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Phase transformations in the central portion of the Nb-Ti-Al ternary system

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Title:
Phase transformations in the central portion of the Nb-Ti-Al ternary system
Creator:
Hoelzer, David Timothy, 1959-
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English
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xvii, 320 leaves : ill. ; 29 cm.

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Subjects / Keywords:
Alloys ( jstor )
Atoms ( jstor )
Conceptual lattices ( jstor )
Cooling ( jstor )
Electrical phases ( jstor )
Heat treatment ( jstor )
Lattice parameters ( jstor )
Parallel plates ( jstor )
Symmetry ( jstor )
Wave diffraction ( jstor )
Dissertations, Academic -- Materials Science and Engineering -- UF ( lcsh )
Materials Science and Engineering thesis, Ph. D ( lcsh )
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bibliography ( marcgt )
non-fiction ( marcgt )

Notes

Thesis:
Thesis (Ph. D.)--University of Florida, 1996.
Bibliography:
Includes bibliographical references (leaves 314-319).
General Note:
Typescript.
General Note:
Vita.
Statement of Responsibility:
by David Timothy Hoelzer.

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University of Florida
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University of Florida
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Copyright [name of dissertation author]. Permission granted to the University of Florida to digitize, archive and distribute this item for non-profit research and educational purposes. Any reuse of this item in excess of fair use or other copyright exemptions requires permission of the copyright holder.
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36683097 ( OCLC )

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PHASE TRANSFORMATIONS IN THE CENTRAL PORTION
OF THE Nb-Ti-Al TERNARY SYSTEM








By


DAVID TIMOTHY HOELZER












A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY


UNIVERSITY OF FLORIDA


1996















ACKNOWLEDGEMENTS


I would like to thank my graduate advisor, Dr. Fereshteh Ebrahimi, for all of

her input and continued support. I am especially grateful to her and to Dr. Michael

Kaufman for the insights and the fruitful discussions we had on this thesis topic. I

also would like to thank the other members of my committee: Dr. Ellis Verink, Jr.,

Dr. Robert Dehoff, and Dr. Anna Brajter-Toth for making my defense of this thesis

truly a memorable occasion for me.

This research work was funded by DARPA under the contract number

N00014-88-J-1100 while I was a full time graduate student at the University of

Florida. However, appreciation for the use of the JEOL 2000FX ASTEM and the

darkroom facilities goes to the NYS College of Ceramics at Alfred, NY where I was

employed during the completion of this thesis.

I would like to thank Pratt and Whitney and in particular Mr. Maloney for so

willingly making the initial samples in the compositions that we specified. I would

also like to thank Mr. Wayne Acree for the microprobe work.

A special thank you goes to all of my friends, which includes at the top of the

list Dr. Wishy Krishnamoorthy. I really appreciated the housing provided by Wishy

while I traveled to UF on thesis related business. Those escapes to Wishys' place and

discussions which we had were physical and mental boosts for me.

Finally, I would like to express my most sincerest thanks to my wife Amy for

her total support and for the sacrifices she made that allowed for me to finish this

thesis. I can now make up for lost time with both her and Rachel, my daughter.
















TABLE OF CONTENTS


Page

ACKNOWLEDGMENTS. ........... ................... ii

LIST OF TABLES .................. ..... ...... .. v

LIST OF FIGURES ................... ........... vii

ABSTRACT.......... ............................ xvi

CHAPTER

1 INTRODUCTION ................. ....... ..... 1

2 LITERATURE SURVEY .................. ..... 5
2.1 The Ternary Phase Equilibria Studies. .......... ...... .... 5
2.1.1 The Survey Studies ............... .. ..... 6
2.1.2 The Ti-Based Alloy Development Studies ............. 18
2.2 The Sigma Phase .................. ....... 22
2.3 The Gamma Phase ................... ....... 24
2.4 The B2 Phase. .................... ....... 26
2.5 The Omega Phase .................. ....... 32
2.6 The Ortho/Hex Phases .. ...... .......... .... 39
2.6.1 Disordered Structures ........... ............ 41
2.6.2 Ordered Structures ........... ............. 43

3 EXPERIMENTAL PROCEDURE ........ .. 53
3.1 Material ................ ...... 53
3.2 Heat Treatments ....................... ....... 56
3.2.1 Long-Term Heat Treatment Experiments. ... 58
3.2.2 Short-Term Heat Treatment Experiments .... 59
3.3 Characterization Techniques. ............... .. 59
3.4 Sample Preparation .................. .. 62
3.4.1 Optical Microscopy. ................ .. .. 62
3.4.2 Transmission Electron Microscopy .... 63

4 EQUILIBRIUM PHASE TRANSFORMATION STUDY .... 65
4.1 Introduction .................. ............ 65
4.2 Results ................. ....... ....... 66
4.2.1 As-Cast Microstructures ....... .. 66
4.2.2 Long-Term Heat Treatments ..... 83









4.2.3 Short-Term Heat Treatments 108
4.3 Discussion ............... .... 120
4.3.1 The p Phase ............. ....... ........ 122
4.3.2 Phase Equilibrium .......... .. ... 127
4.3.3 Phase Transformation Mechanisms ... 131

5 THE OMEGA-RELATED (m-D) PHASE ..... 137
5.1 Introduction ............. ... 137
5.2 Results .... ......... .... ........... 138
5.2.1 Structural Analysis ....... .....138
5.2.2 Effect of Cooling Rate ..... ...... 154
5.2.3 Effect of Low Temperature Heat Treatments .... 158
5.3 Discussion ............. .... ........ .. .... 165
5.3.1 Microstructural Aspects of the a-D Phase. ... 166
5.3.2 Comparison of the c-Related Phases in the Nb-Ti-Al System. 169
5.3.3 Transformation and Site Occupancies of the a Phases 172
5.3.4 Crystallographic Aspects of the a-D Phase Transformation. 187

6 THE ORTH(HEX) PLATES .................. ....... ..194
6.1 Introduction ............... ... 194
6.2 Results ......... ... ..... .......... 195
6.2.1 Structural Analysis ....... .....196
6.2.2 Plate Morphology ......... ... 230
6.2.3 Zig-Zag Plate Morphology ..... .... 233
6.2.4 Defect Structural Analysis of the Plates. .... 236
6.3 Discussion of the Plate Transformation 251
6.3.1 Martensitic Transformation .. ... 251
6.3.2 Structural Analysis of the Plates ... 263
6.3.3 Crystallographic Treatment of the Plate Transformation. 292
6.3.4 Formation of Plate Martensite from the p phase ... 305

7 SUMMARY ................. ....... ....... 309

8 FUTURE WORK. ................... .......... 313

REFERENCES ............ ......... .............. 314

BIOGRAPHICAL SKETCH .................. ........ 320















LIST OF TABLES


Table Page

2.1 Information regarding the phases of the three binary Nb-Ti, Nb-Al, and
Ti-Al phase diagrams.. ................. .. ..... 7

3.1 The nominal compositions of the as-received alloys and the
compositions determined by microprobe analysis of the re-arc melted
alloys .. .. ... .. ... 54

3.2 Long-term heat treatments used for alloys 2 and 4. ... 60

3.3 Short-term heat treatments used for alloys 2 and 4.. ... 61

4.1 Microprobe results of the aged samples of alloy 2. .. 87

4.2 Summary of the phases identified by TEM in the aged samples of
alloy 2. .. ............... .............. 93

4.3 Summary of the phases identified by TEM in the aged samples of
alloy 4........... ....................... .. 109

5.1 Shows the relation between the possible diffraction groups and the
symmetries observed in the Convergent Beam Electron Diffraction
(CBED) patterns at the [0001], zone axis. .. ..... 141

5.2 Shows the relation between the possible diffraction groups and the
symmetries observed in the Convergent Beam Electron Diffraction
(CBED) patterns at the [1T00], zone axis. .. ..... 143

5.3 Shows the crystal point groups that are consistent with the diffraction
groups observed in the CBED whole patterns. .... 144

5.4 Comparison of characteristics between the (-D, the disordered co-Ti,
and the ordered o-B8, phases from the Nb-Ti-Al ternary system. ....... 170

5.5 Proposed site occupancy for the w-D phase with Ti3Al4Nb2
stoichiometry and P63/mcm (193) space group.. ... 182

6.1 Shows the relation between the possible diffraction groups and the
symmetries observed in the Convergent Beam Electron Diffraction
(CBED) patterns of the HCP phase at the [0001]H zone axis. 201









6.2 Shows the relation between the possible diffraction groups and the
symmetries observed in the Convergent Beam Electron Diffraction
(CBED) patterns of the HCP phase at the [1126], zone axis. 204

6.3 Shows the relation between diffraction groups and crystal point groups
for the CBED patterns of the HCP plates. .. 205

6.4 Shows the relation between the possible diffraction groups and the
symmetries observed in the Convergent Beam Electron Diffraction
(CBED) patterns of the ORTH1 phase at the [001]0 and [110]o zone
axis.. .............. ........ ....... .. 218

6.5 Shows the relation between the diffraction groups and crystal point
groups for the CBED patterns of the ORTH1 plates.. ... 220

6.6 Shows the lattice parameters of plates with the ORTH1 structure,
where 4u1 is the angle between the (01T1) and (020)02 reflections. .223

6.7 The imaging conditions of the APDBs observed in the plates. ... 253

6.8 Calculated phase factors, a = 2n(g R), for the APDB vectors in the
HCP phase .......... ............ .......... 268

6.9 Calculated phase factors, a = 2n(g R), for the APDB vectors in the
ORTH1 and ORTH2 phases. ...... ......269

6.10 The proposed atomic site occupancy of the ORTH 1 phase with the
Al2NbTi stoichiometry and the Cmcm (63) space group. 274

6.11 The proposed atomic site occupancy of the ORTH3 phase with the
Al(NbTi) stoichiometry and the Pmma (51) space group. 285















LIST OF FIGURES


Figure Page

2.1 Shows the binary Ti-Al phase diagram. (a) the previously accepted
phase diagram; (b) the modified section of the phase diagram. 9

2.2 Shows the 1200C isothermal section of the ternary Nb-Ti-Al system
determined by Jewett et al.. ................... .... 12

2.3 Shows the liquidus projection of the ternary Nb-Ti-Al system
determined by Perepezko et al.. ................ .... 14

2.4 Shows the 1200C isothermal section of the ternary Nb-Ti-Al system
determined by Das et al. ................... .. ..16

2.5 Shows the partial sections of the isotherms in the ternary Nb-Ti-Al
system determined by Das et al. (a) 12000C; (b) 1150C. ... 17

2.6 Shows the temperature-composition diagram of the Ti3Al to TiAl1Nb
section by Banerjee et al. ...... ........ 21

2.7 Shows the five lattice sites in the projection of the unit cell for the a
phase ........... ......... ......... 25

2.8 Shows the unit cell of the y-TiAl phase. ................. ..27

2.9 Shows the unit cell of the B2 phase with the atomic site occupancy
determined in a Ti-24.5Al-14Nb (at.%) alloy by Banerjee et al. 30

2.10 Shows the {11), plane collapse model of the 0 to ( transformation.
The view is normal to the (110)p planes. ..... 35

2.11 Shows two rotational variants, each one containing three translational
variants, of the a phase formed from the P to a transformation. The
view is normal to the (110)p planes. .................. .. 36

2.12 Shows the transformation paths from the P phase to the w-related
phases using subgroup and symmetry relations from Bendersky et al. 40

2.13 Shows the relationship between the crystal structures of the ac-Ti3Al
phase and the O-Ti2AINb phase. (a) the a2-Ti3Al phase (PG6/mmc space
group); (b) the O-Ti2AINb phase (Cmcm space group) 46









2.14 Shows the transformation paths from the p phase to the O phase using
subgroup and symmetry relations from Bendersky et al. ... 50

4.1 Optical micrographs showing the as-cast microstructures. (a) alloy 2;
(b) alloy 3; (c) alloy 4 .................. 67

4.2 SAED patterns showing the [001] zone axis of the P matrix. (a) alloy 2;
(b) alloy 3; (c) alloy 4. ................... ..70

4.3 TEM micrographs showing the APDBs in the B2 matrix. (a) alloy 2;
(b) alloy 4 ...... .............. .. ........ 72

4.4 SAED pattern from the B2 matrix showing the splitting of diffraction
spots. The specimen was tilted away from the [001] zone axis along
g=(110). .......... ....... .................. 74

4.5 SAED patterns showing the diffuse electron scattering observed in the
B2 matrix of alloy 2 (a and b) and alloy 4 (c and d). (a) [110] zone axis;
(b) [111] zone axis; (c) [110] zone axis; (d) [111] zone axis.. ... 75

4.6 TEM micrographs showing the tweed microstructure in the B2 matrix.... 77

4.7 Shows the microstructure of the as-cast sample of alloy 2. (a) TEM
micrograph showing the grain boundary allotriomorphs and
Widmanstatten laths of the y-TiAl phase; (b) SAED pattern showing
the orientation relationship observed between the y laths and B2
matrix, which was <110], II p and {111), II {(0}p .. 80

4.8 TEM micrographs of the B2 matrix in the as-cast sample of alloy 2.
(a) the small o-related precipitates and lenticular-shaped plates; (b)
the coarse o-related precipitates adjacent to the plate. ... 81

4.9 Shows the acicular microstructure observed in the EM-levitated and
drop quenched sample of alloy 2. (a) Optical micrograph; (b) TEM
micrograph ................... ............ 82

4.10 Optical micrographs showing the microstructures of the long-term
thermally aged samples of alloy 2. (a) 1500'C-4hrs-WQ; (b) 1400"C-
4hrs-WQ; (c) 13000C-4hrs-WQ; (d) 1200C-4hrs-FC. ... 84

4.11 TEM micrographs showing the microstructures of the long-term
thermally aged samples of alloy 2. (a) 14000C-4hrs-WQ;
(b) 1300C-4hrs-WQ. .................. .......... ..88

4.12 TEM micrographs showing the y laths in the o + y microstructure of the
thermally aged 1200C-4hrs-FC sample of alloy 2. (a) longitudinal
view; (b) transverse view.. ...... ..... 91









4.13 SAED patterns of the y phase observed in the 1200C-4hrs-FC sample
of alloy 2. (a) the [110], zone axis; (b) the [101], zone axis. ...... 92

4.14 Optical micrographs showing the microstructures of the long-term
thermally aged samples of alloy 4. (a) 1550C-2hrs-AC;
(b) 1515C-2hrs-AC; (c) 14000C-4hrs-WQ; (d) 13000C-4hrs-WQ;
(e) 1200'C-4hrs-WQ; (f) 1000C-16hrs-AC. .. 95

4.15 TEM micrograph showing the APDBs in the B2 matrix of the sample
from alloy 4 that was heat treated at 15500C and water quenched. ..... 98

4.16 TEM micrographs showing the different morphologies of the retained p
phase observed in the samples of alloy 4 that were heat treated at
1200C, 1300C, and 1400C and water quenched. (a) the irregular and
circular morphologies; (b) the elongated morphology. 100

4.17 TEM micrograph showing the precipitates of the orthorhombic phase
that formed on the APDBs and in the B2 matrix of alloy 4 during air
cooling from 1550C. ................... ... 101

4.18 SAED pattern showing the orientation relationship between the
orthorhombic phase and the B2 phase in the 1550C-2hrs-AC sample of
alloy 4. The OR was consistent with [001]j 11 <110>p and
(110)o 2) ......... ... ................ ............ ... 103

4.19 TEM micrograph showing the orthorhombic plates that formed in the
retained B2 phase of alloy 4 during furnace cooling from the heat
treatment temperatures of 1200C, 1300C, and 1400C. ... 105

4.20 Shows the microstructure observed in the 1000C-16hrs-AC sample of
alloy 4. (a) TEM micrograph showing particles of the orthorhombic
phase observed at the grain boundaries of the a phase; (b) SAED
pattern showing the [110]o zone axis of the orthorhombic phase. ..... 106

4.21 Shows the microstructure observed in the samples of alloy 4 that were
heat treated at 1000C and 1200C. (a) TEM micrograph showing the
colony of a grains; (b) SAED pattern showing multiple [001], zone axes
from the a grains present in the colony. ..... 107

4.22 TEM micrographs showing the microstructure observed in the
12000C-2min-WQ sample of alloy 2. (a) the a and the y phases at a low
magnification; (b) the a grain size. ..... .111

4.23 Optical micrographs showing the microstructure observed in the
12000C-2min-WQ sample of alloy 4. (a) bright field micrograph;
(b) dark field micrograph.. .... 113









4.24 TEM micrographs showing the microstructure observed in the 1200C-
2min-WQ sample of alloy 4. (a) the thin foil specimen was tilted to the
[001], zone axis of the p particles; (b) the thin foil specimen was tilted
to the [001], zone axis of the a matrix. .... 114

4.25 Optical micrograph showing the microstructure observed in the 1000C-
2min-WQ sample of alloy 4.. ................... .. 115

4.26 Shows the microstructure observed in the 1000'C-2min-WQ sample of
alloy 4. (a) TEM micrograph showing the reaction front of a colony that
partially transformed from the B2 matrix; (b) SAED pattern showing
the [100], zone axis of the disordered p phase located between the a
grains in the colony.............. .... 116

4.27 Shows the microstructure observed in the 10000C-2min-WQ sample of
alloy 4. (a) TEM micrograph showing the transverse view of the colony
structure near the interface between the colony and the B2 phase; (b)
SAED pattern showing the orientation relationship between the B2
and a phases, which was <100>2 I [001), and (110)2 I 1110),. ... 118

4.28 Shows the microstructure observed in the 1000C-2min-WQ sample of
alloy 4. (a) TEM micrograph showing the transverse view of the colony
near the center of the colony; (b) CBED pattern showing the
orientation relationship between the a and P phases, which was
<100> II [100]. and (110), 1 {110) ..... 119

4.29 SAED patterns showing two additional orientation relationships that
were observed between the a and P phases in the heat treated samples
of alloy 4. (a) Bz 1<110], and {110}2 1 {110),; (b) B2 I [001],
and {110),2 {110),.................... ............121

4.30 Shows the translational vector for the nearest neighbor (NN) and next
nearest neighbor (NNN) sites in the unit cell of the B2 phase. The
atomic site occupancy shows Nb and Ti atoms randomly occupying the
la Wyckoff site and Al atoms occupying the lb Wyckoff site. ... 126

4.31 Shows the equilibrium phases that formed at the aging temperatures
in alloys 2 and 4. (a) alloy 2; (b) alloy 4 129

5.1 CBED whole patterns showing the 6mm symmetry observed in the
[0001] zone axis of the a-related phase in alloy 2. (a) long camera
length showing the zero order laue pattern; (b) short camera length
showing the faint FOLZ rings. ............... 139

5.2 CBED whole patterns showing the 2mm symmetry observed in the
[1100] zone axis of the o-related phase in alloy 2. (a) long camera
length showing the zero order laue pattern; (b) short camera length
showing the FOLZ rings. .................. ....... ..142









5.3 CBED whole pattern of the [1100] zone axis with the beam tilted
slightly to show the black cross in the kinematically forbidden (0001)
reflection at the Bragg condition. ................. ..146

5.4 SAED pattern showing the orientation relationship that was observed
for the c-related and B2 phases. The OR was determined to be
[0001], I [111], and (TOO1100), (10). ..... 147

5.5 Shows the stereographic projection of the OR relationship that was
observed for the o-related and p phases in alloy 2. The projection
shows the [1lll] and [0001], poles.. 150

5.6 SAED pattern showing the OR between the a-related and p phases at
the [110], zone axis. Two rotational variants with the [1700], and
[1010], zone axes are superimposed on this diffraction pattern. 151

5.7 Shows the microstructure observed in the as-cast sample of alloy 2.
(a) TEM micrograph showing three plates and a coarse o-related grain
observed in the B2 matrix; (b) SAED pattern showing the orientation
relationship observed for the three plates, the a-related grain, and the
B2 phase at the [11] zone axis. .... 152

5.8 Shows the calculated diffraction patterns of the three plates and the
coarse a-related grain at the [111] zone axis of the B2 matrix. (a) the
composite pattern; (b) the [111] zone axis of the B2 matrix; (c) the
[0001], zone axis of the a-related grain; (d) the [110] zone axis of the
orthorhombic plate 1; (e) the [110] zone axis of the orthorhombic plate
2; (f) the [110] zone axis of the orthorhombic plate 3. ... 153

5.9 Optical micrograph showing the microstructure observed in the 1400C-
4hrs-FC sample of alloy 2. ................ ....... 156

5.10 TEM micrographs of the 14000C-4hrs-FC sample of alloy 2. (a) the
in-matrix region consisting of the o-related phase and plates; (b) the
prior grain boundary region consisting of the a and y phases. ... 157

5.11 TEM micrographs of the 400C- 12hrs-WQ sample of alloy 2. (a) shows
the precipitates of the a-related phase observed in the B2 matrix; (b)
shows the APDBs observed in the B2 matrix. ..... 159

5.12 TEM micrographs of the 600C- 12hrs-WQ sample of alloy 2. (a) shows
the fine ca-related domains that formed from the B2 phase; (b) shows
the coarse oa-related domains that formed at the prior B2 grain
boundaries. ................. 160

5.13 SAED patterns from the 600C-12hrs-WQ sample of alloy 2 that show
the diffraction patterns of the former B2 matrix. (a) [111]B zone axis;
(b) [110]82 zone axis; (c) [1001], zone axis; (d) [112]B2 zone axis. ...... .162









5.14 TEM micrograph of the 600C-12hrs-WQ sample of alloy 2 that shows
the APDBs observed in a coarse rotational domain of the o-related
phase. The g = (1120) reflection was used to show the APDBs in the
coarse domain. ............... ..... 164

5.15 Shows the atomic site occupancies of the disordered P phase and the
disordered o phase for Ti.. .................. ...... 174

5.16 Shows the atomic site occupancies of the ordered 0 (B2) phase and the
w-B82 phase from the results of Bendersky et al. [37] .. .... 176

5.17 The (111) projection of the B2 phase with the atomic site occupancy
showing Nb and Ti atoms occupying the la Wyckoff site and Al atoms
occupying the lb Wyckoff site. (a) shows the planes at z = 0.0A,
0.095A, and 0.189A; (b) shows the planes at z = 0.284A, 0.379A, and
0.473A. The dashed lines denote the unit cell of the o-D phase. ..... .179

5.18 The (0001) projection of the o-D phase which is based on the P6,/mcm
space group and A14Ti3Nb, stoichiometry. (a) shows the single layer at
z = 0.0A and double layers at z = 1/4c; (b) shows the single layer at
z = 1/2c and double layers at z = 3/4c. 184

5.19 Shows the atomic site occupancies of the B2 phase and the c-D phase.. 186

5.20 Shows the transformation paths from the p phase to the o-D phase
described by subgroup/symmetry relations. ..... 189

6.1 Shows a thick plate with the HCP structure observed in the 1300C
aged sample. (a) TEM micrograph of the HCP plate; (b) SAED pattern
of the orientation relationship observed between the HCP plate and the
p phase, which was [0001], 1 [011], and (1100) II (2 )p 197

6.2 The stereographic projection of the orientation relationship between
the HCP phase and the p phase which shows the [0001], I [011], poles.. .. 199

6.3 CBED patterns showing the whole pattern symmetry of the [0001],
zone axis observed for the thick HCP plates. (a) a large camera
constant; (b) a small camera constant. 200

6.4 CBED patterns showing the whole pattern symmetry of the [1126],
zone axis observed for the thick HCP plates. (a) a large camera
constant; (b) a small camera constant ..... 203

6.5 CBED pattern showing the diffuse black cross in the (0001) disc
observed in the [1120]n zone axis of the HCP plates. 206









6.6 Shows the boundary observed in a thick HCP plate. (a) TEM
micrograph; (b) SAED pattern of the two regions separated by the
boundary which was consistent with the orientation relationship
observed between the HCP plates and the P phase. 208

6.7 CBED patterns of the two sides (A and B in Figure 6.6). (a) shows the
6mm symmetry that was consistent with the [0001], zone axis of the
HCP phase; (b) shows the 3m symmetry of possibly a different phase. 209

6.8 TEM micrographs of a thick HCP plate. (a) shows the plate inclined,
relative to the beam, at the [TIT], zone axes; (b) shows the plate
edge-on after tilting -110 along the g = (0ll)p I (0001)u reflections. .211

6.9 Shows a medium thick plate with the ORTH1 structure. (a) TEM
micrograph; (b) SAED pattern of the orientation relationship between
the ORTH1 plate and the p phase, which was [001]o0 I [011]p and
(110)0o l(211)p. .............................. 213

6.10 The stereographic projection of the orientation relationship between
the ORTH1 phase and the p phase which shows the [001]o 1H [01 l]
poles ... 215

6.11 CBED pattern showing the whole pattern symmetry of the [001]o, zone
axis observed for the ORTHI plates. 216

6.12 The structural analysis of the medium thick ORTH1 plates. (a) CBED
pattern showing the 2mm whole pattern symmetry of the [1101o zone
axis; (b) SAED pattern obtained by tilting the thin foil specimen from
the [110] zone axis along the g = (001) reflection. 219

6.13 Shows a thin plate with the ORTH2 structure. (a) TEM micrograph;
(b) SAED pattern showing the orientation relationship between the
ORTH2 plate and the 0 phase at the [01l], zone axis. 225

6.14 Enlarged SAED pattern showing the diffuse streaks and diffraction
spots observed for the ORTH2 plates. The diffraction pattern shows
the [l01]p zone axis for the p phase. .... 226

6.15 Shows a thin plate with the ORTH3 structure. (a) TEM micrograph;
(b) SAED pattern showing the orientation relationship between the
ORTH3 plate and the P phase at the [011], zone axis. 228

6.16 Enlarged SAED pattern showing the missing diffraction spots for the
ORTH3 plates. The diffraction pattern shows the [011]p zone axis for
the p phase. ................... ..... 229

6.17 TEM micrographs showing the plate morphology that was determined
from the images of plates at the three orthogonal directions. (a) the
[001]o zone axis; (b) the [110]o zone axis; (c) the [1TO]o zone axis.. 231









