitates are coherent and coplanar with the α-Al matrix (Chapter 5), and

therefore the {111} planes should be continuous through the Al3(Zr1−xTix) precipitate. The re-

constructed precipitate also exhibits a lower atomic density than the surrounding α-Al matrix. At

ambient temperature, the lattice parameters of Al and Al3Zr are similar, 0.405 and 0.408 nm, re-

spectively (Table 1.1), resulting in a 2% difference in molar volume. This small difference would

be visually indiscernible in a correctly reconstructed volume. It is obvious in Figure 4.10, how-

ever, that the reconstructed local atomic density of the α-Al solid-solution is significantly greater

than that of the Al3(Zr1−xTix) precipitate, with a measured density ratio of 3.4. The differences

in atomic densities between the two phases in the reconstruction of Figure 4.3 is even greater,

5.9. This greater disparity is also consistent with the possibly greater fraction of unpositioned

solute species (Zr and Ti) in the Al3(Zr1−xTix) precipitate in explaining the sub-stoichiometric

compositions observed in Figure 4.4.

The difference in evaporation fields between the α-Al matrix and Al3(Zr1−xTix) precipitates

is demonstrated quantitatively by considering the mass spectra in Figure 4.6. As discussed, 95%

of the detected Al ions are field evaporated as Al1+ from the α-Al solid-solution (Figure 4.6(b)),

whereas Al2+ is predominant from the Al3(Zr1−xTix) precipitates (Figure 4.6(c)), indicative of

the higher electric fields required to field-evaporate ions from the Al3(Zr1−xTix) precipitates.

This electric field is estimated from the calculated evaporation fields [320] for the constituent

species (Al, Zr, and Ti) in the precipitates, Table 4.3. Because the Al ions field-evaporated from

the Al3(Zr1−xTix) precipitates are almost all Al2+, the electric field required to field-evaporate

ions then must exceed 35 V nm-1 (E2 = 35 V nm-1, Table 4.3). This is also consistent with the

observed charge states of Zr ions in the precipitates: 93% of the field-evaporated ions are Zr3+

(E3 = 35 V nm-1) rather than Zr2+ (E2 = 28 V nm-1). These results agree with a recent analysis by

Sha and Cerezo [315], who used field ion microscopy and 3-D APT to estimate an evaporation


Table 4.3: Calculated evaporation fields (V nm-1) for singly- (E1),doubly- (E2), and triply- (E3) charged Al, Zr, and Ti ions [320]. Theexpected charge states are underlined.

Element E1 E2 E3

Al 19 35 50

Zr 56 28 35

Ti 41 26 43

field of 36 V nm-1 for Al3Zr precipitates in a commercial 7050 Al alloy.

The effect of local magnification aberrations on the measured compositions of nanometer-

scale precipitates by 3-D APT has been addressed both experimentally and with modeling [317–

319, 321]. Miller and Hetherington [321] showed, for a one-dimensional atom-probe, that while

trajectory overlap leads to significant spatial distortions in the reconstruction, the bulk com-

positions of the phases are generally unbiased. Recent simulations by Blavette and cowork-

ers [317, 319] have similarly concluded that the measured compositions from the inner core of

precipitates are unaffected by ion trajectory overlap effects occurring at the reconstructed ma-

trix/precipitate interface.

The evaporation field of Al3Zr (EAl3Zr = 36 V nm-1) is considerably greater than that of the

α-Al solid-solution (EAl = 19 V nm-1), as measured by Sha and Cerezo [315] and confirmed

presently. The corresponding evaporation field ratio [317], ε = EAl3Zr/EAl = 1.9, represents a

greater disparity in evaporation fields than those modeled by Blavette et al. (0.85≤ ε≤ 1.15) [317,

319]. Sha and Cerezo, however, measured directly the extent of ion trajectory overlap for ex-

perimental reconstructions and showed that, for Al3Zr precipitates ranging from 3.5–6.0 nm in

radius, the experimentally-measured ion trajectory overlaps increased from 1.4–2.0 nm with in-

creasing precipitate size. The reconstructed precipitates in Figures 4.3 and 4.7 are larger than

this, and so the extent of trajectory overlap (and hence the convolution in measured composi-

tion at the interface) may be greater for the precipitates discussed herein.

It is therefore thought that the precipitate composition profiles presented in Figures 4.4, 4.5,

and 4.8 are less accurate than indicated at the matrix/precipitate heterophase interface. Never-

theless, the reported plateau compositions of the proxigrams (and in particular the solute ratios)

at the center of the precipitates are expected to be reliable.


4.4.2 Compositions of the Al3(Zr1−xTix) precipitates

The proxigrams in Figures 4.4, 4.5, and 4.8 are representative of the compositions of all the

Al3(Zr1−xTix) precipitates observed by both LEAP tomography and conventional APT. The pre-

cipitates are invariably Zr-rich, with only a small fraction of Zr replaced by Ti. In all cases, the

absolute Ti concentration in the trialuminide is less than 5 at.%, despite the nominally equal

concentrations of Zr and Ti initially in the alloys (Table 4.1). The proxigrams shown in Fig-

ure 4.5, which indicate the temporal evolution of the Al3(Zr1−xTix) precipitates at 375°C, suggest

that the Zr:Ti ratio is approximately 10 and varies little with aging time. The proxigram in Fig-

ure 4.4, for a precipitate formed after 400 h of aging at 425°C, demonstrates that the Zr:Ti ratio

is approximately 5, roughly twice that at 375°C, corresponding to a precipitate composition of

Al3(Zr0.83Ti0.17). Like the composition of the alloy from which they form (Figure 4.2), however,

the concentrations of solutes in the precipitates are probably not uniform throughout the vol-

ume. Indeed, as shown in Figure 4.8, the measured Zr:Ti ratio can deviate substantially from

those stated above.

The small concentration of Ti observed in the precipitates could be reflective of: (i) the sub-

stantially higher solid-solubility of Ti in α-Al; and (ii) the smaller diffusivity of Ti in α-Al com-

pared to that of Zr. Figure 4.9 indicates that the solid-solubility of Al3Zr (L12) in metastable

equilibrium with α-Al is 0.008 and 0.016 at.% Zr at 375 and 425°C, respectively. The solid-

solubility of Al3Ti (D022) in α-Al is an order of magnitude greater: 0.101 and 0.163 at.% Ti at

375 and 425°C, respectively. This greater solubility of Ti would account for the small partition-

ing of Ti to the Al3(Zr1−xTix) precipitates and the residual Ti observed in α-Al solid-solution in

precipitate-containing 3-D APT analyses (Table 4.2). The calculated ternary Al-Zr-Ti phase di-

agram (Figure 4.9), however, indicates that both solutes are supersaturated with respect to the

Al3(Zr1−xTix) solvus, especially in the solute-rich dendrites. Thus, the small Ti concentration in

the Al3(Zr1−xTix) precipitates, as well as the residual Ti observed in α-Al solid-solution around

them, is not simply a function of the equilibrium solubilities of Zr and Ti.

The measured tracer diffusivity of Ti in α-Al is smaller than that of Zr in α-Al (Table 1.3); kinet-

ics, therefore, are probably the dominant influence dictating the observed small concentration

of Ti in the Al3(Zr1−xTix) precipitates. The diffusivities of Zr and Ti in α-Al are given by an Arrhe-

nius relationship, D = D0 exp (−Q/RgT ), where Q = 242 and 260 kJ mol-1 and D0 = 7.28 × 10−2

and 1.12× 10−1 m2 s-1 for Zr and Ti, respectively (Table 1.3). The root-mean-squared (RMS) dif-


fusion distances,√

6Dt, of Zr and Ti after 1,600 h at 375°C are 279 and 65 nm, respectively; after

400 h at 425°C these distances are 698 and 184 nm, respectively. These calculated RMS diffusion

distances are much smaller than the dimensions of the initial dendritic distribution of solute

(0.5–4 µm in Figure 4.1 and 3–10 µm in Figure 4.2). Achieving a homogenous distribution of Ti

is therefore impossible for reasonable aging treatments. Even at 375°C, the calculated RMS dif-

fusion distance of Ti is, however, greater than the mean edge-to-edge inter-precipitate distance

(ca. 10–50 nm) in the precipitate-rich dendrite centers (Figure 4.1). The diffusion distance of

Ti is also greater than that of a typical 3-D APT reconstruction (ca. 50×50×200 nm3 as shown

schematically in Figure 4.1). The growth of the precipitates may therefore be interface-limited

and not diffusion-limited. Additionally, diffusion of a solute in an ordered compound is slower

than in the matrix because of the effect of correlation on the diffusivity.

There is only one other study where the composition of Al3(Zr1−xTix) precipitates was mea-

sured directly. Using scanning transmission electron microscopy and energy dispersive X-ray

spectroscopy, Sato et al. [271] measured the composition of Al3(Zr1−xTix) precipitates formed

after aging an Al-0.09Zr-0.17Ti (at.%) for 168 h at 500°C. Despite the enhanced kinetics at 500°C

and the greater supersaturation of Ti as compared to the present study, the estimated Zr:Ti ratio

in their precipitates was approximately 10, consistent with our findings that only a small frac-

tion of the available Ti partitions to the Al3(Zr1−xTix) phase. Our results agree also with other

3-D APT studies on Al-Sc-Ti [118] and Al-Sc-Zr [116, 322, 323] alloys, discussed below, where the

slower-diffusing Ti and Zr atoms are generally a minor constituent in the precipitated trialu-

minide.

4.4.3 Residual Ti in the α-Al solid-solution

Table 4.2 demonstrates that the amount of Ti in α-Al observed locally by 3-D APT can be corre-

lated to the proximity of the 3-D APT analysis to the precipitate-rich and therefore solute-rich

regions in the alloy. This is indeed illustrative of how little Ti partitions to the Al3(Zr1−xTix)

precipitates during aging at both 375 and 425°C. That the local Ti concentration in α-Al solid-

solution is sometimes enriched (C/C0 > 1) and other times depleted (C/C0 < 1) relative to the

bulk composition is reflective of the initial inhomogeneous distribution of solute atoms shown

quantitatively in Figure 4.2. In the as-cast alloy the concentration of Ti, as measured by EDS, is

shown to be enriched nearly threefold relative to the bulk alloy composition within the super-


saturated dendrites. This is in reasonable agreement with the over twofold enrichment of Ti in

solid-solution near precipitates measured by 3-D APT (Table 4.2). That there is only a twofold

enrichment of Ti in α-Al in the aged alloys may be explained by the fact that Ti atoms parti-

tion to the precipitated Al3(Zr1−xTix) phase, as shown in the mass spectra of Figure 4.6 and the

proxigrams in Figures 4.4, 4.5, and 4.8, and hence Ti is depleted from the α-Al solid-solution.

As discussed, the solid-solubility of Al3Ti (D022) in α-Al is 0.101 or 0.163 at.% Ti at 375

or 425°C, respectively, which correspond to the Ti composition of the point in the α-Al +

Al3(Zr1−xTix) (L12) + Al3Ti (D022) tie triangle in Figure 4.9. Since all regions of the alloy are

in the α-Al solid-solution or the α-Al + Al3(Zr1−xTix) (L12) two-phase field (Figure 4.9), any tie

line drawn between these two phases must intercept the α-Al solvus at less than 0.101 at.% Ti

at 375°C or 0.163 at.% Ti at 425°C. Table 4.2 shows that the measured Ti concentration in α-Al

solid-solution for three precipitate-containing analyses performed on Al-0.1Zr-0.1Ti(a) aged at

375°C contain more than 0.101 at.% Ti, the equilibrium maximum solubility at this tempera-

ture. Similarly, the α-Al solid-solution from one precipitate-containing analysis performed on

Al-0.1Zr-0.1Ti(b) aged 400 h at 425°C contained 0.215 ± 0.005 at.% Ti (C/C0 = 2.15 ± 0.10), well

beyond the maximum solubility of 0.163 at.% Ti at 425°C. The fact that concentrations beyond

the equilibrium solubility are measured in some precipitate-containing 3-D APT analyses indi-

cates that the alloys have not achieved global thermodynamic equilibrium after 1,600 h at 375°C

or 400 h at 425°C.

As described in the Introduction, improving the thermal stability of Al3Zr with ternary or

quaternary additions of Ti and/or V has generated considerable interest in the scientific liter-

ature [84, 112–114, 264, 271, 288–298]. Most of these studies, however, assumed that the ratio

of solutes in the precipitated trialuminides was the same as that initially in solid-solution (i.e.,

the bulk alloy composition). This is at first glance a reasonable assumption since the transi-

tion elements are generally sparingly soluble in α-Al and all solute atoms should therefore par-

tition to the precipitated phase. In the present study the alloys had initially the same atomic

fraction of solute species (Table 4.1), yet this Zr:Ti ratio of unity is never observed in the pre-

cipitated Al3(Zr1−xTix) phase. Moreover, as evidenced by Table 4.2, the assumption of negligi-

ble solubility of solutes in α-Al is invalid; diffusion kinetics in α-Al, which is anomalously slow

and varies widely among the transition elements (Section 1.2.3), is instead the dominant factor

dictating the relative amounts of solute atoms in the precipitated phase. The precipitate com-


positions cited in previous studies [84, 112–114, 264, 288–298] are therefore highly speculative

(Sato et al. [271], however, verified the compositions of their Al3(Zr1−xTix) precipitates by EDS,

as noted).

4.4.4 Comparison to 3-D APT studies on Al-Sc-Ti/Zr alloys

It is valuable to compare our measured precipitate compositions to those for Al-0.06Sc-0.06Ti

(at.%) alloys reported in a similar 3-D APT study by van Dalen et al. [118]. Like Al3Zr, the or-

dered Al3Sc phase precipitates from a supersaturated solid-solution during aging, and also has

the L12 structure; unlike Al3Zr, however, Al3Sc is a thermodynamically stable L12 ordered struc-

ture. Even after aging at 300°C for 1,536 h (64 days), van Dalen et al. [118] showed that Ti atoms

partition weakly to the Al3(Sc1−xTix) phase, forming precipitates with the average composition

Al3(Sc0.94Ti0.06), corresponding to a Sc:Ti ratio of 16, with Ti remaining primarily in α-Al solid-

solution. The amount of Ti measured in α-Al solid-solution by 3-D APT decreased continuously

with aging time at 300°C, yet still exceeded the calculated equilibrium value in α-Al after 1,536 h,

indicating that global thermodynamic equilibrium had not yet been reached.

The higher coarsening resistance of Al3Zr compared to Al3Sc, due to the slower diffusion

kinetics of Zr (Section 1.2.3), allows higher aging temperatures (up to 425°C vs. 300°C for the

Al-Sc-Ti alloys [118]). This amounts to a significant difference in diffusion kinetics, with the cal-

culated diffusivity of Ti being 1.8× 104 times greater at 425°C than it is at 300°C. Accordingly, the

proxigram shown in Figure 4.4 indicates that the composition of a precipitate formed after 400 h

at 425°C is Al3(Zr0.83Ti0.17). Figure 4.5, for precipitates formed at 375°C, indicates compositions

of Al3(Zr0.91Ti0.09). These both represent a larger fraction of Ti compared to Al3(Sc0.94Ti0.06) ob-

served after 1,536 h at 300°C [118]. This may be explained in terms of diffusion kinetics: the

calculated RMS diffusion distance for Ti in α-Al after 1,536 h at 300°C is 3 nm, while it is 7 nm

after 20 h at 375°C (corresponding to one of the proxigrams in Figure 4.5).

The distribution of Ti in the Al3(Sc1−xTix) precipitates studied by van Dalen et al. [118] is

inhomogeneous, with the precipitates consisting of an Al3Sc core surrounded by a spherical Ti-

enriched outer shell. Similar cored composite precipitates are well-documented for Al-Sc-Zr

alloys, as observed by 3-D APT [116, 322–324], analytical high resolution electron microscopy

[324, 325], small-angle X-ray scattering [324], and atomic-scale simulations [324, 326]. Both Ti

and Zr reduce the lattice parameter of Al3Sc (L12) [120,121], thus reducing the lattice parameter


mismatch with α-Al. The enriched outer shell of the mismatch-reducing solutes (Ti or Zr) should

therefore minimize the elastic strain energy at the coherent matrix/precipitate interface [116,

118,323–325]. It has also been suggested [116,322–326] that, because of the much faster diffusion

rate of Sc compared to Ti or Zr, only Sc atoms are kinetically able to participate in the early stages

of nucleation and growth. These pre-existing clusters then act as heterogeneous nucleation sites

for the less-mobile Ti or Zr atoms, accounting for the observed cored precipitate compositions.

In the Al-Zr-Ti system, Ti is the slowest diffuser and has also been shown to reduce the

lattice parameter mismatch with α-Al of the metastable Al3Zr (L12) phase [292]. By analogy

with the Al-Sc-Ti or Al-Sc-Zr systems, therefore, one might expect segregation of Ti to the ma-

trix/precipitate interface. Nevertheless, our results provide no evidence for such segregation

at the α-Al/Al3(Zr1−xTix) interface after extended aging at 375 or 425°C. The lattice parameter

mismatch of Al3Zr (L12) with α-Al at room temperature, δ = +0.75%, is less than that of Al3Sc,

δ = +1.32% (Table 1.1), which means the elastic strain energy (proportional to the square of the

mismatch [327,328]) of the coherent heterophase interface is also less. There is also the possibil-

ity that interfacial segregation of Ti may be veiled by the significant local magnification for this

system, as discussed.

4.5 Chapter Summary

Atom-probe tomography was applied to measure directly the chemical compositions of

Al3(Zr1−xTix) (L12 structure) precipitates formed in conventionally-solidified Al-0.1Zr-0.1Ti

(at.%) alloys during extended aging times at 375°C (to 1,600 h) or 425°C (to 400 h). The following

results were obtained and discussed:

• Dendritic solidification results in non-uniform distributions of Zr and Ti, which segregate

to the dendrite centers and are enriched by a factor of 2 and 3, respectively. The interden-

dritic regions are concurrently solute-depleted. During subsequent aging, Al3(Zr1−xTix)

precipitates are similarly non-uniformly distributed in a dendritic manner.

• The Zr:Ti atomic ratio in the Al3(Zr1−xTix) precipitates, as measured by atom-probe to-

mography, is 10 (x = 0.09) at 375°C and 5 (x = 0.17) at 425°C, indicating that Ti partitions

weakly to the Al3(Zr1−xTix) phase. The precipitate compositions are almost invariant with

aging time.


• While Ti unambiguously partitions to the Al3(Zr1−xTix) precipitates, residual Ti is ob-

served in analyzed volumes (104–106 nm3) of the α-Al solid-solution near the precipitates,

which is greater than the calculated solid-solubility. The small concentration of Ti in the

Al3(Zr1−xTix) precipitates, as well as the enriched Ti concentration beyond the calculated

solubility in α-Al, indicates that global thermodynamic equilibrium has not been achieved

after 1,600 h at 375°C or 400 h at 425°C. This is consistent with 3-D APT studies on Al-Sc-Ti

alloys aged at 300°C [118].

• There is a correlation between the measured Ti concentration in α-Al solid-solution (within

the volume of a typical 3-D APT analysis, ca. 104–106 nm3) and proximity of the analysis to

the precipitate-rich dendrites, which is reflective of the initial non-uniform distribution of

solute atoms.

• There is no indication of Ti segregation at the α-Al/Al3(Zr1−xTix) heterophase interface.

A significant disparity in evaporation fields exists between the α-Al matrix (19 V nm-1)

and Al3(Zr1−xTix) precipitates (36 V nm-1), leading to ion trajectory overlap at the recon-

structed matrix/precipitate interface. The accuracy of the measured composition at the

interface is therefore questionable. The measured compositions in the center of the pre-

cipitates should, however, be unaffected by trajectory overlap and are expected to be reli-

able.

CHAPTER 5

Microstructure and Mechanical Properties ofAl-Zr and Al-Zr-Ti Alloys during Isothermal

Aging at 375, 400, or 425°C

Precipitation of Al3Zr (L12) and Al3(Zr1−xTix) (L12) is investigated in conventionally-

solidified Al-0.1Zr, Al-0.2Zr, Al-0.1Zr-0.1Ti, and Al-0.2Zr-0.2Ti (at.%) alloys isothermally

aged at 375, 400, or 425°C. Pronounced hardening accompanies precipitation of Al3Zr

or Al3(Zr1−xTix) for all alloys investigated. The magnitude of the peak-aged hardness

is controlled primarily by the Zr concentration. The alloys overage sluggishly, with

no benefit from Ti additions in delaying overaging. After extended aging times (up to

1,600 h) at 425°C, there is no difference in the mean precipitate radii of Al3Zr (L12) or

Al3(Zr1−xTix) (L12) precipitates, indicating that there is no benefit in terms of coars-

ening resistance from ternary alloying additions of Ti. Previous coarsening studies on

Al3(Zr1−xVx) (L12), Al3(Zr1−xTix) (L12), and Al3(Zr1−x−y, Vx, Tiy) (L12) precipitates at

425°C are reviewed and discussed, and are compared to the present results.

5.1 Introduction

AS SHOWN IN PREVIOUS CHAPTERS, THE AL-ZR SYSTEM shows particular promise for develop-

ing thermally-stable precipitation-strengthened Al alloys. In this system, an ordered tria-

luminide, Al3Zr, precipitates during aging from supersaturated solid-solution. While the equi-

librium structure of Al3Zr is tetragonal (D023), decomposition of supersaturated Al-Zr solid-

solutions occurs initially by the formation of nanometer-scale Al3Zr precipitates with a meta-

stable cubic L12 structure, structurally and chemically analogous to the Ni3Al (γ′) phase in the

Ni-based superalloys, which are thermally stable at high homologous temperatures (Chapter 1).

Fine has long argued [1, 37, 38, 84] that even greater stability can be achieved by improving the

lattice parameter mismatch between Al3Zr and the α-Al solid-solution, and showed that the lat-

tice parameters of both the stable D023 and metastable L12 Al3Zr precipitates are reduced by

additions of Ti, Hf, or V [84, 110, 111]. The Al-Zr-V system, in particular, was studied exten-

99

100 CHAPTER 5 MICROSTRUCTURE OF AL-ZR AND AL-ZR-TI ALLOYS DURING AGING AT 375, 400, OR 425°C

sively [84, 112–114]; by substituting V for Zr, Al3(Zr1−xVx) (L12) precipitates were shown to ex-

hibit reduced coarsening compared with that of the binary Al3Zr precipitates [84, 113, 114]. The

improved stability of the Al3(Zr1−xVx) precipitates proved to be effective barriers to dislocation

motion during creep at 425 °C [112–114]. However, the formation of an additional Al10V phase,

particularly at grain boundaries, deteriorated the mechanical properties [112, 113].

Titanium also reduces the lattice parameter mismatch of Al3Zr (L12) with α-Al [292] and,

since Al3Ti is the terminal intermetallic compound (Figure 1.4, Table 1.2), the problem of

additional embrittling intermetallics is mitigated. Furthermore, an initial investigation by

Parameswaran et al. [288] showed that an Al-Zr-Ti alloy containing 1 vol.% Al3(Zr0.75Ti0.25) ex-

hibited much improved coarsening resistance at 425°C compared to Al-Zr and Al-Zr-V alloys of

similar precipitate volume fractions [84].

The present study is an extension of the one by Parameswaran et al. [288]. The Al3(Zr1−xTix)

precipitate compositions, measured by 3-D atom-probe tomography, are discussed in Chapter 4.

This chapter reports on the effects of Ti additions to dilute Al-Zr alloys by examining their am-

bient temperature mechanical properties, utilizing Vickers microhardness measurements, and

correlating these results to their microstructure during isothermal aging at 375, 400, and 425°C.

The effect of Ti additions on the Al3(Zr1−xTix) (L12) coarsening rates at 425°C is also assessed.

In Chapter 6, similar alloys are studied at temperatures in excess of 450°C, where the L12 to D023

transformation is discussed. Elevated-temperature mechanical properties, in the form of creep

properties, are discussed in Chapter 7.


A series of binary Al-Zr and ternary Al-Zr-Ti alloys were investigated; alloy designations, com-

positions, and aging conditions are summarized in Table 5.1. The alloys were prepared employ-

ing non-consumable electrode arc-melting, as described in Sections 3.2 and 4.2. The verified

compositions in Table 5.1 were obtained by bulk chemical analysis performed by direct current

plasma emission spectroscopy (ATI Wah Chang, Albany, OR). As-cast specimens were isother-

mally aged at 375, 400, or 425°C in air, within the range of temperatures shown to produce

a strong precipitation strengthening response for Al-Zr alloys produced by RSP [14, 264, 267–

270] or chill-casting [187, 190, 271]. The microstructures obtained after aging were investigated


Table 5.1: Compositions and aging conditions of the Al-Zr and Al-Zr-Ti alloys investigated.

Nominal comp. (at.%) Verified comp. (at.%) Aging temperatures (°C)Zr Ti Zr Ti

Al-0.1Zr(a) 0.1 — — — 375

Al-0.1Zr(b) 0.1 — 0.085 — 375, 400, 425

Al-0.2Zr 0.2 — 0.186 — 375, 400, 425

Al-0.1Zr-0.1Ti(a) 0.1 0.1 — — 375

Al-0.1Zr-0.1Ti(b) 0.1 0.1 0.089 0.099 375, 400, 425

Al-0.2Zr-0.2Ti 0.2 0.2 0.189 0.198 375, 400, 425

in both alloys by transmission electron microscopy (TEM) and scanning electron microscopy

(SEM), using similar instruments and techniques described in Section 3.2.


5.3.1 As-cast microstructure

Figure 5.1 shows the as-cast macrostructure of the alloys studied. The solidification macrostruc-

ture is typical of cast alloys, with coarse columnar grains, originating at the bottom surface of the

ingot (which was in contact with the chilled copper crucible of the arc-meler), growing upward

toward a zone of equiaxed grains near the center of the ingot. The relative sizes of the columnar

and equiaxed zones is strongly dependent on the solute content of the alloy and, more precisely,

the extent of properitectic Al3M (M = Zr or Zr,Ti) precipitation, since these primary phases are

potent grain refiners, as discussed previously.

Alloy Al-0.1Zr(b) is the most dilute of those studied (Table 5.1), and the extent of the colum-

nar zone, which comprises the entire ingot cross-section in this alloy, is also greatest. The grains

in the columnar zone are large, 0.2–1.0 mm wide, and ca. 5 mm long. The effect of doubling the

solute concentration with an addition of 0.1 at.% Ti is demonstrated for Al-0.1Zr-0.1Ti(b). For

this alloy, the upper half of the ingot exhibits a zone of coarse (0.5–1.0 mm diameter) equiaxed

grains. Alloy Al-0.2Zr exhibits a finer structure still, where the grains in the columnar zone are

typically 0.2 mm wide and the extent of this zone is also less. Consequently, the majority of the

ingot is equiaxed, with finer grain sizes (0.15–0.30 mm) than for alloy Al-0.1Zr(b). Alloy Al-0.2Zr-


Fig. 5.1: Macrostructure of as-cast alloys Al-0.1Zr(b), Al-0.1Zr-0.1Ti(b), Al-0.2Zr, and Al-0.2Zr-0.2Ti (etched using Poultan’s reagent),showing various degrees of grain refinement.

0.2Ti, the most concentrated alloy studied (Table 5.1), exhibits the finest grain size; the entire

ingot cross-section is equiaxed with very fine grains ca. 100 µm diameter.

Figures 5.2 and 5.3 display SEM micrographs of metallographically-polished Al-0.2Zr and Al-

0.2Zr-0.2Ti, respectively. Both alloys exhibit petal-like precipitates of the properitectic phase

(Al3Zr or Al3(Zr1−xTix)), responsible for the grain refinement observed in Figure 5.1. This petal-

like morphology is characteristic of the metastable L12 structure [160], whose cubic structure

is commensurate with fcc α-Al and acts as an effective heterogeneous nucleant of α-Al during

solidification. The presence of these primary precipitates indicates also that, for the conven-

tional casting conditions used in this study, exceeding ca. 0.2 at.% solute (Ti or Zr) results in

primary precipitation of the properitectic trialuminide and hence decreases the amount of so-

lute retained in solid-solution. No primary precipitates are observed by SEM in the more-dilute

Al-0.1Zr and Al-0.1Zr-0.1Ti alloys, consistent with the coarser grain structure in Figure 5.1.

Composition of the Al3(Zr1−xTix) primary precipitates

The composition of one of the primary Al3(Zr1−xTix) precipitates in alloy Al-0.2Zr-0.2Ti (Fig-

ure 5.3) was measured using quantitative energy-dispersive spectrometry (EDS) in SEM, using a


Fig. 5.2: SEM secondary electron images of as-solidified Al-0.2Zr. Panel (a) displays the ingot top. Z-contrast delineates Zr-richdendrite arms (lighter contrast). Panel (b) displays the ingot bottom. Several petal-like primary Al3Zr precipitates (light contrast)are observed.

Fig. 5.3: SEM secondary electron images of as-solidified Al-0.2Zr-0.2Ti. Numerous Al3(Zr1−xTix) primary precipitates are ob-served. Panel (a) displays the ingot top. Numerous Al3(Zr1−xTix) primary precipitates are observed. Panel (b) displays the ingotbottom. Fewer primary precipitates are observed. In Figures 5.2 and 5.3 solute-enriched dendritic cells are also visible.

Hitachi S-3500 operated at 20 kV and equipped with an Oxford Instruments EDS system. X-ray

spectra were acquired for 10 s at a constant sample working distance of 15 mm. The beam was

used at maximum focusing (i.e. with a diameter of a few nm) with a beam current of 0.66 nA,

calibrated with an internal Faraday cup. The recorded spectra were compared quantitatively to

previously-recorded pure reference standards generated under identical beam conditions.

Figure 5.4 shows the analyzed properitectic Al3(Zr1−xTix) precipitate with the approximate


Fig. 5.4: SEM micrograph of a primary Al3(Zr1−xTix) pre-cipitate seen in Figure 5.3(a) for Al-0.2Zr-0.2Ti. Quantita-tive EDS measurements, reported in Table 5.2, were per-formed at the positions indicated.

Table 5.2: Quantitative EDS analysis of the properitecticprecipitate of Figure 5.4 in as-cast Al-0.2Zr-0.2Ti.

Position EDS compositions (at.%) Ti:Zr

Zr Ti Total

1 3.95 6.30 10.25 1.59

2 5.76 8.75 14.51 1.52

3 6.04 8.45 14.49 1.40

4 0.84 2.02 2.86 2.40

positions of the four EDS measurements. The measured compositions are displayed in Table 5.2.

While the data are quantitatively accurate, the compositions are not reflective of the primary

phase since the detected X-rays by EDS originate from an interaction volume (ca. 1 µm3) that

includes the underlying α-Al solid-solution, explaining why the total solute content is less than

the 25 at.% expected for the primary trialuminide. Nevertheless, the measured concentrations

in Table 5.2 show unambiguously that both Ti and Zr segregate to the primary phase, and that

these precipitates are slightly Ti-rich with a Ti:Zr ratio of approximately 1.5.

The halo surrounding the primary phase in Figure 5.4 may be a sub-surface projection of

the properitectic precipitate. This could also represent solute-enriched α-Al solid-solution en-

veloping the primary phase, since the first solid to form is nucleated heterogeneously onto the

Al3(Zr1−xTix) precipitate. Quantitative EDS measurements, position 4 in Table 5.2, indicate

that this halo is enriched in both Zr and Ti, supporting both hypotheses. It is shown in Fig-

ures 5.17 and 5.18 that solid-state spheroidal Al3(Zr1−xTix) (L12) precipitates nucleate in regions

surrounding the petal-like primary phases during post-solidification aging, suggesting that the

halo in Figure 5.4 indicates an enrichment in solute concentration in α-Al solid-solution.

5.3.2 Solute segregation and the precipitated microstructure

In addition to the primary precipitates described above, Figures 5.2 and 5.3 reveal also the solute-

enriched dendrites in the as-solidified alloys. The dendritic regions exhibit lighter contrast, in-


(a) (b)

0 5 10 15 20 25 30 350

0.1

0.2

0.3

0

0.1

0.2

0.3

0.4

C0

C0

Zr

Ti

Com

posi

tion

(at.%

)

Distance (µm)

0 5 10 150

0.2

0.4

0.6

0

0.2

0.4

0.6

0.8

C0

C0

Zr

Ti

Com

posi

tion

(at.%

)

Distance (µm)

Fig. 5.5: Concentration profiles of Zr and Ti, measured by EDS, across three solute-rich dendrite arms in as-cast ternary Al-Zr-Tialloys: (a) Al-0.1Zr-0.1Ti(b); (b): Al-0.2Zr-0.2Ti. Both Zr and Ti segregate to the dendrite centers (k0>1).

dicative of the higher solute concentration in α-Al solid-solution (Z-contrast from backscattered

electrons). This concentration variation between the dendritic and interdendritic regions was

quantified by EDS using a JEOL JSM-7000F SEM equipped with a ThermoNORAN EDS detector

(procedure discussed previously, Section 4.2).

Measured composition profiles spanning three dendrite branches are displayed in Figure 5.5

for alloys Al-0.1Zr-0.1Ti(b) and Al-0.2Zr-0.2Ti. The two composition profiles in Figure 5.5 are

similar in that the dendrite centers are enriched approximately twofold in Zr and threefold in

Ti compared to their respective bulk compositions (Table 5.1). The interdendritic regions are

concurrently solute-depleted.

The mean concentrations of Zr and Ti in Figure 5.5(a) are 0.10 at.% Zr and 0.14 at.% Ti; in

Figure 5.5(b) these values are 0.28 at.% Zr and 0.40 at.% Ti. It is noteworthy that both compo-

sition profiles exhibit a bias in the measured Ti:Zr ratio of approximately 1.4 (as noted also in

Figure 4.9). This apparent enrichment of Ti in the locally sampled region of the alloy (ca. 10 µm)

could represent a macrosegregation of solutes. The fact that both composition profiles exhibit

a similar bias, however, suggests that the apparent enrichment of Ti is an artifact of the EDS

technique. Unlike the data in Table 5.2, the data in Figure 5.5 were not calibrated with respect

to known standards, and are therefore strictly semi-quantitative. Nevertheless, the mean alloy

composition measured by EDS is comparable to the bulk alloy composition, C0 (especially true


Fig. 5.6: TEM and SEM micrographs of Al-0.1Zr-0.1Ti(a) aged at 375°C for 1,600 h. Panels (a) and (b) are centered dark-field TEM im-ages showing the dendritic distribution of nanometer-scale Al3(Zr, Ti) (L12) precipitates. Panels (c) through (f) are SEM secondaryelectron images, which illustrate the inhomogenous nature of the dendritic distribution.


for Zr), indicating that the EDS results are reasonably accurate.

During post-solidification aging, nanometer-scale Al3Zr (L12) or Al3(Zr1−xTix) (L12) precip-

itates form, non-uniformly distributed throughout the alloy, reflecting the dendritic microseg-

regation of Zr and Ti solute atoms in the as-cast alloys. Figure 5.6 shows the precipitated mi-

crostructure of the alloys, obtained after aging alloy Al-0.1Zr-0.1Ti(a) at 375°C for 1,600 h.

5.3.3 Age hardening

Figures 5.7–5.10 display the observed microhardness of the Al-Zr and Al-Zr-Ti alloys after isother-

mal aging at 375, 400, or 425°C. Each data point represents a minimum of 20 measurements, with

the standard deviation of these measurements indicated by the error bars. The position of the

measurement within the ingot (i.e., ingot top, center, bottom) was also recorded since the cool-

ing rate of the melt — and therefore the amount of solute retained in solid-solution — probably

varies with proximity to the chilled copper crucible. A positional-dependence on hardness was

observed only in Al-0.2Zr-0.2Ti, where the ingot bottom is consistently 70 MPa harder than the

rest of the alloy. This increased hardness is probably a result of suppressed properitectic precip-

itation of Al3(Zr1−xTix) near the ingot bottom, as shown in Figure 5.3. To minimize the effect

of solute loss to this primary phase, therefore, only hardness data measured near the bottom

surface of alloy Al-0.2Zr-0.2Ti are reported in Figure 5.10.

There is generally no significant difference in the observed hardness between the Al-Zr and

Al-Zr-Ti alloys for a given Zr concentration (0.1 or 0.2 at.%) and aging temperature. For all alloys

investigated, the peak hardness decreases with increasing aging temperature, attributable to the

natural reduction in volume fraction of the dispersed phase due to the temperature-dependence

of solute solid-solubility. There is a particularly pronounced drop in plateau hardness between

aging at 400 and 425°C in the alloys containing 0.1 at.% Zr (Figures 5.7 and 5.9), suggesting a sig-

nificant decrease in volume fraction of the Al3Zr (L12) or Al3(Zr1−xTix) (L12) precipitates. The

reduction in hardness at 425°C is, however, less for Al-0.1Zr-0.1Ti(b), suggesting that Ti addi-

tions increase the volume fraction of L12 precipitates, which is particularly noticeable at higher

temperatures.1

Differences in age hardening behavior are even less distinct between Al-0.2Zr and Al-0.2Zr-

1In Chapter 4 it is shown by 3-D APT that Ti partitions to the precipitates, forming Al3(Zr1−xTix), and hence anincrease in volume fraction is not unexpected. However, the effect is small since the measured Zr:Ti atomic ratio is5–10, indicating that Ti partitions weakly to the precipitated phase.


0 . 1 1 1 0 1 0 0 1 0 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

3 0

4 0

5 0

6 0

7 0

8 0

�� !��

A l - 0 . 1 Z r

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Fig. 5.7: Vickers microhardness vs. aging time for Al-0.1Zr(a) and Al-0.1Zr(b) at 375, 400, or 425°C.

0 . 1 1 1 0 1 0 0 1 0 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

3 0

4 0

5 0

6 0

7 0

8 0

A l - 0 . 2 Z r

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Fig. 5.8: Vickers microhardness vs. aging time for Al-0.2Zrat 375, 400, or 425°C.

0 . 1 1 1 0 1 0 0 1 0 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

3 0

4 0

5 0

6 0

7 0

8 0

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A l - 0 . 1 Z r - 0 . 1 T i

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Fig. 5.9: Vickers microhardness vs. aging time for Al-0.1Zr-0.1Ti(a) and Al-0.1Zr-0.1Ti(b) at 375, 400, or 425°C.

0 . 1 1 1 0 1 0 0 1 0 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

3 0

4 0

5 0

6 0

7 0

8 0

A l - 0 . 2 Z r - 0 . 2 T i

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Fig. 5.10: Vickers microhardness vs. aging time for Al-0.2Zr-0.2Ti at 375, 400, or 425°C.

0.2Ti. These alloys, however, contain primary (properitectic) Al3Zr or Al3(Zr1−xTix) precipitates,

Figures 5.2 and 5.3, which are especially numerous in alloy Al-0.2Zr-0.2Ti. The solute retained

in α-Al solid-solution is therefore less than the bulk concentrations indicated in Table 5.1, and

the overall volume fraction of dispersed phases precipitated during post-solidification aging is

similarly unknown. Table 5.2 demonstrates that both Ti and Zr segregate to the primary phase

in alloy Al-0.2Zr-0.2Ti. Moreover, these precipitates are more numerous in alloy Al-0.2Zr-0.2Ti

than in Al-0.2Zr, as demonstrated by the pronounced grain refinement in Figure 5.1 and also

shown in Figure 5.3(a). The supersaturation of Zr retained in α-Al solid-solution in alloy Al-


0.2Zr-0.2Ti may indeed be less than that in Al-0.2Zr, and the similar hardness values for these

alloys, Figures 5.8 and 5.10, may indicate a similar benefit from Ti additions as that shown for

Al-0.1Zr-0.1Ti(b), Figure 5.7.

In addition to the noticeable increase in peak strength at 425°C, Ti additions appear to also

accelerate the onset of precipitation hardening in the more-dilute alloys containing 0.1 at.% Zr

(Figures 5.7 and 5.9). According to classical nucleation theory, the incubation time for nucle-

ation [279–281, 329] is inversely proportional to the square of the chemical driving force, which

is dependent on the solute supersaturation (Eq.(3.5)). The increased hardness at higher tem-

peratures and the accelerated incubation time for nucleation in Al-0.1Zr-0.1Ti(b) compared to

Al-0.1Zr(b) are indicative of a larger supersaturation of solute.

The incubation time varies also with aging temperature. For Al-0.2Zr and Al-0.2Zr-0.2Ti, Fig-

ures 5.8 and 5.10, the incubation time decreases with increasing aging temperature. The exact

opposite effect is observed for Al-0.1Zr-0.1Ti(b), Figure 5.9, where the onset of hardening is no-

ticeably accelerated compared to aging at 400 or 425°C. Finally, for alloy Al-0.1Zr(b), Figure 5.7,

there is little effect of aging temperature on the incubation time. The hardness values for short

aging times (t < 6.5 h) in Figures 5.7–5.10 were secured with rather small specimens from var-

ious positions in the ingot cross section. The inconsistent variations in incubation time may

therefore reflect slight variations in solute concentration due to different solidification rates with

respect to the location in the ingot. A more detailed study focussing on the early stages of the

decomposition of these alloys during isothermal aging, accompanied also with electrical con-

ductivity measurements, would be beneficial.

5.3.4 Precipitate morphologies

The SEM and TEM micrographs in Figure 5.6 suggest that the precipitated microstructure is bi-

nary in nature, with fine, spheroidal, coherent Al3Zr (L12) or Al3(Zr1−xTix) (L12) precipitates

confined to the dendritic cells, which are surrounded by precipitate-free interdendritic chan-

nels. In reality, the situation is not so simple due to complicated solute concentration gradients

around the dendrites. Figure 5.11 reveals the precipitated microstructure in an interdendritic

channel between two precipitate-rich dendrite arms, produced in alloy Al-0.2Zr-0.2Ti after ag-

ing at 400°C for 100 h (peak hardness, Figure 5.10). In any precipitation reaction, the precipitate

size decreases with increasing solute content (for a given aging temperature) since the chemical


Fig. 5.11: Centered superlattice dark-field TEM micrograph of an interdendritic channel between two dendrite arms in alloy Al-0.2Zr-0.2Ti aged at 425°C for 100 h. The foil edge is on the right-hand side of the micrograph, with increasing specimen thicknesstowards the left-hand side. Within the dendrites themselves, the Al3(Zr1−xTix) precipitates are small and homogeneously dis-tributed. Because the foil is thicker in these regions (due to differential thinning), the small (<10 nm) precipitates are clearly re-solved only near the foil edge. There is a gradient in the size and number density of the larger interdendritic precipitates, becomingsmaller and more numerous with proximity to the dendrite stalk (towards the left-hand side of the figure).

Fig. 5.12: Centered superlattice dark-field TEM micrographs of heterogeneously-nucleated spheroidal L12 interdendritic precip-itates. Panel (a) displays Al-0.2Zr-0.2Ti aged at 425°C for 100 h (detail of interdendritic region in Figure 5.11). Panel (b) displaysAl-0.2Zr aged at 400°C for 100 h. The precipitate-rich dendritic regions are visible in the upper right and lower left.


Fig. 5.13: Development of growth instabilities with increasing precipitate radius for interdendritic Al3(Zr1−xTix) precipitates inAl-0.1Zr-0.1Ti(b) aged at 425°C for 400 h.

driving force for precipitation increases with supersaturation. This leads to a smaller critical ra-

dius for nucleation and hence smaller precipitates. Within the center of the dendrites, where

the solute supersaturation is largest (Figure 5.5), precipitates are small and homogeneously-

distributed at high number densities. The supersaturation decays with lateral position from the

center of the dendritic cells and correspondingly, the precipitates become progressively larger.

The interdendritic regions contain insufficient solute to effect homogeneous nucleation so

these regions are largely precipitate-free. Precipitates that form here do so heterogeneously,

Figure 5.12, as evidenced by their arrangement in linear arrays. These precipitates are likely

formed on dislocations since the strain energy barrier for precipitation via this mechanism is

significantly reduced [330, 331].

Cauliflower-shaped interdendritic precipitates

In the interdendritic regions, where the spheroidal L12 precipitates are largest and their volume

fraction is smallest, cauliflower-like morphologies are frequently observed, Figure 5.13. These

arise from instabilities during the growth of the Al3Zr or Al3(Zr1−xTix) precipitates, and similar

morphologies have been observed for Al3Sc (L12) [117, 119, 332] and Al3Li (L12) [333] precipi-

tates when nucleated under small supersaturations. Because of their small number density, the

interdendritic precipitates in the present alloys are small enough for them to grow in a super-


saturated matrix before their diffusion fields commence overlapping. The morphology is then

determined by the growth conditions rather than the equilibrium conditions.

The gradients in precipitate radius that develop between the dendritic and interdendritic re-

gions are not uniform. Figure 5.14(a–c) are SEM micrographs of alloy Al-0.2Zr aged at 425°C

for 400 h, and Figure 5.14(d–f) are TEM micrographs of alloy Al-0.2Zr-0.2Ti aged at 425°C for

400 h. The TEM micrograph in Figure 5.14(d) demonstrates that the size gradients occurring

on either side of the same interdendritic channel may vary considerably. On the left-hand side,

Figure 5.14(f), the change in precipitate radius is gradual, similar to that shown in Figure 5.13.

On the right-hand side of the precipitate-free channel, Figure 5.14(d), the gradient in precipi-

tate size is abrupt. These gradients, occurring over the micrometer-scale, are best visualized by

SEM, Figure 5.14(a), which shows similar coexistence of sharp and broad size gradients across

a precipitate-free interdendritic channel. The various gradients in precipitate radius probably

reflect the initial distribution of solute species, and are thus a function of the critical radius for

nucleation as discussed above.

Loss of coherency in Al3Zr and Al3(Zr1−xTix) precipitates

The gradual precipitate size gradients occurring in the interdendritic channels provides an op-

portunity to assess the change in coherency of the Al3Zr and Al3(Zr1−xTix) precipitates. Fig-

ure 5.15 displays three pairs of complementary dark-field and bright-field TEM micrographs

recorded under dynamical two-beam conditions. Under such imaging conditions, coherent

precipitates observed in bright-field exhibit characteristic Ashby-Brown [334] strain contrast,

as shown for the precipitates near the center of Figure 5.15(b) exhibiting distinct lines of no con-

trast normal to the diffraction vector, ~g. For slightly larger interdendritic precipitates, coherency

strain contrast is lost and striped lines of contrast normal to ~g appear in the images of many of

the precipitates. These are interpreted as contrast due to interfacial dislocations, thus signaling

the transition to semicoherency of the Al3Zr or Al3(Zr1−xTix) precipitates. Figures 5.15(d) and

5.15(f) illustrate other examples of this phenomenon in different alloys aged at different temper-

atures. Iwamura and Miura [335] show similar transitions in strain contrast in loss of coherency

for Al3Sc (L12) precipitates.

Coherency across a precipitate-matrix heterophase interface is generally lost when the pre-


Fig. 5.14: SEM micrographs of Al-0.2Zr, Panels (a) through (c), and TEM micrographs of alloy Al-0.2Zr-0.2Ti, Panels (d) through (f),aged at 425°C for 400 h. Gradients in precipitate radius vary considerably throughout the microstructure, even across the same in-terdendritic channel. Panels (c) and (f) illustrate the development of growth instabilities, and demonstrate that the microstructuresof the alloys are indistinguishable.


Fig. 5.15: Complementary dark-field and bright-field TEM micrographs showing loss of coherency in the interdendritic precipitatesize gradients. Panels (a) and (b) display Al-0.1Zr(b) aged at 425°C for 1,600 h. Loss of coherency occurs at R = 35 nm. Panels (c) and(d) display Al-0.2Zr aged at 425°C for 400 h. Loss of coherency occurs at R = 31 nm. Panels (e) and (f) display Al-0.1Zr-0.1Ti(a) agedat 375°C for 1,600 h. Loss of coherency occurs at R = 25 nm.


cipitate size exceeds a critical value, which typically corresponds to

n =1δ, (5.1)

where n is the number of atomic planes in the matrix and δ is the lattice parameter mismatch

between the two phases. The radius of the semicoherent precipitates indicated by the arrows

in Figures 5.15(b), 5.15(d), and 5.15(f) are 35 nm (for Al-0.1Zr(b) aged at 425°C for 1,600 h),

31 nm (for Al-0.2Zr aged at 425°C for 400 h), and 25 nm (for Al-0.1Zr-0.1Ti(a) aged at 375°C for

1,600 h), respectively. These precipitate radii are slightly larger than in Al3Sc precipitates, which

retain coherency to approximately 20 nm radius [43, 119, 335, 336]. This disparity may be un-

derstood considering the reported lattice parameters of the L12 Al3Zr and Al3Sc phases, which

are 4.08 A [79, 94, 105] and 4.103 A [103], respectively. A coherent Al3Zr (L12) precipitate has a

smaller mismatch (δ = +0.75%) than that of Al3Sc (δ = +1.33%) and, by Eq.(5.1), should therefore

retain coherency to larger radii. Marquis and Seidman [119] argued that, for an Al3Sc-α-Al het-

erophase interface Eq.(5.1) predicts n = 80 planes which, considering a 0.2 nm spacing between

{200} planes, corresponds to a distance of 16 nm, which is in good agreement with the observed

coherency loss for Al3Sc precipitates of 20 nm radii. Using similar arguments, a metastable L12

Al3Zr precipitate should be expected to remain coherent to n = 130 planes, corresponding to

radii of 26 nm. This is in reasonable agreement with the observed coherency loss in Figure 5.15.

As discussed by Royset and Ryum [43, 336] for Al3Sc (L12) precipitates, the critical radius for

coherency loss is actually temperature dependent. Since the thermal expansion is different for

the matrix and precipitate phases, coherency is expected to be maintained for larger precipitate

radii at higher temperatures. It is interesting that for the specimens aged at 425°C, Figures 5.15(b)

and 5.15(d), the observed loss of coherency occurs at larger radii (R = 35 nm and 31 nm, respec-

tively) than in the one aged at 375°C, Figure 5.15(f) (R =25 nm).