6.18 TEM micrograph showing plates that were partially enclosed within
the thin foil specimen of the as-cast RAM sample of alloy 2.. ... 234

6.19 TEM micrographs showing the zig-zag morphology of the plates. (a)
bright field micrograph; (b) dark field micrograph.. 235

6.20 Shows two plates connected together along a twin boundary. (a) TEM
micrograph; (b) SAED pattern at the [TIT-p zone axis of the P phase
showing the twin relationship between plate 1 and plate 2 .... 237

6.21 TEM micrographs showing the stacking faults observed in the plates at
the [110]o or [1120]1 II [ T1TJ]B zone axes. (a) shows that the stacking
faults were invisible using g = (002)o or (0002),; (b) shows that the
stacking faults were visible using g=(2T1)o or (201).. .239

6.22 TEM dark field micrographs showing the APDBs observed in the HCP
plates. (a) g = (T210) and (100)p; (b) g = (2110); (c)g=(1120)H.. ..... 241

6.23 TEM micrographs showing the columnar shaped APDBs observed in
the HCP plates at the [1ll0] I [1Tl], zone axes. (a) bright field
micrograph formed with g = (2Z00)H; (b) dark field micrograph formed
with g = (T01) ........ ...... ... ........... 244

6.24 TEM micrographs showing the equiaxed morphology of the APDBs
observed in the HCP plates at the [2T110] zone axis. (a) dark field
micrograph formed with g = (01TO),; (b) dark field micrograph formed
with= (01lTI) ........ .... .. .. .. ...... 245

6.25 TEM dark field micrographs showing the APDBs observed in the
ORTH1 plates. (a) g = (200)0 and (100)p; (b) g = (130)o; (c) g = (190)o. 247

6.26 TEM dark field micrograph showing the columnar shaped APDBs
observed in the ORTH1 plates at the [110]o I [011], zone axes using
g = (11 1)o.. 249

6.27 TEM dark field micrograph showing the equiaxed morphology of the
APDBs observed in the ORTH1 plates at the [110]0 zone axis using
g = (111)o. ... ......... .... ......250

6.28 TEM dark field micrograph showing the columnar shaped APDBs
observed in the thin ORTH2 plates using the diffuse streak. The
specimen was tilted from the [01l1] zone axis of the p phase. ... 252

6.29 Shows the uniaxial stress state of the plate and the three principal
axes of strain which are %,, X2, and .. .260









6.30 Shows the unit cells of the ORTH1 and HCP phases superimposed
upon each other to illustrate how the distortions of the orthorhombic
phase are related to the hexagonal phase. The drawing shows the
[001]o direction of the ORTH1 unit cell and the [0001], direction of the
HCP unit cell ................. ............... 264

6.31. Shows the proposed atomic site occupancy of the ORTH1 phase based
on the Cmcm space group with Al atoms occupying the 8g, Ti atoms
occupying the 4cl, and Nb atoms occupying the 4c2 Wyckoff positions. ... 275

6.32. Shows the calculated CBED patterns of the orthorhombic structures
based on two possible atomic site occupancies. (a) the proposed
ORTH1 phase with the stoichiometry of Al1TiNb; (b) the O-TiAlNb
phase ............. ..... .................. 277

6.33. CBED pattern showing the whole pattern symmetry of the O-Ti2AINb
phase observed in alloy 1.. ........ .....279

6.34. Shows the APDB vectors in the (01 )p planes of the B2 phase and in the
(001)o planes of the ORTH1 phase.. ... 280

6.35. Shows the proposed site occupancy of the ORTH3 phase based on the
Pmma space group with Nb and Ti atoms occupying the 2e Wyckoff
sites and Al atoms occupying the 2f Wyckoff sites.. ... 284

6.36. Shows the calculated SAED pattern of the orientation relationship
between the ORTH3 plate and the B2 phase. The OR showed the
[100]04 and [011], zone axes to be parallel and the (011)04 and (211)p
planes to be parallel ........ 286

6.37. Shows the two transformation paths that led to the formation of the
ORTH1 plates (path 1) and the HCP plates (path 2) from the P phase. 294

6.38. Shows the unit cells of the different structures that led to the
formation of the ORTH1 phase from the B2 phase. (a) the B2 (Pm~m)
structure; (b) the orthorhombic (Cmmm) structure; (c) the ORTH2
(Pmma) structure; (d) the ordered ORTH1 (Cmcm) structure.. ... 297

6.39. Shows the unit cells of the different structures that led to the
formation of the HCP phase from the P phase. (a) the disordered P
(Im3m) structure; (b) the disordered orthorhombic (Fmmm) structure;
(c) the disordered orthorhombic (Cmcm) and the disordered HCP
(P63/mmc) structures; (d) the ordered HCP (P63/mmc) structure. 303


XV















Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy

PHASE TRANSFORMATIONS IN THE CENTRAL PORTION
OF THE Nb-Ti-Al TERNARY SYSTEM

By

David Timothy Hoelzer

December 1996


Chairperson: Dr. Fereshteh Ebrahimi
Major Department: Materials Science and Engineering

Intermetallics in the Nb-Ti-Al ternary system have been considered for high

temperature aerospace applications, and their development requires a thorough

understanding of phase equilibria and phase transformations. In this study,

transmission electron microscopy (TEM) was primarily used to investigate the phase

equilibria and phase transformations of two alloys with compositions of 27Nb-33Ti-

40Al (alloy 2) and 42Nb-28Ti-30A1 (alloy 4). Small arc melted samples were

thermally aged at temperatures between 400"C and 15500C for four to sixteen hours

(long-term), and at 1000C and 1200"C for two to five minutes (short-term), followed

by either water quenching, air cooling, or furnace cooling.

The equilibrium phase study showed that both alloys solidified as the P phase,

which becomes ordered to the B2 phase during solid state cooling. The a-Nb2Al phase

precipitated from the 0 phase slightly below 1400C in alloy 2 and 1550C in alloy 4.

The a phase formed as isolated grains above 13000C in both alloys. Colonies of a

grains formed below 1300C in alloy 4. A eutectoid transformation from 0 to a + y-









TiAl occurred at 12000C in alloy 2. A discontinuous transformation from B2 to a + p

occurred at 1000C in alloy 4.

A metastable ea-D phase formed by the collapse of {111}p planes and chemical

ordering from the B2 phase in alloy 2 during slow cooling. The A-D phase consisted of

the Al4Ti3Nb2 stoichiometry and P6,/mcm space group. A proposed model showed

aluminum and niobium on single layers, and titanium and aluminum on double

layers. The transformation path was described using subgroup and symmetry

relations as: PmJm(B2) -* PYlm(o'") P63/mcm(a-D).

A martensitic transformation of the P phase to plates occurred during fast

cooling in alloy 2. The observed habit plane of the plates agreed with that calculated

using the invariant line theory. The p composition affected the formation of

structurally related orthorhombic (Pmma and Cmcm) and HCP (P63/mmc) plates.

The proposed site occupancy showed the Al2TiNb stoichiometry for the Cmcm phase.

Analysis of domain structures, stacking faults, and electron diffraction suggested two

possible transformation paths: Im3m(p) Cmcm(disordered) -+ P63/mmc(disordered)

- P63/mmc(DO1,) for HCP plates and Pmm(B2) -* Pmma -" Cmcm(ordered) for

orthorhombic plates.















CHAPTER 1
INTRODUCTION


The ternary Nb-Ti-Al system is recognized as a technologically important

system. Alloys based on this system have found many applications, especially in the

aerospace industry. The titanium-based alloys have long been studied and used in

aeronautical applications because of their low densities and high strengths [1,2].

There is currently the need for higher temperature and higher strength materials to

improve the performance of gas turbine engines and structural airframe components

of modern aircraft. This will require different materials than the current nickel-

based superalloys and conventional titanium-based alloys. The nickel-based alloys

are heavy and have reached practical limits imposed by operating temperatures that

are -85% of their melting point [3,4]. The titanium-based alloys do not possess

sufficient mechanical properties, such as creep and oxidation resistance, at high

temperatures [5]. Thus, the current need to improve the properties of these

materials has led to the development of intermetallic titanium-aluminides [5-8].

Additions of niobium to these titanium-aluminides make these alloys even more

attractive by improving the mechanical properties. However, the materials needed

for the next generation of high performance gas turbine engines, etc. still need

further improvements in lower density, better mechanical properties, and higher

operating temperatures. To meet these needs, research is being conducted on

refractory-based alloys, ceramics, composites, and intermetallics.









2

Ordered intermetallic compounds based on refractory metals such as niobium

have been identified as potential materials that may meet the high-temperature

requirements of advanced turbine engines [9]. The attractive properties of refractory-

based intermetallics include combinations of high melting temperature, lower

density, high stiffness, and good creep/strength resistance. However, monolithic

ordered intermetallics typically show poor mechanical properties at low

temperatures, of which the low fracture toughness is the most serious problem. The

current trend to overcome these problems has been through the use of composites and

alloy development based on two-phase and multi-phase systems incorporating

Nb-based intermetallics and ductile second phases. The Nb-Ti-Al system shows

potential phase relationships between the intermetallic a-NbAI phase, which has a

high melting point of 2060C; the BCC p or B2 phases, which have extensive ternary

composition ranges; and other technologically important intermetallics, such as the

a2-Ti3Al, y-TiAl, O-Ti2NbAl, and l-CTi-Nb)A3l phases [10-13]. Thus, the development

of these alloys for the improvement of properties requires a thorough understanding

of the phase relationships in this system.

There have been a number of studies over the past thirty years that have

contributed to the current understanding of the phase equilibria and phase

transformations in the Nb-Ti-Al system. However, a thorough understanding of these

topics is far from being complete. The reason for this lack of understanding can be

attributed to the complex phase relationships arising from a multitude of equilibrium

and metastable phases in this ternary system. In addition, the solidus temperature

over a large portion of this ternary system lies above = 1500C. The high solidus

temperature limits the practical number of isotherms that can be developed in










3

systematic studies using large numbers of alloys with different compositions. Thus,

the systematic studies in the past have concentrated on determining the phase

equilibria at just one or two temperatures, with the most common temperature at

1200C. Surprisingly, the phase equilibria and phase transformations in the central

portion of this system have not been investigated very thoroughly in the past. This

central portion contains ternary solubility extensions of the binary a-NbMAl and y-TiAl

phases and the ternary P/B2 phases, which have attractive properties such as high

melting point, low density, and reasonable oxidation resistance. Therefore, it was the

purpose of this study to provide basic research on the phase relationships in the

central portion of the ternary Nb-Ti-Al system.

A review of previous literature on the phase equilibria in the ternary Nb-Ti-Al

system is given in Chapter 2. From this literature survey and preliminary

experimental results, two alloys were selected based on two-phase microstructures

that contained the a phase and either the B2 or y phases. A third alloy was also

investigated in order to study the influence of aluminum on the ordering in the P

phase to the B2 phase that has been shown to occur in this ternary system. Both long

term and short term heat treatments at high temperatures were employed to study

the high temperature phase equilibria and their evolution. The stability aspect of the

p phase with regard to metastable phase formations was investigated using different

cooling rates in the high temperature heat treatments. Transmission electron

microscopy (TEM), utilizing the imaging, selected area electron diffraction (SAED),

and convergent beam electron diffraction (CBED) capabilities, was selected as the

primary analytical technique to identify and study the phases. The details










4

concerning the alloy compositions, heat treatments, analytical techniques, and

specimen preparation are presented in Chapter 3.

Since the central portion of this ternary system is complicated, an analysis of

the results of the equilibrium phases are presented in Chapter 4. The equilibrium

phases present at high temperatures and the formation of these phases are described

from the analysis of the long term and short term heat treatments. The influence of

cooling rate on the formation of metastable phases from the high temperature P

phase is also introduced in this chapter. Following the analysis of the equilibrium

phases, detailed studies of the metastable phase formations are presented in

Chapters 5 and 6.

Chapter 5 covers the metastable co-related phase that forms from the P phase

in alloy 2. The analysis of the structure and the effects of composition, cooling rate,

and low temperature heat treatments are covered in this chapter. From these

results, the proposed atomic site occupancy of this o-related phase and a description

of the P to a phase transformation using subgroup and symmetry relations is given.

Chapter 6 analyzes the metastable plates that formed from the p phase in

alloy 2. The analysis of the crystal structures and defect structures of these plates is

covered. The influence of the heat treatment and P composition on the structure and

formation of the plates is analyzed. The transformation of the p phase to plates is

shown to be consistent with martensitic transformations using the invariant line

theory. Finally, the structure of the plates is described using subgroup and symmetry

relations to show that two different transition paths lead to different plate

structures.















CHAPTER 2
LITERATURE SURVEY


In this chapter, an overview of the published literature on the phase equilibria

and phase transformations in the ternary Nb-Ti-Al system is presented. Due to the

complexity of this system, this overview will focus only on those phases that were

observed in the alloys investigated in this study. Therefore, this chapter is divided

into six sections: the ternary phase equilibria studies, the sigma (a) phase, the

gamma (y) phase, the B2 phase, the omega (o) phase, and the ortho(hex) phases. The

orth(hex) designation is used in conjunction with the two closely related orthorhombic

and hexagonal close packed (HCP) structures. The first section on the ternary phase

equilibria studies describes the various developments in the overall ternary phase

diagram. This section is divided into two subsections based on the results from the

survey studies and the Ti-based alloy development studies. The remaining sections

on the a, y, B2, omega (o), and orth(hex) phases describes the research pertinent to

each of these specific phases. The sections dealing with the c, y, and B2 phases are

relevant to the equilibrium phase study in Chapter 4; the a phase is relevant to the

o-related phase transformation study in Chapter 5; and the orth(hex) phases are

relevant to the plate transformation study in Chapter 6.


2.1 The Ternary Phase Equilibria Studies

The phase equilibria of the ternary Nb-Ti-Al system has been the subject of

many investigations since the early 1950s. However, most of these studies can be











grouped into two main categories: those that surveyed the phase equilibria of alloys

with compositions covering large regions of the ternary phase diagram at discrete

temperature ranges, and those that concentrated in-depth on just a few compositions

for the commercial development of Ti-based alloys over large temperature ranges.

The survey studies investigated the phase equilibria of alloys at a limited number of

temperatures, which was usually only one temperature and often 12000C. The

Ti-based alloy studies mostly focused on Ti-rich or TiAl + Nb compositions with

constant 25at.%Al. The recent trend in the development of these Ti-based alloys has

been in the ternary region between TiAAl and TiAl. These are binary phases with

ternary additions such as Nb.

2.1.1 The Survey Studies

Most of the phases that have been observed in the ternary Nb-Ti-Al system

are binary phases that showed large ternary solubility ranges. There have been

many studies in the past that have determined the binary phases of the Nb-Ti, Nb-Al,

and Ti-Al systems. These studies and the binary phase diagrams of the Nb-Ti, Nb-Al,

and Ti-Al systems developed from them have been compiled in two main references

[10,11]. The results of these compiled studies have shown that there are at least ten

equilibrium binary phases. A summary of the important information about these ten

phases, such as the phase notation, the crystal structure, the point group, the

reaction type, and the reaction temperature are shown in Table 2.1.

There currently exists some uncertainty concerning the phase equilibria in the

Ti-Al phase diagram. This uncertainty is mainly between the y-TiAl and n-TiAl,

phases and is due to the formation of long-range periodic structures in this part of the

binary system that can complicate the determination of the equilibrium phases












Table 2.1. Information regarding the phases of the three binary Nb-Ti, Nb-Al, and Ti-Al phase diagrams [10,12].


Phase Crystal Space Reaction Temperature
Notation Stoichiometry Structure Group Type (IC)

p Nb BCC Im3m L -p 2467
p Ti BCC Im3m L -p 1670
a Ti HCP P6,/mmc a .p 882
6 Al FCC Fm3m L 1p 660

Pi NbAl Cubic Pm3n L +1 -p, 2060
a Nb2Al Tetragonal P42/mnm L + 13- 1940
T NbAl1 BCT I4/mmm L -T 1680
a, Ti3Al HCP P6,/mmc a at 1180
y TiAl Tetragonal P4/mmm L+a -y 1450
6 LPS --- L+y1 1380
B TiAl2 BCT 14,/amd y+8 1240

l TiAl, BCT I4/mmm L+5 "-, 1350
LPS is a long-period superlattice structure;
p-Ti and p-Nb are isomorphous across the binary Nb-Ti phase diagram; and
ri-NbAl, and il-TiAl, are isomorphous across the ternary phase diagram for constant Al.









8

[10,14]. The binary Ti-Al phase diagram originally developed by Murray [10] has

recently been modified on the Ti-rich side from the results by Valencia et al. [15] in

1987 and by McCullough et al. [16] in 1989. This modification involved the phase

boundaries between the a-Ti and y phases. It had previously been shown that the

a-Ti phase formed from the peritectoid p + y a reaction near 1480C, but the

results of these recent studies showed that the a-Ti phase formed from the peritectic

L + p a reaction. The previously accepted binary Ti-Al phase diagram and the

revised part of the phase diagram are shown in Figure 2.1. The temperature of the

peritectic L + p a reaction was estimated to be 1475C from the study by

McCullough et al. [16]. Further studies by McCullough et al. showed that the single

a-Ti phase was present in the Ti-50at.%Al alloy at a temperature of = 1450C.

Therefore, the results of these two studies indicated that the a-Ti phase in binary

Ti-Al alloys was stable at high temperatures and with compositions containing up to

50at.%Al.

One of the first survey studies of the phase equilibria in the ternary system

was reported by Popov and Rabezova in 1962 [17]. In this study, the Nb rich side of

the ternary phase diagram and the ternary solid solution range of the binary y (TAl)

phase were examined. Isothermal sections were constructed at 1400C, 1200C, and

room temperature from the phase analysis of the heat treated samples. The results

of this study showed the existence of a ternary intermetallic compound that was

given the notation of the y, phase. The y, phase was determined to have a tetragonal

structure with a stoichiometric composition of NbTiAl, (at.%). The lattice parameters

that were determined for the y, phase were a = 3.56A and c = 4.69A. It was also

found that a quasi-binary section existed from NbAl3 to Ti that contained the ternary


































Al, wt -%


Figure 2.1. Shows the binary Ti-Al phase diagram. (a) the previously accepted [10]
phase diagram; (b) the modified section of the phase diagram [51.









10

Yi phase. The NbAl,-Ti quasi-binary section showed that the ternary y, phase

transformed congruently from the liquid phase at a temperature of 1850C. On

opposite sides of the ternary y, phase were the eutectic reactions of L -* + y, and

L y1 + 1. The L p + Y, reaction occurred at = 1550*C, and the L y, + i reaction

occurred at = 1520*C. Although the precise boundaries of the phase equilibria were

not reported, the results of this study indicated that extensive regions of ternary solid

solutions existed for the binary p, (NbAl), a (Nb2Al), 11 (NbAl), and y (TiAl) phases.

In the following year, Wukusick [18] studied the ternary solid solution range

of Nb in the y-TiAl phase. The phase equilibria that were reported in this study were

from X-ray diffraction analysis of alloys that were solution treated at 1425"C and

then water quenched. One of the alloy compositions that Wukusick investigated was

24Nb-26Ti-50Al (at.%), which was close to the stoichiometric composition of the y,

phase (NbTiAl,). However, the results of this study did not support the results of

Popov and Rabezova [17]. In this study, the y phase, showing a large ternary

solubility range for Nb, was observed instead of the y, phase. Thus, the results by

Wukusick contradicted the existence of the ternary y, phase.

A study conducted by Zakharov et al. [19] in 1984 added to the confusion

concerning the existence of the ternary y, phase. This study examined the region of

the ternary Nb-Ti-Al phase diagram where the two quasi-binary sections of NbAl3-Ti

and TiAl-Nb intersected. The alloys investigated in this study were subjected to an

extensive schedule of solutionizing heat treatments with final heat treatments at

1200C, 9000C, and 600C for one hour and then water quenched. The phase

equilibria of the aged samples were determined by X-ray diffraction. The most

significant result of this study was the confirmation that the ternary y, phase, which









11

was reported by Popov and Rebezova [17], existed with the stoichiometry of NbTiAl2.

The study by Zakharov et al. also reported that the y, phase had a tetragonal

structure. However, the lattice parameters that were reported for the y phase by

Zakharov et al. were different from those that were reported by Popov and Rabezova.

The lattice parameters reported by Zakharov et al. were a = 8.418A and c = 4.538A,

while the lattice parameters reported by Popov and Rabezova were a = 3.56A and

c = 4.69A. The reason for the disagreement between the lattice parameters was not

known.

Two comprehensive studies that were published in 1989 by Jewett et al. [20]

and Kaltenbach at al. [21] investigated the phase equilibria of alloys that had

compositions covering the central portion of the ternary phase diagram. The study by

Jewett et al. examined fourteen different alloy compositions formed by heat treating

arc-melted samples at 12000C for up to seven to sixteen days. In a similar manner,

Kaltenbach et al. examined thirty five different alloy compositions that were heat

treated at 1200C for one to seven days. In general, there was good agreement

between the results of these two studies. Both studies indicated that the P phase

existed at 12000C over a substantial part of the ternary phase diagram. The binary

phases were observed to project into the ternary section with large solubility ranges

that loosely followed constant Al compositions. This projection can be seen by the p,

(shown as 8) and cr phases on the binary Nb-Al side and by the y and a phases on the

binary Ti-Al side of the ternary isotherm shown in Figure 2.2. The Tr phase was

determined to be isomorphous between the NbAl3 and TiAlI phases, and was shown to

connect these two phases along the constant 75at.%Al composition line.











12











Ti
1.0

0.9

0.8
0
AC 0.7


r. 3a

0o.




0.3

0.

0. / /^-' / /


1.0 0.9 0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.1 0.0


Nb ATOMIC FRACTION


Figure 2.2. Shows the 12000C isothermal section of the ternary Nb-Ti-Al system
determined by Jewett et al. [201.









13

These results from the studies by Jewett et al. and Kaltenbach et al. helped to

clarify the earlier results, but they also caused future problems concerning the phase

equilibria in the Nb-Ti-Al system. Both of these studies concluded that the ternary y,

phase did not exist in this system. It was found instead that alloys with compositions

close to that reported for the y, phase consisted of the binary y-TiAl phase. These

studies also showed that the y phase had a substantial solubility limit for Nb. Jewett

et al. reported a solubility of up to 30at.%Nb in the y phase, along the 50at.%Al

direction. However, Jewett et al. also claimed that there were two new ternary

phases that existed and these were given the notation of T1 and T2. The location of

the Tl and T2 phases are shown in the 1200C isotherm of Figure 2.2. There was no

structural information reported for the T1 and T2 phases, and only optical microscopy

was used to support the existence of the T1 and T2 phases. It could not be

ascertained whether the TI and T2 phases were present at 12000C, or whether they

were decomposition products of higher temperature phase equilibria.

In recent years, there have been two studies published on the ternary phase

equilibria from the same research group. One study was by Perepezko et al. [12] in

1990 and another study by Das et al. [22] in 1993. The primary purpose of these two

studies was to determine the liquidus projection and to clarify the conflicting results

from the earlier studies of the Nb-Ti-Al system.

The liquidus projection that was determined from the study by Perepezko et

al. [12] is shown in Figure 2.3. This projection was consistent with the recent

modifications to the binary Ti-Al phase diagram [15,16]. It showed that the a phase

formed from the peritectic L + p a reaction at the higher temperature of 1480C,

compared to the y phase which formed from the peritectic L + a y reaction at



















Ti




0.8



0.6
S? (1480
PS 1430)

0.4

1.Ps (137O0)




0 -2/ ,-;/ 4 y \

1.0 0.8 P** f0.6 *e 0.4 0.2 *2
(2449) (2060) (194C) 1 lsC)


Figure 2.3. Shows the liquidus projection of the ternary Nb-Ti-Al system
determined by Perepezko et al. [12].









15

S1450C. There were two important points made from this liquidus projection. The

first point was that the liquidus surface of the p phase covered an extensive portion of

the ternary phase diagram. The second point was that the bivariant L + p + CY and

L + a + y tie-triangles reacted in a class II four phase reaction [23] near the central

portion of the liquidus projection. The product L + p + y tie-triangle then moved

toward the binary Ti-Al side, where it reacted in another class 11 four phase reaction

with the L + p + a tie-triangle. The L + p + a tie-triangle had originally started from

the binary peritectic L + p a reaction. The product of this reaction was the

L + p + a tie-triangle, which terminated at the binary peritectic L + a y reaction.

Thus, these two points indicated that the solidification paths of alloys, that had

compositions near the central portion of the ternary phase diagram, could have

solidified with the P phase or could have complications near the four phase reactions.

Further work by Das et al. [22] indicated that refinements in the 1200C

isotherm had occurred near the composition of Ti4,A]Nb. The revised 1200C

isothermal section, which is currently accepted to be the most accurate

representation of the high temperature phase equilibria in the Nb-Ti-Al system, is

shown in Figure 2.4. The refinement made by Das et al. involved changing the phase

for the region that had the Ti4Al1Nb composition from the T2 phase to the a phase.

The T2 phase, which had been seen in the earlier 1200C isothermal section of Figure

2.2, was identified as the ordered p (B2) phase. However, it was argued that the B2

phase had formed from the specific cooling methods employed, and was not

considered to be an equilibrium phase. Likewise, the T1 phase was also found to

have had the B2 phase, and was disregarded as an equilibrium phase by a similar

argument used for the T2 phase.























































Figure 2.4. Shows the 1200C isothermal section of the ternary Nb-Ti-Al system
determined by Das et al. [22].










i-30AI



~ Ti- 50AI


Figure 2.5. Shows the partial sections of the isotherms in the ternary Nb-Ti-Al
system determined by Das et al. [22]. (a) 12000C; (b) 11500C.