Rod-like and plate-like interdendritic L12 precipitates

A second precipitate morphology observed in the interdendritic channels is that of orthogonal

rods and plates oriented in 〈100〉matrix directions and having the metastable L12 structure, Fig-

ure 5.16. Both rod-like and plate-like morphologies are observed, as confirmed by tilting experi-

ments. A characteristic feature of these precipitates is the appearance of closely-spaced pairs of


Fig. 5.16: TEM micrographs of rod-shaped interdendritic Al3Zr or Al3(Zr1−xTix) precipitates, oriented in 〈100〉 matrix directions.Panel (a) is a centered dark-field image, [100] zone axis, of precipitate-rich dendrites in Al-0.1Zr-0.1Ti(a) aged at 375°C for 1,600 h.Panel (b) is a detailed view of closely-spaced pair of rod-like precipitates shown in Panel (a). Panel (c) is a bright-field image, [111]zone axis, of Al-0.1Zr(a) aged at 375°C for 1,600 h. Arrays of spheroidal precipitates are also visible. Panel (d) displays a centeredsuperlattice dark-field image of several rod-like and plate-like Al3(Zr1−xTix) (L12) precipitates. Panel (e) is a bright-field image,[110] zone axis, of Al-0.1Zr(b) aged at 425°C for 1,600 h. Panel (f) is a centered superlattice dark-field image of rod-like Al3Zr (L12)precipitates and arrays of spheroidal Al3Zr (L12) precipitates.


rods, Figure 5.16(b). These are identical to those described by Nes [79] in a hypoperitectic Al-Zr

alloy (0.05 at.% Zr) aged at 460°C. Ryum [85] observed similar orthogonal rods and plates during

the decomposition of Al-0.027Hf (at.%) alloys after aging at 450°C for 50 h.

In the present alloys, these rod-shaped precipitates are predominantly in the interdendritic

regions where the solute concentration is smallest. Heterogeneous nucleation dominates in

these regions, as discussed, and Nes [79] observed that precipitation of the rod-like and plate-

like precipitate morphologies is closely associated with matrix dislocations, proposing that a

precipitate growth/dislocation climb interaction is responsible for their appearance. Indeed, a

dislocation is seen protruding into the matrix from both ends of the rod-shaped precipitate in

the lower right of Figure 5.16(e).

Near to the rod- and plate-shaped precipitates are linear arrays of spheroidal interdendritic

L12 precipitates, Figures 5.16(c), 5.16(d), and 5.16(f). Nes [79] observed similar arrays of Al3Zr

(L12) precipitates which were described as “rods partly chopped up into rows of spherical parti-

cles.” As with the rods, the arrays of spheroidal precipitates were aligned along 〈100〉 directions.

The arrays observed presently seem to be aligned in particular orientations, Figures 5.16(c) and

5.16(f), but the orientation relationship with α-Al is not strict. Zedalis and Fine [84] observed

similar arrays of Al3Zr (L12) precipitates in Al-0.22Zr (at.%) alloys aged at 450°C, which they at-

tributed to coalescence of spheroidal precipitates.

Precipitation in supersaturated α-Al enveloping primary (properitectic) precipitates

In the more concentrated Al-0.2Zr and Al-0.2Zr-0.2Ti alloys, primary Al3Zr or Al3(Zr1−xTix) pre-

cipitates (Figures 5.2 and 5.3) act as heterogeneous nucleants during solidification of α-Al. Since

k0 > 1, α-Al solid-solution enveloping the primary phase is enriched in solute, as shown by EDS

in the halo surrounding the primary Al3(Zr1−xTix) precipitate in Figure 5.4. This phenomenon

is also observed in aged specimens of alloy Al-0.2Zr-0.2Ti, Figures 5.17 and 5.18, where solid-

state precipitation of Al3(Zr1−xTix) (L12) is observed in α-Al surrounding the petal-like primary

phase.

Figure 5.17 shows SEM micrographs of electropolished TEM foils of alloy Al-0.2Zr-0.2Ti aged

at 425°C for 400 h. The origin of the grain refinement observed in Figure 5.1 is apparent in

Figures 5.17(a) and 5.17(b), where most grains contain a petal-like primary Al3(Zr1−xTix) pre-

cipitate. Closer inspection of these primary phases, Figures 5.17(c)–5.17(f), demonstrate that


Fig. 5.17: SEM secondary electron images of primary Al3(Zr1−xTix) precipitates in alloy Al-0.2Zr-0.2Ti aged at 425°C for 400 h.Panels (a) and (b) display numerous petal-like primary precipitates, similar to those shown in Figure 5.3(a), and the accompanyinggrain refining effect due to heterogeneous nucleation of α-Al during solidification. Panels (c) through (f) reveal the solid-stateprecipitation of Al3(Zr1−xTix) (L12) in supersaturated α-Al surrounding the primary precipitates. Gradients in size and numberdensity vary of the solid-state precipitates varies with initial solute concentration, Panel (f), as discussed previously.


Fig.

5.18

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the α-Al solid-solution enveloping them is initially highly supersaturated, as evidenced by a

concentric layer of homogeneously-distributed nanometer-scale spheroidal Al3(Zr1−xTix) (L12)

precipitates surrounding the petal-like precipitate. Here again we observe similar gradients in

precipitate size and morphology as described in Figures 5.13 and 5.14. Consider Figure 5.17(f).

Near to the primary precipitate where the first α-Al solid-solution was nucleated during solidi-

fication, the solute supersaturation is greatest (k0 >1) and the precipitates are fine and homo-

geneously distributed in high number densities. The solute supersaturation decreases outward

from the primary precipitate periphery and the solid-state spheroidal precipitates are coarser

and in smaller number densities. These variations in precipitate morphologies are illustrated

also in the series of superlattice dark-field TEM micrographs in Figure 5.18.

As shown in Figure 5.18, the primary petal-like precipitates have the metastable L12 struc-

ture, as evidenced by their bright contrast when imaged with an L12 superlattice reflection. This

observation is not unexpected since the petal-like morphology is characteristic of the metastable

L12 structure [160], as discussed previously and also in Section 2.4.2. Figure 5.18 demonstrates

also that the primary precipitates and the spheroidal ones precipitated in solid-solution around

them share the same orientation relationship with the α-Al matrix, indicating that α-Al nucleates

epitaxially on the metastable L12 primary precipitate.

The primary precipitates themselves consist of fine dendritic filaments, Figures 5.17(f) and

5.18, reflective of growth instabilities during their growth in the melt. This observed morphology

is consistent with what has been observed for Al3Ti, Al3Zr, and Al3Hf metastable L12 primary

precipitates, as discussed in Section 2.4.2.

Small dendritic precipitates

The most prevalent precipitate morphology in terms of number density, and certainly the one

exerting the most influence on the observed mechanical properties, is the large population of

small, coherent, homogeneously-distributed spheroidal Al3Zr (L12) or Al3(Zr1−xTix) (L12) pre-

cipitates found in the precipitate-rich dendritic cells (e.g., Figure 5.6). An example of these pre-

cipitates is shown in Figure 5.18(d), where the mean radius, 〈R〉, of Al3(Zr1−xTix) (L12) pre-

cipitates is 6.3 ± 1.9 nm after aging at 425°C for 400 h. This is impressive coarsening resis-

tance indeed, but Al3(Zr1−xTix) (L12) precipitates have been shown [288] to exhibit dramatically-

reduced precipitate radii compared to binary Al3Zr (L12). How different are the precipitate radii


Fig. 5.19: Centered superlattice dark-field TEM micrographs of Al-0.1Zr(b) (Panels (a) and (b)) and Al-0.1Zr-0.1Ti(b) (Panels (c) and(d)) aged at 425°C for 1600 h. Images obtained with a low-index L12 superlattice reflection as indicated in the inset diffraction pat-terns. Panel (a) displays an interdendritic channel separating two precipitate-rich dendritic regions. Panel (b) demonstrates that,within the dendrites, the Al3Zr (L12) precipitates have a mean radius of 〈R〉 = 10.9 ± 1.9 nm. Panel (c) shows precipitate-rich den-dritic regions separated by a 4–5 µm wide precipitate-lean interdendritic channel in Al-0.1Zr-0.1Ti(b). L12-structured orthogonalrods or plates are observed in the channel. The dark spots are holes in the foil, remnant of large spheroidal interdendritic precipi-tates dislodged during sample preparation. Panel (d) displays Al3(Zr1−xTix) (L12) precipitates within the dendrites, 〈R〉 = 11.6 ±2.3 nm, showing no improvement in coarsening resistance as compared to Al3Zr (L12) (Panel (b)).

for Al3Zr (L12) under identical aging conditions? Such a comparison is illustrated in Figure 5.19,

which shows the precipitated microstructures formed in alloys Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b)

after aging at 425°C for 1,600 h.2 Despite the inhomogeneities occurring throughout both alloys,

within the dendrites themselves the alloys are indistinguishable: 〈R〉 = 10.9 ± 1.9 nm for Al3Zr

2It is more appropriate to use the dilute Al-0.1Zr and Al-0.1Zr-0.1Ti alloys for a direct comparison since both alloysare coarse-grained, Figure 5.1, and are free of primary precipitates.


Table 5.3: Mean precipitate radii, 〈R〉, observed in the dendrite centers from six different Al-Zr and Al-Zr-Ti (three of each) speci-mens after extended aging times at 425°C.

Alloy Aging treatment Mean precipitate radius, 〈R〉 (nm) Number of precipitates counted

Al-0.1Zr(b) 1,600 h at 425°C 10.9± 1.9 251

Al-0.1Zr-0.1Ti(b) 1,600 h at 425°C 11.6± 2.3 143

Al-0.2Zr 400 h at 425°C 5.1± 1.7 300

6.7± 1.7 310

Al-0.2Zr-0.2Ti 400 h at 425°C 6.3± 1.9 101

5.9± 1.9 117

and 〈R〉 = 11.6 ± 2.3 nm for Al3(Zr1−xTix). Table 5.3 summarizes results of these and other pre-

cipitate size analyses on more concentrated Al-0.2Zr and Al-0.2Zr-0.2Ti alloys, which indicate

also that the coarsening rates of Al3Zr and Al3(Zr1−xTix) precipitates in these alloys are identical

after extended aging times at 425°C (0.75Tm).

The similarity in dendritic precipitate size between the Al-Zr and Al-Zr-Ti alloys after extend-

ing aging times is also reflected in the comparable hardness of the Al-Zr and Al-Zr-Ti alloys in

Figures 5.7–5.10. At 400°C and below, the alloys exhibit identical precipitation hardening behav-

ior after extended aging times. While the Al-Zr-Ti alloys tend to be harder at shorter aging times

or at temperatures greater than 400°C, this effect may be interpreted in terms of solute super-

saturation influencing the incubation for nucleation and the volume fraction of precipitate-rich

dendritic regions, and is not an effect of precipitate radius or coarsening resistance. The effect

of Ti additions is discussed more thoroughly in Chapter 6, Section 6.5.5.

The data in Table 5.3 illustrate a second point. The two mean precipitate sizes reported for Al-

0.2Zr and Al-0.2Zr-0.2Ti represent different physical specimens from each alloy. In other words,

these mean radii represent the precipitate populations in two different dendrites. The vast in-

homogeneities in precipitate size, growth, and morphology (e.g., Figures 5.14 and 5.16) are not

observed in the dendrites, and the consistency in the observed mean radii suggests that the so-

lute supersaturation (and hence the critical radius for nucleation as well as the local volume

fraction) is more-or-less constant from one dendrite to the next. This contention is supported

also by the measured EDS linescans, Figure 5.5, which indicate that the solute concentration is

similar for all dendritic regions.


Transformation to the equilibrium D023 structure

Despite extended aging times (up to 1,600 h) at 425°C (0.75Tm), equilibrium D023-structured are

not observed by TEM. Indeed, as discussed in Chapter 6, Al3Zr (L12) is kinetically stable at high

homologous temperatures, with significant coarsening and transformation to the D023 structure

commencing at ca. 500°C (0.83Tm), much higher than the temperatures investigated presently.

It is important to point out, however, that the precipitates in Figure 5.18(d) exhibit structural

faults, characterized by sharp lines of no contrast parallel to {100} planes inside the L12 precip-

itates. The faults are attributed to antiphase boundaries (APBs) generated in the transition to

an imperfect tetragonal (D023) structure by the formation of an APB with a displacement vector

a/2〈110〉 on {100} type planes, and hence represent the early stages of the transformation to the

D023 structure. This L12 to D023 transformation is discussed in more detail in Chapter 6.

5.4 Discussion

5.4.1 Comparison to Previous Studies on Al-Zr-V/Ti Alloys

As described in Section 4.1 and the Introduction to this chapter, improving the coarsening resis-

tance of Al3Zr with ternary or quaternary additions of V and/or Ti has generated considerable

interest in the literature. These studies are summarized in Table 5.4, which indicates the al-

loy compositions, volume fractions and purported compositions3 of the L12 precipitates, and

their measured coarsening rates at 425°C. The measured precipitate radii of the L12-structured

Al3(Zr1−xMx) (M = V and/or Ti) precipitates as a function of aging time at 425°C are displayed in

Figure 5.20.

Anomalous behavior in arc-melted alloys

An immediately apparent observation in Figure 5.20 is that the studies on arc-melted alloys [84,

288, 337] (i.e., those containing 1 vol.% precipitates — see also Table 5.4) tend to exhibit larger

precipitates with faster coarsening rates than the subsequent studies on melt-spun alloys [113,

114,289,296–298,338]. Indeed, several researchers [113,289,297,298] noted that the precipitates

they studied in melt-spun Al-Zr-V and Al-Zr-Ti-V alloys were consistently smaller than those

3As pointed out in Section 4.1, with exception to reference [271], the precipitate compositions were not measureddirectly in the studies in Table 5.4.


Table

5.4:Rep

orted

coarsen

ing

rateso

fL12

Al3

Zr-b

asedp

recipitates

inA

l-Zr,A

l-Zr-T

i,and

Al-Z

r-Ti-V

alloysaged

at425°C

Stud

yA

lloyco

mp

s.(at.%)

Meth

od

Cu

bic

L12

ph

aseN

om

.vol.frac.

Pre-age

Rate

con

stant,k

Refs

Zr

Ti

V(%

)(m

3h

-1)

Stud

ieso

nA

l-Zr

alloys

Zed

alisan

dF

ine

(1986)0.2

4—

—A

rc-melted

Al3

Zr

1vo

l.%N

on

e2.4

9×

10−

26

[84,337]

Stud

ieso

nA

l-Zr-V

alloys

Zed

alisan

dF

ine

(1986) a0.0

3—

0.1

9A

rc-melted

Al3

(V0.8

75Zr0.1

25)

1vo

l.%N

on

e1.6

3×

10−

26

[84,337]0.0

7—

0.1

8A

rc-melted

Al3

(V0.7

25Zr0.2

75)

1vo

l.%500°C

1h

2.2

6×

10−

27

Ch

enetal.(1987–1990)

0.3

4—

0.9

8M

elt-spu

nA

l3(V

0.7

5Zr0.2

5)

5vo

l.%500°C

2.5h

1.0

3×

10−

28

[113,114]

Han

etal.(1995)0.6

5—

1.3

5M

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Fig. 5.20: Comparison of L12-structured Al3(Zr1−xMx) (M = V and/or Ti) average precipitate radius (〈R〉, in nm) with aging timeat 425°C. Open symbols represent studies by Fine and coworkers on Al-Zr [84,337], Al-Zr-V [84,113,114,337,338], and Al-Zr-Ti [288]alloys. Shaded symbols are studies by Lee et al.on Al-Zr-V [289] and Al-Zr-Ti-V alloys [289, 297, 298], and half-shaded symbols arefrom a study by Tsau and Chen on Al-Zr-Ti alloys [296]. Specific alloy compositions and reported coarsening rates indicated inTable 5.4.

observed previously by Fine and coworkers on Al-Zr, Al-Zr-V and Al-Zr-Ti alloys produced by

arc-melting. The effect of finite volume fraction of precipitates on coarsening rate, well-studied

theoretically and experimentally [36,339,340], is always increasing coarsening rate with increas-

ing volume fraction. The apparent violation of this rule further supported the belief that ternary

additions of V and/or Ti had dramatically improved the coarsening resistance of Al3Zr, in spite

of the higher volume fractions in the melt-spun studies.

Lee and colleagues [289, 297, 298] attributed the discrepancies in coarsening behavior be-

tween the initial arc-melted studies of Zedalis and Fine [84, 337] and those that followed to dif-

ferences in the pre-aging treatment prior to aging.4 They contended that “the pre-aging treat-

4This is one of the reasons why the specimens were not pre-aged in the present study. Pre-aging was also foundempirically to be detrimental to achievable hardness, Figure B.1.


ment has substantially decreased the amount of supersaturation and slowed down the growth

kinetics of the precipitates.”

A more probable explanation for the discrepancies in Figure 5.20 is that the initial studies

on Al-Zr and Al-Zr-V alloys by Zedalis and Fine [84,337] were biased by an inhomogeneous den-

dritic ditribution of precipitates. Indeed, Zedalis and Fine [84] acknowledged that “the density of

precipitates is not uniform but varies somewhat from place to place” in their alloys, and likened

the microstructure to that reported in a study by Izumi and Oelschlagel [76] who observed Al3Zr

(L12) precipitates “statistically distributed within the dendritic cells of the cast structure.” Fur-

thermore, Zedalis and Fine [84] themselves questioned the validity of their measured coarsening

rates, stating that their calculated value interfacial free energy, σ, driving coarsening “is a hun-

dred times larger or more than expected,” and suspected that “the nonuniform distribution of

precipitates may be playing a role” for this anomalous result.

The (i) anomalously fast coarsening and (ii) anomalously large precipitates observed in the

initial studies by Zedalis and Fine [84, 337] can be easily explained within the context of a den-

dritic distribution of precipitates. Consider the TEM micrographs in Figure 5.21. What is the

mean radius of the Al3Zr precipitates? The most apparent population is the ca. 20 nm spheroidal

Al3Zr (L12) precipitates observed in the interdendritic regions of the specimen, where the foil is

thinnest. The most numerous population is the ca. 5 nm precipitates in the dendrites, which

are not obvious because of their smaller size and the much-reduced contrast from absorption

(incoherent scattering) occurring in the thicker dendritic regions. The precipitate radii Zedalis

and Fine measured, 〈R〉 = 15–20 nm (Figure 5.20), suggest that the precipitates they studied were

interdendritic.5 These tend to be larger, as discussed above, because of an initially smaller so-

lute supersaturation leading to a larger critical radius for nucleation. Moreover, as suggested by

Robson [341], their smaller number density allows them to grow to larger radii before exhausting

their surrounding α-Al solid-solutions of solute, further accounting for their increased size (and

also enhancing the observed gradients in their radius, Figure 5.13). In other words, the interden-

dritic precipitate size is evolving by a growth mechanism rather than a coarsening one, which

would explain the anomalously large radii and anamolously fast coarsening rates observed in

reference [84].

Why are the precipitates in the melt-spun alloys consistently smaller than those observed

5While this theory may seem far-fetch, consider also the study in Appendix D on an alloy prepared byParameswaran et al. [288].


Fig. 5.21: Montage of centered superlattice dark-field TEM micrographs of Al-0.2Zr aged at 425°C for 400 h. The apparent meanradius, 〈R〉, of spheroidal Al3Zr (L12) precipitates is strongly dependent on location in the specimen.

in the alloys produced by arc-melting? As discussed in Section 1.1.1, Al has particular limita-

tions in its capacity to form extensive solid-solutions with all but a few neighboring elements,

so the extension of solid-solubility by RSP will lead to immediately smaller precipitates be-

cause of the larger achievable supersaturation of solutes. A second effect of rapid solidifica-

tion is the stabilization of plane-front solidification, thus minimizing or eliminating microseg-

regation [16, 342, 343]. Because alloys produced by RSP exhibit little or no microsegregation of


solutes, the precipitates that form are not subject to the various inhomogeneities in size and

morphology associated with dendritically-distributed precipitates in arc-melted alloys. The ap-

parently improved coarsening resistance, despite the higher volume fraction of precipitates in

the melt-spun alloys, is also not surprising. Figure 5.6, however, demonstrates that the ap-

parent coarsening kinetics of precipitates in the densely-populated dendrites are much slower

than those in regions of smaller volume fraction, thus seemingly violating volume diffusion-

controlled coarsening theory. Al3Zr precipitates, despite their high number densities in the den-

drites, exhibit impressive stability at 425°C (Table 5.3).

Precipitate compositions

Finally, consider the reported precipitate compositions in Table 5.4; in all cases the implicit as-

sumption is that the ratio of solutes in the precipitated trialuminides is the same as that initially

in solid-solution (i.e., the bulk alloy composition). This is at first glance a reasonable assumption

since the transition elements are generally sparingly soluble in α-Al and all solute atoms should

therefore partition to the precipitated phase, as argued by Lee and colleagues [297,298,344]. It is

shown in the present study, Chapter 4, that for Al-Zr-Ti alloys having the same atomic fraction of

solute species, this Zr:Ti ratio of unity is never observed in the precipitated Al3(Zr1−xTix) phase.

Moreover, as evidenced by Table 4.2, the assumption of negligible solubility of solutes in α-Al is

invalid; diffusion kinetics in α-Al, which is anomalously slow and varies widely among the tran-

sition elements (Section 1.2.3), is instead the dominant factor dictating the relative amounts of

solute atoms in the precipitated phase. Both Ti and V are slower diffusers than Zr in α-Al, Fig-

ure 1.7 (especially true of V), and the precipitate compositions in Table 5.4 are therefore highly

speculative.

Summary

Discounting the anomalous results from the studies on arc-melted alloys, the various

Al3(Zr1−xVx) (L12), Al3(Zr1−xTix) (L12), and Al3(Zr1−x−y,Vx,Tiy) (L12) precipitates in Figure 5.20

have radii 〈R〉 = 3–7 nm after 400 h at 425°C, which are comparable to the radii of the Al3Zr (L12)

and Al3(Zr1−xTix) (L12) dendritic precipitates in the present Al-Zr and Al-Zr-Ti alloys after 400

and 1,600 h at 425°C which range in size from 〈R〉 = 5–7 nm after 400 h and 〈R〉 = 11–12 nm

after 1,600 h (Table 5.3). Binary Al3Zr (L12) precipitates are more coarsening resistant than the


data of Zedalis and Fine [84] indicates, and the apparent improvement by V and/or Ti additions

claimed in the studies of Table 5.4 are probably artifacts of the non-uniform precipitate mor-

phologies associated with an inhomogenous dendritic distribution of precipitates in the initial

studies on arc-melted alloys. As shown presently, ternary additions of solute have little effect on

the coarsening behavior at 425°C (0.75Tm).

5.5 Chapter Summary

Precipitation of Al3Zr (L12) and Al3(Zr1−xTix) (L12) was investigated in conventionally-solidified

Al-0.1Zr, Al-0.2Zr, Al-0.1Zr-0.1Ti, and Al-0.2Zr-0.2Ti (at.%) alloys isothermally aged at 375, 400,

or 425°C. The following results were obtained and discussed:

• Dendritic solidification results in non-uniform distributions of Zr and Ti, which segregate

to the dendrite centers and are enriched by a factor of 2 and 3, respectively (Figure 5.5). The

interdendritic regions are concurrently solute-depleted. During subsequent aging, Al3Zr

Al3(Zr1−xTix) precipitates are similarly non-uniformly distributed in a dendritic manner

(Figure 5.6).

• Pronounced hardening accompanies precipitation of Al3Zr or Al3(Zr1−xTix) at 375, 400, or

425°C for all alloys investigated (Figures 5.7–5.10). The magnitude of the peak-aged hard-

ness is controlled primarily by the Zr concentration. The alloys overage sluggishly, even

after 3,200 h at 425°C, with no obvious benefit from Ti additions in delaying overaging. For

shorter aging times (t < 10 h), Ti accelerates the nucleation kinetics, and for higher temper-

atures (425°C), the peak hardness is increased. These are effects of solute supersaturation

(incubation time and precipitate volume fraction) and is not an indication of coarsening

resistance.

• The precipitated microstructures in the alloys are complicated by the initial segregation

of solute species; because of the locally-varying solute concentration, different precipitate

morphologies are observed. These include: (i) larger interdendritic precipitates, frequently

exhibiting growth instabilities (Figure 5.13); (ii) rods and plates aligned in 〈100〉 matrix di-

rections (Figure 5.16); and (iii) small (< 10 nm) spheroidal L12 precipitates in the centers


of the dendritic cells (Figure 5.19). Even after aging at 425°C (0.75Tm) for 1,600 h, the equi-

librium D023 structure (Al3Zr or Al3(Zr1−xTix)) is not observed.

• The original motivation for adding Ti was to form Al3(Zr1−xTix) precipitates, that exhibit

a reduced lattice parameter mismatch with α-Al, thereby improving the coarsening resis-

tance of the Ti-containing precipitates. Despite the confirmed partitioning of Ti to the

Al3(Zr1−xTix) precipitates by 3-D atom-probe tomography (Chapter 4), there is no bene-

fit in terms of coarsening resistance of the small dendritic Al3(Zr1−xTix) (L12) precipitates

compared to Al3Zr (L12) during extended isothermal aging at 425°C (Table 5.3), consistent

also with the similar overaging behavior observed by Vickers microhardness (Figures 5.7–

5.10).

• The apparent improvement of the coarsening resistance of Al3Zr (L12) by V and/or Ti ad-

ditions claimed in the studies of Table 5.4 may be artifacts of a non-uniform dendritic dis-

tribution of precipitates in the initial studies on arc-melted Al-Zr alloys. Metastable Al3Zr

(L12) precipitates are intrinsically coarsening resistant at 425°C (0.75Tm).

CHAPTER 6

Microstructural Coarsening in Al-Zr andAl-Zr-Ti Alloys, T ≥ 450°C

The transformation of Al3Zr (L12) and Al3(Zr1−xTix) (L12) precipitates to their respec-

tive equilibrium D023 structures is investigated in conventionally-solidified Al-0.1Zr and

Al-0.1Zr-0.1Ti (at.%) alloys isothermally aged at 500°C or isochronally aged 300–600°C.

There is no benefit from Ti additions, both in terms of coarsening resistance of the

metastable L12 precipitates, or in delaying the L12 to D023 transformation. Both alloys

overage at the same rate at or above 500°C, during which spheroidal L12 precipitates

transform to disc-like D023 precipitates ca. 200 nm in diameter and 50 nm in thickness,

exhibiting a cube-on-cube orientation relationship with α-Al. The transformation oc-

curs heterogeneously on dislocations because of a large lattice parameter mismatch of

the equilibrium D023 phase with α-Al. The transformation is extraordinarily sluggish;

even at 575°C, coherent L12 precipitates are also observed. Mechanisms of microstruc-

tural coarsening and strengthening are discussed with respect to the micrometer-scale

dendritic distribution of precipitates.

6.1 Introduction

CHAPTER 5 DISCUSSED the microstructure and ambient temperature mechanical properties

of Al-Zr and Al-Zr-Ti alloys during isothermal aging at 375, 400, and 425°C. Despite ex-

tended aging times (3,200 h) at 425°C (0.75Tm of Al), the alloys exhibit no appreciable overaging

and the precipitates do not transform to the equilibrium D023 structure. This chapter investi-

gates similar alloys during exposure at higher temperatures: (i) isothermally at 500°C (0.83Tm);

and isochronally at 300–600°C (0.94Tm). Microstructural coarsening and strengthening mecha-

nisms, with respect to the micrometer-scale dendritic distribution of precipitates, are developed

and the L12 to D023 transformation is discussed.

131

132 CHAPTER 6 MICROSTRUCTURAL COARSENING IN AL-ZR AND AL-ZR-TI ALLOYS, T ≥ 450°C

Table 6.1: Compositions and aging conditions of the Al-Zr and Al-Zr-Ti alloys investigated.

Nominal comp. (at.%) Verified comp. (at.%) Aging temperatures (°C)Zr Ti Zr Ti

Al-0.1Zr(b) 0.1 — 0.085 — 500

Al-0.1Zr(c) 0.1 — 0.101 — Isochronal aging (300–600)

Al-0.1Zr-0.1Ti(b) 0.1 0.1 0.089 0.099 500

Al-0.1Zr-0.1Ti(c) 0.1 0.1 0.095 0.091 Isochronal aging (300–600)


A series of binary Al-Zr and ternary Al-Zr-Ti alloys were investigated; alloy designations, com-

positions, and aging conditions are summarized in Table 6.1. The alloys were prepared employ-

ing non-consumable electrode arc-melting, as described in Sections 3.2 and 4.2. The verified

compositions in Table 6.1 were obtained by bulk chemical analysis performed by direct current

plasma emission spectroscopy (ATI Wah Chang, Albany, OR).

To investigate the overaging behavior isothermally, peak-aged specimens (aged isothermally

at 375°C for 100 h to ca. 450 MPa, Figures 5.7 and 5.9) of Al-0.1Zr(b) and Al-0.1Zr-0.1ZTi(b) were

overaged isothermally at 500°C. The microstructures obtained after aging for 100 h at 500°C were

investigated in both alloys by transmission electron microscopy (TEM) and scanning electron

microscopy (SEM), using similar instruments and techniques described in Section 3.2.

A second set of alloys, Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c), were aged in a series of multi-step 3 h

isochronal anneals beginning at 300°C and terminating at 600°C, in 25°C increments. Between

each aging step, the specimens were water-quenched and precipitation of Al3Zr or Al3(Zr1−xTix)

was monitored by Vickers microhardness and electrical conductivity measurements. The con-

ductivity measurements were performed using a SIGMATEST 2.069 (Foerster Instruments, Pitts-

burgh, PA) eddy current apparatus at room temperature. Five measurements were recorded,

each corresponding to a different frequency (60, 120, 240, 480, 960 kHz), on each specimen. For

consistency, a single specimen of each alloy was used for conductivity measurements, which

was measured between each isochronal aging treatment. The microstructures obtained after

isochronal aging to 450, 525, or 575°C were investigated in both alloys by TEM.

SECTION 6.3 EXPERIMENTAL RESULTS: ISOTHERMAL AGING AT 500°C 133

1 1 0 1 0 0 1 0 0 0

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6.3 Experimental Results: Isothermal Aging at 500°C

6.3.1 Vickers microhardness

Figure 6.1 displays Vickers microhardness as a function of aging time at 500°C, indicating that

both Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b) alloys overage at the same rate with no improved coarsen-

ing resistance of Al3(Zr1−xTix) (L12) as compared to Al3Zr(L12). These results are consistent with

those of Chapter 5, where the precipitate radii of Al3Zr(L12) and Al3(Zr1−xTix) (L12) precipitates

were statistically identical after extended aging times at 425°C (Table 5.3).

6.3.2 Electron microscopies

Precipitate dissolution and transformation to the equilibrium D023 structure

The overaging behavior exhibited in Figure 6.1 was investigated directly by TEM and SEM, and

is due to dissolution of the small, high number density, dendritic spheroidal Al3Zr (L12) or

Al3(Zr1−xTix) (L12) precipitates which then re-precipitate on grain boundaries or dislocations,

ultimately forming extensive arrays of heterogeneously-nucleated equilibrium D023-structured


Fig. 6.2: Montage of bright-field TEM micrographs of Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h).Imaged with ~g = (020).


Fig. 6.3: Montage of centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at375°C for 100 h). Imaged with the ~g = (010) superlattice reflection.


Fig. 6.4: SEM micrograph of heterogeneously-nucleated Al3Zr (D023) precipitates on dislocations or small-angle grain boundariesin Al-0.1Zr(b) aged at 500°C for 100 h (after aging at 375°C for 100 h). The morphology of the D023 phase is disc-like, 150–200 nm indiameter and 50–60 nm in thickness, and exhibit one of three possible orientation relationships with the α-Al solid-solution.

precipitates. Figures 6.2 and 6.3 display complementary bright-field and dark-field micrographs,

respectively, of Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h). Sev-

eral dislocations are shown, on which numerous disc-like, ca. 150–200 nm diameter and 50 nm

thick, equilibrium D023-structured precipitates have formed (precipitates A in Figures 6.2 and

6.3). The dislocations and linear arrays of precipitates are most apparent in the bright-field mi-

crographs, Figure 6.2.

The arrays of heterogeneously-nucleated D023 precipitates extend for hundreds of microm-

eters and are observable also by SEM, Figure 6.4, which shows similar arrays of disc-shaped


[010]Al

[001]Al

I: c-axis || [100]Al

[100]Al

II: c-axis || [010]Al III: c-axis || [001]Al

c-axisc-axisc-axis

Fig. 6.5: Schematic of the three possible orientation relationships with α-Al for the disc-like equilibrium D023 precipitates. Adaptedfrom Chen [338].

precipitates along several dislocations extending over 50 µm in length. It initially seems unlikely

that such a long dislocation length would randomly exist exactly parallel to the specimen surface

in Figure 6.4 (the probably of observing a dislocation by TEM is much larger since a 3-D volume,

ca. 100 nm thick, is observable). For a metal at room temperature, the dislocation density is of

the order of 1010 m−2 [345], or one dislocation per 10×10 µm2 area. Thus, the area represented

in Figure 6.4(a) can be expected to contain 18 dislocations — many more than what is visually

apparent — intersecting the specimen at various orientations, with one of them parallel to the

specimen surface.

The disc-like D023 precipitates exist in one of three orientations: two orthogonal edge-on

views and one face-on when viewed along [100]Al, indicating a cube-on-cube orientation rela-

tionship with the α-Al solid-solution, shown schematically in Figure 6.5. Only precipitates with

certain orientation relationships, those whose c-axis is parallel to [010]Al, are illuminated in dark-

field when imaged with ~g = (010) in Figure 6.3; this indicates that the precipitates have the D023

structure, as discussed in more detail below.

Transformation mechanism to the equilibrium D023 structure

The equilibrium D023-structured precipitates are not nucleated directly from supersaturated α-

Al solid-solution during aging, but rather form by dissolution and subsequent re-precipitation of

solutes from nearby pre-existing L12 precipitates formed during the prior isothermal aging treat-

ment at 375°C. The consumption of the pre-existing L12 precipitates is unambiguously shown

in the bottom-most precipitate-rich dendrite in Figure 6.3, which initially contained a relatively


high number density of homogenenously-distributed spheroidal L12 Al3(Zr1−xTix) precipitates

in the peak-aged condition (100 h at 375°C). During aging at 500°C, the equilibrium D023 pre-

cipitates form heterogeneously along a dislocation cutting through the dendrite, leaving behind

a precipitate-free cylinder approximately 500 nm in radius around the dislocation (precipitates

A in Figures 6.2 and 6.3). A similar precipitate-free zone extends 1.5 µm to either side of the

heterogeneously-nucleated array of disc-like Al3Zr precipitates in Figure 6.4(c). The widths of

these precipitate-free zones is comparable to the calculated1 root-mean-squared diffusion dis-

tance,√

6Dt, of Zr or Ti after 100 h at 500°C, which is 2.6 or 0.8 µm, respectively.

Arrays of spheroidal L12 precipitates, labeled B in Figures 6.2 and 6.3, are also nucleated

heterogeneously along the dislocation. These precipitates are surrounded by a narrower, 100–

400 nm wide, precipitate-depleted zone, suggesting that the metastable L12 structure repre-

sents an intermediate stage of microstructural coarsening. Figure 6.6 demonstrates that many of

the heterogeneously-nucleated L12 precipitates exhibit preferential growth in 〈100〉 directions,

forming elongated cigar-shaped precipitates. They also contain antiphase boundaries (APBs),

characterized by sharp lines of no-contrast parallel to {100} planes inside the L12 precipitates,

and represent a transition to the equilibrium D023 structure as discussed below.

The precipitates labeled C, most obvious in Figure 6.2, are other examples of the transfor-

mation to the equilibrium D023 structure occurring heterogeneously on dislocations cutting

through a precipitate-rich dendrite, leaving behind precipitate-depleted zones surrounding the

dislocation. Finally, regions marked D in Figures 6.2 and 6.3 indicate a second mechanism of

microstructural coarsening. Here, larger spheroidal L12-structured precipitates at the dendrite

edge have coarsened at the expense of the smaller precipitates immediately on the interior side

of the dendrites, leaving behind a narrow, ca. 100–200 nm, precipitate-free zone on either side

of the ca. 1 µm wide interdendritic channel. This reduction in volume fraction of the small den-

dritic precipitates also contributes to the drop in hardness during exposure at 500°C, Figure 6.1.

Structure and morphology of the D023 precipitates

Figure 6.7 displays a bright-field TEM micrograph of three orthogonal D023-structured precipi-

tates viewed along [100]Al, exhibiting the three orientation relationships depicted schematically

in Figure 6.5. Selected area diffraction patterns (SADPs) from two of the precipitates are also

1Using diffusivities reported in Table 1.3.


Fig. 6.6: Complementary bright-field and dark-field TEM micrographs of heterogeneously-nucleated L12- and D023-structuredAl3(Zr1−xTix) precipitates on dislocations in Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h). Many of theL12-structured precipitates exhibit attributed to antiphase boundaries (APBs) along 〈100〉 directions, indicating transformation tothe equilibrium D023 structure.

shown in Figure 6.7, which agree with kinematically-simulated diffraction patterns in Figure 6.9,

indicating that the precipitates in Figure 6.7 have the equilibrium D023 structure of Al3Zr.

The disc-like morphology of the equilibrium Al3Zr precipitates reflects the tetragonal sym-

metry of the D023 unit cell. The lattice parameters of Al3Zr (D023) are a = b = 4.014 and c = 17.321 A,

resulting in lattice parameter mismatches with α-Al (a = 4.0496 A) of δa = δb = -0.88% and δc = 6.92%

(Table 1.1). Since the lattice parameter mismatch is most severe in the c-direction, the energetic

preference is for precipitates that are thin along c, explaining the disc-like morphology. The

large overall lattice parameter mismatch with α-Al, δ = 2.89% (Table 1.1), also explains why the


Fig. 6.7: Bright-field TEM micrograph of disc-like D023 precipitates in three possible orientations and corresponding selected areadiffraction patterns (SADPs) of the equilibrium D023 phase in Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for100 h).The patterns are indexed in Figure 6.9.

equilibrium phase is only observed heterogeneously-nucleated along dislocations.

Orientation relationships and diffraction contrast of the D023 precipitates

The disc-like D023-strcutured precipitates exhibit a cube-on-cube orientation relationship with

α-Al, with the three possible orientations, designated I, II, or III as viewed along [100]Al (Fig-

ure 6.5). The equilibrium precipitates are observed only on dislocations, and appear to be at

least partially coherent with α-Al, as evidenced by strong strain contrast when imaged under

two-beam conditions, Figure 6.8(a). Because of their large size (ca. 200 nm in diameter), the

disc-like precipitates are most likely semicoherent with α-Al, which is supported also by the ap-

pearance of striped lines of contrast normal to ~g in the images of many of the precipitates when

viewed face-on (i.e., those in orientation I). These are interpreted as contrast due to interfacial

dislocations, indicating semicoherency with the α-Al matrix.

The dark-field micrographs in Figure 6.8 illustrate a second diffraction contrast phenomenon.

As observed previously (e.g. in Figures 6.3 and 6.6), and shown in Figure 6.8(b), only the disc-

like precipitates whose c-axes are parallel to [010]Al (orientation II) are illuminated when imaged

with the~g = (010) L12 superlattice reflection. Similarly, only precipitates whose c-axes are parallel

to [100]Al (orientation I) are illuminated when imaged with ~g = (011), Figure 6.8(d). These diffrac-


Fig. 6.8: Complementary bright-field and dark-field TEM micrographs of heterogeneously-nucleated D023 precipitates on dislo-cations in Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h). The strong contrast and presence of interfacialdislocations in bright-field indicates that the disc-like precipitates are semicoherent with α-Al. In dark-field, only certain orienta-tions are illuminated, depending on the g vector that is operating.

tion contrast phenomena may be understood considering the calculated diffraction patterns for

the D023 and L12 structures of Al3Zr in Figure 6.9. For the D023 precipitates in orientation II,

the 004T reflections nearly coincide with the 010C reflections from the L12 precipitates, explain-

ing why only those tetragonal precipitates in orientation II are illuminated in dark-field when

imaged with the ~g = (010) L12 superlattice reflection as in Figure 6.3 and 6.8(b). With ~g = (011)

operating, Figure 6.8(d), only the disc-like precipitates in orientation I are illuminated, since the

011T reflections from these precipitates overlap with the 011C superlattice reflections from the

L12 precipitates.


II: c-axis || [010]Al III: c-axis || [001]Al

011C

012C

001C

021C

010C

008T

004T

011T

013T

015T

017T

019T

024T

028T

020T

000 020Al

022Al

002Al

220T

011C

012C

001C

021C

010C

020T

110T

200T

000 020Al

022Al

002Al

011C

012C

001C

021C

010C

020T

004T

011T013

T015

T017

T019

T

008T

024T

028T

000 020Al

022Al

002Al

I: c-axis || [100]Al

Fig. 6.9: Calculated electron diffraction patterns from the equilibrium tetragonal D023 (closed circles, T subscript) and metastablecubic L12 (crosses, C subscript) Al3Zr phases in three possible orientation relationships with α-Al (open circles). The spots notindexed (e.g. 002T, 006T) are kinematically-forbidden.

Small dendritic precipitates

Figure 6.10 shows that the precipitates in both Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b) are L12-structured,

coherent with α-Al, and exhibit impressive coarsening resistance at 500°C. The mean radii of the

dendritic precipitates in Figures 6.10(a) and 6.10(c) are displayed in Table 6.2; after aging at 500°C

for 100 h (preceded by aging at 375°C for 100 h), the mean radii of the Al3Zr and Al3(Zr1−xTix)

precipitates are statistically identical: 12.3± 2.9 and 12.0± 4.0 nm, respectively, which is consis-

tent also with the identical drop in hardness exhibited by both alloys during exposure at 500°C,

Figure 6.1.

Figure 6.11 shows that the larger dendritic precipitates (R & 20 nm) exhibit planar faults,

indicating a partial transition to the D023 structure. Nevertheless, as shown in Figure 6.10, the

vast majority of the dendritic precipitates are L12-structured and fully coherent, indicating that

the L12 to D023 transformation does not occur spontaneously even after 100 h at 500°C. The

transformation instead requires pre-exisitng defects and other heterogeneous nucleation sites.


Fig. 6.10: Complementary dark-field and bright-field TEM micrographs of dendritic precipitates in Al-0.1Zr(b) (panels a and b) andAl-0.1Zr-0.1Ti(b) (panels c and d) aged at 500°C for 100 h (after aging at 375°C for 100 h). The Al3Zr or Al3(Zr1−xTix) precipitateshave the metastable L12 structure and are coherent with the α-Al solid-solution, with a mean radius ca. 12 nm.

Table 6.2: Mean precipitate radii, 〈R〉, of precipitates observed in the dendrite centers of Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b) aged at500°C for 100 h (after aging at 375°C for 100 h).

Alloy Aging treatment Mean precipitate radius, 〈R〉 (nm) Number of precipitates

Al-0.1Zr(b) 500°C for 100 h (after 375°C for 100 h) 12.3± 2.9 206

Al-0.1Zr-0.1Ti(b) 500°C for 100 h (after 375°C for 100 h) 12.0± 4.0 168


Fig. 6.11: Centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for100 h). Many of the spheroidal L12-structured precipitates exhibit APBs along 〈100〉 directions, indicating a partial transformationto the equilibrium D023 structure.

6.4 Experimental Results: Isochronal Aging, 300–600°C

6.4.1 Vickers microhardness and electrical conductivity

Figure 6.12 indicates the precipitation behavior of Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c) during 3 h

isochronal aging, as monitored by Vickers microhardness and electrical conductivity. Precipi-

tation of Al3Zr commences between 350 and 375°C in the binary alloy Al-0.1Zr(c), as evidenced

by the increase in strength and the accompanying change in electrical conductivity. For the

Ti-containing alloy, Al-0.1Zr-0.Ti(c), precipitation of Al3(Zr1−xTix) begins at temperatures lower

by about 25°C, consistent with the accelerated isothermal nucleation kinetics (shorter incuba-

tion time) for Al-0.1Zr-0.1Ti(b), Figure 5.9, compared with Al-0.1Zr(b), Figure 5.7. The maxi-

mum strength achieved during isochronal aging is also slightly greater (ca. 25 MPa) in the Ti-

containing alloy indicating a finer population, or larger volume fraction, of Al3(Zr1−xTix) pre-

cipitates.

The amplitude of the change in conductivity between as-cast and peak-aged condition is,

however, less for Al-0.1Zr-0.1Ti(c) (2.7 MS m-1) than it is for Al-0.1Zr(c) (3.7 MS m-1), suggesting

that less solute is precipitated from solid-solution in the ternary alloy despite having twice the

initial total solute concentration. The most obvious effect of Ti additions is the reduced con-

ductivity initially in Al-0.1Zr-0.1Ti(c) compared to that of Al-0.1Zr(c), which might explain the

reduced absolute change in conductivity for the ternary alloy. Barghout et al. [346] showed that

SECTION 6.4 EXPERIMENTAL RESULTS: ISOCHRONAL AGING, 300–600°C 145

3 0 0 4 0 0 5 0 0 6 0 0

2 6

2 7

2 8

2 9

3 0

3 1

3 2

3 3

3 4

2 0 0

2 5 0

3 0 0

3 5 0

4 0 0

4 5 0

5 0 0

3 0 0 4 0 0 5 0 0 6 0 0

Ele

ctr

ica

l c

on

du

cti

vit

y

(MS

m-1

)

��

A l - 0 . 1 Z r - 0 . 1 T i ( c )

A l - 0 . 1 Z r ( c )

A s - c a s t

A l - 0 . 1 Z r - 0 . 1 T i ( c )

Vic

ke

rs

mic

ro

ha

rd

ne

ss

(M

Pa

)

A l - 0 . 1 Z r ( c )

��

Fig. 6.12: Vickers microhardness and electrical conductivity evolution during isochronal aging (3 h at each temperature) of Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c).

the presence of second phase precipitates reduces the electrical conductivity in Al-Fe and Al-Mg

alloys. The smaller change in conductivity for Al-0.1Zr-0.1Ti(c) may therefore indicate a greater

volume fraction of Al3(Zr1−xTix) precipitates compared to Al3Zr precipitates in Al-0.1Zr(c), con-

sistent with the increased hardness of the ternary alloy over the binary.

The Vickers microhardness and electrical conductivity data in Figure 6.12 indicate that Ti

additions have a slight effect on the hardness of the alloys in the peak-aged condition, and a

negligible influence on the total amount of solute precipitated from solid-solution. These obser-

vations are consistent with the results in Chapter 4, where only a small amount of Ti is precipi-

tated from α-Al solid-solution in Al-Zr-Ti alloys, and also the hardness data in Figures 5.7–5.10,

where the effect of Ti is small compared to that of Zr in dictated the hardness of the alloys during


Fig. 6.13: Bright-field and centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(c) isochronally aged to 450°C. (a)Precipitate-rich dendrite are clearly defined. Contrast in bright-field is due to thickness contrast. (b) Dark-field micrograph dis-playing several precipitate-rich dendrites. (c) Arrays of coalescing precipitates near a dendrite edge. (d) Dendritic L12 precipitateswith a mean radius 〈R〉 = 3.4±0.5 nm.

isothermal aging.

6.4.2 Transmission electron microscopy

Microstructural coarsening of the alloys, which occurs for T > 450°C in Figure 6.12, was investi-

gated directly by TEM in both alloys after isochronal aging to 450, 525, and 575°C.

Figure 6.13 displays the microstructure in Al-0.1Zr-0.1Ti(c) isochronally aged to 450°C, which

corresponds to the onset of coarsening in Figure 6.12. Distinct, well-defined precipitate-rich

dendrites are observed in Figures 6.13(a) and 6.13(b); the contrast observed in bright-field, Fig-


Fig. 6.14: Centered superlattice dark-field TEM micrographs of Al-0.1Zr-0.1Ti(c) isochronally aged to 525°C. (a) Precipitate-richdendrite (outline dotted). D023-structured precipitates are nucleated heterogeneously on a dislocation, leaving behind a ca. 400 nmwide precipitate-free zone. (b) Magnified region from panel (a). (c) Arrays of coalescing precipitates in an earlier stage of coarsening.(d) Dendritic L12 precipitates with a mean radius 〈R〉 = 6.6±1.9 nm.

ure 6.13(a), is from thickness differences arising from preferential thinning of the interdendritic

channels during electropolishing, discussed also in Appendix C. Figure 6.13(c) displays precip-

itates near a dendrite edge, where precipitate coalescence along ca. 100 nm linear arrays are

observed within the precipitate-rich dendrite. It is believed that these precipitates are coalesc-

ing on dislocations, ultimately leading to vast arrays of heterogeneously-nucleated D023 precip-

itates during the latter stages of coarsening. Figure 6.13(d) shows that within the dendrites, the

precipitates are small and spheroidal, with a mean radius 〈R〉 = 3.4± 0.5 nm.

Figure 6.14 shows the microstructure in Al-0.1Zr-0.1Ti(c) isochronally aged to 525°C, which


Fig. 6.15: Complementary bright-field and dark-field TEM micrographs of Al-0.1Zr-0.1Ti(c) isochronally aged to 575°C. Most of themetastable L12 precipitates are transformed to the disc-like D023 structure.

corresponds to an intermediate stage of coarsening. Figures 6.14(a) and 6.14(b) display disc-like

D023 precipitates heterogenenously nucleated on a dislocation cutting through a precipitate-

rich dendrite. A surrounding precipitate-free cylinder, up to 400 nm in radius, is formed as a

result of the D023 formation. The calculated cumulative root-mean-squared diffusion distance,√

6Dt, of Zr or Ti atoms isochronally aged to 525°C is 1.8 or 0.5 µm, respectively, which is com-

parable to the radius of the precipitate-free cylinder in Figure 6.14(b); this is consistent with the

mechanism of microstructural coarsening described previously in Figures 6.3 and 6.4. Other re-

gions of the alloy, Figure 6.14(c), exhibit earlier stages of microstructural coarsening with smaller

precipitate-free zones surrounding other arrays of coalescing precipitates. Within the dendrites,

most precipitates are spheroidal and have the L12 structure, with a mean radius of 6.6± 1.9 nm.

By 575°C, the majority of the precipitates have transformed to the equilibrium D023 structure,

demonstrated in Figure 6.15. The disc-like D023 precipitates exhibit a cube-on-cube orientation

relationship with α-Al, as discussed previously in Figure 6.5. There are, however, other regions of

the alloy containing spheroidal L12 precipitates, indicated in Figure 6.15(b). These precipitates


Fig. 6.16: Complementary dark-field and bright-field TEM micrographs of dendritic precipitates in Al-0.1Zr(c) (panels a and b)and Al-0.1Zr-0.1Ti(c) (panels c and d) isochronally aged to 575°C. The Al3Zr or Al3(Zr1−xTix) precipitates have the metastable L12

structure and are coherent with the α-Al solid-solution, with a mean radius ca. 16.5 nm.

exhibit impressive coarsening resistance despite exposure at 575°C (0.91Tm), with a mean radius

of 16.3± 3.4 nm.