10 at. %









18

The a phase that was studied by Das et al. [22] was assumed to have the HCP

structure since no structural results were reported in the study. Nevertheless, the

phase boundaries near the Ti4AlNb composition of the a phase were shown to be

complicated. The tie-lines surrounding this composition were significantly changed

after just a 50C drop in temperature from 12000C. These are shown in the partial

sections of the 1200C and 1150C isotherms in Figure 2.5. One of the most dramatic

changes in the boundaries of the phase equilibria between 1200C and 1150C

involved the presumed class II four phase reaction of the p + y + a* and p + y + a

tie-triangles (Figure 2.5a) to form the a* + a + p and a* + a + y tie-triangles (Figure

2.5b). This reaction was shown to have shifted the tie-lines in the two-phase a* + a

fields by nearly 90 from the two-phase p + y fields.

The combined results of Perepezko et al. [12] and Das et al. [22] also indicated

that the B2 phase covered an extensive composition range in the central portion of

the Nb-Ti-Al system. The extent of the B2 phase at 1200C was shown previously in

Figure 2.4. However, the compositional boundaries of the B2 phase were not

precisely known.

2.1.2 The Ti-Based Alloy Development Studies

The phase equilibria and the microstructures that can be produced from

various heat treatments of alloys in the Ti-rich part of the ternary Nb-Ti-Al system

have been reviewed by Williams [1] and Rhodes [2]. These alloys show that a variety

of microstructures can be produced by processing or heat treatments to form different

morphologies, distributions and combinations of the equilibrium a and P phases, the

metastable a' and a" martensites, and the metastable a phase. The reason the

different microstructures were produced was due to the existence of the p phase at









19

higher temperatures that could be stabilized down to room temperature with

sufficient Nb additions.

The renewed interest during the 1980s in the development of Ti-aluminides

prompted several studies of the phase equilibria in a2-TiAl based alloys that

contained up to 30at.%Nb additions. In 1988, a study by Banerjee et al. [13] showed

that an orthorhombic phase (O-phase) formed in a Ti-25A1-12.5Nb (at.%) alloy that

was heat treated at 1100C and furnace cooled. This heat treatment produced

equiaxed a2 grains in an ordered p (B2-CsCI structure) matrix at 1100C. The slow

cooling rate resulted in the growth of the a2 grains, which caused the formation of the

O-phase at the triple points of the impinged a2 grains. The O-phase was studied by

selected area electron diffraction (SAED) and convergent beam electron diffraction

(CBED) to show that the structure was consistent with the Cmcm space group and

that the lattice parameters were a = 4.50A, b = 5.88A, and c = 9.60A. The

channelling enhanced microanalysis technique was used to determine the atomic site

occupancy of the O-phase. This technique indicated that the composition of this

phase was based on the TiMAINb stoichiometry.

The presence of the O-phase was subsequently confirmed by Kaufman et al.

124] in 1988. However, the lattice parameters determined for the O-phase by

Kaufman et al. were different than those determined by Banerjee et al. [13]. The

O-phase was observed in Ti-25Al + Nb (at.%) alloys that contained at least 12at.%Nb

and were heat treated between 800C and 1000C. The CBED analysis showed the

O-phase to have the same Cmcm space group, but the lattice parameters were

a = 6.2A, b = 9.4A, and c = 4.7A.









20

The temperature and composition range of the O-phase were investigated in

several studies immediately following the discovery [25-35]. These studies

contributed mainly to the understanding of the phase equilibria in the section of the

phase diagram that connects the binary ao-Ti3Al phase to the O-Ti2NbAl phase. The

Ti-27.5Al + Nb section shown in Figure 2.6 was developed from the results by

Banerjee et al. [32], and is the currently accepted phase equilibria for this part of the

ternary phase diagram.

Evidence showing the existence of a second ternary phase was reported in the

study by Strychor et al. [36] in 1988. This evidence was obtained from low

temperature heat treatments on alloys based on the composition of Ti.Al alloys with

ternary additions from 5 to 17at.%Nb. The alloys containing more than 5at.%Nb

showed a rapid drop in the M, temperature of the a' martensite and retention of the 3

phase, which had ordered to the B2 phase during cooling. Then, the metastable B2

phase was shown to have decomposed to an o-type phase during the low temperature

heating. The SAED analysis of the o-type phase indicated that the structure was

ordered and was consistent with the B82 structure, which was a HCP-Zr2Al prototype

structure.

The w-B82 phase was confirmed in a 1990 study by Bendersky et al. [37]. In

this study, an alloy with the Ti4AlaNb composition was heat treated at 700C for 26

days. This caused the prior B2 matrix to transform completely to the o-B82 phase.

Thus, the stoichiometric composition of the (o-B82 phase was determined to be

'I., \I. N The structure of the a-B82 phase was investigated using SAED and CBED

analysis to show that the space group was P6,/mmc, and the lattice parameters were
































1200


Ti-27 5A10 20
at% Nb


Figure 2.6. Shows the temperature-composition diagram of the TiAl to Ti'AINb
section by Banerjee et al. [32].









22

a = 4.580A and c = 5.520A. The mechanism of this transformation was also studied,

but will be discussed later in section 2.3 on the omega phases.

There was a follow up study by Bendersky et al. [38] in 1990 that reported the

possibility of another ordered derivative of the o-phase. In this study, an alloy with

the composition of Ti-37.5A1-20Nb (at.%) was heat treated at 700C for 18 days. The

results showed that small precipitates formed in the matrix of the o-B82 grains. The

SAED analysis indicated that the lattice parameters of the precipitates were

a = 7.93A and c = 5.52A. These lattice parameters of the precipitates were larger

than those of the a-B8, phase. Two possible structures of this ternary phase were

proposed. One possible structure was based on the Pearson symbol hP18, which had

the P6/mcm space group and Ga4Ti5 prototype structure. The other possible

structure was based on the Strukturbericht D8, structure, which had the P6,/mcm

space group and MnSSi, prototype structure.


2.2 The Sigma Phase

The sigma (a) phase occurs in many transition-metal alloy systems that are of

technological interest in alloy development, such as the stainless steels and the

superalloys. However, the occurrence of the a phase in these alloys has usually been

avoided because of its generally hard and brittle properties at room temperature.

These properties have had a deleterious effect on the mechanical properties, since the

a phase usually forms along the grain boundaries of affected alloys, where cracks can

nucleate and propagate. Conversely, the a phase has a high melting point and a

complex ordered structure, that may provide some strength and creep resistance at

high temperatures. A full review of the structure and properties of the a phase in

various binary and ternary alloy systems is described by Hall and Algie [39].









23

The a phase has been determined in the binary Nb-Al system to form by the

peritectic L + p, 3 a reaction at = 19400C [40]. From Table 2.1, the stoichiometries of

the P, and the a phases were determined to be NbAl and Nb2Al, respectively. The

maximum solubility of Al in binary compositions of the a phase was 12at.%Al at

S1600C. The range of ternary compositions of the a phase in this system is shown in

Figure 2.4 [12,20-22]. The ternary a phase region was shown at 1200C to extend

from the binary Nb-Al side into the ternary Nb-Ti-Al system, along the constant Al

composition section. The solubility range for Ti, in the ternary compositions of the a

phase, was estimated to be =35at.%Ti at 12000C. The L + P, *y a liquidus valley,

which extends from the binary peritectic reaction into the ternary section, decreases

with temperature, as shown in Figure 2.3.

The structure of the a phase, based on the Nb2A stoichiometry, has been

investigated by Wilson and Spooner [41,42]. This phase has been described as a

topologically close packed (TCP) structure, since it can be viewed as consisting of

distorted, hexagonal close packed, layers of atoms that are rotated by 900 between

each alternating layer. Like the a phases in other various alloy systems, the a phase

in the Nb-Al system has a tetragonal structure with the P4Jmnm space group. The

lattice parameters of the stoichiometric Nb2Al composition were measured to be

a = 9.935A and c = 5.169A [11]. However, the lattice parameters of the unit cell for

this a phase do change with the Al content. There are 30 atoms associated with the

unit cell of the a phase. These atoms are arranged with varying degrees of order, on

five different lattice sites in the unit cell. The schematic shown in Figure 2.7 shows

the five different lattice sites in the projection of the unit cell along the c axis of the a

phase. These five sites with their equivalent Wyckoff designations are A (2a), B (4f),









24

C (8i1), D (8i), and E (8j) [43]. The occupation of these sites by the Nb and Al atoms

was found to consist of predominantly Nb atoms on the B, C, and E sites; mostly Al

atoms on the A site; and a mixture of both Nb and Al atoms on the D site [42]. The

coordination number does have an effect on this particular site occupancy. The B site

has the largest coordination number (CN = 15), the A and D sites have the smallest

(CN = 12), while the C and E sites have an intermediate value (CN = 14). Therefore,

both the atomic size and electronic behavior can affect the specific site occupancy in

the c structure.


2.3 The Gamma Phase

The gamma (y) phase from the binary Ti-Al system has received extensive

research during the last ten years. This surge in research has been attributed to

several attractive properties of the y phase, such as low density, good oxidation

resistance, and high temperature strength retention. The development of potential

alloys based on this phase has resulted from studies on the phase relations;

microstructural formation by processing; and microstructure-property relationships

relating to oxidation, deformation, and fracture. Further information on these topics

for the y phase is provided in the review articles by Kim [7,44].

The formation of the y phase in the binary Ti-Al system was recently modified

to show that it solidified in the peritectic L + a y reaction at 14500C, as shown in

Figure 2.1. The solubility range of this phase extends more to the Al-rich side, rather

than to the Ti-rich side, from the exact TiAl stoichiometry. In studies of the ternary

Nb-Ti-Al system, the y phase was found to have an extensive solubility range for Nb

at temperatures in the 1200C range [12,20-22]. These studies indicated that up to

26at.%Nb was soluble in the y phase at 12000C.












































O Z-i
0 '-0


Figure 2.7. Shows the five lattice sites in the projection of the unit cell for the
a phase [41].









26

The structure of the y phase was determined to be the face-centered tetragonal

L10 structure and the P4/mmm space group [10]. The unit cell of this structure has a

layered arrangement of Ti and Al atoms on alternating (001) planes, as shown in

Figure 2.8. The Ti atoms occupy the la and Ic Wyckoff sites at (0, 0, 0) and (%, %, 0),

while the Al atoms occupy the 2e Wyckoff site at (0, %, %) and (%, 0, %). The lattice

parameters, for this equiatomic TiAl composition, were determined to be a = 3.976A

and c = 4.049A (45]. However, these lattice parameters also depended on the Al

content. Measuring these changes in the lattice parameters by the c/a ratio showed

that the c/a ratio increased from = 1.02 to =1.03 with increasing Al content [46].

The atomic site occupancy of the y phase, containing ternary additions of Nb,

was studied by Kronitzer et al. [47] and Jackson [48]. Kronitzer et al. showed that

Nb atoms randomly occupied the la and le Wyckoff sites with the Ti atoms, for small

additions of Nb to the y-TiAl phase. In comparison, Jackson reported the existence of

a new phase that was based on the tetragonal L60 structure and P4/mmm space

group. This structure showed the Ti atoms occupied the la Wyckoff site, the Nb

atoms occupied the Ic Wyckoff site, and the Al atoms occupied the 2e Wyckoff site.

Thus, the new structure reported by Jackson showed that ordering had occurred

between the Nb and Ti atoms on the la and Ic Wyckoff sites. The occurrence of this

ordering reaction from the y-Llo structure was supported by the observation of

APDBs in the y-L6, structure.


2.4 The B2 Phase

The B2 phase has been shown to exist in alloys of the Nb-Ti-Al system at both

high temperatures and over large composition ranges. The composition range of

alloys that had been found to contain the B2 phase included compositions in the




































* la and Ic WyckoffSites: Ti

) 2e WyckoffSite: Al


Figure 2.8. Shows the unit cell of the y-TiAl phase.









28

central portion of the ternary phase diagram [12], compositions close to Ti3Al

containing Nb additions in the range of 5Nb to 30Nb (at.%) [13,24-35], the Ti4Al3Nb

composition [37], and binary Nb-Al compositions with additions of 13.5A1 to 16.9A1

(at.%) [49]. The 1200C isotherm by Das et al. [22] showed that the solid solution

range of the B2 phase was fairly small and was located close to the center of this

ternary system, as shown in Figure 2.5. The temperature of the p to B2 transition

was determined to increase with an increase in the Nb content of alloys based on the

approximate TiAAl composition. More specifically, the p to B2 transition temperature

was found to increase from 11000C for 12.5at.%Nb additions to over 14000C for

25at.%Nb [13,24-26,28-31,33-35]. Likewise, a high transition temperature of more

than 14000C was suggested by Bendersky et al. [37] for an alloy with a Ti4AlNb

composition. The B2 phase was also recently observed in binary Nb-13.5A1 and

Nb- 16.9A1 (at.%) alloys after being heat treated at only 800C for 10 hours [49]. This

B2 phase was determined to be a metastable phase that had formed, instead of the

equilibrium Nb3Al (p,) phase, at the low aging temperature. The formation of the B2

phase, instead of the p, phase, was possible because the activation energy of the

second order p to B2 transition was lower than that of the first order p to P,

transition. Therefore, the diffusion length that was required to form the B2 phase

was shorter and the formation of the B2 phase occurred, rather than the p, phase.

The B2 phase has the ordered BCC (CsCI) structure and the PmTm space

group. The atomic site occupancy of the B2 phase in a Ti-25Al-10Nb (at.%) alloy was

determined by Banerjee et al. [50] using channelling enhanced microanalysis. The

composition of the B2 phase in the two phase a2 + B2 microstructure of this alloy was

found to be Ti-24.5A- 14Nb (at.%). The channelling results of the B2 phase indicated









29

that the two distinct sublattices in the B2 unit cell were occupied mostly by Ti on the

la Wyckoff site at (0, 0, 0) and by a mixture of Ti, Nb, and Al on the lb Wyckoff site

at (%, %, %). The exact atomic percent values that were determined for the two sites

were 48Ti and 2Nb for the la site, and 14Ti, 12Nb, and 24A1 for the lb site. The

schematic shown in Figure 2.9 shows the B2 unit cell using the atomic site occupancy

determined by Banerjee et al.

A 1988 study by Strychor et al. [36] showed that the B2 phase in Ti3Al + Nb

alloys exhibited diffuse scattering and extra maxima in the selected area electron

diffraction (SAED) patterns, and tweed microstructures in the TEM images. The

diffuse scattering and extra maxima in the SAED patterns was interpreted to show

an instability in the B2 lattice that suggested two possible modes of transformation.

The first mode was determined from the observation of diffuse scattering at the

1/3<111> positions in the SAED patterns of the B2 matrix. This diffuse scattering

was attributed to 2/3<111> longitudinal displacement waves, or phonons, that led to

the formation of an ordered o-type phase during low temperature aging of TisAl + Nb

alloys. The second mode was determined from the observation of diffuse streaks that

ran parallel to the <110> directions and extra maxima at 1/2(110} positions in the

SAED patterns of the B2 matrix. These observations were attributed to

1/2<110>{lT0} type transverse lattice displacement waves, that caused localized

strain in the B2 lattice, which led to the formation of the tweed microstructure.

These shear strains were shown to be consistent with the lattice deformation of the

martensitic transformation of the 2H pseudo ortho/hex (orthorhombic/hexagonal)

structure from the B2 phase. The diffuse streaking and extra maxima observed in






































* la WyckoffSite: Ti

lb WyckoffSite: Nb and Al


Figure 2.9. Shows the unit cell of the B2 phase with the atomic site occupancy
determined in a Ti-24.5Al-14Nb (at.%) alloy by Banerjee et al. [50].









31

the SAED patterns were formed from the tweed microstructure, and were explained

as rel-rods that intersected the Ewald sphere.

This tweed microstructure has been observed in the B2 phase of many alloy

systems that exhibited martensitic transformations [51]. This relationship between

the tweed microstructure and the martensitic transformation in these types of alloys

has been investigated by Robertson and Wayman [52-54], Tanner et al. [55], and

Schryvers and Tanner [56]. Robertson and Wayman showed that the tweed

microstructure was formed by <110>{110) static displacement waves, which resulted

in the softening of the elastic constant C', where C' = Y(C,1 C12), in the B2 phase of a

63Ni-37Al (at.%) alloy. Tanner et al. and Schryvers and Tanner showed that the

tweed microstructure in the B2 phase of the 63Ni-37Al (at.%) alloy was composed of a

fine-scaled mosaic assembly of non-uniformly distorted and micromodulated domains,

which were coined inhomogeneously strained domains (ISDs). The ISDs were

examined by high resolution electron microscopy (HREM). The ISDs were shown to

lie parallel to {110) plane traces in the B2 matrix and to have a size of =40-60A in

length, with an average spacing of 13A thick. The computer simulations by

Schryvers and Tanner confirmed that the atomic structure of the ISDs was consistent

with small transverse shuffles, plus shear distortions of the <110>{110) type. The

displacements in the ISDs were correlated with the low energy transverse

Z4 (50)-TA2 phonon mode that was found to have an anomolous, temperature-

dependent, incomplete softening at C a 0.16. It was theorized that the ISDs were

formed by the dynamic softening of the C' constant, which resulted from the TA,

phonon that became coupled to the static strain fields of defects in the B2 lattice. It









32

was suggested that the nucleation site of the martensitic phase was determined by

the strength of the strain field that was associated with the defect.


2.5 The Omega Phase

The omega (c) phase has been an intriguing phase that was first discovered in

thermally aged, p-stabilized, titanium alloys in 1954 [57]. Since then, it has been

observed in the group IV transition elements of Ti, Zr, and Hf; in numerous alloys

that consisted of the group IV elements; and in the elements immediately to the right

of the group IV elements (the d-rich transition elements of Nb, V, and Mo) [58]. The

a phase has also been observed in many alloys that were not based on the group IV

elements, such as: Cu-Zn [59,60], Cu-Sn [59,611, Cu-Zn-Al [62], Cu-Zn-Si [63],

Cu-Al-Ni [64], Ag-Al [65], Ag-Mg [66], Fe-Al [67], and Ni-Al [68]. One of the main

reasons for investigating the a phase was because of the harmful effect this phase has

on the mechanical properties of these alloys. The formation of the a phase has been

shown to embrittle the alloy by decreasing the ductility and increasing the hardness.

Extensive studies have been conducted on the morphology; the effects on the physical

properties; the nature of the diffuse x-ray, electron, and neutron scattering; the

transformation kinetics; and the transformation mechanism of the a phase. The

results of these studies were summarized in a review article by Sikka et al. [58].

The o phase that forms in Ti has been suggested to be the low-temperature

high-pressure modification of the p phase [69]. Normally, the a-HCP phase is the

stable phase in the unary Ti system below the allotropic transition of -8820C at

atmospheric pressure. However, the a phase has been shown at high pressures to be

the stable phase, instead of the a phase, based on the equilibrium unary P-T diagram

that was developed for Ti [58]. The structure of the ideal o phase in pure Ti was









33

found to be the HCP structure, with the P6/mmm space group. The unit cell of this

ideal phase contains three atoms: one atom located on the la Wyckoff site at (0, 0, 0)

and two atoms located on the 2d Wyckoff site at (1/s, %2, %) and (%, /s, %). However,

the ( phase has also been shown to have the trigonal structure with the Pmml space

group in Ti-based alloys. The difference in the symmetry between the o-trigonal and

a--HCP phases was due to the small atomic displacement of the two atoms on the 2d

Wyckoff site in the aO-HCP phase. Thus, the atomic site occupancy of the (c-trigonal

phase was shown to have one atom located on the la Wyckoff site at (0, 0, 0) and two

atoms located on the 2d site at (/s, %, z) and (%, Va, z). The z parameter indicates

the magnitude of the atomic displacement and that the displacement occurs in the

direction of the c-axis in the trigonal unit cell.

The p to a transformation has been described by a plane collapse mechanism

that involves three (111}) planes. In this mechanism, one pair of planes collapses

together to an intermediate position, while the third plane remains unaltered [58].

The atomic displacements that were required for this transformation have been

described by a soft-mode transformation mechanism. This mechanism involves a

longitudinal sinusoidal wave with atomic displacements, U = Wsin[qx + i(x)]; a wave

vector, q. = 2/3<11 l>2/ap; a phase for the three possible variants, O(x) = 0, 2i/3, 4n/3;

and an amplitude, W = ap/6 for w-HCP or a value less than this for e-trigonal. The

{111)} plane collapse mechanism and the soft mode mechanism of the p to o

transformation are shown in the schematic of Figure 2.10. This schematic shows the

projection of both the (110), planes for the P phase and the (1100), planes for the a

phase. If the collapse of the two {111}1 planes is incomplete, then the atoms on the B









34

plane are displaced. This results in a rumpled B plane and the trigonal symmetry for

the a phase.

The transformation of the 1 phase to the a phase has been shown to form

twelve variants of the a phase. These variants were formed from the orientation

relationship that the a phase has with the p phase:

(0001)o I(11)l}p and <110>(o 1<1TO>p.

The twelve variants were formed from four rotational variants that each had three

translational variants, as shown in the schematic of Figure 2.11. The two rotational

variants were formed by aligning the (0001). plane parallel to either the (111), or

(1iT ) planes of the P phase. The three translational variants were determined from

the (11)}p plane that remained unaltered during the plane collapse. Since the

stacking sequence of the (111), planes in the BCC p phase is ...ABCABC..., then the

(0001), plane of the A-variant is formed from the A plane, the B-variant from the B

plane, and the C-variant from the C plane. Therefore, six out of the twelve total

variants are illustrated in the schematic of Figure 2.11.

There have been several studies in the past that have investigated the

influence of Nb on the formation of the o-type phase in binary Ti-Nb alloys [70-73].

These studies have shown that the a phase can transform both athermally and

isothermally from the P phase in the Nb-Ti alloys. The primary difference between

the two formation mechanisms was that the composition of the athermal o phase was

the same as that of the p matrix, while the composition of the isothermal a phase was

different from that of the p matrix. However, the isothermal a phase was determined

to form from the p matrix by the same mechanism as the athermal a phase, which

involved the collapse of alternating pairs of {111} planes.








35








(111)p trace Oft
A 0

A
C .. -O "










[0001]o
A (0001)c- trace

B 00
A
B 00 '
A--------------
A ----------@----Q----A


Figure 2.10. Shows the ( 111), plane collapse model of the P to m transformation. The
view is normal to the (110)p planes.














































(111) Plane Trace


(111) Plane Trace




Figure 2.11. Shows two rotational variants, each one containing three translational
variants, of the c phase formed from the P to (o transformation. The
view is normal to the (110), planes.









37

The most systematic studies of the a phase in the binary Nb-Ti system were

by Moffat and Larbalestier [72,73]. In these studies, the effect of cooling rate and

isothermal aging were investigated in alloys that contained 20Nb to 35Nb (at.%). The

results of these studies showed that there was a competition between the formation

of the o phase, the metastable a" martensite, and the equilibrium a phase depending

on the cooling rate from 1000C and the isothermal aging temperature. It was shown

that water quenching caused the formation of the a" martensite, while air cooling and

furnace cooling caused the formation of the a precipitates. The aging experiments

indicated that the equilibrium a phase was formed at the higher aging temperatures,

such as 400C to 500C, while the lower aging temperatures, such as 200C and 300C,

caused the formation of the a phase. The competition between the a" martensite and

the a precipitates in the alloys was determined to have resulted from the lowering of

the M, martensitee start) temperature of the a" martensite by the addition of

>20at.%Nb [72]. It was also suggested by Moffat and Larbalestier that since the ao

and a" phases were never observed together in the microstructure, then which ever

phase formed first excluded the other phase from forming.

There have been a only a few studies that have investigated the formation of

the c-related phases in ternary Nb-Ti-Al based alloys [36-38]. The study by Strychor

et al. [36] was the first to show that ordered derivatives of the ao phase were formed

in Ti3Al + 5Nb to 17Nb (at.%) alloys. The a-related phase that formed in these water

quenched samples was detected only in the SAED patterns of the B2 matrix. These

diffraction patterns showed diffuse streaking and extra diffraction maxima at

1/3(111), positions that were attributed to an ordered derivative of the c-related

phase. The effect of isothermal aging for different times at 400C and 500C was









38

found to slowly degrade the tweed pattern of the B2 matrix and to intensify the

reflections of the ordered a-related phase in appropriate SAED patterns. Strychor et

al. determined that the ordered o-related phase had the B8, structure, which

consisted of the P6,/mmc space group. The o-B82 structure was found to have A2B

ordering, where it was assumed that the A represented a mixture of Ti and Nb atoms

and the B represented the Al atom.

The transformation mechanism of the o-B82 phase from the B2 phase in a

'1, \1 ,M alloy was investigated by Bendersky et al. [37]. This study showed that

furnace cooling from 14000C resulted in the transformation of the B2 matrix

completely to the trigonal o"-type phase. This )"-type phase was determined by

CBED to have the P-ml space group. It was deduced from microstructural evidence

that the B2 phase was stable at 14000C, and that the subsequent heat treatment of a

furnace cooled sample for 26 days at 700C resulted in the transformation of the

o-B82 phase from the (" matrix. The site occupancies of the a" and 0-B82 phases

were determined by X-ray diffraction. This analysis showed that the Nb atoms

flowed out of sites on the collapsed double layers and into sites on the single layers.

The vacated sites on the double layers were then preferentially occupied by Ti and Al

atoms, and the sites on the single layers were enriched by a mixture of Nb, Ti, and Al

atoms. The driving force for the chemical exchange between these atoms was shown

to be from the maximization of the Ti-Al bonds on the double layers. This conclusion

was based on the fact that the sites on the double layers had a greater coordination

number of nearest neighbors than those on the single layers in the a" and a-B82

phases. The Ti and Al atoms were shown to have a preference for the sites on the









39

double layers, since the Ti-Al bonds were the most stable and had the shortest bond

length compared to the other possible types of bonds.