Figure 6.16 displays complementary dark-field and bright-field TEM micrographs of den-

dritic L12 precipitates in Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c) isochronally aged to 575°C. The pre-

cipitates have the same radius, 16.9 ± 4.4 nm and 16.3 ± 3.4 nm, respectively, demonstrating

that there is no difference in coarsening of spheroidal Al3Zr(L12) and Al3(Zr1−xTix) (L12) precip-

itates at 575°C. Moreover, the dendritic L12 precipitates are coherent with α-Al, demonstrated by

the Ashby-Brown strain contrast exhibited in the bright-field micrographs in Figure 6.16. Con-


Table 6.3: Mean precipitate radii, 〈R〉, observed in the dendrite centers in Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c) isochronally aged at 450,525, and 575°C.

Alloy Last aging treatment (°C) Mean precipitate radius, 〈R〉 (nm) Number of precipitates counted

Al-0.1Zr(c) 450 3.7± 0.5 188

525 6.8± 1.8 180

575 16.9± 4.4 75

Al-0.1Zr-0.1Ti(c) 450 3.4± 0.5 99

525 6.6± 1.9 205

575 16.3± 3.4 85

sidering the size of these precipitates, their coherency is not too surprising since their radii are

below the 〈R〉 = 25–35 nm threshold for coherency loss discussed in Figure 5.15.

Figures 6.13, 6.14, and 6.15 are for Al-0.1Zr-0.1Ti(c) aged isochronally at 450, 525, and 575°C,

respectively, but are representative of the binary Al-0.1Zr(c) as well. Indeed, Table 6.3 demon-

strates that measured precipitate radii after isochronal aging at 450, 525, and 575°C are statis-

tically identical for Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c), consistent with their similar overaging be-

havior demonstrated in Figure 6.12. As found in Chapter 5, the addition of Ti has little effect on

the coarsening of Al3Zr (L12) precipitates.

6.5 Discussion

6.5.1 Transformation to the equilibrium D023 structure

The L12 to D023 transformation is summarized schematically in Figure 6.17. Spheroidal L12 pre-

cipitates coarsen, eventually developing planar faults parallel to {100} planes inside the L12 pre-

cipitates. This strain contrast is attributed to antiphase boundaries (APBs) generated in the tran-

sition to an imperfect D023 structure. As coarsening progresses, the precipitates elongate along

these faults into a cigar-shaped morphology, eventually transforming into D023-structured pre-

cipitates exhibiting a disc-like morphology with a cube-on-cube orientation relationship with

α-Al.

Examples of spheroidal L12 precipitates exhibiting APBs are provided in Figures 5.18(d)

and 6.11. Zedalis and Fine [84] observed similar faults in Al3Zr (L12) precipitates formed af-


[010]Al

[001]Al

[100]Al

Fig. 6.17: Schematic of the L12 to D023 transformation in Al3Zr precipitates. Spheroidal L12 precipitates coarsen, eventuallydeveloping antiphase boundaries (APBs) parallel to {100} planes. Preferential coarsening occurs along the APBs, resulting in cigar-shaped precipitates. These eventually transform into disc-shaped D023-structured precipitates like those in Figure 6.5.

ter aging Al-0.22Zr (at.%) at 600°C for 12 h, which they postulated were APBs. This was con-

firmed by Chen et al. [114, 338] studying Al3(Zr0.25V0.75) (L12) precipitates formed in melt-spun

Al-0.34Zr-0.98V (at.%) alloys aged at 500°C. Using high resolution electron microscopy, these

researchers showed that the APBs are associated with a displacement of a/2〈110〉 on {100}

planes [338]. Lee et al. [297] showed similar planar faults in L12 precipitates (purportedly

Al3(Zr0.40,V0.40,Ti0.20)2) formed in Al-0.8Zr-0.8V-0.4Ti (at.%) alloys aged at 425°C for 200–400 h.

Ryum [85] demonstrated that the transformation from the cubic L12 structure to an imperfect

tetragonal structure involves the introduction of an APB with a displacement vector a/2〈110〉 on

{100}-type planes in D022-structured Al3Hf. The observed APBs in the spheroidal L12-structured

precipitates (Figures 5.18(d) and 6.6) therefore represent the early stages of transformation to the

equilibrium D023 structure. As indicated in Figure 6.17, a preferential growth in the precipitate

along the APB represents a later stage of transformation. Figure 6.6 displays examples of such

cigar-shaped precipitates, which have also been observed by others [84, 114, 297].

Temperature of the transformation

The L12 to D023 structural transformation seems to occur at approximately 500°C, which corre-

sponds also to the onset of overaging as observed by Vickers microhardness in Figures 6.1 and

6.12. In Chapter 5, D023 precipitates were never observed during extended aging times (1,600 h)

at 425°C. In the present isochronally aged alloys, D023 precipitates are not observed at 450°C,

2The precipitate compositions were not measured directly.


Figure 6.13, and are observed to a limited extent at 525°C, Figure 6.14. After extended aging

times (100 h) at 500°C, Figures 6.2 and 6.3, the transformation is extensive along dislocations,

although most precipitates are still L12-structured. Even at 575°C, L12 precipitates are observed

in Figure 6.15, indicating that the L12 to D023 transformation is not spontaneous. Additional fac-

tors, including aging time and dislocation density in the specimen, will also naturally influence

on the kinetics and the apparent transformation temperature. Nevertheless, the transformation

process seems to occur at 500°C within reasonable aging times.

This conclusion concurs with results of several studies in the scientific literature. Zedalis and

Fine [84] observed transformation to the equilibrium D023 structure in Al-0.24Zr (at.%) aged at

450°C. Their specimens, however, were initially cold rolled 95%; the high dislocation density

therefore explains why these researchers observed D023 precipitates at a relatively low tempera-

ture. Izumi and Oelschlagel [76] reported a “fine dispersion of small round particles (L12-crystal

structure), which are statistically distributed within the dendritic cells of the cast structure” in

an Al-0.33Zr (at.%) aged at 450°C. Even after extended aging times (218 h), these researchers re-

ported that the Al3Zr precipitates had not transformed to the stable D023 structure and the L12

precipitates were resistant to coarsening with 〈R〉 = 10.9 nm. Nes [79] aged an Al-0.05Zr (at.%)

alloy at 460°C up to 700 h and observed small, 〈R〉 < 35 nm Al3Zr (L12) precipitates and, like

Izumi and Oelschlagel [76], never observed the equilibrium D023 structure at 460°C. Ryum [78]

studied precipitation of Al3Zr (L12) in a chill-cast Al-0.15Zr (at.%) alloy aged at 500°C and at this

temperature, the equilibrium D023 structure was observed only after long aging times (>120 h),

primarily at grain boundaries. Rystad and Ryum [190] studied recrystallization in an Al-0.15Zr

(at.%) alloy annealed in the range 400–600°C; L12-structured Al3Zr precipitates were observed

at all temperatures except at 600°C, when plate-shaped equilibrium D023structured Al3Zr pre-

cipitates were formed after 8 h. Chaudhury and Suryanarayana [83] observed the precipitation

behavior of Al3Zr in an Al-3.0Zr (at.%) alloy prepared by RSP. The highly supersaturated solid

solution was stable beyond 300°C, with discontinuous precipitation of L12 Al3Zr observed after

aging for 1 h at 320°C. Metastable L12-structured Al3Zr was still observed after aging aging for 1

h at 500°C, followed quickly by the formation of the equilibrium D023 structure after 3 h at the

same temperature. Finally, Srinivasan et al. [94] formulated single-phase Al3Zr (L12) through

mechanical alloying of pure elemental powders, and found that the L12 to D023 transformation

occurs at 550°C, as monitored by X-ray diffraction and differential scanning calorimetry.


Stage I Stage II

Stage III Stage IV

Fig. 6.18: Schematic of four stages of microstructural coarsening occurring on the nanometer-scale of the precipitates and themicrometer-scale of the dendrites.

6.5.2 Mechanism of microstructural coarsening

The mechanism of microstructural coarsening described in this chapter may be summarized

schematically in Figure 6.18, in which four stages of coarsening are defined. Stage I coarsen-

ing corresponds to the peak-aged condition, where well-defined precipitate-rich dendrites con-

tain a dense distribution of small (R < 5 nm) coherent L12 precipitates. Larger interdendritic

spheroidal L12 precipitates like those in Figures 5.11 and 5.12 are also depicted, many of them

heterogeneously nucleated on dislocations.


Stage II is the initiation of microstructural coarsening, where dendritic L12 precipitates

coarsen by Ostwald ripening and the heterogeneously-nucleated precipitates coarsen rapidly by

dislocation pipe diffusion. Within the precipitate-rich dendrites, preferential coarsening occurs

heterogeneously on dislocations. During Stage II there is a macroscopically-observed decrease

in strength.

Stage III coarsening signals the beginning of the L12 to D023 structural transformation, lim-

ited to dislocations and probably other heterogeneous nucleation sites such as grain boundaries.

The heterogeneously-nucleated L12 precipitates grow into elongated cigar-shaped precipitates,

exhibiting preferential growth in 〈100〉 directions. Antiphase boundaries, characterized by sharp

lines of no-contrast parallel to {100} planes inside the L12 precipitates, signal the beginning of

the transition to the D023 structure. The transformation to the disc-like D023 phase occurs at the

expense of the surrounding L12 precipitates, leaving behind precipitate-free cyclinders up to ca.

1 µm in diameter surrounding the dislocation.

By Stage IV coarsening, most of the metastable L12 precipitates have transformed to the equi-

librium D023 structure. There are, however, remnants of the precipitate-rich dendrites that still

contain coherent, spheroidal L12 precipitates. Despite their presence, virtually all strength is lost

on the macroscopic scale.

The dark-field TEM micrographs in Figure 6.19 displays actual examples of the four stages

of coarsening described in Figure 6.18. Figure 6.19(a) displays Al-0.1Zr-0.1Ti(a) aged at 375°C

(0.69Tm) for 1,600 h. The microstructure contains well-defined dendrites full of small (〈R〉< 5 nm)

coherent L12 precipitates. Stage II coarsening is reached during extended aging at 425°C (0.75Tm),

as evidenced by Figure 6.19(b) displaying Al-0.1Zr-0.1Ti(b) aged at 425°C (0.75Tm) for 1,600 h.

The microstructure consists of coherent L12-structured precipitates, 〈R〉 ca. 12 nm, but precip-

itate coalescence in heterogeneous arrays is also widespread. Stage III coarsening is character-

ized by the beginning of the L12 to D023 transformation, which is occurring along a dislocation

in Figure 6.19(c) for Al-0.1Zr-0.1Ti(b) aged at 500°C (0.83Tm) for 100 h (after aging at 375°C for

100 h). The transformation occurs at the expense of the surrounding L12 dendritic precipitates,

forming a precipitate-free cylinder 400 nm in radius in Figure 6.19(c). In Stage IV coarsening, the

majority of the precipitates have transformed to the D023 structure, as shown Figure 6.19(d) for

Al-0.1Zr-0.1Ti(d) isochronally aged to 575°C (0.91Tm).


Fig. 6.19: Centered superlattice dark-field TEM micrographs displaying the four stages of microstructural coarsening shownschematically in Figure 6.18. Panel (a) is Al-0.1Zr-0.1Ti(a) aged at 375°C for 1,600 h. Well-defined dendrites contain small (〈R〉ca. 5 nm) coherent L12-structured precipitates. Panel (b) is Al-0.1Zr-0.1Ti(b) aged at 425°C for 1,600 h. Heterogeneously-nucleatedspheroidal L12-structured precipitates, fed by rapid pipe diffusion, grow at the expense of nearby dendritic precipitates. Panel(c) is Al-0.1Zr-0.1Ti(b) aged at 500°C for 100 h (after aging at 375°C for 100 h). Transformed D023 precipitates are observedheterogeneously-nucleated on a dislocation, forming a precipitate-free cylinder around the dislocation. Panel (d) is Al-0.1Zr-0.1Ti(c) isochronally aged to 500°C. The microstructure consists of fully-transformed D023 precipitates exhibiting a disc-like mor-phology (ca. 200 nm in diameter and 50 nm in thickness) and a cube-on-cube orientation relationship with α-Al.

6.5.3 Mechanism of strengthening

Theories of precipitation strengthening have been thoroughly reviewed and assessed, originally

by Brown and Ham [347] and more recently by Ardell [20, 23]. Precipitate shearing, precipitate

bypass by dislocation looping, or a combination of these two mechanisms can generally explain

ambient temperature strength in coarse-grained, non-strain-hardened, precipitate-strengthened

alloys in the absence of other strengthening mechanisms (e.g., Hall-Petch or solid-solution strength-


ening). For large precipitate radii, R, the strength is controlled by Orowan dislocation looping,

defined in Eq.(1.1) [25]:

∆σor = M · 0.4 ·Gb

π√

(1− ν)·ln

(2Rb

)λ

, (1.1)

where M = 3.06 [26], b = 0.286 nm [27], G = 25.4 GPa [27], and ν = 0.345 [26] for Al. The mean

planar radius, R, and effective inter-precipitate distance λ, are, for a monodispersed assembly,

given by [20, 22, 23]:

R =π

4〈R〉 , (1.2)

and

λ =(√

2π

3f− 2

π

4

)〈R〉 , (1.3)

where f is the precipitate volume fraction. Equations (1.2) and (1.3) are also good approxima-

tions for polydispersed arrays [23].

Deformation by dislocation shearing is expected to occur at small R, as observed experimen-

tally in Al-Li [348, 349] and Al-Sc [115, 180] alloys with shearable, coherent precipitates. For the

shearing mechanism, the increase in yield strength results from three contributions: (i) modulus

hardening; (ii) coherency strengthening; and (iii) order strengthening. The strengthening due to

modulus mismatch, ∆σms, results from differences in the shear moduli of the precipitate and

matrix phases and is given by [20]:

∆σms = M · 0.0055(∆G)3/2

(f

Γ

)1/2

b

(〈R〉b

)3m/2−1

, (6.1)

where ∆G is the modulus mismatch between the matrix and precipitate, Γ = 12Gb2 is the line

tension of the dislocation, and m = 0.85 is a constant [20]. The elastic constants of Al3Zr (L12)

are not clear at present. Nakamura and Kimura [350] measured the elastic constants of Al3Zr

(D023) single crystals and report G = 85.1 GPa, which is in good agreement with the extrapo-

lated results of Zedalis et al. [351] for D023-structured Al3Zr precipitates in α-Al solid-solution.

George et al. [352] studied the elastic constants of an Fe-stabilized Al3Zr L12 trialuminide,

(Al0.67,Fe0.08)3Zr, and measured G = 68 GPa as reported also in references [49, 50]. George et

al. [49,352] cautioned that this value of G may be affected by Kirkendall voids in their specimen,


perhaps accounting for the differences in the reported values for the D023- and L12-structured

Al3Zr trialuminides. Their measured value of 68 GPa for Fe-modified Al3Zr (L12) is, however,

identical to that of Al3Sc (L12) [353–355]. Taking the smaller value, G = 68 GPa, results in a value

of ∆G = 42.6 GPa entering Eq.(6.1); this has the added benefit of providing a conservative esti-

mate for the magnitude of ∆σms, thereby over-predicting the critical R at which the transition

from shearing to looping occurs.

Coherency strengthening, ∆σcs, arises through the elastic strain-field interactions between a

coherent precipitate and a dislocation and is given by [20]3:

∆σcs = M · χ(εG)3/2

(〈R〉 fb

Γ

)1/2

, (6.2)

where χ = 2.6 [20], ε is a mismatch parameter approximated by 23δ; δ = 0.0075 is the lattice

parameter mismatch at room temperature defined in Table 1.1 for Al3Zr (L12).

Order strengthening is due to the formation of an antiphase boundary (APB), which occurs

when a matrix dislocation shears an ordered precipitate. At peak strength, ∆σos is given by [20,

23]:

∆σos = M · 0.81γAPB

2b

(3πf

8

)1/2

, (6.3)

where γAPB ≈ 0.5 J m-2 is an average value of the Al3Sc APB energy for the (111) plane taken from

several reported values [49, 352, 353, 355, 357]. A comparable APB energy is assumed for Al3Zr

(L12).

Equations (1.1), (6.1), (6.2), and (6.3) each depend on the precipitate volume fraction, f . For

a homogeneous alloy containing 0.1 at.% Zr, the equilibrium volume fraction of precipitated

Al3Zr would be f = 0.004 from the lever rule, assuming negligible solubility at the aging tem-

perature. The solutes in the present alloys are microsegregated, however, with the dendrite cen-

ters enriched approximately twofold in Zr based on the measured EDS linescans in Figure 5.5.

The equilibrium volume fraction of Al3Zr precipitates in the dendrites is therefore approaching

3There is some confusion regarding the expression for ∆σcs, Eq.(6.2), used in previous studies on Al-Sc alloys byMarquis et al. [180,356] and Fuller et al. [115,322]. In those studies, the effective value of Γ entering Eq.(6.2) is 0.18Gb2,attributed to Ardell [20]. Ardell, however, does not use this expression for Γ; indeed, as he discusses in detail, Γ is notconstant but instead depends on the character of the dislocation; Γ = 1

2Gb2 is a reasonable approximation used by

Brown and Ham [347]. The inconsequential discrepancy between Γ = 12Gb2 used presently, and Γ = 0.18Gb2 used

previously [115, 180, 322, 356], can easily be accommodated by the value of χ chosen, which varies between 2 and 3for various theories [20].


f = 0.008. There is also the assumption of negligible solid-solubility. For the isochronally aged

specimens considered presently, the aging temperature is progressively raised to 600°C, with a

concomitant (and significant) decrease in equilibrium volume fraction of precipitates (assuming

equilibrium is reached during the 3 h aging treatment). The volume fraction of Al3Zr (L12) is fur-

ther reduced during the transformation to the equilibrium D023 structure, which commences by

525°C, Figure 6.14(b). Finally, a small amount of Ti is incorporated into the precipitates, shown

in Chapter 4, and Ti also influence the metastable Al-Zr solvus, discussed in Figure 6.22. While

these effects are small, as evidenced by the isothermal age hardening curves in Figures 5.7–5.10

where Ti has little effect on the strength, ternary additions will nevertheless influence f .

Which volume fraction, f , to use for modeling the theoretical contributions to strength

is therefore not trivial. For isochronally aged specimens, the assumption of negligible solid-

solubility is reasonably valid, at least in the peak-aged condition, since precipitates are nucle-

ated at the lowest possible temperature. As coarsening commences at higher temperatures, the

equilibrium volume fraction of L12 precipitates is reduced by increasing solubility of solute(s)

and the transformation to D023. An advantage of isochronal aging, however, is that the relative

fractions of dendritic and interdendritic regions remains relatively constant (this is not true of

isothermal aging, where the initial aging temperature has a strong influence on the volume frac-

tion of the precipitate-rich dendrites, as discussed in Figure 6.21). As a point of reference, a value

of f = 0.004 is assumed, which is approximately the total volume fraction (including precipitate-

free regions) of Al3Zr or Al3(Zr1−xTix) precipitates in the alloy at peak strength on the isochronal

aging curve. This f will under-predict the strengthening in the precipitate-rich dendrites, but

will over-predict the strengthening at higher temperatures because of increased solid-solubility

and L12 to D023 transformation.

Figure 6.20 displays the theoretical and measured increases in strength of isochraonally-

aged Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c) as a function of the precipitate radius, R. The measured

strength increment is compared to the value for the as-cast, unaged specimens using a conver-

sion factor of 3 between Vickers microhardness and strength [358]. Calculations for the shearing

mechanisms, ∆σms + ∆σcs and ∆σos, are applicable for R . 3 nm, confirming that shearing is

not operative for the data in Table 6.3. These conclusions are in agreement with prior results

on Al-0.18Sc [180] and several Al-Sc-Zr alloys [115], for which Orowan dislocation looping is the

main strengthening mechanism for R & 2 nm.


0 5 1 0 1 5 2 0 2 5

0

5 0

1 0 0

1 5 0

2 0 0

A l - 0 . 1 Z r ( c )

A l - 0 . 1 Z r - 0 . 1 T i ( c )

f = 0 . 0 0 4

∆σm s

+ ∆σc s

∆σo s

∆σo r

Str

en

gth

en

ing

in

cre

me

nt

(M

Pa

)

M e a n p r e c i p i t a t e r a d i u s ( n m )

Fig. 6.20: Vickers microhardness yield stress increment vs. mean precipitate radius 〈R〉 for Al-0.1Zr(c) and Al-0.1Zr-0.1Ti(c)isochronally aged at 450, 525, and 575°C. The curves represent predictions of Eqs.(1.1), (6.1), (6.2), and (6.3) for f = 0.004.

Table 6.4: Mean precipitate radii, 〈R〉, observed in the dendrite centers after different thermal histories.

Alloy Aging treatment Vickers microhardness (MPa) Mean precipitate radius, 〈R〉 (nm)

Al-0.1Zr(b) 425°C for 1600 h 319± 12a 10.9± 1.9e

Al-0.1Zr-0.1Ti(b) 351± 10b 11.6± 2.3e

Al-0.1Zr(b) 500°C for 100 h (after 375°C for 100 h) 325± 17c 12.3± 2.9f

Al-0.1Zr-0.1Ti(b) 319± 12c 12.0± 4.0f

Al-0.1Zr(c) Isochronally aged to 525°C 327± 13d 6.8± 1.8g

Al-0.1Zr-0.1Ti(c) 341± 8d 6.6± 1.9g

a Figure 5.7. b Figure 5.9. c Figure 6.1. d Figure 6.12. e Table 5.3. f Table 6.2. g Table 6.3.


For the 〈R〉 values measured in Table 6.3, the experimentally-measured strength is much less

than that predicted by Eq.(1.1), indicating that the observed strengthening is not directly con-

trolled by the Orowan bypass mechanism. This is illustrated more dramatically by the data in Ta-

ble 6.4, which compares experimentally-measured precipitate radii and Vickers microhardness

from Al-0.1Zr and Al-0.1Zr-0.1Ti (at.%) alloys reported previously in Tables 5.3, 6.2, and 6.3. The

alloys in Table 6.4 exhibit similar hardness values, yet the precipitate radii differ substantially de-

pending on the alloy thermal history. Specimens isochronally aged to 525°C exhibit smaller radii

than those aged isothermally at 425°C or in a sequential 375°C + 500°C aging treatment, despite

the similar strength. The data in Figure 6.20 and Table 6.4 indicate that strengthening cannot be

predicted from the measured 〈R〉, assuming precipitate bypass by Orowan dislocation looping.

Precipitation strengthening is instead probably occurring on multiple length scales, as might

be anticipated considering the nanometer-scale of the precipitates and the micrometer-scale

of the precipitate-rich dendrites. Table 6.4 also demonstrates that thermal history has a strong

influence on alloy strength.

6.5.4 Effect of aging temperature

Increasing the aging temperature (decreasing the undercooling) in any precipitation reaction

reduces the chemical driving force for nucleation leading to: (i) larger precipitates (because the

critical radius for nucleation is larger); and (ii) reduced equilibrium volume fraction of the dis-

persed phase. Thus, precipitation-strengthening (at ambient temperature) is generally reduced

because the precipitates are at a smaller number density and volume fraction. Higher aging

temperatures therefore lead to reduced precpitation hardening, as observed in Figures 5.7–5.10.

The effect of decreasing strength with decreasing precipitate volume fraction must also be in-

terpreted within the context of the micrometer-scale dendritic distribution of Al3Zr or Al3(Zr,Ti)

precipitates. Figure 6.21(a) shows schematic isoconcentration contours of the as-solidified α-

Al solid-solution for a peritectic system (k0 > 1) in which the cores of the dendritic cells are

enriched in solute concentration compared to regions at the periphery (C1 < C2 < C3). Fig-

ure 6.21(b) shows the expected precipitation behavior for such a cored solid-solution at three

temperatures T1 < T2 < T3. At the lowest aging temperature, T1, all compositions C1, C2, and C3

are supersaturated with respect to the solvus and hence these regions of the original dendritic

α-Al solid-solution may be expected to exhibit precipitation, as indicated by the dots in Figure


T1

T2

T3

Te

mp

era

ture


( a ) ( b )

C1 < C

2 < C

3

S o l v u s

C1 C

2 C

3

C1

C2

C3

C3

C2

C1

Fig. 6.21: Schematic isoconcentration contours and the expected precipitation behavior for a dendritically-solidified α solid-solution in which k0 > 1. (a) Isoconcentration contours corresponding to C1 < C2 < C3 within the original α dendrites. (b)The expected precipitation behavior from the α solid-solution in (a) for three aging temperatures T1 < T2 < T3. As aging tem-perature is increased, fewer isoconcentration contours have sufficient supersaturation to allow precipitation, with those that doindicated by dots. Therefore, for higher aging temperatures, the relative amount of interdendritic precipitate-free zones is expectedto increase.

6.21(b). At the higher temperature, T2, the region corresponding to C1 is no longer supersatu-

rated with solute and therefore precipitation should not occur. At the highest aging temperature,

T3, only the concentration C3 is sufficiently rich to effect precipitation.

Higher aging temperatures, therefore, not only result in smaller volume fractions of the pre-

cipitated phase (typical for any precipitation reaction), but because of the dendritic distribution

of precipitates, the volume fraction of precipitate-free interdendritic regions is also increased.

This also explains why during isochronal aging, Figure 6.12, strengths of the order of 415 MPa

may be obtained for Al-0.1Zr after aging up to 450°C, whereas during isothermal aging, Fig-

ure 5.7, the attainable strength at 425°C is much less: ca. 300 MPa. A similar effect was observed

by Sato et al. [271] during aging of an Al-0.17Zr (at.%) alloy at 450°C, whereby ramp heating at a

controlled rate of 100 °C h-1 produced a nearly 200 MPa increase in peak hardness compared to

the same alloy aged directly at 450°C (540 MPa vs. 350 MPa).

6.5.5 Effect of Ti additions

The isothermal (Figures 5.7–5.10 and 6.1) and isochronal (Figure 6.12) hardness curves, as well

as the mean precipitate radii in Tables 5.3, 6.2, and 6.4 indicate that there is no benefit, in terms

of coarsening resistance, of Al3(Zr1−xTix) (L12) precipitates as compared to Al3Zr (L12). Even


0 0 . 0 5 0 . 1 0 0 . 1 5 0 . 2 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

0 . 2 a t . % T i

0 . 1 a t . % T i

M e t a s t a b l e L 1 2 S o l v u s

6 6 0 . 8 °C

( A l )

( A l ) + A l3Z r

L + A l3Z r

��

��

��

�

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

L

Fig. 6.22: Al-rich equilibrium binary Al-Zr phase diagram (solid lines) [123], and calculated (dotted lines) metastable L12 Al3Zrand Al3(Zr1−xTix) solvus by Murray [316]. Additions of Ti reduce the solid-solubility of Zr.

at 425°C (0.75Tm), the Al-Zr or Al-Zr-Ti alloys overage very sluggishly and the precipitates in the

dendrites remain small, 〈R〉. 10 nm, after extended aging times (1,600 h or 9.5 weeks, Table 5.3).

Only at or above 500°C (0.83Tm) do the alloys significantly overage within a 3 h treatment, Fig-

ure 6.12, and after 100 h at this temperature the microstructures are again indistinguishable,

〈R〉 ≈ 12 nm (Table 6.2). After 3 h at 575°C (0.91Tm), the Al3Zr (L12) or Al3(Zr1−xTix) (L12) pre-

cipitates still exhibit impressive coarsening resistance, 〈R〉. 17 nm (Table 6.3), with no improve-

ments in coarsening from Ti additions.

While Ti has no effect on the long-term stability of the Al3(Zr1−xTix) precipitates, there is

a noticeable influence on the incubation time for nucleation, as evidenced in Figures 5.7, 5.9,

and 6.12. Moreover, the peak hardness of Al-0.1Zr-0.1Ti(b) is greater than that of Al-0.1Zr(b)

at 425°C (Figures 5.7 and 5.9) and that of Al-0.1Zr-0.1Ti(c) is similarly greater than that of Al-

0.1Zr(c) during isochronal aging (Figures 6.12). The same effect was observed by Sato et al. [271]


in chill-cast Al-0.17Zr and Al-0.12Zr-0.15Ti (at.%) alloys aged at 450°C for up to 1,200 h. The

Al-0.12Zr-0.15Ti alloy had a reduced incubation time for nucleation and also exhibited a larger

peak strength (400 MPa) compared to Al-0.17Zr (350 MPa).

Murray [316] has calculated that additions of Ti reduce the solid-solubility of Al3Zr (L12),

Figure 6.22. The resulting increased supersaturation (for a given Zr concentration) accounts for

the reduced incubation time for nucleation in the Al-Zr-Ti alloys and also explains the larger

peak strength, since the equilibrium volume fraction of precipitate-rich dendrites is greater.

Titanium additions thus have a similar effect as isochronal aging in Al-Zr-based alloys. The

effective supersaturation of Zr is enhanced during nucleation, increasing the volume fraction of

the precipitate-rich dendrites and reducing the width of the precipitate-free interdendritic chan-

nels. Unfortunately, the variation in apparent dendritic cell size varies considerably throughout

the alloy, Figure 5.6, and there is no practical way to quantify this phenomenon using traditional

2-D microscopy.

6.6 Chapter Summary

The transformation of Al3Zr (L12) and Al3(Zr1−xTix) (L12) precipitates to their respective equi-

librium D023 structures was investigated in conventionally-solidified Al-0.1Zr and Al-0.1Zr-0.1Ti

(at.%) alloys isothermally aged at 500°C or isochronally aged 300–600°C. The following results

were obtained and discussed:

• There is no benefit from Ti additions, both in terms of coarsening resistance of the

metastable L12 precipitates, or in delaying the L12 to D023 transformation. Both alloys

overage rapidly at the same rate at or above 500°C, during which metastable L12 precipi-

tates transform to their equilibrium D023 structure.

• The L12 to D023 transformation is limited primarily to dislocations and other defects, where

heterogeneously-nucleated L12-structured precipitates, fed by rapid pipe diffusion, coarsen

by dissolution of surrounding L12 precipitates within a ca. 500 nm radius of the dislocation.

The coarsening precipitates often exhibit antiphase boundaries (APBs) parallel to {100}

planes, representing a transition to an imperfect equilibrium D023 structure. Preferential

growth is similarly observed along 〈100〉 directions, resulting in elongated spheroids. The

cigar-shaped precipitates eventually transform to disc-like D023 precipitates, ca. 200 nm in


diameter and 50 nm in thickness, exhibiting a cube-on-cube orientation relationship with

α-Al.

• The disc-like morphology of the equilibrium phase reflects the tetragonal symmetry of the

D023 unit cell. The lattice parameter mismatch of Al3Zr (D023) is δa = δb = -0.88% and

δc = 6.92% (Table 1.1); the mismatch is most severe in the c-direction, favoring disc-like

precipitates that are thin along c. The equilibrium D023-structured precipitates form het-

erogeneously on dislocations because of a significant lattice parameter mismatch with the

α-Al solid-solution, δ = 2.89% (Table 1.1).

• The transformation does not occur spontaneously throughout the alloy, as evidenced in

Figure 6.6 displaying spheroidal L12 precipitates, partially-transformed imperfect L12 pre-

cipitates with APBs, and disc-like D023 precipitates in close proximity. The transformation

is also extraordinarily sluggish; even after aging at 575°C (0.91Tm) for 3 h, L12-structured

precipitates are observed, Figure 6.15.

• Precipitation strengthening in segregated Al-Zr alloys occurs on multiple length scales: (i)

on the nanometer-scale scale by an Orowan strengthening mechanism and (ii) on the

micrometer-scale related to the volume fraction of the precipitate-rich dendrites. Mi-

crostructural coarsening similarly occurs on several length scales: (i) by Ostwald ripen-

ing (volume diffusion-controlled coarsening) of spheroidal nanometer-scale precipitates

within the dendrites; and (ii) by local dissolution of the micrometer-scale precipitate-rich

dendrites during the D023 transformation.

• Because strengthening depends on the volume fraction of the precipitate-rich dendrites,

thermal history has a strong influence on the observed alloy strength. During isochronal

aging, Figure 6.12, strengths of the order of 415 MPa may be obtained for Al-0.1Zr (at.%)

after aging up to 450°C; during isothermal aging, Figure 5.7, the attainable strength at 425°C

is only ca. 300 MPa. Unlike isothermal aging, isochronal aging nucleates precipitates at

the lowest possible aging temperature where the supersaturation is greatest. Because of

the corresponding large chemical driving force, the critical radius for nucleation is small

and the equilibrium volume fraction of the precipitate-rich dendrites is large, promoting a

larger number density of fine precipitates.


• Ternary additions of Ti have no effect on the coarsening resistance of Al3(Zr1−xTix) (L12) as

compared to Al3Zr (L12); Ti does, however, accelerate the incubation time for nucleation

and also increases the peak hardness achieved during aging (Figures 5.7, 5.9, and 6.12).

Titanium reduces the solid-solubility of Al3Zr (L12), Figure 6.22, and these are effects of

solute supersaturation (chemical driving force and precipitate-rich dendrite volume frac-

tion) and is not an indication of coarsening resistance.

CHAPTER 7

Creep of Al-Zr and Al-Zr-Ti Alloys

Creep behavior at 300, 350, or 400°C of Al-0.1Zr and Al-0.1Zr-0.1Ti (at.%) alloys are pre-

sented and discussed. At all temperatures investigated, creep threshold stresses are ob-

served and found to be much smaller than the ambient temperature yield stress, indica-

tive of a climb-controlled bypass mechanism. For a given temperature, the ternary Al-

Zr-Ti alloy exhibits a smaller threshold stress than the binary Al-0.1Zr alloy. This may be

attributable to: (i) differences in texture arising from a reduced grain size in the ternary

Al-Zr-Ti alloy; (ii) reduction in the lattice parameter mismatch of the Al3(Zr1−xTix) pre-

cipitates with the α-Al matrix; or (iii) finer precipitates. Comparisons are made with

Al-Sc and Al-Sc-M (M = Zr or Ti) alloys with similar radii and volume fractions of precip-

itates.

7.1 Introduction

IN CHAPTER 5 the microstructure and ambient temperature mechanical properties of Al-0.1Zr

and Al-0.1Zr-0.1Ti (at.%) alloys, during isochronal aging at 375, 400, or 435°C, are discussed.

This chapter investigates the elevated mechanical properties of similar isothermally aged alloys

by creep experiments performed at 300, 350, or 400°C.

The creep behavior of precipitation- or dispersion-strengthened materials generally follows

a power-law equation of the form [359]:

ε = Aap · σnap · exp(− Q

RgT

), (7.1)

where ε is the minimum, secondary strain rate, Aap is a constant, σ is the applied stress, nap is the

apparent stress exponent, Qap is the apparent activation energy, Rg is the universal gas constant,

and T the absolute temperature. When the apparent stress exponent significantly exceeds that

of the matrix, an athermal threshold stress, σth, is assumed, below which creep is negligible. This

167

168 CHAPTER 7 CREEP OF AL-ZR AND AL-ZR-TI ALLOYS

Table 7.1: Compositions and aging conditions of the Al-Zr and Al-Zr-Ti alloys investigated by creep.

Alloy Nominal comp. (at.%) Verified comp. (at.%) Aging conditions Creep temperature (°C)Zr Ti Zr Ti

Al-0.1Zr(d) 0.1 — 0.079 — 100 h at 400°C 300, 350, 400

Al-0.1Zr-0.1Ti(d) 0.1 0.1 0.081 0.092 100 h at 400°C 300, 350, 400

leads to a modified power-law equation [360, 361]:

ε = A · (σ − σth)n exp(− Q

RgT

), (7.2)

where A is a constant, n is the matrix stress exponent, and Q is the matrix creep activation en-

thalpy, which is usually equal to the activation enthalpy for self-diffusion. The threshold stress,

σth, arises from the increase in length of the dislocation during the climb bypass process for

metals with non-shearable, coherent precipitates [30]. Since creep deformation is negligible for

stresses below this threshold, σth is analogous to the yield stress describing deformation at am-

bient temperature. It is thus desirable to have large values of σth. Indeed, an upper bound for σth

is the Orowan stress, σor, above which dislocations bypass dispersoids by bowing within their

glide plane, and hence σor controls the yield stress.


A series of binary Al-0.1Zr and ternary Al-0.1Zr-0.1Ti alloys were investigated; alloy designa-

tions, compositions, aging conditions prior to creep, and creep temperatures are summarized

in Table 7.1. The alloys were prepared employing non-consumable electrode arc-melting, as de-

scribed in Sections 3.2 and 4.2. The verified compositions in Table 7.1 were obtained by bulk

chemical analysis performed by direct current plasma emission spectroscopy (ATI Wah Chang,

Albany, OR).

Creep specimens in the shape of parallelepipeds (ca. 4×4×8 mm3) were electrode-discharge-

machined (EDM) with their axes perpendicular to the surface of the ingot which had been in

contact with the water-cooled copper crucible during arc-melting. The height of the creep spec-

imen spans nearly all of the button ingot height. Constant-load compression creep tests, corre-


Table 7.2: Vickers microhardness of the aged Al-Zr and Al-Zr-Ti alloys before and after creep.

Alloy Vickers microhardness (MPa) Creep temperature and timeAs-aged Post-creep

Al-0.1Zr(d) 335± 7 354± 7 411 h at 300°C

366± 18 712 h at 400°C

Al-0.1Zr-0.1Ti(d) 401± 15 439± 10 420 h at 400°C

sponding to compressive stresses in the range of 4–30 MPa, were performed in air at 300–400°C.

A superalloy creep cage translated tensile loads in the pull-rods to compressive stresses on the

specimen. Frictional effects on the end-loaded specimens were minimized by using boron ni-

tride coated alumina platens in the creep cage. Specimen temperature was measured in the

three-zone furnace with a temperature stability of± 1°C after a 1 h soak at the test temperature.

Specimen strains were calculated from extensometric displacements of cage plattens measured

using a linear variable differential transducer (LVDT) with a resolution of 2.5 µm. Continuous

measurements were made on each specimen by monitoring the strain rate during the test; the

data were collected in ca. 1–30 s increments depending on the strain rate. Once steady-state

deformation was achieved, the load was increased, resulting in three to five data points suitable

for determining a stress exponent from one specimen. Total specimen strain never exceeded

ε = 0.10.

The alloys were age hardened isothermally at 400°C for 100 h prior to creep, which produces

a peak-aged hardness of ca. 420 MPa in alloys Al-0.1Zr(b) and Al-0.1Zr(b)-0.1Ti(b) (Figures 5.7

and 5.9). The hardness of the present alloys prior to creep is slightly less, as indicated in Ta-

ble 7.2: 335± 7 MPa for Al-0.1Zr(d) and 401± 15 MPa for Al-0.1Zr-0.1Ti(d). The reduced strength

is consistent with the slightly smaller solute concentration in Al-0.1Zr(d) and Al-0.1Zr-0.1Ti(d)

(Table 7.1) compared to Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b) (Table 5.1). The greater strength of the

ternary alloy is, however, inconsistent with what is observed for Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b)

(Figures 5.7 and 5.9), where the peak strengths of the alloys are indistinguishable at 400°C and

below. The higher strength of Al-0.1Zr(d) may indicate: (i) a larger volume fraction of precip-

itates (or precipitate-rich dendrites); or (ii) smaller precipitates leading to increased Orowan

strengthening at ambient temperature. Based on the results in Chapters 5 and 6, however, the

latter explanation seems unlikely since at all aging temperatures investigated, Al-Zr and Al-Zr-


0 1 0 0 2 0 0 3 0 0 4 0 0 5 0 0 6 0 0

0 . 0 0

0 . 0 1

0 . 0 2

0 . 0 3

0 . 0 4

0 . 0 5

0 . 0 6

0 . 0 7

0 . 0 8

0 . 0 9

0 . 1 0

σ = 1 9 M P a

ε = 4 . 8 x 1 0- 6 s

- 1

σ = 1 7 M P a

ε = 6 . 3 x 1 0- 7 s

- 1

σ = 1 5 M P a

ε = 2 . 1 x 1 0- 7 s

- 1

σ = 1 4 M P a

ε = 1 . 8 x 1 0- 8 s

- 1

Str

ain

��

��

Fig. 7.1: Variation of strain, ε, with time for alloy Al-0.1Zr(d) crept at 300°C for four levels of stress. These data correspond to theright-most curve in Figure 7.2. The strain rate measured at steady-state is also indicated for each load.

Ti alloys are observed to have statistically identical precipitate radii. Moreover, strength is not

controlled directly by the Orowan mechanism, Chapter 6, and therefore hardness is not a direct

indication of precipitate radius.


A primary creep regime, where the strain rate decreases continuously with time, always precedes

steady-state creep. As shown in Figure 7.1, the strain associated with primary creep, even near

the threshold stress, can be large (of the order ε = 0.02–0.03). This may represent unimpeded

dislocation glide through the precipitate-free interdendritic channels, not unlike the deforma-

tion through γ matrix channels between γ′ precipitates in Ni-based superalloys [28,29]. Primary

creep is eventually exhausted as deformation accumulates in the interdendritic regions, at which

point steady-state creep ensues and dislocation climb around the precipitates in the dendritic

regions controls the deformation.

The minimum strain rate from the steady-state regime is plotted as a function of applied


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

1 0- 9

1 0- 8

1 0- 7

1 0- 6

1 0- 5

1 0- 4

1 0- 3

1 0- 2

��

��

��

��

��

��

�

�

Min

imu

m s

tra

in r

ate

(s

-1

)

A p p l i e d s t r e s s ( M P a )

��

��

Fig. 7.2: Double logarithmic plot of minimum creep rate at 300, 350, or 400°C vs. applied stress, for Al-0.1Zr(d) and Al-0.1Zr-0.1Ti(d)aged for 100 h at 400°C. Each curve represents data from a single specimen; two specimens each of Al-0.1Zr(d) and Al-0.1Zr-0.1Ti(d)were tested at 400°C. Data for pure Al calculated from reference [27].

Table 7.3: Threshold stress, σth, of the Al-Zr and Al-Zr-Ti alloys crept at 300, 350, or 400°C.

Alloy Creep temperature (°C) σth (MPa)

Al-0.1Zr(d) 300 12

350 9

400 67

Al-0.1Zr-0.1Ti(d) 300 9

350 7

400 66


- 0 . 5 0 . 0 0 . 5 1 . 0 1 . 5 2 . 0

- 2 0

- 1 9

- 1 8

- 1 7

- 1 6

- 1 5

- 1 4

- 1 3

- 1 2

- 1 1

- 1 0

- 9

- 8

��

��

��

��

�

ln (

ε)

l n ( σ- σt h)

Fig. 7.3: ln(ε) vs. ln(σ − σth) for the data in Figure 7.2.

stress, Figure 7.2, for tests performed at 300, 350, or 400°C. Despite the non-uniform precipitate

distribution and the associated precipitate-free interdendritic channels described in Chapter 5,

the Al-Zr and Al-Zr-Ti alloys show remarkably high creep resistance at 300–400°C. The apparent

stress exponent, n, of both alloys is much higher than for pure Al (n = 4.4), indicative of the

existence of a threshold stress. The value of this threshold stress is determined from linear fits

of (ε)1/n vs. σ [362] using n = 4.4 [27] for dislocation creep of Al, and are given in Table 7.3 for

the alloys studied. The validity of this approach is demonstrated in Figure 7.3, where the data at

each temperature fall on straight lines with slopes near n = 4.4, the assumed value of the stress

exponent.

7.4 Discussion

7.4.1 Transformation to the equilibrium D023 structure

Chapter 6 discusses that the transformation to the equilibrium D023 structure of Al3Zr or

Al3(Zr1−xTix) occurs on dislocations and heterogeneous nucleation sites. Dislocations gener-


ated during creep may therefore hasten the transformation of the strengthening precipitates,

thereby reducing creep resistance. Post-creep Vickers microhardness measurements for creep

experiments performed at 400°C for hundreds of hours, Table 7.2, indicate no loss in strength

compared to the as-aged condition prior to creep, suggesting that the mechanism of microstruc-

tural coarsening on dislocations is not active at 400°C. This is consistent with the results in Chap-

ter 6, which indicate that the L12 to D023 transformation occurs at ca. 500°C. Indeed, the hard-

ness is actually increased after creep, which may indicate additional precipitation during the

test.

7.4.2 Differences in creep behavior between Al-Zr and Al-Zr-Ti alloys

As discussed in Section 6.5.5, additions of Ti to Al-Zr alloys accelerates the nucleation kinet-

ics and increase the peak strength at aging temperatures greater than 400°C, attributable to a

reduction in the solid-solubility of Zr. This decreases the incubation time for nucleation and in-

creases the equilibrium volume fraction of the precipitated Al3(Zr1−xTix) (L12) phase. There is

no improvement, however, in coarsening resistance of Al3(Zr1−xTix) precipitates as compared

to Al3Zr during isothermal aging at 375–425°C; even after extending aging times at 425°C (75%

of the absolute melting temperature of Al), the Al-Zr or Al-Zr-Ti alloys investigated overage very

sluggishly and the precipitates in the dendrites remain small, 〈R〉. 10 nm, Table 5.3.

Because of the indistinguishable microstructure, it might be expected that the Al-Zr and Al-

Zr-Ti alloys would behave similarly during creep. Figure 7.2, however, indicates that at 300, 350,

or 400°C, the ternary Al-Zr-Ti alloy consistently creeps faster than the binary Al-Zr, despite the

higher concentration of solutes (Table 7.1). This may be attributable to: (i) the reduced grain size

in the ternary Al-Zr-Ti alloy; (ii) reduction in the lattice parameter mismatch of the Al3(Zr1−xTix)

precipitates with the α-Al matrix; or (iii) smaller precipitates. Each of these effects is discussed

in the following.

Crystallographic texture

Figure 5.1 displays the as-cast macrostructure of alloys Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b), which

are similar to those studied presently. The relative extents of equiaxed and columnar zones in the

ingot cross-section varies between the two alloys: Al-0.1Zr(b) is comprised entirely of columnar

grains whereas half of the ingot cross-section in Al-0.1Zr-0.1Ti(b) is equiaxed. The grain struc-


ture of these alloys is shown more clearly in Figures 7.4 and 7.5, which display a series of optical

micrographs for alloys Al-0.1Zr(b) and Al-0.1Zr-0.1Ti(b), respectively.

It is conceivable that the poorer creep performance of the Al-Zr-Ti alloy is attributable to a

macroscopic texture effect. It is well-known [175–177] that dendrites, in the absence of hetero-

geneous nucleation sites, grow in crystallographically-preferred 〈100〉 directions, explaining the

columnar structure shown in Figure 7.4. Because primary precipitates of Al3(Zr1−xTix) act as

heterogeneous nucleants of α-Al, the as-cast structure for Al-0.1Zr-0.1Ti(b) (Figure 7.5) is pre-

dominantly equiaxed and is therefore no longer textured along 〈100〉. This difference in crys-

tallographic texture may explain the poorer creep performance of the ternary alloy. The refined

grain structure in Figure 7.5 would also be more susceptible to grain boundary sliding, but this

deformation mechanism seems unlikely given the existence of a threshold stress in the Al-Zr-Ti

alloy.

Reduced lattice parameter mismatch

Marquis and Dunand [33] showed that elastic interactions due to precipitate-matrix modulus

and lattice parameter mismatches can influence significantly the creep threshold stress for co-

herent precipitates, with σth increasing with increasing lattice parameter mismatch, δ. This the-

ory has been confirmed by numerous creep experiments on Al-Sc [180, 181], Al-Mg-Sc [182],

Al-Sc-Zr [115], Al-Sc-Ti [118], and Al-Sc-RE (RE = Y, Dy, or Er) [363] alloys.

It has been shown [292] that the lattice parameter of the metastable L12 Al3Zr phase may

be reduced by additions of Ti, Figure 7.6. The reduced creep strength of the Ti-containing al-

loys, therefore, may also be attributable to a reduced lattice parameter mismatch between the

Al3(Zr1−xTix) precipitates and the α-Al matrix.

Assuming Vegard’s Law holds, the lattice parameter of Al3(Zr1−xTix) (L12) may be estimated,

Figure 7.6, with δ = 0 (at room temperature) occurring at x = 0.25. In Chapter 4 the compositions

of Al3(Zr1−xTix) precipitates formed at 375 or 425°C were measured directly by 3-D atom-probe

tomography (3-D APT) to be Al3(Zr0.91Ti0.09) and Al3(Zr0.83Ti0.17), respectively, resulting in cal-

culated lattice parameter mismatches with α-Al of δ = +0.43% and +0.22%, respectively, which

is less than that of Al3Zr (L12), δ = +0.75%. It is important to note, too, that these mismatches

are reduced further at elevated temperature due differences in coefficient of thermal expansion

for Al3M and α-Al, as shown experimentally by Harada and Dunand [121] for Al3Sc-based tria-


Fig. 7.4: Series of optical micrographs illustrating themacrostructure of alloy Al-0.1Zr(b). The ingot cross-section is composed entirely of columnar grains (etchedusing Poultan’s reagent).

Fig. 7.5: Series of optical micrographs illustrating themacrostructure of alloy Al-0.1Zr-0.1Ti(b). The grain sizeis more equiaxed and refined than that of Al-0.1Zr(b), Fig-ure 7.4 (etched using Poultan’s reagent).


0 0 . 2 5 0 . 5 0 0 . 7 5 1 . 0 0

0 . 3 9 6

0 . 3 9 8

0 . 4 0 0

0 . 4 0 2

0 . 4 0 4

0 . 4 0 6

0 . 4 0 8

0 . 4 1 0

L 12 A l

3T i , A l

3Z r S r i n i v a s a n e t a l . ( 1 9 9 1 )

L 12 A l

3( Z r , T i ) M a l e k e t a l . ( 1 9 9 9 )

A l3T i ( L 1

2)

A l3Z r ( L 1

2)

Al 3

(Zr

1-xT

i x)

latt

ice

pa

ra

me

ter

(nm

)

S t o i c h i o m e t r i c p a r a m e t e r , x

L a t t i c e p a r a m e t e r o f p u r e A l

Fig. 7.6: Dependence of the lattice parameter (at ambient temperature) of the metastable Al3(Zr1−xTix) (L12) phase on the sto-ichiometric parameter x. Measured lattice parameters of Al3(Zr1−xTix) are from Malek et al. [292]. The lattice parameters of themetastable L12 Al3Zr (4.077 A) and Al3Ti (3.967 A) trialuminides are from Srinivasan et al. [94].

luminides.