In the study by Bendersky et al. [37] the transformation of the parent B2

phase to the final (o-B82 phase was described by a series of structural changes that

involved subgroup/symmetry relations in crystallography [43]. The transitions that

described the B2 to o-B82 phase transformation are shown in the schematic of Figure

2.12. These individual transitions were connected together by subgroup/supergroup

relations that indicated how many variants of the product phase were formed from

the parent phase, whether the symmetry in the product phase was increased or

decreased relative to the parent phase, and whether there was chemical ordering or

whether distortions had occurred. The symmetry and atomic site occupancy results

were then used to show that the transformation path from the B2 phase to the o-B82

phase traversed the trigonal o" phase. This trigonal o" phase had the lowest

symmetry as compared to either the B2 or a-B82 phases. The transitions from the B2

phase to the a" phase occurred during furnace cooling, since these transitions

involved only the partial collapse of the double layers and the incomplete exchange of

atoms between double and single layers. However, the transition from the a" phase

to the a-B82 phase required thermal energy, since it involved the full collapse of the

double layers, the complete chemical exchanges that led to the disordered single

layers, and the fully populated sites on the double layers by Ti and Al atoms.


2.6 The Ortho/Hex Phases

The ortho/hex phases are the phases with the closely related HCP and

orthorhombic structures that have been observed in the Ti-rich alloys of the Nb-Ti-Al

system. These phases are formed from alloys with compositions that show the P or





















Chemically Chemically
Disordered Ordered

Im3m
A2: bce

Chemical Ordering


Homogeneous Distortion 12
[4 1m3m
1 B2: CsC1


R3m Homogeneous Distortion
[41
Chemical Ordering
[2]

P6/c'2m 4
P6/mo
u-Ti -Collapse R3m

Completion of Collapse
(2) Tin- -1 1
P3ml w-Collapse P6/mmc
Trigonal-o [3) B8,

Chemical Ordering | Disordering
[2] between Single Layers
P3ml-
c=. 2e w 2


Figure 2.12. Shows the transformation paths from the P phase to the
on-related phases using subgroup and symmetry relations from
Bendersky et al. [37].









41

B2 phase to be stable at high temperatures. This stable P or B2 phase can then be

quenched and retained at room temperature, or it can undergo a variety of

transformations to different HCP and orthorhombic phases depending on the

composition, heat treatment, and cooling rate. Thus, the studies of these different

HCP and orthorhombic phases are divided into two different groups: the disordered

structures of the a, a', and a" phases in binary Nb-Ti alloys and the ordered

structures of the a2, a2', and O (orthorhombic) phases in ternary TiAl + Nb alloys,

with constant Al content. This division facilitates a comparison of the influences of

Nb and Al on the structure and transformation of these phases in the Ti-rich alloys.

2.6.1 Disordered Structures

The HCP and orthorhombic phases that were observed in the binary Nb-Ti

alloys consisted of the equilibrium a phase and the metastable a' and a" martensite

phases. Both equilibrium and metastable have disordered site occupancies between

the Nb and Ti atoms. The structure of the equilibrium a phase was determined to be

the HCP structure and the P63/mmc space group [10]. The metastable a' phase was

determined to have formed instead of the a phase when a martensitic transformation

occurred by rapid cooling from the high temperature P phase in Ti-based alloys

containing up to "7at.%Nb [72,74]. The structure and the lattice parameters of the a'

phase were identical to the a phase. For binary alloy compositions greater than

S7at.%Nb, the a" phase was determined to have formed instead of the a' phase by a

similar martensitic transformation from the P phase [721. The a" phase was found to

have the C-centered orthorhombic structure and Cmcm space group [75].

The orthorhombic structure of the a" phase resulted from distortions between

the Nb-Ti bonds. These distortions became more noticeable in binary alloys that









42

contained more than s7at.%Nb [72]. The range in compositions for the distortions in

the a and b lattice parameters, as measured by the a/b ratio, increased with the Nb

content from 0.578 for 7at.%Nb to 0.654 for 20at.%Nb [74]. It was also determined

that the orthorhombic structure of the a" phase was produced from the a'-HCP

structure by varying the positions of the atoms on the Wyckoff sites in the a" phase

[72]. There were four lattice sites in the Cmcm space group of the a" structure that

could be described by the 4c Wyckoff site [43]. The atomic positions of these 4c sites

were (0, y, %); (0, y, %); (%, y+%, %); and (%, y+%, %). The distortion in the

orthorhombic structure was indicated by the y parameter, which was -0.2 for Ti-Nb

alloys [76]. The y parameter for the HCP structure in Ti-Nb alloys was 0.1667. This

value for the y parameter of the HCP structure was based only on the a/b ratio, and

would produce the ideal a/b ratio of 0.578 for the undistorted HCP structure.

The M, temperature of the martensitic transformation for the a' and a" phases

was determined by Jepson et al. [74] and Moffat and Larbalestier [72] to be

dependent on the Nb content of the binary Nb-Ti alloys. The study by Jepson et al.

showed that the M, temperature dropped rapidly from -850C in Ti to =3000C in the

Ti-17.at.%Nb alloy. The study by Moffat and Larbalestier showed that the Ms

temperature fell below room temperature in alloys that contained more than

s30at.%Nb. This conclusion was based on the observation that the retained P phase

was the only phase that was observed in alloys of these compositions after water

quenching from 1000C.

Jepson et al. also showed that the M, temperature was a function of cooling

rate in the binary Nb-Ti alloys. It was found that rapid cooling rates could suppress









43

the martensitic transformation of the a" phase. From this result, it was suggested

that the a" structure could form both isothermally and martensitically.

2.6.2 Ordered Structures

The equilibrium a2 and O phases and metastable a,' phase form in ternary

TiAl + Nb alloys with ordered structures. It was determined that the as and a2'

phases formed in TiAl + Nb alloys with less than llat.%Nb [24,26,271. The a2

phase has the IICP structure with the Ti3AI stoichiometry and P6,/mmc space group

[10]. The ordering between Ti and Al in the a2 structure causes the a-axis lattice

parameter to be twice the a-axis lattice parameter of the disordered a-Ti structure.

The a2' phase has the same HCP structure and lattice parameters as the a, phase,

but has been shown to form by a martensitic transformation during rapid cooling

from the high temperature p phase [24,26,27]. The O phase forms in TiA1 + Nb

alloys that contain more than 12at.%Nb [13,24-32]. The O phase has the C-centered

orthorhombic structure with the Ti2AINb stoichiometry and Cmcm space group [13].

In Ti3Al + 0-1 lat.%Nb alloys, the a2 and a2' phases can both form from the

high temperature P or B2 phase during cooling. However, only the equilibrium a2

phase can form during isothermal heating of alloys containing the retained P phase.

In fact, the a, and a2' phases are so cooling rate dependent, that it is possible to

retain the high temperature disordered P phase that exists over these composition

ranges by rapid quenching. This disordered P phase then orders to the B2 phase at

lower temperatures in alloys that contained more than =5at.%Nb [24]. The

metastable a2' phase usually forms by martensitic transformation for most cooling

rates. In comparison, the equilibrium as phase forms by nucleation and growth

processes only at high temperatures and for extended periods of time [27,32,35]. This









44

equilibrium a2 phase can form from the p phase as grain boundary allotriomorphs,

intragranular plates, and equiaxed grains depending on the alloy composition and

thermomechanical history.

The M, temperature of the martensitic transformation has been shown to

decrease rapidly in the Ti3Al + Nb alloys with increasing Nb content [24,26,27,36]. In

the binary Ti3AI alloy, the martensitic transformation was found to be impossible to

suppress by rapid cooling [77]. The a,' phase that formed in this martensitic

transformation occurred by first forming the a' phase and then ordering to the a2'

phase. In ternary Ti3AI + Nb alloys, the a2' martensitic transformation could be

suppressed completely, provided that the cooling rate was rapid enough. A very high

cooling rate, such as by splat quenching, was necessary to suppress the martensitic

transformation in a TiAl + 5at.%Nb alloy, but a lower cooling rate could achieve the

same result in alloys with higher Nb content [24]. For the alloy compositions and

cooling rates that showed the occurrence of the a2' martensitic reaction, it was

suggested that plates formed with the a' phase first and then later ordered to the a2'

phase [27]. The development of this transformation sequence was based on the

observation of midribs and anti-phase domain boundaries (APDBs) in the plates.

These APDBs were the same as those that were observed in the a to as ordering

reaction of binary Ti2Al based alloys [78].

The formation of the O phase in 'Ti2Al + Nb alloys with Nb contents of 12 to

30at.%Nb requires thermal activation, since water quenching these alloys from high

temperatures retains the P phase with the B2 structure at room temperature [13,24-

35]. Therefore, the O phase observed in these studies has been found to form only by

slowly cooling from the high temperature p or B2 phase, or by isothermally heating









45

the retained p phase. The main points concerning the 0 phase that were determined

from these studies were the O phase was structurally related to the a2 phase [13], the

O phase could form from the B2 phase by a shearing mechanism that involved

thermally activated processes [27,291, the O phase could form from the B2 phase by a

composition invariant transformation [31], the transformation of the O phase could be

described crystallographically [33,34], and the O phase could exist with two different

atomic site occupations [31,35]. The details of these points are covered in the

following discussion.

The O-TiAINb phase was found to be structurally related to the a2-Ti3Al

phase and to involve additional ternary ordering that caused small orthorhombic

distortions [13]. The structural relationship between these two phases is shown in

the schematic of Figure 2.13. The atomic site occupancy of the a2-Ti3Al phase shown

in Figure 2.13a consists of Ti atoms on the 6h Wyckoff sites and Al atoms on the 2d

Wyckoff sites [43]. Kronitzer et al. [47] showed that the addition of Nb to the a2

phase, in amounts that remained in solid solution, preferentially occupied the 6h sites

with Ti atoms. However, in TiAl + Nb alloys that contained more than 12at.%Nb,

the O-Ti2AINb phase was found to form instead of the a2 phase [13,24-33]. The

O-Ti2A1Nb structure shown in Figure 2.13b involved further ternary ordering of the a2

phase that caused the unit cell to be distorted in the direction of the a-axis and b-axis

lattice parameters which are contained in the (001) planes. In the O-TiAINb

structure, the Ti, Nb, and Al atoms were determined to predominantly occupy three

different lattice sites. These were the 8g Wyckoff site by Ti, the 4cl Wyckoff site by

Nb, and the 4c2 Wyckoff site by Al [13,28]. Mozer et al. [28] performed a structural

refinement using a sample that had the composition of Ti-25at.%Al-25at.%Nb to











(b)



I* l




0 -0


b


a


* 6h WyckoffSite: Ti

* 2d WyckoffSite: Al


W 4c, WyckoffSite: Nb

O 8gWyckoffSite: Ti

S4c, WyckoffSite: Al


Figure 2.13. Shows the relationship between the crystal structures of the a2-Ti3Al
phase and the O-Ti2AINb phase. (a) the ac-Ti3Al phase (P63/mmc space
group); (b) the O-Ti2AINb phase (Cmcm space group). The dark shaded
atoms are at z = 0 and the light atoms are at z = %c.









47

determine the relative atomic occupancies of the three different Wyckoff sites in the

O-Ti2AlNb structure. The results of this study showed that the Nb atoms occupied

18% of the 8g sites which were Ti-rich, and that the Ti atoms occupied 18% of the

4cl sites which were Nb-rich. The 4c2 site was found to be occupied by only Al atoms.

The effect of the Nb content on the distortions in the O-TiAINb phase was

studied by Kestner-Weykamp et al. [26]. These distortions were measured by the a/b

ratio and were found to increase from -0.632 for the O-Ti2AlNb phase containing

20at.%Nb to =0.645 for the O-Ti2ANb phase containing 30at.%Nb.

The O-phase has been found to form from the 0 phase as plates by a lattice

shear mechanism in the studies by Kestner-Weykamp [27] and Bendersky et al. [29].

Kestner-Weykamp examined plates that formed during air cooling from the p phase,

which was present at 12500C, in a Ti3Al + 20at.%Nb alloy [27]. These plates

contained defect structures that consisted of midribs, columnar APDBs, and {110}

twins aligned roughly parallel to the midrib. From the analysis of these defects, the P

to O transformation was determined to be by a lattice invariant shear mechanism,

that initially formed the orthorhombic a" shear product with the plate shape and

then later ordered to the 0 structure. A subsequent diffusional growth mechanism

was used to explain the thickening of the plates and the formation of the columnar

shaped APDBs.

The formation of the 0 phase from the retained B2 phase during isothermal

heating of two alloys was investigated by Bendersky et al. [29]. The two alloys

examined in this study had compositions of Ti-12.2Al-37.2Nb and Ti-23.9AI-25Nb

(at.%). Following the heat treatment at 7000C for 26 days, the 0 phase was observed

as plates in the Ti-12.2Al-37.2Nb alloy and as equiaxed grains in the Ti-23.9A1-25Nb









48

alloy. The plates were determined by energy dispersive X-ray spectroscopy (EDX)

analysis to have the TiAlNb composition. This composition of the plates was

different than the composition of the B2 matrix, which was Ti-10Al-45Nb (at.%)

composition. However, the transformation of the plates was still described using the

minimization of elastic strain energy approach which was consistent with the

phenomonological theory of martensitic transformations [79]. This theory accurately

predicted the habit plane and rigid body rotation of the plates using the measured

lattice parameters of the B2 and O phases in this alloy. It was determined for the

Ti-23.9A1-25Nb alloy that the formation of the equiaxed grains involved two steps:

initially the retained B2 phase transformed completely to a highly faulted O phase

and then later recrystallized into fault free grains.

The formation of the O phase from the B2 phase in a Ti-24Al-15Nb (at.%) alloy

was described by a composition invariant transformation mechanism in the study by

Muraleedharan et al. [31]. This alloy was heat treated for short times that lasted

from one to sixty minutes at temperatures of 800C, 900C, and 950C. The shorter

aging times and lower temperatures favored the complete transformation of the

retained B2 phase to the O phase without a change in composition. The O phase

formed as plates, which contained complex defect structures determined to be coarse

APDBs with the displacement vector of 1/4[110], fine APDBs with the displacement

vector of 1/2[100], and stacking faults with the displacement vector of 1/10[025]. The

coarse APDBs observed in the O plates were shown to have the same size and a

related displacement vector to the APDBs that were observed in the retained B2

matrix prior to aging. The presence of these APDBs with no variance in composition

supported the conclusion in this study that the transformation from the B2 phase to









49

the O phase occurred by a shear type mechanism. However, this transformation

mechanism did require short-range diffusion, since the fine APDBs observed in the

O plates were formed from an ordering reaction that required atomic exchanges

between lattice sites.

Bendersky et al. [33,34] described the formation of the 0 phase from the high

temperature P phase using a crystallographic model based on a sequence of structural

changes related by subgroup and symmetry relations. Three alloys with compositions

of Ti-25A1-12.5Nb, Ti-25A1-25Nb, and Ti-28Al-22Nb (at.%) were examined after

heating at 1100C for four days. This heat treatment caused the partitionless

transformation of plates in these three alloys. The analysis showed different defect

structures in the plates of the Ti-25A1-12.5Nb composition as compared to those in

the two Nb-rich alloys. From the identification of these defect structures, such as

APDBs and stacking faults, the transformation of the plates in the Ti-25A1-12.5Nb

alloy was shown to have occurred from the disordered p phase, while the

transformation of the plates that formed in the two Nb-rich alloys occurred from the

B2 phase. These results were then used to show that the plates present in these

alloys formed along two different transformation paths, as shown in Figure 2.14. The

plates in the Ti-25Al-12.5Nb alloy formed from the disordered 1 (Imam) phase and

then followed the path that passed through the intermediate HCP structures before

reaching the final 0 structure. In comparison, the plates observed in the two Nb-rich

alloys formed from the B2 phase and then followed the path through the B19

structure to the final O structure. There was no structural confirmation of the

intermediate transitional structures in each of these two paths. However, evidence of



























Imlm (A2)

3
PmIm (B2) 14/mmm
S 2 2

P4/mmm Fmnmm P64/mmc (A3)

2 2 2 4

Cmmm Cmcm (A20) P63/mmc (D0,)

2
Pmma (819)


CLcm i(A;BC)


Figure 2.14. Shows the transformation paths from the p phase to the O phase using
subgroup and symmetry relations from Bendersky et al. [33).









51

the individual transitions was obtained from the analysis of the defect structures in

the plates and from the partitionless nature of the transformation in these alloys.

The O phase has been found to exist in two structural forms with different

atomic site occupancies according to studies by Muraleedharan et al. [31,35]. In the

first study by Muraleedharan et al. in 1992, the O phase was formed at the lower

temperature of 800C in the Ti-24Al-15Nb alloy and was determined to have the

commonly observed site occupancy of Ti on the 8g Wyckoff site, Al on the 4cl Wyckoff

site, and Nb on the 4c2 Wyckoff site. However, the 0 phase that formed at the

slightly higher temperature of 900C showed a different site occupancy. This O phase

was determined to have a random occupancy of Nb and Ti on the 8g and 4c2 Wyckoff

sites, while Al still occupied the 4cl Wyckoff site. In a recent study by

Muraleedharan et al. in 1995, the same results showing two different site occupancies

for the O phase were obtained from a series of heat treatments that were conducted

on Ti-27.5at.%Al alloys with up to 25at.%Nb additions. In both studies, the site

occupancies of the two different O phases were determined from intensity variations

between reflections in the CBED patterns and by channelling enhanced

microanalysis. The CBED analysis was conducted in thin regions of the O phase to

minimize the dynamical scattering effects. The order parameter (S) was defined in

terms of the Ti and Nb site occupation of the 8g and 4c2 Wyckoff sites. This

parameter was calculated by thermodynamic analysis and showed that a random site

occupancy between Ti and Nb atoms on these two lattice sites was stabilized at

higher temperatures. Thus, it was suggested in both of these studies that Al

stabilized the disordered a" martensite structure. This disordered a" martensite

structure is a metastable phase that forms in the binary Nb-Ti system and then










52

orders into the ordered O phase, which is an equilibrium phase that forms in the

ternary Nb-Ti-Al system.















CHAPTER 3
EXPERIMENTAL PROCEDURES


3.1 Material

The compositions of the alloys used in this investigation were specified to

Pratt and Whitney, who then manufactured the alloys. Out of a total of ten alloy

compositions originally manufactured, the results of three alloys were used in this

study. The nominal compositions of these three alloys (referred to as alloys 2, 3, and

4) are given in Table 3.1.

The main criterion for selecting the three alloy compositions was that each

contain the BCC p phase in the microstructures at high temperatures. Therefore, the

alloy compositions were chosen based on the 12000C isotherm shown in Figure 2.2,

which was developed from previous studies of the ternary Nb-Ti-Al phase diagram

[17-20]. Thus, it was expected that at 1200C the microstructures of alloy 3 should

consist of a single P phase and alloys 2 and 4 should consist of the 0 + x phases,

where x is the y phase (alloy 2) or the a phase (alloy 4). One of the purposes of this

investigation was to then construct the phase equilibria for the alloys at higher

temperatures than 1200C.

The alloys were supplied by Pratt and Whitney in the form of 200 gram

arc-melted samples. The as-received samples had been arc melted a total of four to

six times to ensure complete chemical mixing. This was verified by composition line

profiles performed on cross sections of the as-received samples using the Electron

Microprobe Analyzer (EMPA). These profiles did not show significant chemical

53

























Table 3.1. The nominal compositions of the as-received alloys and the compositions
determined by microprobe analysis of the re-arc melted alloys.


Analysis Composition (at.%)
Alloy
Method Nb Ti Al

Nominal 27 33 40
2
Microprobe 26.8 ( 0.2) 33.8 ( 0.2) 39.3 ( 0.2)
Nominal 50 40 10
3
Microprobe 49.8 ( 1.0) 40.6 ( 0.7) 9.57 (a 0.3)
Nominal 42 28 30
Microprobe 41.4 0.8) 29.5 0.3) 2
Microprobe 41.4 (+ 0.8) 29.5 (: 0.3) 29.1 (a 0.5)










55

inhomogeneities from top to bottom or from center to outer edge. However, the

microstructural characterization of the as-received alloy samples did show

inhomogeneous microstructures that were observed primarily along the thickness

direction of the cast samples, i.e. from top to bottom. This inhomogeneity was most

probably due to uneven solidification rates between the surface making contact with

the water cooled Cu plates of the arc melter and the untouched top surface of the

samples.

Due to the significant nonuniformities that were observed in the as-received

microstructures, a procedure was adopted that involved re-arc melting fragments of

the 200 gram samples into smaller 3 gram samples. The re-arc melting was

performed at the University of Florida using a non-consumable tungsten electrode

under pressurized flowing argon gas. The fragmented pieces were placed in cavities

on a water cooled copper base plate and re-arc melted (RAM) a total of at least 4

times to ensure complete mixing, in case there were some chemical inhomogeneities

between the fragments. The molten samples took approximately 2 to 3 seconds to

solidify and reached a temperature that gave them a metallic lustre. This procedure

reduced the microstructural inhomogeneities of alloys 2, 3, and 4. The compositions

of the re-arc melted samples of alloys 2, 3, and 4 were analyzed by electron

microprobe and are given in Table 3.1. The results obtained by the microprobe

analysis for the re-arc melted alloys are close to the nominal compositions of the

as-received alloys.

The interstitial oxygen and nitrogen content of the as-received alloy 2 was

determined by wet chemical analysis at Teledyne Wah Chang Albany (TWCA). The










56

results of this analysis showed that the oxygen content was -490 parts per million

(ppm) and the nitrogen content was -48 ppm.


3.2 Heat Treatments

The heat treatment experiments were conducted with two types of furnaces

depending on whether a fast or slow cooling rate was needed. The two furnaces used

were a CM model 1600 vertical tube furnace and a Vacuum Industries high vacuum

furnace.

The vertical tube furnace was used when fast cooling rates were needed, such

as during water quenching. This furnace consisted of a mullite tube that was

surrounded by MoSi2 heating elements. The mullite tube was sealed at the top and

bottom with water-cooled removable fixtures. When the fixtures were closed at both

ends, the tube could be pressurized slightly with a flowing stream of argon gas that

entered at the top and exited at the bottom fixtures. The top fixture was designed to

allow both a sample and a type B thermocouple to be positioned within the heating

zone of the furnace. The thermocouple was used to monitor the temperature during

the heat treatment and was also used to calibrate the heating zone. The calibration

showed that the heating zone was constant to within :5C over the tube length of

6cm from the center of the furnace. The sample was placed on an alumina boat

which was suspended in the heating zone with molybdenum or tungsten wire. The

temperature difference between the sample and the thermocouple was estimated to

be less than 5C, since the heating zone of the furnace was radially uniform due to

the cylindrical design. Following is the typical experimental procedure:

(1) Ramp the furnace at -100C/min to the heat treatment temperature.

(2) Insert the sample into the heating zone and seal the top fixture.










57

(3) Hold the sample under flowing argon for the duration of the heat treatment.

(4) Open the bottom fixture, cut the wire, and allow the sample to fall into the

quenching media.

Water was used as the quenching media in this study in order to obtain rapid cooling

rates. However, there were a few heat treatments that were performed in the

vertical tube furnace in which the sample was simply dropped from the heating zone

onto the bottom fixture. This permitted the samples to be air cooled, which

represented a cooling rate that was intermediate between water quenching (fast) and

furnace cooling (slow).

The vacuum furnace was used in experiments that needed a slow cooling rate.

This furnace consisted of an alumina crucible that was surrounded by a tantalum

resistive heating cage. A type R thermocouple was used to monitor the temperature

and was positioned 3cm from the sample in the heating zone. The furnace

incorporated a diffusion pump capped by a water baffle, so that the furnace was

capable of obtaining a pressure of 1 to 4 x 10'6 Torr depending on the heat treatment

temperature. Following is the typical procedure for using the vacuum furnace:

(1) Place the sample on the alumina crucible.

(2) Place the bell jar over the sample and pump to the base vacuum pressure.

(4) Ramp the furnace to the heat treatment temperature at ~100C/min.

(5) Hold at the heat treatment temperature for the duration of the experiment.

(6) Turn the power off to the transformer and allow the sample to cool to room

temperature inside the furnace.

(7) Vent to atmospheric pressure when the temperature, monitored by the

thermocouple, is below -200C.










58

This procedure involved a slow heating rate for the sample up to the aging

temperatures since a high vacuum had to be obtained prior to heating the sample.

The temperature was monitored with the thermocouple during furnace cooling and

followed a parabolic curve for the heat treatments.

3.2.1 Long-Term Heat Treatment Experiments

Table 3.2 lists the long-term heat treatment conditions of samples that were

used in this study. The list includes the initial cast condition, the heat treatment

parameters, and the cooling method that was employed.

Most of the heat treatments that were used in this study were conducted with

3 gram RAM samples using the vertical tube furnace. These samples were heat

treated in the as-cast RAM condition for 2 to 12 hours, depending on the

temperature, and then water quenched. Heat treatments conducted above ~1200C

varied from 2 to 4 hours in duration while those that were conducted below this

temperature were either 12 or 16 hours.

Three of the heat treatments were performed with material that was cut from

the as-received 200 gram arc-melted sample of alloy 4. The three samples used in

these heat treatment experiments were first heated to 1550C for 2 hours in the

vertical tube furnace and then air cooled. Further heat treatments were then

conducted on two of these samples: one was heat treated at 1515C for 2 hours and

air cooled and the other was heat treated at 10000C for 16 hours and air cooled.

There were five heat treatment experiments that were performed with the

vacuum furnace. With the exception of the alloy 2 sample which was heat treated at

1200C and furnace cooled, these experiments were designed to investigate the effect

that a slow cooling rate had on the microstructures of alloys 2 and 4. The 3 gram









59

RAM samples were used in all of these heat treatments and all of these heat

treatments lasted for a duration of 4 hours.

3.2.2 Short-Term Heat Treatment Experiments

Table 3.3 lists a set of 3 gram RAM samples that were heat treated for very

short times lasting from 2 to 5 minutes. These heat treatments were conducted in

the vertical tube furnace and water quenched. These heat treatments were designed

to investigate the evolution of the high temperature phase equilibria.


3.3 Characterization Techniques

The microstructures of the cast and heat treated samples were investigated

using optical microscopy and transmission electron microscopy (TEM) techniques.