Precipitate size

There is also a strong effect of precipitate size on σth, as predicted by Marquis and Dunand [33]

and verified experimentally in numerous studies on Al reinforced with Al3Sc (L12) precipitates [115,

118, 180–182, 363]. The threshold stress is essentially a trade-off between the Orowan stress

(which decreases with precipitate size) and repulsive interactions from modulus and lattice pa-

rameter mismatches (which increase with precipitate size), resulting in an optimum precipitate

size for creep. A difference in the Al3Zr (L12) and Al3(Zr1−xTix) (L12) precipitate radii in the Al-

0.1Zr(d) and Al-0.1Zr-0.1Ti(d) alloys studied presently may account for their disparate behavior

in creep. Based on the microstructural study of similar alloys, Chapter 5, this seems unlikely

since the microstructures of those alloys are shown to be identical after extended isothermal

aging times at 425°C (Table 5.3). Moreover, as discussed in Chapter 6, the strengthening by the

dendritically-distributed precipitates cannot be predicted completely by the Orowan dislocation

looping mechanism; precipitate radius, therefore, is not indicated by the strength.


1 0 1 5 2 0 2 5

1 0- 9

1 0- 8

1 0- 7

1 0- 6

1 0- 5

1 0- 4

1 0- 3

1 0- 2

A l - 0 . 0 6 S c - 0 . 0 6 T i

R = 7 . 4 n m

A l - 0 . 0 6 S c - 0 . 0 6 T i

R = 1 0 . 8 n m

A l - 0 . 0 6 S c - 0 . 0 6 T i

R = 5 . 8 n m

A l - 0 . 0 9 S c - 0 . 0 5 Z r

R = 4 . 8 n m

A l - 0 . 0 9 S c - 0 . 0 5 Z r

R = 8 . 1 n m

A l - 0 . 1 2 S c

R = 3 . 0 n m

A l - 0 . 0 6 S c

R = 8 . 5 n m

A l - 0 . 1 Z r ( d )

A l - 0 . 1 Z r - 0 . 1 T i ( d )

A l - Z r A l - Z r - T i

A l - S c A l - S c - T i A l - S c - Z r

�

Min

imu

m s

tra

in r

ate

(s

-1

)


��

n = 4 . 4

Fig. 7.7: Minimum creep rate at 300°C vs. applied stress,comparing data in Figure 7.2 to Al-Sc [180, 181], Al-Sc-Zr [115], and Al-Sc-Ti alloys [118]. Mean precipitate ra-dius, 〈R〉, is indicated for the Al-Sc alloys.

6 7 8 9 1 0 1 1 1 2 1 3 1 4 1 5 1 6 1 7 1 8 1 9 2 01 0

- 9

1 0- 8

1 0- 7

1 0- 6

1 0- 5

1 0- 4

1 0- 3

1 0- 2

��

��

��

��

�

�

Min

imu

m s

tra

in r

ate

(s

-1

)


� �

��

��

��

Fig. 7.8: Minimum creep rate at 350 or 400°C vs. ap-plied stress, comparing data in Figure 7.2 to Al-Sc-Ti al-loys [364].

7.4.3 Comparison to Al-Sc alloys

Figures 7.7 and 7.8 compare the present creep results on Al-0.1Zr and Al-0.1Zr-0.1Ti alloys with

those secured in the same laboratory on Al-Sc [180,181], Al-Sc-Zr [115], and Al-Sc-Ti alloys [118]

crept at 300°C, and Al-Sc-Ti alloys crept at 350 and 400°C [364], for alloys of comparable precipi-

tate volume fraction and radius (〈R〉 ca. 5–10 nm).

At 300°C, Figure 7.7, the Al-Sc–based alloys generally outperform the Al-Zr–based ones stud-

ied presently. The poorer performance of the Zr-containg alloys might reflect the inhomoge-

neous precipitate distributions discussed in Chapter 5. The lattice parameter mismatch with

α-Al of Al3Zr, δ = +0.75%, is also less than that of Al3Sc (L12), δ = +1.32% (Table 1.1), and so mis-

match effects might also be contributing to the reduced σth for the Zr-containing alloys. It is also

conceivable that the present Zr-containing alloys are not aged to the optimum precipitate size

for creep. Figure 7.7 indicates that Al-0.1Zr is comparable to Al-0.12Sc (〈R〉 = 3.0 nm) [180, 181].

Despite having a smaller volume fraction, Al-0.06Sc exhibits a much-improved threshold stress,

attributable to the larger precipitate radii (〈R〉 = 8.1 nm). The alloys crept presently correspond

approximately to the ambient temperature peak-aged strength, which is typically not optimum

for creep, as discussed. The present Al-0.1Zr(d) and Al-0.1Zr-0.1Ti(d) alloys were not investigated


by TEM; the precipitate sizes are assumed to be similar to those in Table 5.3.

Figure 7.8 compares the current results to those of van Dalen et al. [364] on Al-0.06Sc-0.06Ti

alloys crept at 350 and 400°C. Here again the Zr-containing alloys exhibit smaller values of σth;

Al-0.06Sc-0.06Ti at 400°C is comparable to Al-0.1Zr at 350°C.

7.5 Chapter Summary

Al-0.1Zr and Al-0.1Zr-0.1Ti (at.%) alloys crept at 300, 350, or 400°C exhibit creep threshold stresses

σth ca. 6–12 MPa, indicative of a climb-controlled bypass mechanism. For a given temperature,

the ternary Al-Zr-Ti alloy exhibits a smaller threshold stress than the binary Al-0.1Zr alloy. This

disparity is probably attributable to: (i) a different texture due to the reduced grain size in the

ternary Al-Zr-Ti alloy; or (ii) reduction in the lattice parameter mismatch of the Al3(Zr1−xTix)

precipitates with the α-Al solid-solution. Comparisons are made with Al-Sc and Al-Sc-M (M = Zr

or Ti) alloys with similar radii and volume fractions of precipitates.

CHAPTER 8

Final Remarks andSuggestions for Future Work

8.1 Introduction

THIS THESIS has demonstrated that Al-Zr–based alloys offer tremendous potential for devel-

oping novel high-strength, high-temperature Al-based alloys. Coherent nanometer-scale

precipitates, Al3Zr (L12), are remarkably stable at temperatures up to 500°C (0.83Tm), well be-

yond the capabilities of existing creep-resistant Al alloys. There are, however, significant limita-

tions related to the peritectic nature of Al-Zr alloys and the associated: (i) high casting tempera-

tures necessary to melt the alloy above the liquidus; (ii) fast cooling rates necessary to suppress

primary Al3M; and (iii) inability to homogenize the alloys when the dendrites are supersatu-

rated beyond the maximum solid-solubility. This chapter offers new directions for developing

improved creep-resistant alloys by considering: (i) judicious aging treatments for maximizing

strength and creep resistance of Al-Zr alloys; (ii) an alternate system to Al-Zr; or (iii) ternary and

higher-order alloying additions to Al-Zr alloys.

8.2 Improvements with Judicious Aging

Chapter 6 indicates that precipitation strengthening in segregated Al-Zr alloys occurs on mul-

tiple length scales: (i) on the nanometer-scale scale by an Orowan strengthening mechanism

and (ii) on the micrometer-scale related to the volume fraction of the precipitate-rich dendrites.

The thermal aging history consequently has a strong influence on the observed alloy strength.

Compared to isothermal aging, isochronal aging nucleates precipitates at the lowest possible

aging temperature where the solute supersaturation and chemical driving force is greatest, re-

sulting in: (i) smaller precipitates since the critical radius for nucleation is reduced; and (ii) a

larger volume fraction of precipitate-rich dendrites since the supersaturation is greater at lower

179

180 CHAPTER 8 FINAL REMARKS AND SUGGESTIONS FOR FUTURE WORK

temperatures (Figure 6.21).

In Chapter 7, the maximum temperature investigated by creep was 400°C, primarily because

beyond this temperature a precipitous drop in hardness was observed in the isothermally aged

specimens (Figures 5.7–5.10). Isochronal aging, Figure 6.12, indicates that Al-Zr alloys retain

strength to ca. 450°C, beyond which precipitate coarsening occurs along with transformation to

the equilibrium D023 structure. Furthermore, alloys isochronally aged to 450°C (Figure 6.12) are

stronger than similar alloys aged isothermally at 400°C (Figures 5.7 and 5.9) at ambient tempera-

ture. Because of the larger Orowan stress, isochronally aged creep specimens should also exhibit

a larger threshold stress in creep than the alloys studied in Chapter 7. Moreover, isochronally ag-

ing creep specimens to 450°C (0.77Tm) permits creep experiments at this temperature, where

the microstructures in Al-Zr alloys are still intrinsically stable.

8.3 Alloys Containing Ta

Based on the criteria established in Chapter 1, Al-Ta alloys seem particularly intriguing for devel-

oping new high-strength high-temperature Al-based alloys. With exception to the Al-Sc system,

Al-Ta has the most attractive solvus for precipitation hardening of the systems in Figure 1.4, with

a relatively large maximum solid-solubility (0.24 at.% Ta) that decreases significantly with tem-

perature. The maximum solid-solubility of Ta in α-Al is approximately three times that of Zr,

suggesting that Al-Ta alloys may be more amenable to solutionizing prior to precipitation aging

(Al-Zr alloys cannot be homogenized as discussed in Appendix B). This would mitigate many of

the problems associated with inhomogeneous precipitate distributions that occur from aging

as-cast Al-Zr alloys.

As a Group 5 element, Ta should be an extremely slow diffuser in α-Al (Figure 1.7), and as with

Al3V and Al3Nb, trialuminides based on the other Group 5 elements, Al3Ta could conceivably

exist as a metastable L12 structure. Table 1.1 indicates, however, that Al3Ta (L12) has not been

reported in the scientific literature.

The Al-Ta system is peritectic, however, which introduces complications with respect to casta-

bility, discussed in Chapters 1 and 2. Nevertheless, peritectic Al-Ta alloys may be conventionally

cast, provided one appreciates the necessity for high casting temperatures and relatively fast

solidification rates to suppress primary precipitation of Al3Ta.

SECTION 8.4 ALLOYING ADDITIONS TO AL-ZR ALLOYS 181

There is much to glean from a study on Al-Ta alloys. Tracer diffusivity studies of Ta in α-

Al are unknown to the author. As mentioned above, no information is available concerning the

existence of Al3Ta (L12). There are also no reported studies investigating this system with respect

to precipitation hardening.

Since the lattice parameter of Al3Ta (L12) is unknown and, anticipating extremely sluggish

diffusion kinetics, a prudent approach might initially involve Al-Ta alloys produced by rapid

solidification processing (RSP), resulting highly-supersaturated solid-solutions and bypassing

many of the challenges in conventionally casting peritectic alloys. The concomitant large chem-

ical driving force for nucleation will furthermore help to overcome any significant elastic strain

energy hindering nucleation, should one exist if the lattice parameter mismatch of Al3Ta (L12)

with α-Al is large (this is a limiting factor in developing conventionally-cast alloys based on Al-

Ti, Chapter 3). Assuming Al3Ta (L12) precipitates form during aging, coarsening studies may be

performed to estimate the diffusivity of Ta in α-Al, and transmission electron microscopy (TEM)

may elucidate the lattice parameter of Al3Ta (L12), and its coherency with α-Al. From these re-

sults, the Al-Ta system may then be evaluated from the standpoint of developing a conventionally-

solidified alloy.

Tantalum may also be beneficial as a ternary addition to Al-Sc– and Al-Zr–based alloys, where

the anticipated slower diffusion of Ta may confer improved coarsening resistance of Al3Sc (L12)

or Al3Zr (L12). Partitioning of Ta to these precipitates might also alter the lattice parameter mis-

match with α-Al, thereby further affecting coarsening resistance as well as influencing the creep

threshold stress.

8.4 Alloying Additions to Al-Zr Alloys

Inhomogeneously dendritically-distributed Al3Zr (L12) precipitates are a significant problem in

commercial wrought alloys, where Zr is added as a recrystallization inhibitor [78, 168, 188–191].

During solutionizing, which is typically performed at ca. 500°C, coherent Al3Zr (L12) precipitates

form, which inhibit recrystallization by pinning migrating grain boundaries. Because precipita-

tion of Al3Zr (L12) occurs at such a high temperature, the interdendritic precipitate-free chan-

nels are larger than those observed in the present study. As in the present study, the precipitate-

free interdendritic regions are deleterious since there is locally a lower resistance to recrystal-


lization [365]. As a result, a number of recent studies have emerged investigating the effect of

ternary alloying additions in mitigating the effects of an inhomogeneous distribution of Zr in

commercial wrought Al alloys.

8.4.1 Al-Zr-Mg Alloys

Calculations by Sigli [366] indicate that the metastable L12 Al-Zr solvus is strongly shifted to

smaller solubilities by additions of Zn, Cu, Mg, and Li (in increasing order of efficacy). Rob-

son and Prangnell [367] similarly demonstrated that additions of Zn, Cu, and Mg to commercial

7xxx alloys influences the supersaturation of Zr in α-Al, thereby accelerating the precipitation

kinetics of Al3Zr (L12) and ultimately influencing the recrystallization behavior. By decreasing

the solid-solubility of Zr in α-Al, these ternary additions should increase the volume fraction of

precipitate-rich dendrites during post-solidification aging, as schematically illustrated in Fig-

ure 6.21. Robson and Prangnell [367] predict that Mg additions have the greatest influence on

minimizing the precipitate-free interdendritic regions. In addition to the increased volume frac-

tion of preciptiate-rich dendrites created by a decreased solid-solubility of Zr, Mg additions also

confer solid-solution strengthening, further improving creep performance as observed in Al-Sc-

Mg alloys [182].

8.4.2 Al-Zr-Sc Alloys

There is currently strong interest in adding Sc to 7xxx and other commercial Al alloys to im-

prove the Al3M (L12) precipitate distribution, thereby inhibiting recrystallization [341, 368, 369].

Robson [341] has shown that by combining a Sc, eutectic-forming solute (k0 < 1), with Zr, a

peritectic-forming solute (k0 > 1), the precipitate-free regions associated with a Sc-free alloy

may be eliminated. His argument, supported by experimentally-measured solute concentra-

tion profiles across the dendrites, is that during solidification Zr and Sc solute atoms segregate

to the dendrite interiors and exteriors, respectively, effectively “filling-in” the interdendritic re-

gions with Sc, which forms Al3Sc (L12) precipitates during post-solidification aging. A similar

effect was observed by Lieblich and Torralba [370,371] with Al-Li-Ti alloys, where Ti (a peritectic-

forming solute) segregates to the dendrite interiors and Li (a eutectic-forming solute) segregates

to the periphery. A further advantage of combining Zr and Sc is the improved coarsening re-

sistance of Al3(Sc1−xZrx)compared to Al3Sc (L12) [115–117]. Finally, Sc has been calculated to

SECTION 8.4 ALLOYING ADDITIONS TO AL-ZR ALLOYS 183

2 0 0 3 0 0 4 0 0 5 0 0 6 0 0

2 6

2 7

2 8

2 9

3 0

3 1

3 2

3 3

3 4

3 5

3 6

2 0 0

2 5 0

3 0 0

3 5 0

4 0 0

4 5 0

5 0 0

5 5 0

6 0 0

6 5 0

7 0 0

7 5 0

8 0 0

8 5 0

2 0 0 3 0 0 4 0 0 5 0 0 6 0 0

A l - 0 . 1 S c

A l - 0 . 1 Z r - 0 . 1 S c ( a )

A l - 0 . 1 Z r - 0 . 1 S c ( b )

A l - 0 . 1 Z r

A l - 0 . 1 S c

( h o m o g e n i z e d )

Ele

ctr

ica

l c

on

du

cti

vit

y

(MS

m-1

)

�� A s - c a s t

A l - 0 . 1 Z r - 0 . 1 S c ( a )

A l - 0 . 1 S c

A l - 0 . 1 Z r - 0 . 1 S c ( b )

A l - 0 . 1 Z r

Vic

ke

rs

mic

ro

ha

rd

ne

ss

(M

Pa

)

A l - 0 . 1 S c


��

Fig. 8.1: Vickers microhardness and electrical conductiv-ity evolution during isochronal aging (3 h at each temper-ature) of Al-0.1Zr, Al-0.1Sc and Al-0.1Zr-0.1Sc (at.%) [373].

2 0 0 3 0 0 4 0 0 5 0 0 6 0 0

2 6

2 7

2 8

2 9

3 0

3 1

3 2

3 3

3 4

3 5

3 6

2 0 0

2 5 0

3 0 0

3 5 0

4 0 0

4 5 0

5 0 0

5 5 0

6 0 0

6 5 0

7 0 0

7 5 0

8 0 0

8 5 0

2 0 0 3 0 0 4 0 0 5 0 0 6 0 0

A l - 0 . 0 6 S c

A l - 0 . 0 6 Z r - 0 . 0 6 S c ( a )

A l - 0 . 0 6 Z r - 0 . 0 6 S c ( b )

A l - 0 . 0 6 Z r

A l - 0 . 0 6 S c


Ele

ctr

ica

l c

on

du

cti

vit

y

(MS

m-1

)

�� A s - c a s t

A l - 0 . 0 6 S c

A l - 0 . 0 6 Z r - 0 . 0 6 S c ( b )

A l - 0 . 0 6 Z r

A l - 0 . 0 6 Z r - 0 . 0 6 S c ( a )

Vic

ke

rs

mic

ro

ha

rd

ne

ss

(M

Pa

)

A l - 0 . 0 6 S c


��

Fig. 8.2: Vickers microhardness and electrical conduc-tivity evolution during isochronal aging (3 h at eachtemperature) of Al-0.06Zr, Al-0.06Sc and Al-0.06Zr-0.06Sc(at.%) [373].

stabilize the L12 structure of Al3Zr [372].

This thesis is closed by returning to the conclusions posited in Chapter 1. Of all the so-

lutes capable of forming L12-structured Al3M precipitates, Sc and Zr were shown to exhibit

the most promise for developing castable, creep-resistant Al-based alloys. Scandium is singu-

lar among the transition metal solutes in that it forms a thermodynamically stable cubic L12

trialuminide, providing the highest increment of strengthening per atomic percent of any alloy-

ing element added to Al [178]. Zirconium is an extremely slow diffuser in α-Al, Figure 1.7, and

forms metastable L12-structured Al3Zr precipitates that exhibit impressive thermal stability to

ca. 500°C, discussed in Chapters 5 and 6.


Figures 8.1 and 8.2 present results of a preliminary study [373] comparing the precipitation

behavior in several binary Al-Zr and Al-Sc alloys during 3 h isochronal aging, as monitored by

Vickers microhardness and electrical conductivity. Precipitation of Al3Sc (L12) commences be-

tween 200 and 250°C, with peak hardness occurring at 325°C; precipitation of Al3Zr (L12) is com-

paratively delayed, commencing between 350 and 375°C with a peak at 400–500°C, depending

on the supersaturation. Figures 8.1 and 8.2 summarize, in a nutshell, the merits of the two sys-

tems: Al-Sc providing improved precipitation hardening and Al-Zr conferring improved thermal

stability.

Ternary Al-Zr-Sc alloys containing equiatomic fractions of Zr and Sc (2:1 Zr:Sc by weight)

are also studied in Figures 8.1 and 8.2, where a pronounced synergistic effect is observed when

both Zr and Sc are added to Al alloys. At 325°C, where peak hardness occurs in the binary Al-Sc

alloys, Zr additions result in a secondary increase in strength that is attributed to precipitation

of Al3Zr (L12) at higher temperatures. The ternary alloy reaches peak strength at 400°C, delaying

overaging by over 100°C compared to the Zr-free alloy.

These results are preliminary and are not entirely understood at present. The increased

strength of the Al-Zr-Sc alloys compared to their respective binaries may be attributable to the

phenomenon observed by Robson [341], where the precipitate-free interdendritic channels are

filled with Al3Sc (L12) precipitates (the precipitate-free interdendritic regions limit the precip-

itation strengthening of Al3Zr (L12) in Al-Zr alloys, Chapter 6). Preliminary 3-D atom-probe

tomographic analyses [373] on Al-0.1Zr-0.1Sc (at.%) alloys indicate that the increase in strength

comes from segregation of Zr to pre-existing Al3Sc precipitates, thereby decreasing the interpre-

cipitate distance, λ, and increasing the yield stress of the alloy. Perhaps it is a combination of

both phenomena.

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[340] A. D. Brailsford and P. Wynblatt, “The Dependence of Ostwald Ripening Kinetics on Particle Volume

Fraction,” Acta Metall, 27(3), pp. 489–497 (1979).

[341] J. D. Robson, “A New Model for Prediction of Dispersoid Precipitation in Aluminium Alloys Con-

taining Zirconium and Scandium,” Acta Mater, 52(6), pp. 1409–1421 (2004).

[342] R. Mehrabian, “Rapid solidification,” Int Met Rev, 27(4), pp. 185–208 (1982).

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tion to Aluminum-Alloys,” Int Mater Rev, 37(1), pp. 1–44 (1992).

[344] H. M. Lee and H. Lee, “Site Fraction Behavior of Transition-Elements in Al3(Ti,V,Zr) Phase,” CAL-

PHAD, 16(1), pp. 19–24 (1992).

[345] R. W. Baluffi, S. M. Allen, and W. C. Carter, Kinetics of Materials, John Wiley & Sons, Hoboken (2005).

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[347] L. M. Brown and R. K. Ham, Strengthening Methods in Crystals, Elsevier Publishing Company, Lon-

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[348] S. M. Jeon and J. K. Park, “Transition Behavior of Deformation Mode from Shearing to Looping in

Al-Li Single Crystals,” Phil Mag A, 70(3), pp. 493–504 (1994).

[349] S. M. Jeon and J. K. Park, “Precipitation Strengthening Behavior of Al-Li Single Crystals,” Acta Mater,

44(4), pp. 1449–1455 (1996).

[350] M. Nakamura and K. Kimura, “Elastic Constants of TiAl3 and Zr Al3 Single Crystals,” J Mater Sci,

26(8), pp. 2208–2214 (1991).

[351] M. S. Zedalis, M. V. Ghate, and M. E. Fine, “Elastic Moduli of Al3Zr,” Scripta Metall, 19(5), pp. 647–

650 (1985).

[352] E. P. George, J. A. Horton, W. D. Porter, and J. H. Schneibel, “Brittle Cleavage of L12 Trialuminides,”

J Mater Res, 5(8), pp. 1639–1648 (1990).

[353] C. L. Fu, “Electronic, Elastic, and Fracture Properties of Trialuminide Alloys: Al3Sc and Al3Ti,” J

Mater Res, 5(5), pp. 971–979 (1990).

[354] R. W. Hyland, Jr. and R. C. Stiffler, “Determination of the Elastic Constants of Polycrystalline Al3Sc,”

Scripta Metall Mater, 25(2), pp. 473–477 (1991).

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[356] E. A. Marquis, Microstructural Evolution and Strengthening Mechanisms in Al-Sc and Al-Mg-Sc Al-

loys, Ph.D. thesis, Northwestern University (2002).

[357] K. Fukunaga, T. Shouji, and Y. Miura, “Temperature Dependence of Dislocation Structure of L12–

Al3Sc,” Mater Sci Eng A, 239-240, pp. 202–205 (1997).

[358] D. Tabor, “The Physical Meaning of Indentation and Scratch Hardness,” Br. J. Appl. Phys., 7(5), pp.

159–166 (1956).

[359] A. K. Mukherjee, J. E. Bird, and J. E. Dorn, “Experimental Correlations for High-Temperature Creep,”

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[360] J. H. Gittus, “Theoretical Equation for Steady-State Dislocation Creep in a Material Having a Thresh-

old Stress,” Proc R Soc Lond, A, 342(1629), pp. 279–287 (1975).

[361] J. C. Gibeling and W. D. Nix, “The Description of Elevated Temperature Deformation in Terms of

Threshold Stresses and Back Stresses: A Review,” Mater Sci Eng, 45(2), pp. 123–135 (1980).

[362] R. Lagneborg and B. Bergman, “The Stress/Creep Rate Behavior of Precipitation-Hardened Alloys,”

Metal Science, 10, pp. 20–28 (1976).

[363] R. A. Karnesky, D. N. Seidman, and D. C. Dunand, “Creep of Al-Sc Microalloys with Rare-Earth Ele-

ment Additions,” Mater Sci Forum, 519-521, pp. 1035–1040 (2006).

[364] M. van Dalen, D. C. Dunand, and D. N. Seidman, “Creep Data for Al-Sc-Ti Alloys, Personal Commu-

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[365] J. D. Robson and P. B. Prangnell, “Dispersoid Precipitation and Process Modelling in Zirconium

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[366] C. Sigli, “Zirconium Solubility in Aluminum Alloys,” in International Conference on Aluminum Al-

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[367] J. D. Robson and P. B. Prangnell, “Modelling Al3Zr Dispersoid Precipitation in Multicomponent Alu-

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[368] Y. W. Riddle and T. H. Sanders, “A Study of Coarsening, Recrystallization, and Morphology of Mi-

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[369] B. Forbord, H. Hallem, N. Ryum, and K. Marthinsen, “Precipitation and Recrystallisation in Al-Mn-

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[370] M. Lieblich and M. Torralba, “Cellular Microstructure and Heterogeneous Coarsening of δ′ in

Rapidly Solidified Al-Li-Ti Alloys,” J Mater Sci, 26(16), pp. 4361–4368 (1991).