Optical microscopy was used primarily for the macroscopic characterization of the

samples due to its low magnification capabilities. A Nikon microscope and Leica

microscope were used in this study. The TEM was used extensively for the

microscopic analysis of the samples due to its combined diffraction and image

analysis capabilities while complimented with composition analysis using Energy

Dispersive Spectroscopy (EDS) on a submicron microstructural scale. Three

microscopes were used over the course of this investigation: a JEOL 200CX ASTEM

(Analytical Scanning Transmission Electron Microscope) and a JEOL 4000FX TEM

located at the University of Florida, FL and a JEOL 2000FX ASTEM located at the

New York State College of Ceramics at Alfred University, NY.

The analysis of materials using the TEM has been widely employed over the

past 30 years. The methods that were predominantly used in this investigation

included selected area electron diffraction (SAED), convergent beam electron

diffraction (CBED), and amplitude contrast image formation. The SAED analysis was















Table 3.2. Long-term heat treatments used for alloys 2 and 4.


Initial Cooling Furnace
Alloy Condition Temp. Time Method Type

2 RAM 1500C 2 hours Water Quench Tube
2 RAM 1400C 4 hours Water Quench Tube
2 RAM 1400C 4 hours Furnace Cool Vacuum
2 RAM 1300C 4 hours Water Quench Tube
2 RAM 1200C 4 hours Furnace Cool Vacuum
2 RAM 6000C 12 hours Water Quench Tube
2 RAM 400C 12 hours Water Quench Tube
4 AR 1550C 2 hours Air Cool Tube
4 AR 1550C 1 hour Water Quench Tube
4' AR 1515C 2 hours Air Cool Tube
4 RAM 1400C 4 hours Water Quench Tube
4 RAM 1400C 4 hours Furnace Cool Vacuum
4 RAM 1300C 4 hours Water Quench Tube
4 RAM 1300C 4 hours Furnace Cool Vacuum
4 RAM 1200C 4 hours Water Quench Tube
4 RAM 1200C 4 hours Furnace Cool Vacuum
4' AR 1000C 16 hours Air Cool Tube

note: AR As-Received 200 gram arc-melted sample.
RAM Re-Arc Melted 3 gram sample.
solutionized at 1550C for 2 hours and air cooled.























Table 3.3. Short-term heat treatments used for alloys 2 and 4.


Initial Cooling Furnace
Alloy Temp. Time
Condition Method Type
2 RAM 1200C 2 min. Water Quench Tube
2 RAM 1200C 5 min. Water Quench Tube
4 RAM 12000C 2 min. Water Quench Tube
4 RAM 12000C 5 min. Water Quench Tube
4 RAM 1000C 2 min. Water Quench Tube









62

used for routine phase identification and crystallographic information, such as the

analysis of twins, orientation relationships, and stacking faults. The convergent

beam electron diffraction (CBED) technique was used for detailed phase analysis

where symmetry information was required. This method enabled the crystal point

group and the space group of the phase to be determined. These images were

obtained using amplitude contrast methods based on two-beam and multiple-beam

(i.e. near Laue zone axes) conditions. This method was also used to obtain either

bright field or dark field images depending on the beam tilt conditions.

The SAED patterns that were used for phase identification were measured

with a Starrett measuring table. This instrument permitted both linear and angular

measurements to be made with accuracies of +0.025mm for linear measurements and

+0.083 for angular measurements.


3.4 Sample Preparation

3.4.1 Optical Microscopy

The sample preparation that was used for optical microscopy consisted of

standard metallographic methods. The 3 gram RAM samples were mechanically

sectioned into 0.5 to 1.0mm thick samples using a diamond edged cutting wheel.

These samples were then mounted in lin diameter molds using phenolic powder. A

polished surface of 0.03jm was obtained using standard grinding and polishing

methods. The grinding steps were done using 240 to 600 grit SiC paper. The

polishing steps were performed with A1203 powder on appropriate cloths. Finally, the

samples were etched using Kroll's etchant.












u.t.A i rauIUILluaIbuh ithtjCtVIl IVIUI1UVBUIJU


The sample preparation for TEM used the jet-polishing method. This method

consisted of a mechanical preparation step and an electro-chemical polishing step.

The mechanical preparation step involved cutting a wafer from the bulk

sample, cutting a 3mm diameter disc from the wafer, and finally reducing the

thickness of the 3mm disc to about 0.2mm (200um). The starting bulk sample was

either a 3 gram RAM sample or roughly a 1cm x Icm x Icm as-received (AR) sample.

Care was taken in this procedure to select TEM samples that were representative of

the bulk sample. In general, wafers that were cut near the mid-section of the bulk

sample were selected for the next step of obtaining the 3mm disc. It was necessary to

obtain the 3mm disc from the center of the wafer in order to avoid the heat affected

zone (HAV) that was observed by optical microcopy in several heat treated samples.

The method that was employed in obtaining the 3mm disc was different than that

which is customarily used, since the heat treated samples often fractured in a brittle

manner during the preparation when using the hole punch or the ultrasonic disc

cutter. The alternative procedure consisted of mounting the wafer using superglue on

the end of a -2.8mm diameter stainless steel rod. The edges of the wafer were then

smoothed down by mechanical grinding on 600 grit SiC paper until a circular 3mm

disc was obtained. After obtaining the 3mm disc, the thickness was then reduced by

mechanical grinding methods to a thickness of -200 to 300pm using SiC grit paper.

The electro-chemical polishing step used the jet-polishing technique in order to

obtain an electron transparent region near the center of the 3mm disc sample. This

procedure was conducted with a Struers Tenupol jet-polisher and an electrolytic

solution. The electrolyte that gave the best polishing results was based on









64

4%Hf + 10%H2SO4 + 86%Methanol. The following polishing parameters gave

consistently good results: -500C to -400C temperature range, 25 volts DC, and

moderate to low flow rate. After a hole developed in the sample, the sample was

quickly removed from the polisher and rinsed in two successive methanol baths.















CHAPTER 4
EQUILIBRIUM PHASE TRANSFORMATION STUDY


4.1 Introduction

The literature review in Chapter 2 indicated the complexity of the phase

equilibria in the Nb-Ti-Al system. As was previously stated, the foremost reason for

the complexity of this system is that there are a multitude of equilibrium binary

phases and at least two ternary phases which exist in this ternary system. This

combined with the fact that the solidus temperature over most of the ternary system

lies above 1500oC, means that most of the comprehensive phase equilibria studies

have concentrated on one or two temperatures for heat treatments. The most

common temperature has been 1200C. A further complication is that a number of

metastable phases, such as the martensitic a', a2', and a" and the
form in competition with the equilibrium phases in this system. The presence of

these metastable phases can complicate the development of the equilibrium

microstructures and may cause confusion in the equilibrium phase analysis.

Therefore, the purpose of chapter 4 is two-fold: to determine the equilibrium

phases and to describe their development into the equilibrium microstructures in

alloys 2 and 4. The as-cast microstructure of a third alloy (alloy 3) with a composition

of 50Nb-40Ti- 10A (at.%) was investigated in order to compare properties of the BCC

p phase to that of alloys 2 and 4. This background information on the P phase will

then provide the basis for investigation of the metastable phases in chapters 5 and 6.









66

The results of this chapter are divided into three sections: (1) the as-cast, (2)

the long-term heat treatment, and (3) the short-term heat treatment microstructures.

Using these results, the discussion of this chapter is then divided into three main

parts: (1) high temperature p-phase, (2) equilibrium microstructures at the aging

temperatures, and (3) equilibrium phase transformation mechanisms.


4.2 Results

4.2.1 As-Cast Microstructures

The analysis of as-cast microstructures consisted of re-arc melted (RAM)

samples of alloys 2, 3, and 4. The cooling rate associated with these samples was

relatively fast due to their small size which minimized inhomogeneous

microstructures in the RAM samples. However, the RAM sample of alloy 2 showed a

microstructure with an inhomogeneous distribution of precipitates. Therefore, a

RAM sample of alloy 2 was also analyzed which was electromagnetically (EM)

levitated and drop quenched in order to suppress the solid-state precipitation by

rapid solidification.

4.2.1.1 Re-arc Melted

The as-cast microstructures that were observed in the RAM samples of alloys

2, 3, and 4 are shown in Figure 4.1. These microstructures consisted of large primary

grains with a coarse dendritic structure. The size of the primary grains was typically

>300pm (0.3mm). In alloy 2, an inhomogeneous distribution of acicular precipitates

was observed near the grain boundaries and within the interdendritic regions (Figure

4. la). Alloy 4 (Figure 4. Ic) occasionally showed a second phase at the grain

boundaries, however, alloy 3 (Figure 4. lb) showed no additional phases.




















(a)


50pim
















(b)


50pm


Figure 4.1. Optical micrographs showing the as-cast microstructures. (a) alloy 2;
(b) alloy 3; (continued)









68

















(c)


50um























Figure 4.1. (continued) (c) alloy 4.











4.2.1.1.1 The Primary Phase

The primary phase in the RAM samples of alloys 2, 3, and 4 was examined by

TEM and was determined to have the BCC structure, known as the P phase. Selected

area electron diffraction (SAED) patterns showing the [001]p zone axis from the P

matrix in each alloy are shown in Figure 4.2. The patterns show diffraction spots at

the {100} positions for alloys 2 and 4 (Figures 4.2a and 4.2c, respectively); however,

the one for alloy 3 does not show these spots (Figure 4.2b). The {100) spots denote

that the P phase in alloys 2 and 4 has an ordered BCC structure, or B2 (CsC1)

structure, and are referred to as superlattice reflections. The fact that no

superlattice reflections are observed in alloy 3 indicates that the p phase has a

disordered structure.

Micrographs that were obtained using a two-beam condition with a (100)p

superlattice reflection are shown in Figure 4.3. These micrographs show the presence

of large anti-phase domain boundaries (APDBs) in the matrix of alloys 2 and 4. The

APDBs are formed during a disorder to order transition and indicate that ordering

occurred during solid state cooling.

The B2 phase that was observed in alloys 2 and 4 exhibited several diffuse

scattering anomalies. The results showing these anomalies are grouped into three

categories: diffuse streaking, splitting of diffraction spots, and localized diffuse

intensity maxima. A tweed structure that correlated with the diffuse streaking and

spot splitting was also observed.

The diffuse streaking and the splitting of diffraction spots are best observed at

the [001]p zone axis, as shown in Figure 4.2a for alloy 2 and 4.2c for alloy 4. The

streaking was continuous in the <110> directions and intersected both fundamental




















































Figure 4.2. SAED patterns showing the [001] zone axis of the 0 matrix. (a) alloy 2;
(b) alloy 3; (continued)






















































Figure 4.2. (continued) (c) alloy 4.





















(a)


0.2pm

















(b)


0.2fim


Figure 4.3. TEM micrographs showing the APDBs in the B2 matrix. (a) alloy 2;
(b) alloy 4.









73

and superlattice reflections. The splitting of diffraction spots was observed as

satellite reflections that were displaced from the reciprocal lattice position in <100>

directions. The separation of the split reflections increased as the order of reflection

increased. This is demonstrated in Figure 4.4 using a SAED pattern that was tilted

off the [001]p zone axis along the g=(110) reflection. Splitting is observed for

reflections along both the [100] and [010] directions and is greater for the (400) spot

as compared to the (200) spot. Also, notice that for the (220) spot there are four

satellite reflections. These are comprised of two pairs of reflections that are split

along orthogonal [100] and [010] directions.

The diffuse electron scattering is best observed in SAED patterns of the

<110>, and zone axes, as shown in Figure 4.5. The scattering consists of

localized segments of diffuse intensity that extend in <112> or <110> directions and

have a maxima located between the BCC diffraction spots at fractional coordinates.

The diffuse intensity is superimposed on the continuous streaking that is also

observed in the <112> and <110> directions. The fractional coordinates for intensity

maxima in alloy 2 were at 1/s<1T0>,>,<10>s< 2>, %<1 12>, and 1/

positions in the <110>, zone axis (Figure 4.5a) and at %s< TO> and 1/% positions

in the <1ll1> zone axis (Figure 4.5b). The intensity maxima positions observed in

alloy 4 were the same as those observed in alloy 2, with the exception of the Vs<<110>

positions that were not observed at either the <110>, zone axis (Figure 4.5c) or the

zone axis (Figure 4.5d).

The tweed structure consisted of striations that lied parallel to (110} traces of

the ordered p phase, as shown in Figure 4.6. The micrograph was obtained near the

[001], zone axis to show the tweed striations along two orthogonal (110) and (110)





















































Figure 4.4. SAED pattern from the B2 matrix showing the splitting of diffraction
spots. The specimen was tilted away from the [001] zone axis along
g=(110).






















































Figure 4.5. SAED patterns showing the diffuse electron scattering observed in the
B2 matrix of alloy 2 (a and b) and alloy 4 (c and d). (a) [110] zone axis;
(b) [111] zone axis; (continued)




















































Figure 4.5. (continued) (c) [110] zone axis; (d) [111] zone axis.






























O.1 Im


Figure 4.6. TEM micrographs showing the tweed microstructure in the B2 matrix.









78

planes. The image characteristics of the tweed structure was investigated using

two-beam amplitude contrast conditions, and was found to obey the g-R=0 invisibility

criterion (where g is the reflection vector and R is a general description of a

displacement vector) and to depend on the deviation parameter, s. For example, a

two-beam condition using g=(100) caused both (110) and (110) striations to be visible,

but using g=(110) caused only the (110) striation to be visible. The (110) striations

are invisible using g=(110) since g-R=0, assuming R=(1TO). Likewise, the deviation

parameter, s, affected the tweed image, causing it to have a coarse appearance for a

small s magnitude and a fine appearance for a large s magnitude.

4.2.1.1.2 The Precipitates in Alloy 2

The inhomogeneous distribution of precipitates observed in the as-cast

microstructure of alloy 2 were divided into three representative areas: the in-matrix,

the grain boundary, and the interdendritic regions. The primary difference between

the in-matrix and the interdendritic regions was the presence of acicular shaped

precipitates which were observed by optical microscopy, as seen in Figure 4.1.

The in-matrix region consisted of a high number density of very small

precipitates, as can be seen in Figure 4.3b. The precipitates were homogeneously

distributed within the matrix and yet were not affected by the APDBs that had

formed during the p phase disorder/order transition. The analysis of the small

precipitates identified them as being related to the class of o-phases, a close-packed

hexagonal structure [58] which will be discussed further in chapter 5.

A TEM micrograph of the grain boundary region is shown in Figure 4.7. This

region consisted of B2 grains with the y phase, based on the LI, tetragonal structure

with TiAl stoichiometry, distributed along the grain boundaries. Two morphologies









79

were observed for the y grains: a blocky type that formed along the B2 grain

boundaries and a lath type that extended from the blocky grains into the B2 matrix.

The formation of Widmanstatten laths from grain boundary allotriomorphs [80]

resembles this type of microstructure. The laths were observed to have an

orientation relationship with the B2 matrix which was determined to be as follows:

<1<110], l and {11, II{l10}p

and is shown in the SAED pattern in Figure 4.7b. The grain boundary allotriomorphs

contained stacking faults that formed on the (111), planes.

Micrographs that are representative of the interdendritic regions are shown in

Figure 4.8. These regions consisted of B2 matrix with large acicular shaped

precipitates as well as small in-matrix o precipitates. The analysis of the acicular

shaped precipitates showed that they were plates and that they had an orthorhombic

structure (referred to from this point on as plates). Their size was typically observed

to be >10pm in length (1) and <0.5pm in thickness (t), giving them an aspect ratio

>20(l/t). The small o precipitates were observed to be homogeneously distributed

about the plates. However, occasionally large a precipitates were observed in contact

with the plates, as shown in Figure 4.8b. A detailed analysis of the plates will be

covered in chapter 6.

4.2.1.2 EM Levitated and Drop Quenched

The levitated and drop quenched sample of alloy 2 was found to consist of an

acicular microstructure, as shown in Figure 4.9. The optical micrograph in Figure

4.9a showed that the microstructure had a basket weave appearance. The TEM

micrograph in Figure 4.9b revealed that a high number density of lenticular-shaped

plates had formed in the ordered p matrix. The ordered p matrix was also found to




















(a)


0.2pm


















(b)











Figure 4.7. Shows the microstructure of the as-cast sample of alloy 2. (a) TEM
micrograph showing the grain boundary allotriomorphs and
Widmanstatten laths of the y-TiAI phase; (b) SAED pattern showing the
orientation relationship observed between the y laths and B2 matrix,
which was <110], II <1ll> and {IT1(, I {T10p.




















(a)


1.O .m
















(b)


1000A









Figure 4.8. TEM micrographs of the B2 matrix in the as-cast sample of alloy 2.
(a) the small a-related precipitates and lenticular-shaped plates; (b) the
coarse a-related precipitates adjacent to the plate.























-- -"'.>. *,, -, "BS .,'.















(b)












Figure 4.9. Shows the acicular microstructure observed in the EM-levitated and
drop quenched sample of alloy 2. (a) Optical micrograph; (b) TEM
micrograph.










83

have contained APDBs. The analysis of these plates showed that they were the same

as those that were observed in Figure 4.8 for the RAM sample.

4.2.2 Long-Term Heat Treatments

4.2.2.1 Analysis of Alloy 2

The long-term heat treatments used for alloy 2 were shown in Table 3.2. The

1500C heat treatment lasted 2 hours while the 1400C, 1300C, and 1200C heat

treatments were for a duration of 4 hours. All of the heat treatments used a water

quench to cool the samples to room temperature except for the 1200C treatment

which used a furnace cool.

4.2.2.1.1 Optical Microscopy

Micrographs that are representative of the microstructures observed in

samples aged at 13000C, 1400C, and 15000C and subsequently water quenched are

shown in Figure 4.10. Note that the microstructures do not show any evidence of the

prior dendritic structure and that the heat treatments applied were sufficient to

remove chemical inhomogeneities.

The microstructures of the samples aged above 1300oC all showed a matrix

that resembled the acicular microstructure observed in the interdendritic regions of

the as-cast sample. These microstructures showed basket weave morphologies, as

shown in Figures 4. 10a to 4. 10c. The grain boundaries observed in these

microstructures indicated that the matrix consisted of a single phase at high

temperatures and had a grain size of 1 to 2mm. The entire microstructures of 1400C

and 1500oC aged samples consisted of the acicular microstructure. However, the

sample aged at 13000C also showed large blocky-shaped second phase particles that

formed within the matrix and at grain boundaries. A two-phase microstructure was




Full Text

PAGE 1

PHASE TRANSFORMATIONS IN THE CENTRAL PORTION OF '11IE Nb-Ti-Al TERNARY SYSTEM By DAVID TIMOTHY HOELZER A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVE RSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 1996

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ACKNOWLEDGEMENTS I would like to thank my graduate advisor, Dr. Fereshteh Ebrahimi, for all of her input and continued support. I am especially grateful to her and to Dr. Michael Kaufman for the insights and the fruitful discussions we had on this thesis topic. I also would like to thank the other members of my committee: Dr. Ellis Verink Jr., Dr. Robert Dehoff, and Dr. Anna Brajter-Toth for making my defense of this thesis truly a memorable occasion for me. This research work was funded by DARPA under the contract number N00014-88-J-1100 while I was a full time graduate student at the University of Florida. However appreciation for the use of the JEOL 2000FX ASTEM and the darkroom facilities goes to the NYS College of Ceramics at Alfred, NY where I was employed during the completion of this thesis. I would like to thank Pratt and Whitney and in particular Mr Maloney for so willingly making the initial samples in the compositions that we specified. I would also like to thank Mr. Wayne Acree for the microprobe work. A special thank you goes to all of my friends, which includes at the top of the list Dr. Wishy Krishnamoorthy. I really appreciated the housing provided by Wishy while I traveled to UF on thesis related business. Those escapes to Wishys' place and discussions which we had were physical and mental boosts for me. Finally, I would like to express my most sincerest thanks to my wife Amy for her total support and for the sacrifices she made that allowed for me to finish this thesis. I can now make up for lost time with both her and Rachel, my daughter. 11

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TABLE OF CONTENTS ACKNOWLEDGMENTS . . . . . . . . . . . . . . . LIST OF TABLES . . . . . . . . . . . . . . . . Pa ge 11 V LIST OF FIGURES . . . . . . . . . . . . . . . . Vll. ABSTRA C T. C HAPTER . XVl 1 INTRODUCTION 1 2 LITERATURE SURVEY . . . . . . . . . . . . . 5 2.1 Th e T e rnary Phas e Equilibria Studi es . . . . . . . . . 5 2 1.1 Th e Surv ey Studi es . . . . . . . . . . . . 6 2 1.2 Th e Ti Bas e d All oy D e v e lopm e nt Studi es . . . . . . 18 2. 2 Th e Sigm a Phas e . . . . . . . . . . . . . . 22 2 3 Th e Gamma Phas e . . . . . . . . . . . . . . 24 2.4TheB2Phase ............. . ... .. .. .. ... .. 26 2. 5 Th e Omega Phas e . . . . . . . . . . . . . . 32 2. 6 Th e Ortho/H ex Phase s . . . . . . . . . . . . . 39 2. 6.1 Di sor d ere d Structur es . . . . . . . . . . . 41 2.6 2 Ordered Structur es . . . . . . . . . . . . 43 3 EXPERIMENTAL PROCEDURE . . . . . . . . . . 53 3.1 Material . . . . . . . . . . . . . . . . 53 3.2 H e at Tr ea tm e nts . . . . . . . . . . . . . . 56 3.2 1 Long-T e rm H e at Tr eat m e nt Exp e riments. . . . . . . 58 3.2 2 Short-T e rm H e at Tr eat m e nt Experim e nts . . . . . . 59 3. 3 Characterization T ec hniqu es. . . . . . . . . . . . 59 3.4 Sampl e Pr e paration . . . . . . . . . . . . . 62 3. 4.1 Optical Micr osco py . . . . . . . . . . . . 62 3.4 2 Transmission Electron Mi c r osco py . . . . . . . . 63 4 EQUILIBRIUM PHASE TRANSFORMATION STUDY. . . . . . 65 4.1 Introdu ctio n . . . . . . . . . . . . . . . 65 4. 2 R es ult s . . . . . . . . . . . . . . . . . 66 4.2.1 As-Cast Mi c rostructur es . . . . . . . . . . . 66 4.2.2 Long-T e rm H ea t Tr eat m e nts . . . . . . . . . 83 ill

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4.2 3 Short-T er m H eat Tr eat m e nts ..... .. .. ... ....... 108 4. 3 Di scussion . . . . . . . . . . . . . . . . 120 4. 3 .1 Th e f3 Pha se . . . . . . . . . . . . . . 122 4.3.2 Pha se Equilibrium . . . . . . . . . . . . 127 4. 3. 3 Pha se Transf o rmati on M ec hanism s . . . . . . . . 131 5 THE OME GA -RELATED (co -D ) PHASE . . . . . . . . . 137 5. 1 Introdu c tion . . . . . . . . . . . . . . . 1 3 7 5. 2 R es ult s . . . . . . . . . . . . . . . . . 1 38 5.2.1 Stru ct ural Analy s i s . . . . . . . . . . . . 138 5. 2. 2 Effect of Coo ling Rat e . . . . . . . . . . . 154 5.2.3 Effect of Low T e mp e ratur e H ea t Tr eatme nts . . . . . 158 5. 3 Dis c u ss i on . . . . . . . . . . . . . . . . 1 6 5 5. 3.1 Micro st ru ct ural Aspects of th e co -D Pha se . . . . . . 166 5.3.2 Co mpari so n o f the co -Relat e d Phas es in th e Nb-Ti-Al Sy ste m ... 169 5.3.3 Tran s f o rmation and Sit e Occupancies of th e co Pha ses . . . 172 5.3 4 Cryst all og raplu c A spects of th e co-D Phas e Transformation. . 187 6 THE ORTI-I(HEX) PLATES . . . . . . . . . . . . 194 6.1 lntr o du ctio n . . . . . . . . . . . . . . . 194 6.2 Results . . . . . . . . . . . . . . . . . 195 6.2.1 Stru c tural Analy s is . . . . . . . . . . . . 196 6 2. 2 Plate M o rphol ogy . . . . . . . . . . . . 230 6.2 3 Zig-Zag Plat e M o rphol ogy . . . . . . . . . . 233 6.2.4 D efect Stru c tural Analy s i s of the Plat es. . . . . . . 236 6. 3 Dis c ussion o f th e Plat e Transformati on . . . . . . . . 251 6. 3.1 Mart e n s itic Tran s f orma tion . . . . . . . . . . 251 6.3.2 Stru ct ural Analy s i s of the Plat es . . . . . . . . 263 6. 3. 3 C rystallographi c Tr ea tm e nt of the Plat e Transf o rmation . . 292 6 .3 .4 Formation of Plat e Mart e nsit e from th e f3 pha se ...... .. 305 7 S ... .,Y ..... ..... ............ 8 FUTURE WORK ........ . . . . 309 . . . . 313 REFEREN CE S ..... .. ......... .................. 314 BIOGRAPHICAL SKETCH . . . . . . . . . . . . 320 lV

PAGE 5

LIST OF TABLES Pag e 2 .1 Informati o n regarding the pha ses o f th e thr ee binary Nb Ti Nb-Al and Ti-Al phas e diagrams ..... .. ...... .. ............... 7 3.1 Th e nominal compositions of the as-r ece iv e d alloys and the compositions d e t e rmin e d by mi c roprob e analysis of th e re-arc m e lt e d alloys. . . . . . . . . . . . . . . . . . . 54 3.2 Long-t e rm h e at tr ea tm e nts us e d for alloys 2 and 4. . . . . . . 60 3.3 Short-t e rm h e at tr e atm e nts us e d for alloy s 2 and 4 . . . . . . . 61 4.1 Mi c roprob e r es ults of th e ag e d samples of alloy 2 . . . . . . . 87 4 2 S11mmary of t h e phas es id e ntifi e d by TEM in the ag e d sampl e s of all oy 2. . . . . . . . . . . . . . . . . . . 93 4. 3 Summary of the phas es identifi e d by TEM in the aged samples of alloy 4. . . . . . . . . . . . . . . . . . . 109 5.1 Shows the r e lation b et w ee n the possibl e diffraction groups and th e symmetries observed in the Co nv e rg e nt B e am El ec tron Diffraction (CBED) patt e rns at th e [0001] 00 zone axis .................. 141 5 2 Shows th e r e lation b e tw ee n the possibl e diffraction groups and th e symmetries observed in the Co nverg e nt B ea m El ec tron Diffraction (CBED) patt e rn s at th e (1100] 00 zone axis. . . . . . . . . . 143 5. 3 Shows th e crystal point groups that ar e consistent with the diffraction groups obs e rv e d in th e CBED whol e patt e rns. . . . . . . . . 144 5.4 Co mpari so n o f c hara c t e ristics b e tw ee n the co-D, the disorder e d co-Ti, and the ordere d co-B8 2 pha ses from the Nb-Ti-Al t e rnary system. . . . 170 5.5 Pr o p ose d site occ upancy f or the co-D phas e with Ti 3 Al 4 Nb 2 sto i c hiom etry and P6 3 / m c m (193) space group ... ... 6.1 Shows th e r e lation b et w ee n the po ss ibl e diffraction groups and th e sym metri es obse rv e d in the Co nv e rg e nt B ea m El ec tron Diffra ction . . 182 (C BED ) patt er ns of the HCP phase at th e [OOOl]H zone axis. . . ... 201 V

PAGE 6

6.2 Shows th e r e lation b e tw ee n th e possibl e diffraction groups and th e s ymmetri es o bs e rv e d in th e Convergent Beam Electron Diffraction (CBED) patt e rns of th e HCP phas e at th e [1126]H zone axis ......... 204 6 3 Shows the r e lation betw ee n diffraction groups and crystal point groups for the CBED pattern s of th e I-ICP plates. . . . . . . . . . 205 6.4 Shows the r e lation b e tw ee n th e possibl e diffraction groups and th e symmetri es observed in th e Convergent B e am Electron Diffraction (CBED) patt e rns of tl1 e ORTHl phas e at th e l001] 0 and [110] 0 zone axis. . . .. . . . . . . . . . . . . . . . . 218 6.5 Shows the r e lation b e tw ee n the diffraction groups an d crystal point groups for th e CBED patt e rns o f th e ORTH I plat e s.. . . . . . . 220 6.6 Sh o w s the latti ce param e t e rs o f plat es with th e ORTHl structure wh e re 4>1 i s th e angle b e tw ee n th e (OT1) 13 and (020) 02 refl ec tions ....... 223 6 7 Th e imaging co nditions of the APDBs observed in the plat es. . . . . 253 6.8 C al c ulat e d phas e factors a= 2n(g R ), for th e APDB vectors in th e HCP phas e . . . . . . . . . . . . . . . . . 268 6 9 Calculated phase factors a= 2n(g R) for the APDB vectors in th e ORTHl and ORTH2 phas e s. . . . . . . . . . . . . 269 6.10 Th e propos e d atomic site occupancy of th e ORTHl phas e with the Al 2 NbTi stoic hiom e try and th e Cmcm (63) space group ............ 274 6.11 Th e propo se d atomic site occ upancy of th e OR'Til3 phas e with the Al(NbTi) stoichiometry and th e Pmma ( 51) space group. ... .. .. 285 Vl

PAGE 7

LlST OF FIGURES Figure Page 2.1 Shows the binary Ti Al phase diagram. (a) the previously accepted phase diagram; (b) the modified section of the phase diagram. ....... 9 2.2 Shows the 1200C isothermal section of the ternary Nb-Ti Al system determined by Jewett et al.. . . . . . . . . . . . . 12 2 3 Shows the liquidus projection of the ternary Nb Ti-Al system determined by Perepezko et al. . . . . . . . . . . 14 2 4 Shows the 1200C isothermal section of the ternary Nb Ti Al system determ i ned by Das et al. . . . . . . . . . . . . . 16 2 5 Shows the partial sections of the isotherms in the ternary Nb-Ti-Al system determined by Das et al (a) 1200; (b) 1150. . . . . . 17 2.6 Shows the temperature-composition diagram of the Ti 3 Al to Ti 2 A1Nb section by Banerjee et al. . . . . . . . . . . . . . 21 2 7 Shows the five lattice sites in the projection of the unit cell for the cr phase. . . . . . . . . . . . .. . . . . . . 25 2.8 Shows the unit cell of the y-TiAl phase. . . . . . . . . . 27 2. 9 Shows the unit cell of the B2 phase with the atomic site occupancy determined in a Ti 24.5Al-14Nb (at.%) alloy by Banerjee et al. . . . 30 2.10 Shows the {111} 13 plane collapse model of the J3 to ro transformation. The view is normal to the (l lO)p planes. . . . . . . . . . 35 2.11 Shows two rotational variants each one containing three translational variants of the ro phase formed from the J3 to ro transformation. The view is normal to the (110) 13 planes. . . . . . . . . . . 36 2.12 Shows the transformation paths from the J3 phase to the ro-related phases using subgroup and symmetry relations from Bendersky et al. . 40 2.13 Shows the relationship between the crysta l structures of the a 2 Ti 3 Al phase and the 0-TizA].Nb phase. (a) the a 2 -Ti 3 Al phase (P6 3 /mmc space group); (b) the O Ti 2 A1Nb phase (Cmcm space group). . . . . . . 46 Vll

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2.14 Show s th e tr a n s f o rmation p a ths fr o m th e J3 pha se t o th e O ph ase u s in g s ubgroup and s ymm e try r e lati o ns fr o m B e nd e r s ky e t al. . . . . . 5 0 4. 1 Opti c al mi c r o graph s s h o wing th e a s c a s t mi c ro s tructur e s. ( a ) all o y 2 ; (b ) alloy 3 ; (c) all o y 4. . . . . . . . . . . . . . . 67 4 2 SAED pa t t e r ns s h o win g t h e [001] zo n e a xi s o f t h e J3 matri x. (a) a ll o y 2; (b ) all oy 3 ; (c) all o y 4 . . . . . . . . . . . . . . . 70 4 .3 ( a ) all o y 2 ; TEM mi c r og r a ph s s h o wing th e APDB s in th e B2 matrix (b ) all o y 4 . . . . . . . . . . . . . . 72 4 .4 SAED p at t e rn fr o m th e B2 m at rix s h o wing th e s plitting o f diffr act i o n s p o t s Th e s p ec im e n w as tilt e d a w a y from th e (00 l] zon e axi s al o ng g -( 110 ). . . . . . . . . . . . . . . . . . . 7 4 4 5 SAED pa t t e r ns s h o wing t h e diffus e e l ec tr o n sc att e ring o bs e rv e d in th e B2 matrix o f al l o y 2 ( a a nd b ) and all o y 4 (c a nd d ). ( a ) [110] zo n e axis; (b) [111] zo n e axi s; (c) [110] zo n e axi s; ( d ) [111] zo n e axi s . . . . . 75 4.6 TEM mi c r o gr a phs sh o wing th e tw ee d mi c r os tru c tur e in th e B2 matr ix . . 77 4 7 Sh o ws th e mi c r os tru c tur e of th e as c a s t s ampl e o f alloy 2 ( a ) TEM mi c r o graph s h o wing th e grain b o undary all o tri o morphs and Widman s tat te n l a th s o f th e r -TiAl pha se; (b ) SAED patt e rn s h o wing th e o ri e nt a ti o n r e l at i o n s hip o b se rv e d b e tw ee n th e 'Y laths and B2 m atrix, whi c h w as < lIO] r II < lll >p a nd {l T l} r II {TIO} p . . . . . . 80 4 .8 TEM mi c r og r a p hs o f t h e B2 m at rix in t h e as c a s t s ampl e o f all oy 2 (a) t h e s mall ro -r e l ate d pr ec ipi ta t es an d l e nti c ul ar s h a p e d pl a t es; (b ) th e coarse ro re l a t e d pr ec ipit a t e s adja ce n t t o th e pl a t e. . . . . . 81 4.9 Sh o w s th e acic ul ar mi c r os tru ct ur e o b se rv e d in th e EM-l e vitat e d and dr o p qu e n c h e d sa mpl e o f all o y 2. (a) Opti ca l mi c r o graph ; (b ) TEM mi c rograph. . . . . . . . . . . . . . . . . 82 4. 10 O pti c al mi c r o graph s s h o wing th e mi c r os tru c tur es o f th e l o n gt e rm t h e rmall y a ge d s ampl es o f all oy 2 ( a ) 1500 -4br s -WQ ; (b ) 1400 4hr s -WQ ; (c) 1300 -4hr s -WQ ; ( d ) 1200 -4hr s F C . . . . . . . 84 4 11 TEM mi c r o gr a phs sh o wing th e mi c r os tru c tur es o f th e l o ng-t e rm th e rmally ag e d s ampl es of all oy 2 ( a ) 1400 -4brs-WQ ; (b ) 1300 -4br s -WQ . . . . . . . . . . . . . .. 4. 1 2 TE M mi crog r ap h s s h o win g th e y l at h s i n t h e cr + r mi cros tru c tur e of th e t h e rm a ll y age d 1200 -4br s -F C sa mpl e o f al l oy 2 (a) l o ngitudinal 88 vie w ; (b) t r ans v e r se vi e w .. . . . . . . . . . . . . . 9 1 Vlll

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4 13 SAED pa t t e rns o f th e y pha se o b se rv e d in th e 1200 -4hrs-F C s ampl e o f a ll o y 2 ( a ) t h e (110] 1 zo n e axi s; (b ) th e [101] 1 zo n e axi s. . . . . 92 4.14 Optical mi c r o graphs sh o wing th e micro s tru c tur e s o f th e l o ng-t e rm t h e rmally a g e d sa mpl es o f all o y 4 ( a ) 1550 2hr s -A C; (b ) 1515 2hr s -A C; (c) 1400 C4hrs-WQ ; ( d ) 1300 -4hr s -WQ ; (e) 1200 -4br s -WQ ; ( f) 1000 -16hr s A C. . . . . . . . . . 95 4 15 TEM mi c r o graph showing th e APDBs in th e B2 matrix o f th e sampl e fr o m alloy 4 t h a t was h e at tr e at e d at 1550 C and wat e r qu e nch e d . . . 98 4 16 TEM mi c r ogr aphs sh o wing th e diff e r e nt m o rph o l o gi es o f th e r e tain e d phas e ob se rv e d in th e sampl es of alloy 4 that w e r e h e at tr e at e d at 1200 C, l300 C, and 1400 C and wat e r qu e n c h e d ( a) th e irr e gular and c ir c ular m o rph o l o gi e s ; (b ) th e e l o ngat e d morph o logy . . . . . . 100 4 17 TEM mi c r o gr a ph s h o wing th e pr ec ipitat es of th e o rthorhombic phas e t hat f o rm e d o n t h e APDB s and in th e B2 m a trix o f alloy 4 during a ir coo lin g fr o m 15 5 0 C. . . . . . . . . . . . . . . 101 4.18 SAED p a tt e rn s h o wing th e o ri e nt a tion r e lati o n s hip b e tw ee n th e o rth o rh o mbi c pha se and t h e B2 pha se in th e 1550 -2hr s -A C sampl e o f all o y 4 Th e OR was co nsi s t e nt with [001] 0 11 < 110 > 13 and ( 110 ) 0 II {1 T2} 13 103 4.19 TEM mi c r o graph s howing th e o rth o rh o mbi c plat es that form e d in th e r et ain e d B2 pha se of all o y 4 during furna ce co oling fr o m th e h e at tr e atm e n t t e mp e ratur es o f 1200 1 1300 C, and 1400 C. . . . . . 105 4.20 Sh o w s th e mi c r o stru c tur e o b se rv e d in th e 1 000 -16hr s -A C s ampl e o f all o y 4 ( a ) TEM mi c r o gr a ph s h o wing parti c l es o f th e o rth o rh o mbi c pha se o b se rv e d a t th e grain b o undari es of th e cr pha se; (b ) SAED patt e rn s h o wing th e [110] 0 z o n e axis o f th e orthorhombi c phas e . . . 106 4 21 Show s th e mi c r os tru ct ur e o b ser v e d in th e s ampl es of all o y 4 tha t w e r e h ea t tr ea t e d a t 1000 c a nd 1200 C. ( a ) TEM mi c r o graph s howing th e co l o ny o f cr gr ain ~; (b ) SAED p atte rn s h o wing multipl e [OOll a zo n e axes fr o m th e cr gr ain s pr ese n t in th e co l o ny. . . . . . . . . . 107 4 22 TEM mi c r o gr a ph s sh o wing th e mi c r os tru c tur e o b se rv e d in th e 1200 -2min WQ s ampl e of all o y 2 ( a ) th e cr and th ey pha ses a t a l o w magnifi ca ti o n ; (b ) th e cr grain s i ze. . . . . . . . . . . . 111 4 2 3 Opti c al mi c r o gr a phs s h o wing th e mi c r ost ru c tur e obs e rv e d in th e 1200 -2min-WQ s a mpl e o f alloy 4. ( a ) bright fi e ld mi c r o graph ; (b ) dark fi e ld mi c r o graph .. . . . . . . . . . . . . . 113 IX

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4.24 TEM micrographs showing th e microstru c tur e observed in th e 1200 2min-WQ sample of alloy 4 (a) th e thin foil sp ec imen was tilted to the [001] 13 zon e a.xis of th e p particles; (b) th e thin foil specjm e n was tilt e d to th e [OOl]a zone axis of th e cr matrix. . . . . . . . ..... 114 4.25 Optical micrograph showing th e microstructur e observed in the 1000C2min-WQ sample of all o y 4 .. . . . . . . . . . . . . 115 4.26 Sh o ws the mi c r ostr u ct ur e obse rv e d in t h e 1000 -2min WQ sample o f all o y 4. (a) ,.fEM mi c r o graph showing the r eac tion front of a co lony that, partially tran s f o rm e d from the B2 matrix ; (b) SAED pattern showing th e [100) 13 zone axis of th e disorder e d p phas e locat e d betwe e n the cr grains in the co lony.. . . . . . . . . . . . . . . 116 4.27 Shows th e microstructur e observed in th e 1000 -2rnin-WQ sample o f alloy 4. (a) TEM micrograph showing th e transvers e view of the colony s tructur e n ear th e int e rface b et we e n th e colony and the B2 phase ; (b ) SAED patt e rn showing the orientation r e lationship between th e B2 and cr phas es, which was <100>B 2 II [OOlla and {110}B 2 II {llO)a ......... 118 4.28 Shows th e mi c rostructur e observed in th e 1000 -2min-WQ sampl e of alloy 4. (a) TEM micrograph showing th e transverse view of' th e colony n e ar the center of th e colony; (b) CBED patt e rn showing the orientation r e lationship b e tween th e cr and p phases which was < 100> 13 II [lOOJ a and {110} 13 II {1 lO)a. . . . . . . . . . . . 119 4.29 SAED patterns showing tw o additi o nal orientation r e lati o nship s that w e r e observed between the cr and p phas es in th e heat treated samples of alloy 4 (a) < 110> 82 II B 2 II [001J a and {l 10} B 2 II {l lO)a . . . . . . . . . . . . . . . 121 4.30 Shows th e translational v ec tor for th e n ea r es t n e ighbor (NN) and n ext n e arest neighbor (NNN) sit es in th e unit cell of th e B2 phase. Th e atomic site occupancy shows Nb and Ti atoms randomly occupying the la Wyckoff s it e and Al atoms occupying th e lb Wyckoff site. . . . . 126 4 31 Shows th e eq uilibrium phas es that form e d at the aging t e mperatur e s in alloys 2 and 4 (a) alloy 2 ; (b) alloy 4 . . . . . . . . . . 129 5.1 CBED whol e p a tt e rns s howing th e 6mm symmetry observed in th e [0001] zone axis o f th e co-related phas e in alloy 2 (a) long camera l e ngth showing th e z e ro order lau e patt e rn ; (b ) s hort camera l e ngth s h o wing th e f a int FOLZ rings. . . . . . . . . . . . . 139 5.2 CBED whol e patterns showing th e 2mm symmetry observed in th e (1100] zone axis of th e co-related phas e in alloy 2 ( a) long c am e ra l e ngth showing the zero order lau e patt e rn ; (b) short ca mera l e ngth s howing the FOLZ rings. . . . . . . . . . . . . . 142 X

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5.3 C BED wh o l e pa t t e rn o f th e [1100] z on e axis with th e b e am tilt e d s lightly t o s h o w th e bla c k c ro ss in th e kin e m a ti c ally f o rbidd e n ( 0001 ) r e fl ec tion at t h e Bragg co ndi tio n . . . . . . . . . . . 146 5 4 SAED patt e rn showing th e ori e ntati o n r e lationship that wa s o bs e rv e d f o r th e co -r e l a t e d and B2 pha se s Th e OR wa s d e t e rmin e d t o b e (0001] 00 II (11 l] p and ( 1100 ) 00 II ( llO )p ... . .. ... ... .. .. ..... 147 5. 5 Shows th e s t e r e ographi c pr o j ec tion of th e OR r e lationship that wa s o bs e rv e d for th e co r e lat e d and J3 ph ase s in all o y 2 Th e proj ec t i on s h o w s th e [lll] p and [0001] 00 p o l es . . . . . . . . . . . 150 5.6 SAED patt e rn s howing th e OR b e tw ee n th e co r e lat e d and J3 phas e s at th e [1 lO] p z o n e axis Tw o rotational variant s wi.th th e [ 1 100] 00 and [lOTOl w zo n e ax es ar e s up e rimp ose d o n t hi s diffra c ti o n patt e rn . . . 1 5 1 5. 7 Sh o w s th e mi c r o stru c tur e o b se rv e d in th e a s c a s t sampl e o f all o y 2 ( a ) TEM mi c r o graph s h o wing t hr ee plat e s and a co ars e co -r e lat e d grain o b se rv e d in th e B2 matri x; (b ) SAED patt e rn s h o wing th e ori e ntati o n r e lati o nship o b se rv e d f o r th e thr ee pl a t es, th e co -r e lat e d grain and th e B2 pha se a t t h e (11 l] p zo n e axis. . . . . . . . . . . . 1 5 2 5 8 Sh o ws th e c al c ulat e d diffra c ti o n patt e rns of th e thr ee plat es and th e co ars e cor e la te d grain at th e [111] zo n e axis of th e B2 matrix ( a ) th e co mp o sit e pa t t e rn ; (b ) th e [111] zo n e axi s o f th e B2 matrix ; (c) th e [OOOl l w zo n e axi s o f th e co -r e lat e d grain ; ( d ) th e [110 ] z o n e axi s o f th e o rthorh o mbi c plat e 1 ; (e) th e [110] zo n e axi s o f th e o rthorhombi c plat e 2 ; (f) th e (110] zon e axis o f th e o rthorh o mbi c plat e 3 . . . . . . 1 5 3 5 9 O pti c al mi cro graph s h o win g th e mi c r os tru c tur e o b se rv e d in th e 1400 4hr s -F C s ampl e o f all o y 2 . . . . . . . . . . . . . 1 5 6 5 10 TEM mi crog r a phs o f th e 1400 -4lu~ s -F C s ampl e o f all o y 2 ( a ) th e in-matrix r e gi o n co n s i s ting o f t h e ro -r e lat e d pha se and plat es; (b ) th e pri o r grain b o undary r e gion co nsi s ting o f th e cr and y pha ses . . . . 157 5.11 TEM mi c r og raphs o f th e 400 -12hr s -WQ sampl e o f alloy 2 ( a ) s h o w s th e pr ec ipit a t es of th e co -r e lat e d ph ase o bs e rv e d in th e B2 m a trix ; (b ) s h o w s th e APDBs o bs e rv e d in th e B2 matrix. . . . . . . .... 159 5.12 TEM mi c r ogr aphs o f th e 600 -12hr s -WQ s ampl e of alloy 2. ( a ) sh o w s th e fin e co -r e lat e d domains that form e d from th e B2 phas e; (b ) show s th e c oar se cor e lat e d d o main s that f o rm e d at t h e pri o r B2 gr a in b o undari es. . . . . . . . . . . . . . . . . . 160 5 .13 SAED patt e rn s fr o m th e 600 -12hrs WQ sampl e o f alloy 2 that sh o w th e diffr ac ti o n patt e rn s o f th e form e r B2 matrix. ( a) [lll]B 2 z on e axis ; (b ) [110] B 2 zo n e axis ; (c) [100] B 2 zo n e axi s; ( d ) [112] B 2 z o n e axis. . . .. 1 6 2 Xl

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5 14 TEM micrograph of th e 600C-12hrs-WQ sample of alloy 2 that shows th e APDB s obse rv e d in a coarse r ota tional domain of the co-related phase. The g = ( 1120 ) r e fl ectio n was u se d to show the APDBs in the coarse domain. . . . . . . . . . . . . . . . . 164 5.15 Shows the ato mi c s it e occupanc i es of the disordered P phase and the disordered ro phase for Ti.. . . . . . . . . . . . . . 17 4 5. 16 Sh ows the atomic site occupanc i es of the or d ered P (B2) phas e and the roB8 2 phase from tl1e r es ult s of B e nd ers ky et al. l 37]. . . . . . . 176 5. 17 The ( 111 ) projection o f the B2 phase with the atomic site occupancy s h o wing Nb and Ti atoms occ upying th e la Wy c k o ff site and Al atoms occ upying the lb Wyckoff site. (a) shows the plan es at z = O.OA 0 095.A and 0.189A; (b) shows the planes at z = 0.284A 0.379A, and 0 .473A. Th e dashed lin es denote the unit ce ll of the ro D phase. ..... 179 5.18 Th e ( 0001 ) projection of the roD phas e whi c h i s bas e d on the P6 3 / m c m space group and Al 4 Ti 3 Nb 2 stoichiometry. (a) s h o w s the s ingl e lay er at z = O.O A and double lay ers at z = 1/4c ; (b) shows th e single lay e r at z = 1 /2c and double lay ers at z = 3/4c. . . . . . . . . . . 184 5. 19 Shows the atomic site occupancies of the B2 pha se and the ro-D phase. . 186 5.20 Shows the transformation paths from the p phase t o the roD phase d escri b ed by subgro up/ sy mm etry relations. . . . . . . . . 189 6.1 Sh o w s a thick plate with the H CP str u cture observed in the 1300 C aged sa mpl e. (a) TEM mi crograp h o f th e H C P plat e; (b) SAED pattern o f the orie ntati on r e lationship o b se rv ed between th e HCP plat e and th e p phase, which wa s [OOOl]a II [011] 13 and ( lTOO )H I I (21 1 ) 13 197 6.2 Th e stereographic proj ect ion of the o ri e ntati o n relationship betw een the HCP phase and the p phase which s h o w s the [OOOl]g II [Oll] p poles.. . 199 6.3 CBE D patterns showing the wh o l e pattern s ymm etry of' the [0001] 8 zo n e axis observed for the thick H C P plates. (a) a larg e ca m era consta nt ; (b) a s mall camera constant. . . . . . . . . . . 200 6.4 CBE D patterns s h owing the whole pattern s ymm etry of the [1126J a zo n e axis observed for the thick H C P plates. (a) a larg e ca m e ra co n stant; (b) a s mall ca m e ra co n stant . . . . . . . . . . 203 6.5 CBE D pattern sho wing the diffuse bla c k c ros s in the ( 0001 ) disc observed in the [1120]H zone axis of the H C P plates. ..... .. ... 206 Xll

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6.6 Shows the boundary observed in a thick HCP plate. (a) TEM micrograph; (b) SAED pattern of the two regions separated by the boundary which was consistent with the orientation relationship observed between the HCP plates and the J3 phase . . . . . . . 208 6. 7 CBED patterns of the two sides (A and Bin Figure 6.6). (a) shows the 6mm symmetry that was consistent with the [OOOl]H zone axis of the HCP phase; (b) shows the 3m symmetry of possibly a different phase .... 209 6.8 TEM micrographs of a thick HCP plate. (a) shows the plate inclined, relative to the beam, at the [TIT]p zone axes; (b) shows the plate edge-on after tilting ~11 along the g = (Oll)p II (OOOl)n reflections ...... 211 6.9 Shows a medium thick plate with the ORTHl structure. (a) TEM micrograph; (b) SAED pattern of the orientation relationship between the ORTHl plate and the J3 phase, which was [001] 01 II [011] 13 and (110) 01 II (211) 13 213 6.10 The stereographic projection of the orientation relationship between the ORTHl phase and the J3 phase which shows the [001] 01 II [011] 13 poles. . . . . . . . . . . . . . . . . . . 215 6.11 CBED pattern showing the whole pattern symmetry of the [001] 0 1 zone axis observed for the ORTHl plates. . . . . . . . . . .. 216 6.12 The structural analysis of the medium thick ORTHl plates. (a) CBED pattern showing the 2mm whole pattern symmetry of the [110] 0 1 zone axis; (b) SAED pattern obtained by tilting the thin foil specimen from the [110] 0 1 zone axis along the g = (001) reflection. . . . . . .. 219 6.13 Shows a thin plate with the ORTH2 structure. (a) TEM micrograph; (b) SAED pattern showing the orientation relationship between the ORTH2 plate and the J3 phase at the [011] 13 zone axis ... ......... 225 6.14 Enlarged SAED pattern showing the diffuse streaks and diffraction spots observed for the ORTH2 plates. The diffraction pattern shows the [Oll]p zone axis for the J3 phase. . . . . . . . . . . 226 6.15 Shows a thin plate with the ORTH3 structure. (a) TEM micrograph; (b) SAED pattern showing the orientation relationship between the ORTH3 plate and the J3 phase at the [Oll]p zone axis. . . . . . .. 228 6.16 Enlarged SAED pattern showing the missing diffraction spots for the ORTH3 plates. The diffraction pattern shows the [011] 13 zone axis for the J3 phase. . . . . . . . . . . . . . . . . . 229 6.17 TEM micrographs showing the plate morphology that was determined from the images of plates at the three orthogonal directions. (a) the [001] 0 zone axis; (b) the [110] 0 zone axis; (c) the [110] 0 zone axis ....... 231 XJJJ