[371] M. Lieblich and M. Torralba, “Aging Characteristics of Rapidly Solidified Al-Li-Ti Alloys at 473 K,” J

Mater Sci, 27(13), pp. 3474–3478 (1992).

[372] J. H. Xu and A. J. Freeman, “Phase Stability and Electronic Structure of ScAl3 and ZrAl3 and of Sc-

Stabilized Cubic ZrAl3 Precipitates,” Phys Rev B, 41(18), pp. 12,553–12,561 (1990).

[373] C. P. Lee, K. E. Knipling, D. C. Dunand, and D. N. Seidman, In preparation (2006).

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of Ultradispersed Phases in Rapidly Solidified Aluminium Alloys,” Nanostructured Materials, 6(1-4),

pp. 477–479 (1995).

[375] I. Brodova, V. M. Zamyatin, P. Popel, V. O. Esin, B. A. Baum, A. I. Moiseyev, I. P. Korshunov, A. L.

Topchii, Y. G. Tikhomirov, and I. V. Polents, “Conditions of Formation of Metastable Phases During

Crystallization of Al-Zr Alloys,” Rasplavy, 2(6), pp. 23–27 (1988).

[376] I. G. Brodova, D. V. Bashlykov, A. B. Manukhin, T. I. Yablonskikh, and N. A. Zolotova, “Effect of the

Time-Temperature Melt Treatment on the Structure and Phase Composition of a Rapidly Quenched

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[377] I. Brodova, D. Bashlykov, A. Manukhin, E. Rozhicyna, P. Popel, and V. Manov, “Disperse Structure

Forming in Rapidly Quenched Al-Hf Alloy,” Mater Sci Eng A, 304, pp. 544–547 (2001).

[378] E. Nes and N. Ryum, “On the Formation of Fan-Shaped Precipitates During Decomposition of a

Highly Supersaturated Al-Zr Solid Solution,” Scripta Metall, 5(11), pp. 987–990 (1971).

[379] E. Nes and H. Billdal, “The Mechanism of Discontinuous Precipitation of Metastable Al3Zr Phase

from an Al-Zr Solid-Solution,” Acta Metall, 25(9), pp. 1039–1046 (1977).

[380] O. Beeri, D. C. Dunand, and D. N. Seidman, “Effect of Fe and Si Contamination on Preciptiation in

Dilute Al-Sc Alloys, Personal Communication,” (2006).

APPENDIX A

Alloy Preparation and Castability

Peritectic alloys are not amenable to conventional casting because of the necessary liquidus

elevation and the potential for properitectic precipitation in the melt, as discussed in Chap-

ter 2. The alloys must be superheated well beyond the melting point of Al, and the cooling

rate and/or solute concentration must be controlled so as to avoid primary Al3M. Initially,

however, these difficulties were not appreciated and efforts to conventionally cast1 Al-Ti– and

Al-Zr–based alloys were invariably met with failure. Large primary Al3Ti or Al3Zr precipitates

formed, Figure A.1, and the alloys exhibited no appreciable precipitation hardening during post-

solidification aging.

The presence of primary precipitates, like those in Figure A.1, was initially (and erroneously)

believed to indicate that the α-Al solid-solution surrounding the needle-like primary precipi-

tates is supersaturated in solute. In fact, the exact opposite is true — the presence of primary

precipitates ensures that the α-Al solid-solution has a composition following the Al3M liquidus,

meaning the last solid to form has a composition CLp, Figure 2.1, that is necessarily less than

maximum solid-solubility.

A.1 Advantages of Arc-Melting

It was revealed, largely through trial and error, that arc-melted alloys consistently exhibit stronger

precipitation hardening than those of similar composition that were conventionally-cast.1 The

reasons for this are manifold (in order of decreasing significance):

1. High melt temperatures The temperature of the arc is much greater than that obtained in a

resistively-heated furnace. High temperatures are required to melt into the single-phase

melt since the liquidus of a peritectic alloy increases rapidly with solute content.1By melting in a resistively-heated furnace (in air), and casting into graphite molds.

213

214 APPENDIX A ALLOY PREPARATION AND CASTABILITY

Fig. A.1: Optical micrograph of a conventionally-solidified hyperperitectic Al-Ti alloy. Numerous plate-like Al3Ti primary precipi-tates are observed.

2. Inert atmosphere The arc-melter chamber has a gettered argon atmosphere, which mini-

mizes oxidation of the solutes and maintains accurate alloy compositions.

3. Accelerated solidification rates Solidification in the arc-melter is effected by an underlying

continuously water-cooled copper hearth, which might accelerate the solidification of the

alloys. Accelerated solidification cooling rates are required to suppress properitectic pre-

cipitation, as discussed in Chapter 2.

A.1.1 High melt temperatures

The interrelation between solidification rate, solute concentration, and solidified microstruc-

ture in peritectic alloys was discussed in Chapter 2. In addition to these factors, Brodova et al.

investigated the effect of melt superheat (above the liquidus) on the microstructures of dilute

Al-Ti [225,374], Al-Zr [218,374,375], Al-Hf [376,377] alloys. It was shown that increasing the melt

temperature and/or hold time above the liquidus has the same effect as increasing the cool-

ing rate: i.e., primary Al3M precipitates become smaller and more equiaxed; the likelihood for

primary precipitation of metastable phases is increased; and, for sufficiently large superheat-

ing, primary precipitation of Al3M may be suppressed completely. Brodova et al. argued that

APPENDIX A ALLOY PREPARATION AND CASTABILITY 215

significant superheating is necessary to transform the liquid from a colloidal state (consisting

fine Al3M primary precipitates dispersed in the melt) to a homogenous single phase. This same

phenomenon is responsible for so-called ‘fade’ in commercial Al casting alloys, where the grain

refining efficacy of Ti additions falls off with holding time of the melt [156, 157, 161, 163, 168].

A.1.2 Inert atmosphere

Alloys produced by conventional means, melting in air, were generally found to be significantly

depleted in solute compared to their nominal initial compositions. This was attributed to oxi-

dation of Zr and/or Ti during melting. Ingot melting in the arc-melter occurs in an inert argon

atmosphere, mitigating solute loss from oxidation.

A.1.3 Accelerated solidification rates

Solidification in the arc-melter is effected by the water-chilled copper hearth, which might re-

sult in an enhanced solidification rate since the mold is continuously cooled after melting. In

actuality, the effect is probably not significant since, unlike in a conventional mold where heat

is removed through multiple surfaces, heat is extracted only through the bottom of the ingot in

the arc-melter. Indeed, solidification of the button ingot can be observed by eye to occur over

the course of several seconds, indicating that the solidification rate is at most 101–102°C s-1.

APPENDIX B

Homogenization and Other Prior AnnealingTreatments

The alloys in this thesis were not homogenized because, empirically, homogenization was found

to be deleterious to the attainable peak strength achieved during the ensuing isothermal aging

treatment of Al-Zr–based alloys. Other annealing treatment (e.g., pre-aging at 500°C) prior to the

precipitation aging heat treatment had a similar effect. Figure B.1 shows the observed hardness

of an Al–0.2 at.% Zr subjected to various initial annealing treatments, prior to isothermal aging

at 375°C for 100 h. There is a systematic decline in peak hardness attained by aging at 375°C with

increasing thermal exposure prior to aging.

B.1 Homogenization (or Solution) Anneals

A homogenization, or solution, heat treatment involves heating the alloy in the single phase α-Al

solid-solution and allowing diffusion to eliminate the solute microsegregation incurred during

solidification, thereby homogenizing the solute distribution in the alloy. For peritectic alloys

like Al-Zr with k0 > 1, however, the dendrites can be supersaturated well beyond the maximum

solid-solubility of Zr in α-Al. During homogenization, there will be a tendency for precipita-

tion to occur since, locally, the alloy composition does not lie in the single phase α-Al solid-

solution. Because most of Zr atoms are precipitated into coarse (likely equilibrium D023) Al3Zr

precipitates during homogenization, the achievable hardness during the ensuing isothermal ag-

ing treatment is reduced dramatically as demonstrated in Figure B.1.

Figure B.2 supports this hypothesis of precipitation of Al3Zr during homogenization. Dis-

played are two scanning electron microscope (SEM) micrographs of a conventionally-solidified

Al-2.36Zn-0.26Zr (at.%) alloy after homogenization at 640°C for 650 h. As discussed previously,

conventional solidification often results in large primary Al3Zr or Al3Ti precipitates; an example

of a primary Al3Zr precipitate is visible in Figure B.2(a). Smaller plate-like precipitates, ca. 2–

217

218 APPENDIX B HOMOGENIZATION AND OTHER PRIOR ANNEALING TREATMENTS

100

200

300

400

500

600

700

800

20

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Fig. B.1: Observed Vickers microhardness of Al-0.2 at.% Zr subjected to various annealing treatments prior to isothermal aging at375°C for 100 h. Pre-aging was performed at 500°C for 1 h. Homogenization was carried out at 640°C for 240 h.

Fig. B.2: SEM micrographs of a conventionally-solidified Al-2.36Zn-0.26Zr (at.%) alloy following homogenization at 640°C for 650 h.

APPENDIX B HOMOGENIZATION AND OTHER PRIOR ANNEALING TREATMENTS 219

4 nm in length, are visible in Figure B.2(b). Unlike the large primary phase, these smaller pre-

cipitates exhibit one of three unique orientations in Figure B.2(b), suggesting that they formed

in the solid-state and have a cube-on-cube orientation relationship with α-Al. These precipi-

tates are probably disc-like equilibrium D023-structured Al3Zr precipitates, like those described

in Chapter 6, formed in locally supersaturated regions of the alloy during homogenization at

640°C. This conclusion concurs also with that of Dahl et al. [82], who observed orthogonal D023-

structured precipitates after homogenizing alloys containing 0.03–0.36 at.% Zr for 25 h at 640°C.

Dahl et al. also concluded that Al-Zr alloys are incapable of precipitation hardening following

homogenization treatments, consistent with the experimental evidence in Figure B.1.

B.2 Pre-Aging at 500°C

Pre-aging at 500°C similarly results in a significant decrease in hardness attained during subse-

quent aging at 375°C (Figure B.1). The purpose of the 500°C pre-age is to suppress discontinuous

precipitation, which can occur in Al-Zr alloys [378,379]. Pre-aging has been employed in several

studies on Al-Zr, Al-Zr-V, Al-Zr-Ti, and Al-Zr-Ti-V alloys as discussed in Table 5.4.

It is known that Al3Zr (L12) precipitates are nucleated during pre-aging at 500°C, as shown

by Tsau and Chen [296] using transmission electron microscopy (TEM). As discussed in Fig-

ure 6.21, higher isothermal aging temperatures yield a greater volume fraction of precipitate-free

interdendritic regions because nucleation is limited to the inner-most regions of the solute-rich

dendrites. Since precipitates are nucleated at 500°C, the volume fraction of precipitate-free in-

terdendritic regions is greater than if aging had been carried out isothermally at 375°C, thus

explaining the reduced peak hardness for the pre-aged specimen.

Zedalis and Fine [84] claimed that the necessity of pre-aging is dependent on the temper-

ature of the ensuing aging treatment. For their coarsening studies on Al-Zr and Al-Zr-V alloys

at 375, 400, and 425°C, pre-aging was necessary only at lower aging temperatures; the discon-

tinuous precipitation mechanism was not observed at 425°C. Discontinuous precipitation, het-

erogenous nucleation on a moving grain boundary, is prevalent at lower aging temperatures

where sluggish diffusion kinetics may preclude homogenous nucleation. Since their arc-melted

alloys were not pre-aged prior to aging at 425°C, the coarsening study of Al3Zr (L12) precipitates

by Zedalis and Fine [84] is directly comparable to the data for Al3Zr (L12) and Al3(Zr1−xTix) (L12)

220 APPENDIX B HOMOGENIZATION AND OTHER PRIOR ANNEALING TREATMENTS

precipitates presented in Chapter 5.

APPENDIX C

Length Scales of Analytical Techniques andPotential Biases

The alloys studied in this thesis are highly segregated, as discussed an Chapter 5 and shown in

Figure 5.6. During aging, the initially solute-rich dendrites produce homogeneously-distributed

nanometer-scale Al3Zr (L12) or Al3(Zr1−xTix) (L12) precipitates which demarcate the dendritic

and interdendritic regions of the alloy. The precipitate-rich dendritic regions, when observed

in 2-D projection by transmission electron microscopy (TEM) or scanning electron microscopy

(SEM), are of the order of 1–10 µm across and vary considerably in size and morphology across

length scales of the order of 100 µm. The width of the interdendritic precipitate-free channels

similarly varies.

C.1 Length Scales of Analytical Techniques

Figure C.1 is a visual depiction of the primary analytical techniques used in this study. These are

(in increasing length scale): 3-D atom-probe tomography (3-D APT), transmission electron mi-

croscopy (TEM), scanning electron microscopy (SEM), Vickers microhardness, creep, and elec-

trical conductivity. Because of microsegregation of solutes occurring on the 1–100 µm scale,

what is observed on the sub-micrometer-scale cannot be straightforwardly correlated to what is

measured on the macroscopic scale. As discussed in Chapter 4, the measured solute concentra-

tions within a 3-D APT analysis volume varies considerable from analysis to analysis, depending

on the proximity to the solute-rich dendrites. The observed precipitate size and morphology by

TEM, discussed in Chapter 5, is similarly strongly dependent on where one looks. Figure C.2

is an SEM micrograph displaying several precipitate-rich dendrites in relation to the hole in a

TEM foil. The microstructure observable by TEM, which is limited to electron-transparent ar-

eas, is a small sample of the actual microstructure varying on the micrometer-scale. Indeed,

montages of several micrographs, e.g. Figures 5.21, 6.2, and 6.3, are usually required to study the

221

222 APPENDIX C LENGTH SCALES OF ANALYTICAL TECHNIQUES AND POTENTIAL BIASES

10-9 10-8 10-7 10-6 10-4 10-310-5 10-2

Length scale (m)

ElectricalconductivityAtom-probe

tomographyVickersmicrohardness Creep

Transmissionelectron microscopy

Scanningelectron microscopy

10-10

Dendrites

Fig. C.1: Length scales of the analytical techniques used. Those techniques sampling regions on a scale larger than the dendrites(> 10−4 m) are unbiased by the dendritic distribution of precipitates, whereas those techniques limited to the submicron scale(< 10−6 m) are highly biased by the inhomogeneous distribution of solutes and precipitates.

micrometer-scale of the dendrites by TEM.

Fortunately, the larger-scale techniques in Figure C.1 are largely unbiased by the dendritic

distribution of solutes and precipitates. Figure C.3 indicates that the tip-to-tip length of a Vickers

microhardness pyramidal indent is of the order of 100 µm for the alloys studied. Considering

the scale of the microsegregation, Figure 5.6, a single hardness measurement therefore samples

several precipitate-rich dendritic regions and precipitate-free interdendritic channels, meaning

that the measured hardness value is an average of the strong and weak regions of the alloy and is

generally not subject to sampling biases. This applies also to creep and electrical conductivity.

C.2 Biases Associated with Sample Preparation

The dendritic distribution of precipitates was unrecognized for some time, primarily due to pref-

erential chemical etching during electropolishing of TEM foils. Figure C.4 displays low mag-

nification TEM micrographs illustrating preferential electropolishing around the precipitate-

rich dendrites. The small dendritic precipitates are consequently hard to image, and are eas-

ily missed, because of absorption from incoherent scattering occurring in the thicker regions

of the foil. The most apparent population of precipitates observed by TEM is the larger (〈R〉

ca. 20 nm) spheroidal interdendritic precipitates, whereas the most numerous population is the

smaller (〈R〉 ca. 5–10 nm) dendritic precipitates. The increased foil thickness can also accentuate

the apparent volume fraction of precipitates in the dendrites.

APPENDIX C LENGTH SCALES OF ANALYTICAL TECHNIQUES AND POTENTIAL BIASES 223

Fig. C.2: SEM micrograph displaying precipitate-rich dendrites around the specimen edge of an electropolished TEM foil.

5 0 7 5 1 0 0 1 2 5 1 5 0

1 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

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9 0 0

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2 0

3 0

4 0

5 0

6 0

7 0

8 0

9 0

1 0 0

Vic

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ne

ss

(M

Pa

)

I n d e n t d i a g o n a l l e n g t h ( µm )

Ra

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ard

ne

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C o r r e s p o n d i n g i n d e n t s i z e

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(H

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00)

d

Fig. C.3: Vickers microhardness value as a function of indent diagonal length for a 200 g load (used on all hardness measurements).

224 APPENDIX C LENGTH SCALES OF ANALYTICAL TECHNIQUES AND POTENTIAL BIASES

Fig. C.4: Bright-field and centered superlattice dark-field TEM micrographs of Al-0.1Zr and Al-0.1Zr-0.1Ti (at.%) alloys, respectively,after aging at 375°C. The foils preferentially electropolish around the precipitate-rich dendrites.

As in TEM foil preparation, preferential electropolishing in the precipitate-free interdendritic

regions probably also occurs during preparation of 3-D APT specimens since the electropolish-

ing procedures are similar. There is therefore a tendency for necking in the interdendritic re-

gions, resulting in 3-D APT specimens whose apexes (the region analyzed) are largely devoid

of precipitates. This theory explains why even very large 3-D APT analysis volumes are often

precipitate-free, as discussed in Chapter 4.

APPENDIX D

Study on an Al-Zr-Ti Alloy Prepared byParameswaran et al.

A brief study, summarized here, was carried out on an arc-melted Al-Zr-Ti alloy investigated

previously by Parameswaran et al. [288]. Prof. M. E. Fine kindly provided as-cast arc-melted but-

tons from that study, which demonstrated a dramatic improvement in coarsening resistance of

Al-Zr-Ti alloys compared to Al-Zr or Al-Zr-V alloys [84] prepared by similar arc-melting proce-

dures. The reported improved coarsening resistance is displayed in Figure D.1, reproduced from

reference [288]. The as-cast button provided by Prof. Fine, labeled “AZT 2,” had a composition of

Al-0.19Zr-0.18Ti (at.%) as confirmed presently by Bodycote Materials Testing (Skokie, IL). This

is comparable to the alloy composition reported in [288], Al-0.20Zr-0.17Ti (at.%), indicating that

the alloy described presently is comparable (if not identical) to that studied in [288].

The present author repeated the aging procedures in [288], which involved pre-aging the al-

loys at 500°C for 1 h prior to isothermal aging at 425°C. Pre-aging had been reported to suppress

discontinuous precipitation, as discussed in Appendix B. Precipitation hardening was mon-

itored by Vickers microhardness, Figure D.2, which was not performed by Parameswaran et

al. [288]. During the 500°C pre-age, Al3(Zr1−xTix) (L12) precipitates are nucleated, as indicated

by the increase in hardness and shown elsewhere by Tsau and Chen using transmission electron

microscopy (TEM) [296]. During the subsequent aging treatment at 425°C, there is negligible

additional strengthening1 and the L12 precipitates overage sluggishly.

The microstructure from specimen “AZT 2” aged at 425°C for 400 h was investigated using

TEM, offering a direct comparison with data in Figure D.1. The dark-field TEM micrograph of the

Al3(Zr1−xTix) precipitates observed are displayed in Figure D.3. The mean radius, 〈R〉, of the 104

precipitates in Figure D.3 is 19.2±3.6 nm, which is inconsistent with the measured 〈R〉 = 6.9 nm

1After pre-aging, the alloy reaches a peak strength of only ca. 420 MPa during the ensuing isothermal aging treat-ment, consistent with the results in Figure B.1. Comparable alloys that are not pre-aged are much harder, Figures 5.8and 5.10.

225

226 APPENDIX D STUDY ON AN AL-ZR-TI ALLOY PREPARED BY PARAMESWARAN ET AL.

Fig. D.1: Variation of spheroidal L12 precipitate radius vs. time for Al3(Zr1−xTix) [288] (curve A) and Al3(Zr1−xVx) [84] (curve B)precipitates at 425°C. From Parameswaran et al. [288].

reported by Parameswaran et al. [288] (curve A in Figure D.1). The precipitates in Figure D.3

are, however, quite similar to those measured by Zedalis and Fine, 〈R〉 = 21.7 nm, on Al-Zr and

Al-Zr-V alloys [84] (curve B in Figure D.1).

Because of the disparity in observed precipitate size from what Parameswaran et al. [288] re-

ported, further work on this alloy was curtailed. In hindsight, and armed with the understanding

of the various precipitate morphologies discussed in Chapter 5, it is obvious by their ca. 20 nm

radius and their slightly lobed morphology that the precipitates observed by the present author

in Figure D.3 are interdendritic, with the dendrite center out of view toward the bottom right of

the micrograph (see also Figure 5.13). It is emphasized that the alloy whose microstructure is

displayed in Figure D.3 was not made by the present author ; dendrite formation, and the atten-

dant microsegregation of solutes, is ubiquitous in conventionally-solidified alloys.

The difference in the curves of Figure D.1 — and hence the apparent improved coarsening

resistance of Al3(Zr1−xTix) precipitates compared to Al3Zr and Al3(Zr1−xVx) (L12) — is there-

fore simply an artifact of two inidividuals viewing the same heterogeneos microstructure and

reporting two different precipitate radii. After 400 h at 425°C, Parameswaran et al. [288] reported

APPENDIX D STUDY ON AN AL-ZR-TI ALLOY PREPARED BY PARAMESWARAN ET AL. 227

0 . 1 1 1 0 1 0 0 1 0 0 0

1 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

7 0 0

8 0 0

2 0

3 0

4 0

5 0

6 0

7 0

8 0

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Vic

ke

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Fig. D.2: Vickers microhardness vs. aging time for Al-0.19Zr-0.18Ti (at.%) at 425°C.

Fig. D.3: Centered superlattice dark-field TEM micrograph of Al3(Zr1−xTix) (L12) precipitates observed after aging at 425°C for400 h (preceded by a 500°C 1 h pre-age). The precipitates have a mean radius 〈R〉 of 19.2±3.6 nm.

228 APPENDIX D STUDY ON AN AL-ZR-TI ALLOY PREPARED BY PARAMESWARAN ET AL.

〈R〉 = 6.9 nm precipitates. In Figure D.3, the precipitate radius observed by the present author,

on the same alloy, is 〈R〉 = 19.2 nm. It is a classic case of “missing the forest for the trees,” or in

this case, the dendrites for the precipitates.

APPENDIX E

Potential Contamination of Fe and/or Si, andits Effects on Precipitation

As discussed in Sections 3.2 and 4.2, one source of Al used in preparing the alloys in this the-

sis was supplied by Atlantic Equipment Engineers (Bergenfield, NJ). Though nominally 99.99%

pure, independent chemical analysis [380] revealed that the Al was contaminated with 260 and

257 at.ppm of Si and Fe, respectively. These are common impurities in Al, that can produce

dramatically accelerated nucleation kinetics and enhanced precipitation strengthening in Al-

Zr [81, 267], Al-Hf [89, 91, 92], and Al-Sc [380] alloys.

Figure E.1 displays scanning electron microscope (SEM) micrographs of as-cast Al-0.2Zr-

0.2Ti (compositions in Table 5.1), displaying numerous petal-like properitectic Al3(Zr1−xTix)

precipitates (discussed previously in Figure 5.4 and Table 5.2) causing pronounced grain re-

finement in this alloy. Decorating the grain boundaries are other precipitates, confirmed to be

Fe-rich by energy dispersive X-ray spectroscopy (EDS) in SEM. The precipitates exhibit a lamel-

lar structure, characteristic of a eutectic constituent, and are believed to be eutectic α-Al/Al3Fe

formed at the end of solidification.

It is unknown as to what degree the various alloys in this thesis are contaminated with Fe

and/or Si, and efforts are currently underway to verify the concentrations of these impurities.

Nevertheless, Figure E.1 clearly demonstrates that there is a sensible amount of Fe contamina-

tion in the richest Al-Zr-Ti alloy studied; it is perhaps most apparent in this alloy because this has

the finest grain structure, Figure 5.1. Since all alloys were made by similar methods, it is reason-

able to suspect that they are also equally contaminated, meaning that the comparisons between

the various alloys are still valid. Moreover, as shown in Figure E.1(a), the interaction of Fe and/or

Si with Zr and/or Ti may be limited since the solutes segregate opposite to one another during

solidification; Fe and Si are eutectic-forming constituents that segregate to the grain boundaries,

whereas Zr and Ti are peritectic-forming constituents that segregate to the dendrite centers.

229

230 APPENDIX E POTENTIAL CONTAMINATION OF FE AND/OR SI, AND ITS EFFECTS ON PRECIPITATION

Fig. E.1: SEM secondary and backscattered electron images displaying Fe-rich precipitates decorating the grain boundaries in as-solidified Al-0.2Zr-0.2Ti. Panel (a) is a secondary electron image. Z-contrast delineates solute-rich dendritic cells (lighter contrast).Numerous petal-like Al3(Zr1−xTix) primary precipitates are also observed, leading to pronounced grain refinement. Panel (b) isa backscattered electron image displaying a detailed view of the primary Al3(Zr1−xTix) precipitates and Fe-rich precipitates atthe grain boundaries. Panel (c) is a backscattered electron image, showing a grain boundary triple junction decorated by lamellarα-Al/Al3Fe eutectic.

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