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6.18 TEM micrograph showing plat es that w e r e partially e nclos e d within the thin foil s p ec imen of th e as-cast RAM sample of alloy 2.. . . . . 234 6.19 TEM mi c rographs showing the zig-zag morphol o gy of' th e plat es. (a) bright field mi c rograph ; (b) dark fi e ld micrograph.. . . . . . . 235 6.20 Shows two plates connected tog e th e r along a twin boundary (a) TEM micrograph ; (b) SAED patt e rn at th e [I1T] 13 zone axis of th e f3 pha se showing th e twin relationship betw ee n plate 1 and plat e 2 . . . . . 237 6.21 TEM micrographs showing the stacking faults observed in the plat es at th e [110] 0 or [112"0Jn II [ T1T]B 2 zone ax es (a) shows that th e stacking faults w e r e invisibl e using g = (002) 0 or ( 0002)fl ; (b) shows that the stacking faults w e r e visibl e using g = (2T1) 0 or (2201)H .......... 239 6 22 TEM dark fi e ld micrographs showing th e APDBs observed in th e HCP plat es. (a) g = (I2TO)H and ( 100) 13 ; (b) g = (2)a; (c) g = (1120)H ...... 241 6.23 TEM mi crogr aphs s h o wing the co lumnar s hap e d APDBs observed in the HCP plat es at the [1120] H II [111] 13 zone axes. ( a) bright fi e ld mi cro graph formed with g = ( 2200)H ; (b) dark fi e ld micrograph form e d with g = ( lTOl) g ... . .. . ............. ........ 244 6.24 TEM micrographs showing the e quiax e d morphology of th e APDBs observed in the HCP plat es at the [2 1 1 O ] H zone axis. (a) dark fi e ld micrograph formed with g = ( OlTO )H; (b) dark fi e ld micrograph form e d with g = (Olll)H. . . . . . . . . . . . . . . . 245 6 25 TEM dark fi e ld micrograph s showing th e APDBs observed in the ORTHl plat es (a) g = (200) 0 and (100) 13 ; (b) g = ( 130) 0 ; (c) g = (130) 0 247 6.26 TEM dark fi e ld micrograph showing th e columnar shaped APDBs observed in the ORTHl plat es at th e [110] 0 II [011] 13 zone axes using g = ( 111) 0 .. . . . . . . . . . . . . . . . . 249 6.27 TEM dark fi e ld mi c rograph s h o wing the e quiax e d morphology of th e APDBs o bs e rv ed in th e OR,..fHI plat es at th e [110] 0 zone axis using g = ( 111 ) 0 ........ . ... .... .. .. ... . .. . ..... 250 6 28 TEM dark fi e ld micrograph showing th e columnar shaped APDBs observed in the thin ORTH2 plates using the diffus e s tr e ak. Th e spec im e n wa s tilted from the [011 ] 13 zone axis o f th e f3 phas e. . . . . 252 6.29 Sh o ws th e uniaxial stress state of th e plat e and th e three principal axes of strain which are A 1 )... 2 and A 3 . . . . . 260 XIV

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6.30 Show s th e unit ce lls o f the ORTHl and HCP phas es superimposed upon e a c h ot h er t o illtlStrat e how th e distortions of th e orthorhombic phas e ar e r e lat e d to th e h e xagonal p h ase. Th e drawing shows the [ 001 ] 0 dir ectio n of th e ORTHl unit cell and the [OOO I] H direction of th e HCP 11ni t ce ll. . . . . . . . . . . . . . . . . 264 6.31. Shows th e propos e d atomic site occupancy of the ORTHI phas e bas e d on th e C m c m space gr o up with Al atoms occupying the 8g Ti atoms occupying the 4cl and Nb atoms o cc upying the 4c2 W yckoff positions. . 275 6.32 Shows th e calc ulat e d C BED patt e rns of th e o rthorhombic structures based o n two possible atomic site occupancies. (a) th e propos e d ORTHl pl1 ase with th e sto i c hiom e try of A1 2 TiNb ; (b) the O Ti 2 A1Nb pha se. . . . . . . . . . . . . . . . . . . 277 6.33. C BED pattern showing the whol e patt e rn symmetry of the O Ti 2 A1Nb phas e o b se rv e d in alloy 1 .. . . . . . . . . . . . . . 279 6.34. Sh o w s the APDB v ec tor s in th e (011) 13 plan es o f th e B2 phas e and in the ( 001 ) 0 plan es of th e ORTHl phas e. . . . . . . . . . . . 280 6. 35. Sh o ws th e pr o pos e d sit e occ upan cy o f th e ORTH3 phas e bas e d on the Pmma space group with Nb and Ti atoms occupying the 2 e Wyckoff sites an d Al atoms occupying th e 2f W yckoff sit es.. . . . . . . 284 6. 36. Sh o ws th e calc ulat e d SAED pattern o f the o ri e ntation r e lati o nship b e twe e n the ORTH3 plat e and the B2 phase. Th e OR showed th e (100] 04 and (011] 13 zone axes to b e paralle l and th e (011) 04 and (211) 13 planes to be parallel. . . . . . . . . . . . . . . 286 6.37. Shows the two transformation paths that l ed to th e formati o n of the ORTIIl plat es (path 1 ) and the H C P plat es (path 2 ) from the p phas e . . 294 6. 38. Sh o ws th e unit ce lls of th e diff e r e nt s tru c tur es that l e d t o th e formation of the ORTHl phas e from the B2 phas e. ( a) th e B2 (Pm3m ) structure; (b) the orth o rh o mbi c (C mmm ) structure; (c) the ORTH2 ( Pmma) s tru c tur e; ( d) th e or d e r e d ORTHl (Cmcm) structure.. . . . 297 6. 39. Shows th e unit cells of th e diff e r e nt s tructur es that led to th e formation of the HCP phas e from the p phas e. (a) the disord e r e d p (lm3m ) structure; (b) th e disord e r e d orthorhombic (Fromm) structure; (c) th e dis o rd e r e d orthorh o mbi c (Cmcm) and th e disorder e d HCP ( P6 3 / mm c) str u c tur es; (d) th e ordered HCP (P6 3 /mmc) structure. . . . 303 xv

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Abstra c t of Diss e rtation Pr e s e nt e d to the Graduat e S c hool of th e Univ e rsity of Florida in Partial Fu1fillment of th e R e quirem e nts for th e D e gr ee of Doctor of Philosophy PHASE TRANSFORMATIONS IN THE C ENTRAL PORTION OF THE Nb-Ti-Al TERNARY SYSTEM By David Tim o thy Ho e lz e r D ece mb e r 1996 Chairperson: Dr. F e r e shteh Ebrahimi Major D e partm e nt: Mat e rials Sci e nc e and Engin ee ring Int e rm e talli cs in the Nb-Ti-Al ternary system hav e be e n considered for high temperature a e rospace applications and th e ir d e v e lopm e nt requir es a thorough und e rstanding of phase equilibria and phase transformations. In this study transmission e l ec tron micro sco py (TEM) was primarily used to inv es tigat e tl1 e phas e e quilibria and pha se transformati o ns o f two alloys with co mpositions of 27Nb-33Ti40Al ( alloy 2) and 42Nb-28Ti-30Al (alloy 4). Small arc melted samples wer e thermally ag e d at temperatures b e tw ee n 400 C and 1550 C for four to sixteen h o ur s (long-term), and at 1000 C and 1200 C for two to fiv e minutes (short-term), follow e d by e ith e r water qu e nching air cooling, or furnac e cooling. The equilibri11m phas e study showed that both alloys solidifi e d as th e f3 phas e, whi c h b eco m es ordered to th e B2 phas e during solid s tat e cooling. Th e cr-Nb 2 Al phase pr ec ipitat e d from the p phas e slightly below 1400 C in alloy 2 and 1550C in alloy 4. Th e cr phas e form e d as isolated grains above 1300 C in both alloys. Colonies o f cr grains formed b e low ::: 1300 C in all oy 4. A e ut ecto id transformati o n from p to cr + yXVI

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TiAl o c curred at 1200 C in alloy 2. A discontinuous transformation from B2 to cr + P occ urr e d a t 1000 C in all o y 4. A m e ta s t a bl e ro -D pha se form e d by th e co llap se o t {111} 13 pl a n es and c h e mi c al o rd e ring from th e B2 phas e in alloy 2 during slow cooling Th e ro D phas e c on s ist e d o f th e Al 4 Ti 3 Nb 2 s t o i c hi o m e try and P6 3 / m c m s pa ce gr o up. A prop o s e d mod e ] s h o w e d aluminum and niobium on singl e lay e r s, and titanium and alt1minum o n d o ubl e l a y e r s Th e transf o rmati o n path was d esc rib e d using subgroup and symm e try r e lati o n s a s: Pm "3" m(B2 ) P 3 1m (co '" ) P6 3 / m c m (ro -D ) A mart e nsiti c transformation of th e p phas e to plat e s o ccurr e d during fa s t coo ling in alloy 2 Th e o bs e rv e d habit plan e of th e plat es agr ee d with that c al c ulat e d u si n g t h e invarian t lin e th eo ry Th e p co mp os iti o n aff ec t e d th e f o rmation o f s tru c turally r e lat e d o rthorl1ombi c (Pmma and C mcm ) and H C P (P6 3 /mm c) pla tes. Th e prop ose d s it e occ upan c y s h o w e d th e Al 2 TiNb s toi c hiom e try for th e C m c m ph ase. Analy s i s o f domain s tru c tur es s ta c king faults and e l ec tr o n diffr ac tion s ugg e s te d tw o p o ssibl e tran sformati o n paths : lm3m ( P ) C m c m(di so rd e r e d ) P6 3 /mm c( di so rd e r e d ) P6 3 /mm c (D0 1 9 ) f o r H C P plat es and Pm 3 m(B2 ) Pmm a C mcm ( ord e r e d ) f o r o rth o rhombi c pl a t es. XVll

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CHAPTER 1 INTRODUCTION The te rn ary Nb-Ti-Al s y ste m is r eco gnized as a t ec hnologically important system. Alloys bas e d on this system hav e found many applications, es p ec ially in the aerospace industry Th e titani11m-based alloys hav e long been studied and us e d in aeronautical applications b eca us e o f their low d e nsiti es and high strengths [1 ,2 ]. Th e r e is currently the need for high e r temperature and high e r str e ngth materials to improve the performance of gas turbine e ngin es and str u c tural airfram e components of modern aircraft. This will require diff e r e nt mat e rials than the current ni c k e based s up era ll oys and co nv e ntional titanium-based alloys. Th e nick e l-bas e d alloys are h eavy and hav e reached practi cal limit s impos e d by operating temperatures th at are ::::; 85% of their melting point [3 4] The titanium-based alloys do n ot poss ess s uffi cie nt m ec hani ca l prop e rti es, such as c r ee p and oxidation resistanc e, at high temperatures [5]. Thus, the current n ee d t o improv e the prop e rti es of these mat e rials ha s l e d to the developm e nt of int e rmetallic titanium-al11minide s (5-8] Additions o f niobium t o these titanium-aluminid e s mak e these alloys eve n m ore attract iv e by impr oving th e m ec h a ni cal prop e rti es. H o w eve r th e mat e rials n ee d e d for the n ext generation o f high p e rformanc e ga s turbine e ngines etc. still need further improv ements in low e r d ensity, b etter mechanical pr o p e rti es, and high er o p e rating temperatures. To m eet these n eeds, research is b e ing conducted on refractory-based alloys, ceramics, co mposit e s and int e rm et allics. 1

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2 Ordered int e rmetallic co mpounds bas e d on r e fractory m e tals such as niobium hav e b ee n identifi e d as pot e ntial mat e rials that may me e t th e high-temperatur e r eq uir e m e nts of advanced turbin e e ngin es [9]. Th e attractive prop e rti e s of r e fract o ry bas e d int e rmetalli cs include co mbinations of high m e l ting temp e rature low e r density high stiffn ess, and good c r ee p /s tr e ngth r es istan ce. How e v e r monolithi c or d ere d intermetallics typically show poor m ec hani c al prop e rti es at low temperatures, o f whi c h th e low fra ct ur e t o ughness is th e most serious problem. Th e curre nt trend to overcome th ese problems has be e n thr o ugh th e use of composites and alloy d e v e lopm e nt based on tw o -phas e and multi-phase systems in co rporating Nb-ba se d int e rm etallics and du ct il e second phas es. Th e Nb-Ti-Al system shows potential phas e r e lationships b e tw ee n th e int e rm e tallic cr Nb 2 Al phas e, which has a high m e lting p o int of 2060 ; th e B CC J3 o r B2 phas es, which hav e e xt e nsiv e ternary co mp os iti o n rang es; and other t ec hnologically important int e rm e tallics such as th o a 2 -1 :ia Al y -TiAl O-Ti 2 NbAl and 11-(Ti-Nb ) Al 3 phases [10-13] Thus th e d eve l op m e nt of th ese alloys f o r the impr ove m e nt of prop e rti es r eq uir es a thorough und e rstan
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3 syste mati c studies using larg e numb e rs of all o ys with diff e rent compositions. Thu s, the systematic s tudi es in th e past hav e concentrated on determining the pha se e quilibria at just one or two t e mp e ratur es, with th e most common t e mp e ratur e at 1200 C. Surprisingly the phas e e quilibria and phas e transformations in the ce ntral portion of this system hav e not b ee n inv es tigat e d v e ry thoroughly in th e pa s t. Thi s ce ntral portion contains t e rnary solubility e xt e nsions of the binary cr-NbzA-1 and y-TiAl pha ses and th e ternary ~/B2 phas es, which have attractive prop e rti es such as high m e lting p o int l o w density and r easo nabl e oxidation r es istanc e Th e r e for e, it was the purpos e o f this s tudy to provid e ba s i c r ese ar c h o n th e phas e relati o nship s in th e ce ntral p o rti o n o f the ternary Nb Ti-Al system. A r e vi e w of pr e vious lit e ratur e on th e phas e e quilibria in th e ternary Nb Ti-Al syste m is giv e n in C hapt e r 2. From this lit e ratur e s urv ey and pr e liminary e xp e rim e ntal r es uJt s, two alloys w e r e s e l ecte d bas e d on two phas e microstru ct ur cs that co ntain e d th e cr phas e and e ith e r the B2 or 'Y phas es A third alloy wa s also inv es tigat e d in o rd e r t o study th e influ e n ce of al11minum on th e ord e ring in th e ~ phas e to th e B2 pha se that has b ee n shown t o occur in this ternary system. Both l o ng te rm a nd s hort t erm h e at tr e atm e nt s at high t e mp e ratur es w e r e e mpl o y e d t o s tudy t h e high temperature phas e e quilibria and th e ir e volution. Th e stability aspect of the phas e with r e gard to metastabl e phas e f o rmations was inv es tigat e d using diff e r e nt coo ling rat es in th e high t e mp e ratur e h eat tr ea tm e nt s. Transmission e l ect r o n mi crosco py (TEM) utilizing th e imaging se l ecte d ar e a e l ec tron diffra c tion (SAED) and co nv e rg e nt b ea m e l ec tron diffra c tion (C BED) capabilities was selected a s th e primary analytical t ec hniqu e to id e ntify and study the phases. Th e details

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4 co n ce rning th e all oy c ompositions h ea t tr e atm e nts analytical t e chniqu e s and s p ec im e n pr e par a ti o n ar e pr ese nt e d in C hapt e r 3 Sin ce th e ce ntral p o rtion of this t e rnar y sy s t e m i s co mpli c at e d an analy s i s of t h e r es ult s o f th e eq uilibrium pha ses ar e pr ese nt e d in C hapt e r 4 Th e e quilibrium pha ses pr e s e nt at hi g h t e mp e ratur es and t h e f o rmati o n o f th ese phas es ar e d esc rib e d fr o m th e analy s i s o f th e l ong t e rm and s hort t e rm h e at i--r e atm e nts Th e influ e n ce of c o o ling rat e on th e f o rmation of m e ta s tabl e phas es fr o m th e high te mp e ratur e f3 pha se i s al s o intr o du ce d in this c hap te r. F o llowing th e analy s i s o f th e e quilibrium ph ases, d e tail e d s tudi e s of th e m e tastabl e pha se formations ar e pr e s e nt e d in C h a pt e r s 5 and 6 Chapt e r 5 co v e rs th e m e tastabl e co r e l at e d pha se that forms from th e f3 pha se in a ll o y 2. Th e anal ys is o f th e s tru c tur e and th e e ff ects o f co mp os i t i o n coo ling r a t e, an d l o w te mp e ratur e h e at tr e atm e nt s ar e co v e r e d in this c hapt e r. Fr o m th ese r es ul ts, th e pr o p ose d at o mi c s it e occ upan c y of t hi s co r e l at e d phas e and a d esc rip t i o n o f t h e f3 t o co ph ase transformati o n u s ing subgroup and symm e try r e lati o ns i s gi ve n C hapt e r 6 analyz es th e m e tastabl e p l at e s that f o rm e d from th e f3 pha se in all oy 2 Th e analy s i s o f th e c ry s tal s tru c tur es and d e f ec t s tru c tur es of th e s e pla tes i s co v e r e d. Th e influ e n ce of th e h e at tr e atm e nt and f3 c omp o sition on th e s tru ct ur e and f o rmati o n of th e pl a t es is analy ze d Th e t ransformation o f th e f3 pha se t o pl a t es is s h o wn t o b e co n s i ste n t with mari e nsi t i c tr a nsf o rma t i o ns using th e invariant lin e th eo r y. Finally th e st ru c tur e o f th e plat es i s d esc rib e d using s ubgr o up and sy mm et ry r e l at i o n s t o s h o w th a t two diff e r e n t transiti o n path s l e ad t o diff e r e nt plat e s tru c tur es

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C HAPTER2 LITERATURE SURVEY In this c hapt er, an overview of th e publish e d lit e ratur e on the phas e eq uilibria and ph ase transformations in the ter n ary Nb Ti-Al syste m is pr ese nt e d Du e to tl1e comp l exity of this sys t e m this overvie w will fo c us o nly o n those phases that were o bs e rv e d in the alloys inv es tigat e d in this study Th e r e for e, this chapter i s divid e d into six sections: the t e rnary phase eq uilibria st udi es, the sigma ( cr ) phase th e ga mma (y) phas e, the B2 phas e, th e o m ega (ro) phas e, and th e ortho(hex) pha ses. Th e o rth(h ex) d es ignation is us e d in co njun ct i o n with th e two c los e ly r e lat e d orthorhombic and h exago nal c l ose packed (H C P ) str u ct ur es. Th e first section o n the ternary phase eq uilibria s tudi es d escr ib es th e vari ous d e v e lopm e nts in th e overall t e rnary pha se diagram. Thi s sectio n is divid ed int o two s ub sect i ons ba se d o n the r es ul ts fr o m th e survey st udi es a nd t h e Ti-ba se d alloy development stu di es. Th e r e maining sectio n s o n tl1e cr y, B2 o m ega (ro), and o rth(h ex) phas es d esc rib es the r ese arch p e rtin e nt to eac h o f these specific phases. Th e sectio ns d ea ling with th e cr y, and B2 phas es are r e l evant to th e e quilibrium pha se s tudy in C hapt e r 4 ; the ro pha se is relevant to th e co -r e lat e d pha se transformation study in C hapt e r 5 ; and th e orth(hex) pha ses are re l evant to th e plate transformation st udy in C h a pt er 6 2.1 Th e T er nary Phas e Equilibria Studi es Th e phase eq uilibria of the ternary Nb-Ti-Al syste m ha s been th e s ubj ect o f many inv esti gati o n s s in ce the ea rly 1950 s. How eve r m ost of th ese studies can be 5

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6 gr o up e d int o two main c at e g o ri es : tho se that s urv e y e d th e phas e e quilibria o f alloy s with compositions c ov e ring larg e r e gi o ns o f th e t e rnary phas e diagram at dis c r e t e t e mp e ratur e rang es, and tho se that co n ce ntrat e d in-d e pth on just a f e w c omp o siti o ns f o r th e c omm e r c ial d e v e lopm e nt of Ti-bas e d alloys o v e r larg e t e mp e ratur e rang e s. Th e s urv e y s tudi es inv e stigat e d th e phas e e quilibria o f alloys at a limit e d numb e r o f te mp e ratur e s whi c h was usually only o n e t e mp e ratur e and oft e n 1200 C Th e Ti-b ase d all o y s tudi es m os tl y f oc u se d o n Ti-ri c h o r Ti 3 Al + Nb c omposition s with co n s tant 25at. % Al. Th e r ece nt t r e nd in th e d e v e l o pm e nt o f th ese Ti-ba se d all o y s ha s b ee n in th e t e rn a ry r e gion b e tw ee n Ti 3 Al and TiAl. Th e s e ar e binary phas e s with te rnar y additi o ns s u c h as Nb 2 1.1 Th e Surv e y Studi e s M os t o f th e phas e s that hav e b ee n o b se rv e d in th e t e rnary Nb-Ti-Al sy s t e m ar e binary pha ses that show e d larg e t e rnary so lubility rang e s. Th e r e hav e b ee n many s tudi e s in th e past that hav e d e t e rmin e d th e binary phas e s of th e Nb-Ti Nb-Al and Ti-Al s y s t e m s. Th e s e st udi es a nd th e bin a ry pha se diagram s o f th e Nb-Ti Nb-Al a nd Ti-Al s y s t e m s d eve l o p e d from t h e m hav e b ee n co mpil e d in tw o main r e f e r e n ces [10 11]. Th e r es ul ts o f th ese co mpil e d s tudi es hav e s h o wn that th e r e ar e at l eas t te n e quilibrium binary phas es A summary o f th e important informati o n ab o ut th ese te n pha ses, s u c h a s th e phas e n o tati o n th e c ry s tal stru ct ur e, th e p o int gr o up th e r eac t io n typ e, and th e r e acti o n t e mp e ratur e ar e s h o wn in Tabl e 2 1 Th e r e c urr e ntly exist s so m e un ce rtrunty c on ce rnin g th e phas e e quilibria in th e Ti-Al ph ase diagr a m This 1m ce rtainty i s mainly b et w ee n th e y -1'i.Al and 11-TiA1 3 pha ses and i s du e t o th e formati o n o f long rang e p e ri o di c structur es in this part o f th e b inary sys t e m th at c an c ompli ca t e th e d e t e rminati o n o f t h e e quilibrium pha ses

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Tabl e 2.1. Inf ormation regarding the phas es of th e thr ee binary Nb-Ti Nb-Al and Ti-Al phas e diagrams [10 12]. Pha se Stoi c hiom etry C rystal Spac e Reaction T e mperatur e No tat i o n Stru c ture Group Typ e (C) p Nb B CC Im3m L .. p 2467 p Ti B CC Im3m L .. p 1670 ex, Ti HCP P6 3 / rnm c ex, .. p 882 e Al FC C Fm3m L .. p 660 P1 Nb 3 Al Cubic Pm3n L + p .. pl 2060 cr Nb 2 Al T e tragonal P4 2 / mnm L + P1 .. cr 1940 11 NbAl 3 B CT 14/rnmm L .. ,, 1680 CX.2 TiAl 3 H C P P6 3 / mm c ex, CX. 2 1180 y TiAl Tetragonal P4/rnrnm L + ex. y 1450 8 --LPS -L + r .. P 1 1380 E TiAl 2 BCT 14 1 / amd y + 0 -,g 1240 ,, TiAl 3 BCT 14 /rnm m L + 8 .. ,, 1 350 LPS is a l ong-period super lat tice str u cture; P-Ti a nd P-Nb are i somo rphous ac r oss t h e binary Nb-Ti phas e diagram; and 11-NbA1 3 and 11-TiA1 3 are isomorphous across th e t er nary phas e diagram for constant Al.

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8 [10 14]. The binary Ti-Al phase diagram originally dev e loped by Murray [IO] has recently b ee n modifi e d on the Ti-rich side from the results by Valencia et al. [15] in 1987 and by Mc C ullough et al. [16) in 1989. Thi s modification involved th e phas e boundaries b e tw een the a. Ti and y phases It had pr e viously b ee n shown that the a.-Ti phase form e d from the peritectoid f3 + y "'"" a. r eac tion near ~ 1480 C, but the results of these r ecent studies s h o w ed that the a.-Ti phas e form e d from the peritectic L + f3 "'""a. r eactio n. Th e pr e viously accepted binary Ti-Al phas e diagram and the revised part of the phase diagram are shown in Figur e 2.1 Th e t e mp era tur e of the peritectic L + f3 "'""a. r eac tion was est imat e d to b e 1475 C from th e study by M cC ullough et al. [16] Furth e r studies by McCullough e t al. showed that th e single a.-Ti phase was present in th e Ti-50at %Al alloy at a temperature of~ 1450 C. Th erefore, the r es ult s of th ese two studies indicat e d that th e a.-Ti phas e in binary Ti-Al alloys was stable at high t e mp era tur es and with co mpositions containing up to 50at. %Al. One of the first survey studies o f th e phas e e quilibria in the t e rnary system was r e port ed by P opov and Rab ezova in 1962 [17 ] In this s tudy the Nb ri c l1 side of the ternary phas e diagram and the t e rnary solid solution rang e of the binary y (l'iAl) phas e wer e ex amin e d. Isoth e rmal sec tions w e r e constructed at 1400 C, 1200 C, and room temp e rature from the phas e analysis of the h ea t tr e at e d samples Th e r es ult s of this study showed th e e xist e n ce of a ternary interm et alli c compound that was given th e n ota ti o n of th e y 1 phas e The y 1 phas e was d ete rmined to hav e a t etra gon a l str u c tur e with a stoichiometric co mposition o f NbTiA1 2 (at.%). Th e lattic e param ete r s th at w e r e deter min ed for the y 1 pha se w ere a = 3. 56A. and c = 4. 69A. It was also found that a quasi-binary sect ion e xisted from NbA1 3 to Ti that contained the ternary

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u 0 V L ::, _,J t;I L. cu a. E cu fFigur e 2 1. 9 A l o m1 c Pe r ce n t A l umi n um 0 10 20 :JO ~O 50 60 70 60 90 100 1600 1.... -r-''--'r--_,_..,..._~ ........ -....L-..,----4-....,.---J....---.-L---_.,_ 1700 187 0C 1600 1500 140 0 (PTi) l:JOO 1200 1100 1 0 00 900 9a2o c 800 (aT I ) 700 600 5 0 0 0 10 T i 0 1800 1 7 00 16 70 1600 1.500 u 0 1400 w er 1300 :, 1200 UJ a. 1 100 UJ 1 000 900 8 00 a Ti 7 00 600 .5 00 0 10 T i L ; ..... ; I '---, ' 1:z:e, I I I '6 I I I I TiAl ~~ ; ' , h ' I I I I ~--'-... , I ' I I Ti3Al I ' ' ' I ' I I I I I I I 2 0 10 ~Ti 20 I I I I I I I I I I I I I I I g ' I ' I ' t ' 0 ' ' t TiAl3 ., ., tt I l: .. ' ft I II ti .. I II i I II tr-------,,-TIAlz ------------------------' I It I aTIAl:, 30 4 0 5 0 6 0 70 Wei g ht P e r ce nt A l uminum Al, wt -lo 2 0 30 40 .50 60 I I I I I , /, , , , ,, , ,, T i3 AI I ' ' I ' I I I I I I I I I I I 30 40 Al L , o I I I -- I A \ t Ti AI 1 It I I ,' ~ t.. , , , , I I I '' I I I I I I ' I I f I I I I I I I I I I f f I I I f I I I I I I f I I I -----' I : :Tl A '1 ' I I I I I I I I I I I I I I I I I I I I 50 60 70 at .-lo (Al ) e o 90 70 8 0 90 100 665 ----------80 90 1 00 Al eao 4&i 0 c 100 Al ( a ) (b) Show s th e binary Ti-Al phas e diagram. (a ) th e pr e viously ac ce pt e d (10] phas e diagram; (b ) th e modifi e d s e ction of th e phas e diagram [5J

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10 y 1 phase. Th e NbA1 3 -Ti qua s i-binary sect ion s how e d that the t e rnary y 1 phas e transformed congrt1ently from th e liquid pha se at a temperature of~ 1850C On o pp os it e sides o f the ternary y 1 pha se w e r e the e ut ect i c r eactio n s oi' L ,.. f3 + y 1 and L,.. y 1 + rtTh e L .... f3 + y 1 r eac ti o n occ urr e d at~ 1550C and th e L,.. y 1 + rt r eactio n occ urr e d at~ 1520 C. Alth o ugh the pr ec is e b o undari es of th e phas e e quilibria w e r e not r e port e d the r es ults of thi s s tudy indi ca t e d that e xt e nsiv e r e gions of ternary solid so luti ons exis t ed f or the binary f3 1 (Nb 3 Al) cr (Nb 2 Al ) 1 rt (NbA1 3 ), and y (TiAl ) pha ses. In th e f o ll o wing y ear, Wukusi ck [18] st udi e d th e t e rnary so lid so lution rang e of Nb in th e y -TiAl phase. Th e ph ase e quilibria th a t w ere r e port e d in this s tudy w ere from X-ra y diffr actio n analysis o f all oys tha t w e r e so luti o n t-r e at e d at 1425 C and then w a t e r qu e n c h e d. On e o f the alloy co mpo s ition s that Wuku sick inv es tigat e d wa s 24Nb-26Ti-50Al (at.%), whi c h w as c l ose t o the sto i c hi o m etr i c co mp osi ti on o f the y 1 phase (NbTiAl ;). H o w e v e r the r es ults o f thi s s tudy did n o t supp o rt th e r es ult s o f Popov and Rab ezova [17] In thi s s tudy th e y phas e, s howing a larg e ternary so lubility rang e for Nb wa s o b se rv e d in s t ea d o f the y 1 pha se. Thu s, the r es ult s by Wuku s i c k co ntradi cte d th e existe n ce o f th e t e rnary y 1 phas e. A stu dy co ndu cte d by Zakb arov et al. [19] in 1984 a dd e d to the co nfu sio n co n ce rning the existe n ce of the t e rnary y 1 phas e Thi s st udy exa min e d the r egio n of the ternary Nb-Ti -Al phas e diagram wh e r e th e tw o qua s i-binary sec ti o ns o f NbA1 3 -Ti a nd Ti 3 Al-Nb int e r secte d. Th e all oys inv esti gat e d in this s tudy w e r e s ubj ected to an extensi v e schedule o f solutionizing h eat treatments with final h eat tr ea tm e nt s at 1200 C, 900 and 600C for o n e hour and then w ate r qu e n c h e d Th e pha se eq uilibria of th e aged s ampl es w e r e d ete rmin e d by X ray diffra ct ion Th e m ost sig nifi ca nt r es ult of t his study wa s the co nfirmati o n that the ternary y 1 pha se whi ch

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11 was r e port e d by Pop o v and Reb e zova [17] e xist e d with th e stoi c hiom e try o f NbTiA1 2 Th e study by Zakharov e t al. als o r e p o rt e d that th e y 1 phas e had a t e trag o nal s tru c tur e How e v e r th e latti ce param e t e rs that w e r e r e ported for th e y 1 phas e by Zakhar o v e t al. w e r e diff e r e nt fr o m th ose tha t w e r e r e p o rt e d by P o p o v and Rab ezo v a. Th e lat t i ce param e t e r s r e port e d by Zakhar o v e t al. w e r e a= 8 418A and c = 4 538A whil e th e latti ce param e t e rs r e p o rt e d by P o pov and Rab ez ova w e r e a = 3. 56A and c = 4 69A. Th e r e a so n f o r th e dj s agr ee m e nt b e tw ee n th e latti ce param e t e r s wa s n ot known. Two co mpr e h e nsiv e studi e s that w e r e publish e d in 1989 by J e w e tt e t al [20] and Kalt e nba c h e t al [21] inv es tigat e d th e ph ase e quilibria of a ll o y s that had c omp os iti o ns c ov e ring th e central portion of th e t e rnary phas e diagram. Th e s tudy by J e w et t e t al exa min e d f o urt ee n diff e r e nt all o y co mpo s iti o n s f o rm e d by h e at tr e ating a r c -m e l te d s ampl es at 1200 C f o r up t o se v e n t o s ixt ee n day s In a similar mann e r Kalt e nba c h et al e x a min e d thirty fiv e diff e r e nt alloy c omp o sitions that w e r e h e at tr e at e d a t 1200 C f o r o n e t o se v e n day s 1n g e n e ral th e r e wa s g oo d agr ee m e nt b e tw ee n th e r es ult s of th e s e two studi e s. Both s tudi es indi c at e d that th e p pha se e xi s t e d at 1200 C o v e r a s ub s tantial part o f th e t e rnary phas e diagram Th e binary pha ses w e r e o b se rv e d to proj ec t int o th e t e rnary sectio n with lar ge solubility r a ng es that l o os e ly follow e d c onstant Al co mpo s itions. This proj e ction can b e s ee n by th e P 1 (s hown as 8) and cr pha ses o n th e binary Nb-Al s id e and by th ey a nd a. phas es o n t h e bin ary Ti-Al s id e o f th e t e rnar y iso th e rm s h o wn in Figur e 2.2 Th e 11 pha se w as d ete rmin e d t o b e i s om o rphou s b e tw ee n t h e NbA1 3 and Ti.Al 3 pha ses, and was s h o wn to co nn ec t th e s e tw o ph ases along th e co n s tant 7 5a t. % Al co mpo s iti o n lin e.

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12 Ti Nb A.1 1.0 0 9 0.8 0. 7 0 6 0 5 0 4 0.3 0.2 0 1 0 0 ATOMIC FRACTION Figur e 2 .2. Show s the 1200 C isoth e rmal section of th e t e rnary Nb-Ti-Al system d ete rmjn e d by J e w e tt et al. (20]

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13 Th e s e r es ults fr o m th e s tudi es by J e w e tt e t al and Kalt e nba c h e t al. h e lp e d t o c larify th e e arli e r r es ult s, but th e y als o c aus e d n1tur e probl e ms c on ce rning th e pha se e quilibria in th e Nb-Ti-Al syst e m B o th o f th e s e studi e s c onclud e d that th e t e rnary y 1 pha se did not e xi s t in thi s s y s t e m. It wa s found inst e ad that alloy s with co mp os iti o ns c lo se t o that r e p o rt e d f o r th e y 1 ph ase c onsist e d o f' th e binary y-TiAl pha se. Th ese s tudi es al so sh o w e d that th e y phas e had a sub s tantial s o lubility limit f o r Nb J e w e tt et al r e p o rt e d a so lubility o f up to 3 0 a t % Nb in th ey pha se, al o ng th e 50at. % Al dir ec ti o n How e v e r J e w e tt e t al al so c laim e d that th e r e w e r e tw o n e w t e rnary pha ses th a t e xi s t e d and th e s e w e r e giv e n th e n o tation of Tl and T2. Th e l o cation o f th e Tl and T2 pha ses ar e s h o wn in th e 1200 C i so th e rm of Figur e 2 2 Th e r e w a s n o s tru c tural informati o n r e port e d f o r th e Tl and T2 pha se s and o nly opti c al mi c r osco py w as use d t o supp o r t th e e xi s t e n ce o f t h e Tl and T2 pha ses. It co uld n o t b e asce rtain e d wh e th e r th e Tl and T2 pha ses w e r e pr e s e nt a t 1200 C, or wh e th e r th e y w e r e d eco mp o siti o n pr o du c ts o f high e r t e mp e ratur e phas e e quilibria. In r e c e n t y e ar s, th e r e h a v e b ee n tw o studi es publi s h e d o n th e t e rnary pha se e quilibri a fr o m th e s am e r esea r c h gr o up On e s tudy wa s by P e r o p e zko e t al. [12] in 1990 and an o th e r s tudy by Da s e t al. (22] in 1993 Th e primary purpos e o f th e s e tw o s tudi es wa s t o d e t e rmin e th e liquidu s pr o j ec ti o n and t o c l arify th e co nfli c ting r es ult s from th e e arli e r s tudi e s o f th e Nb-Ti-Al s yst e m Th e liquidu s pr o j ec tion that w a s d e t e rmin e d from th e study by P e r e p ez k o et al. [12] i s s h o wn in Figur e 2. 3 This proj ec ti o n wa s c onsist e nt with th e r ece nt m o difi ca ti o ns t o th e binary Ti Al pha se diagram [15 16]. It show e d th a t th e a. pha se f o rm e d fr o m th e p e rit ec ti c L + J3 a. re a c ti o n at th e high e r t e mp e ratur e o f :::: 1480 co mpar e d t o th e r pha se whi c h f o rm e d fr o m th e p e rit ec ti c L +a.~ y r e a c ti o n at

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0.2 14 Ti I. o c1s7o> 0.6 0.4 0.8 ,. Po:0. 6 (2060) (194C) /3 I I I ,,,, P7 Cl480J ~a , 5 c,.,o, I ,,, I/ I ,' j' P, (13702 y :I 11 7') ,, \ / ,, 0.2 Fi g ur e 2. 3 Sh o w s th e liquidu s pr o j ec ti o n o f th e t e rnary Nb-Ti Al s y s t e m d e t e rmin e d by P e r e p ez k o et a l [12] p 3 AI e1

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15 1450 C. Ther e w e r e two important points made from this liquidus projection Th e first point was that the liquidu s surface of th e P phas e co ver e d an e xtensiv e p o rtion of t h e ternary pha se diagram. Th e seco nd point wa s that th e bivariant L + P + cr and L + cr + y ti e -triangl es reacted in a c lass II four phas e r e a c tion [23] n e ar the central p o rti o n o f th e liquidus proj ec tion. Th e produ c t L + P + y ti e -triangl e then m ove d t o ward th e binary Ti-Al side, wh e r e it r e a c t e d in an o th e r c l ass II four phas e r eac ti o n with the L + p + n tie -triangl e. Th e L + p + n ti e -triangl e had originally started from the binary p e rit ectic L + p .n r e a c tion. 1 h e product of this r eac tion was th e L + p + n tie-triangl e, which t e rminat e d at th e binary p e rit ec tic L + n .y r eact i o n Thus th ese two point s indi ca t e d that th e so lidification paths of alloys that had co mp osi tions n ea r th e ce ntral portion of the t e rnary phas e diagram could hav e so lidifi e d with the p phas e or co uld hav e co mplications near th e four phas e r eac tions. Furth e r work by Da s et al. [22] indi ca t e d that r e fin e m e nt s in th e 1200 C isotherm had occ urr e d n e ar th e co mp os ition of Ti 4 Al 3 Nb. Th e r e vi se d 1200 C i so th e rmal section, which is c urr e ntly acc e pt e d to b e the most accurat e r e pr ese ntation of th e high t e mp e rature phas e e quilibria in the Nb-Ti-Al system, is shown in Figur e 2 4 Th e r e fin e m e nt mad e by Das e t al. involv e d c hanging the ph ase for th e r e gion that had the Ti 4 Al 3 Nb composition from th e T2 phas e to the n phas e. Th e T2 phas e whi c h h a d b ee n see n in th e e arli e r 1200 C isothermal section o f Figur e 2.2 wa s id e ntifi e d as the ordered p (B2 ) pha se. H o w e v e r it was argu e d that th e B2 pha se had f o rm e d from th e specific cooling m e thod s e mploy e d and was not co n.sid e r e d to b e an e quilibrium phas e. Lik e wis e th e Tl phas e wa s also f o und t o hav e had th e B2 phas e, and wa s disr e gard e d as an e quilibrium phas e by a similar ar gum e nt us e d for th e T2 phas e.

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, , ,,. 16 Ti 82 Figur e 2.4. Shows the 1200 C isoth e rmal section of the t e rnary Nb-Ti-Al system d e t e rmin e d by Das et al. [22].

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10 at. o/o Ti-30Al-45Nb "'1 ~,._.._.,.-, I , I I I, I I ,, 17 Ti-30 Al Ti50AI
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18 Th e ex. pha se that was studi e d by Das e t al. [22] was a s sum e d t o hav e th e H C P s tru c tur e s in ce n o s tru c tural r es ult s w e r e r e port e d in th e study. N e v e rth e l e ss th e phas e boundari e s n e ar th e Ti 4 Al 3 Nb co mpo s ition o f th e ex. pha se w e r e s hown t o b e co mplicat e d. Th e ti e -lin e s surrounding this co mpo s ition w e r e signifi c antly c h a ng e d aft e r just a 50 C dr o p in t e mp e ratur e from 1200 Th e s e ar e sh o wn in th e partial sec ti o n s o f th e 1200 C and 1150 C i so th e rm s in Figur e 2 5 On e o f th e m os t dr am a t ic c hang es in th e b o undari es o f th e pha se e quilibria b e tw ee n 1200 C and 1150 C inv o lv e d th e pr es um e d c la ss II f o ur phas e r e a c ti o n of th e p + y + ex.* and p + y + cr tie -triangl es (Figur e 2.5a ) to form t h e ex.*+ cr + p and ex.*+ cr + y t i e -triangl es (Fi g ur e 2 5b ) Thi s r e a c ti o n was shown to hav e shift e d th e ti e -lin e s in th e two ph a s e ex.*+ cr fi e ld s by n e arly 90 fr o m th e two-pha se p + y fi e ld s Th e co mbin e d r e sult s of P e r e p ez k o e t al. [12] and Das e t al [22] al so indi c a te d that th e B2 phas e co v e r e d a n ex t ens iv e co mp o sition rang e in th e ce ntral p o r t i o n of th e Nb-Ti Al s y ste m. Th e e xt e nt o f t h e B2 ph ase a t 1200 C was s h o wn pr e vi o u s ly in F i gur e 2 4 H o w e v e r th e co mp osi ti o nal b o undari es o f th e B2 pha se w e r e n o t pr ec i se ly kn o wn 2 1 2 Th e Ti-Ba se d All o y D e v e l o pm e nt Studi es Th e phas e e quilibria and th e micr os tructur e s that c an b e produ ce d from vari o u s h ea t tr e atm e nts of alloy s in th e Ti-rich part of th e t e rnary Nb-Ti-Al syst e m hav e b ee n r e vi e w e d by Williams [l] and Rhod e s [2]. Th e s e alloys show tha t a vari ety o f mi c r os tru c tur e s c an b e produ ce d by pro ces sing or h e at tr e atm e nt s t o f o rm diff e r e nt m o rphol o gi e s di s tributi o n s and co mbinati o n s o f th e e quilibrium a. and p pha ses, th e m e t as tabl e ex.' a nd ex." mart e n s i tes 1 and th e m e ta s t a bl e ro pha se Th e r easo n t h e diff e r e nt mi c r os tru ct ur es w e r e produ ce d wa s du e t o th e e xi s t e nc e o f th e p ph ase at

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19 higher temperatures that could be stab iliz ed down to room temperature with sufficient Nb additions. The renewed int e r es t during the 1980s in th e d e v e lopm e nt of Ti-aluminid es prompted se v e ral studies of the pha se eq uilibri a in a. 2 -1'i 3 Al ba se d alloys that co ntain e d up to 30at. %Nb additions. In 1988 a study by Ban e rj ee et al. [13] showed that an orthorhombic phase (0 -phas e) form e d in a Ti-25A1-12.5Nb (at.%) alloy that wa s h eat treated at 1100 C and furnace coo l ed. This h eat treatment produced eq uiax e d a. 2 grains in an ordered 13 (B2-CsCl structure) matrix at 1100 C. Th e slow coo ling rate resulted in th e growth of the a. 2 grain s, whi c h ca us e d the formati o n o f the 0-phase at t h e triple p o ints of the imping e d a. 2 grains. The 0-phase was studied by se l ecte d area e l ectro n diffraction (SAED) and co nv e rg e nt b ea m e l ec tron diffra ctio n (CBED) to s h o w that the structure wa s cons i ste nt with the C m c m space group and that the lattic e parameters w e r e a= 4.50A, b = 5 88A and c = 9.60A. Th e c l1ann e lling e nhan ced microanalysi s technique wa s used to d e t e rmin e th e atomic site occ upan cy of the 0-phase. Thi s technique indi cated that the co mp os ition o f this phas e wa s bas e d on the TizA].Nb stoichiometry. Th e presence of the 0-phase w as s ubs eq u e ntly co nfirm e d by Kaufman et al. [24] in 1988 How ever, the latti ce parameters determined f or the 0-pbase by Kaufman et al. w ere different than those d e t e rmin e d by Banerj ee et al [13]. Th e 0 -pha se wa s observed in TI-25Al + Nb (at.%) alloys that co ntain e d at l east 12 at. % Nb a nd w e r e h eat treated between 800 C and 1000 C. Th e C BED analysis showed th e 0-phase to hav e the sa m e C m c m space group, but the lattic e param e t e rs w e r e a = 6. 2A, b = 9. 4A., and c = 4 7 A.

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20 Th e temp e rature and composition rang e of the 0-phase w e r e investigat e d in several st udi es imm e diately following th e dis co v e ry [25-35]. Th ese st udi es cont-ri but e d mainly to the und e r sta nding of the pha se e quilibria in the section o f the pha se diagram tha t co nn ec ts th e binary a. 2 -Ti 3 Al pha se to th e O-Ti 2 NbAl pha se. Th e Ti-27. 5Al + Nb sect i o n shown in Figur e 2. 6 w as developed fr o m th e r es ults by Ban e rj ee et al (32] and is th e c urr e ntly accepted pha se e quilibria for this part o f the ternary phase diagram Evidence showing the existe nc e of a seco nd t e rnary phas e wa s r e port ed in th e st udy by Strychor et al. [36] in 1988. This e videnc e was obtained from low temperature h eat treatments on alloys ba se d o n th e composition o f Ti 3 Al alloy s with t e rnary additions from 5 to l 7at %Nb. Th e alloys co ntaining mor e than 5at. %Nb s h o w ed a r a pid dr op in the t e mp erat ur e of t h e a.' mart e nsit e and r e t e nti on o f th e J3 phase, whi c h had orde r e d to the B2 phase during coo ling Th e n th e m etas t ab l e B2 phase w as shown to hav e d eco mpo se d to an ro -typ e ph ase during th e l o w temperature h eat ing. Th e SAED analysis of the ro -typ e pha se indi ca t ed that th e structure wa s o rd e r e d and wa s cons ist e nt with the B8 2 structure, whi c h was a l-I C P-Zr 2 Al pr ototype str u ct ur e. Th e co-B8 2 phas e was confirmed in a 1990 study by Bend e rsky e t al. [37]. In this st L1dy an alloy with th e Ti 4 Al 3 Nb composition was h ea t tr eated at 700 C for 26 d ays. This ca us ed the pri o r B2 matrix to transform co mpl ete l y to the co B8 2 pl1a se. Thu s, the sto i c hiom et ri c co mp os iti on o f the co -B8 2 pha se w as d ete rmin e d to be Ti 4 Al 3 Nb. The st ru ct ur e of the coB8 2 phase w as investigated u s ing SAED and C BED analysis to s how th a t th e space group was P6 3 / mm c, and th e lattic e param e t e r s w e r e

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21 1200 IP (82)j CJ .. C) 1100 '::s Cll+B2 ... ta fl 'C) -.. a. -\ _.,. .. \ E --\ 1000 C) ... ... 3 pha 0+8 2phaN pha [QJ 900 Ti-27 SAI 10 20 30 at 0 /o Nb Figur e 2.6 Show s the temperature-composition diagram of th e Ti 3 Al to Ti.zA].Nb section by Ban e rj ee et al. [32).

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22 a= 4.580A and c = 5.52 0A The mechanism of this transformation was also studied, but will be discuss e d later in sect i on 2. 3 on th e om e ga phas e s. There wa s a follow up study by B e nd e rsky et al. [38] in 1990 that r e p o rt e d the possibility of anoth e r ordered d e rjvative of th e ID-phas e In this st udy an alloy with th e co mposition of Ti-37. 5Al-20Nb ( at.%) was h eat treated at 700C for 18 day s Th e r es ults sho w ed that small pr ec ipitat es form e d in the matrix of th e co -B8 2 grains. Th e SAED analysis indi ca t e d that th e l attice param e ters o f th e pr ec ipitat es w e r e a= 7 93.A. and c = 5.52.A. Th ese l attice param e ter s of th e pr ec ipitat es wer e l arger than those of th e IDB8 2 pha se. Tw o possibl e str u ctures of this ternary phase w e re proposed. On e p ossi bl e structure was based on the P ea rson symbo l hP 18 whi c h had th e P6 3 / mcm space group and Ga 4 Ti 5 prototyp e structur e Th e ot h e r possibl e structure was based o n th e Strukturb e ri c ht D8 8 structur e, whi c h h ad th e P6 3 / m c m space group and ~Si 3 prototyp e structure. 2 2 Th e Sigma Phas e The sigma ( cr ) phas e occurs in many transition-m e tal alloy syste m s that are of tec hn o l ogical int e r est in a ll oy d e v e lopm e nt s uch as the s tainl ess stee l s and the s up e ralloys. How eve r the occ urr e n ce of th e cr phas e in th ese alloys has usually b ee n avoided b ec aus e of its gen e rally hard and brittle prop e rties at room t e mp e ratur e. Th ese prop e rti es hav e had a d e l e t e rious e ff ect on th e m ec hanical prop e rti es, since the cr phas e usually forms along th e grain boundaries of aff ec ted alloys, wh e r e cracks can nucleate and propagat e Converse l y, th e cr phas e has a high melting p o int and a complex o rder e d structure that may provide so m e strengt h and creep r es istan ce at high temperatures. A full r e vi e w of the structur e and properties of th e cr phas e in various binary and t e rnary all o y syste ms i s d escr ib e d by Hall and Al gie [39]

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23 Th e cr pha se ha s b ee n d e t e rmin e d in th e binary Nb-Al syst e m t o form by t h e p e rit ec ti c L + f3 1 cr r e a c ti o n at ~ 1940 C [40]. Fr o m Tabl e 2 1 th e stoichiom e tri es o f th e f3 1 and th e cr ph ases w e r e d e t e rmin e d t o b e Nb 3 Al and Nb 2 Al r es p ec tiv e l y. Th e maxim11m s olubility o f Al in binary co mpo s itions of' th e cr phas e wa s 12at %Al a t 1600 C. Th e ran ge o f t e rnary co mp os iti o n s of th e cr phas e in thi s s yst e m i s s h o wn in Figur e 2 4 [12 20-22] Th e t e rnary cr pha se r e gion wa s shown at 1200 C to e xt e nd fr o m th e binary Nb-Al s id e int o th e t e rnary Nb-Ti-Al s y s t e m al o ng th e co nstant Al c ompo s ition sec ti o n Th e so lubility rang e f o r Ti in t h e t e rnary co mp o sition s o f th e cr pha se, was e stimat e d to b e ~ 35at %Ti at 1200 C Th e L + f3 1 cr liquidus vall e y whi c h e xt e nds fr o m th e binary p e rit ec ti c r e a c tion int o th e t e rnary sec tion d ec r e a ses with t e mp e ratur e, as s h o wn in Figur e 2. 3 Th e s tru c tur e o f th e cr pl1as e, ba se d o n th e Nb 2 Al s t o i c hiom e try ha s b ee n in ve stigated by Wil so n and Sp oo n e r [41 42] This pha se ha s b ee n d esc rib e d a s a t o p o l o gi c ally c l ose pa c k e d (T C P ) s tru c tur e, s in ce it c an b e vi e w e d a s co nsi s ting o f distor te d h e xagonal c los e pack e d lay e r s o f at o ms tha t ar e rotat e d by 90 b e tw ee n e a c h a lt e rnating lay e r. Lik e th e cr phas e s in o th e r various alloy sy s t e ms th e cr phas e in th e Nb-Al sys te m ha s a t et rag o nal s tru c tur e with th e P4 2 / mnm s pa ce gr o up Th e latti ce param e t e rs o f th e s toi c hiom e tric Nb~ compositi o n w e r e m e asur e d to b e a= 9.935A and c = 5. 169A [11] H o w e v e r t h e latti ce param e t e r s o f th e unit ce ll f o r thl s cr pha se do c han ge wi t h th e Al co nt e nt Th e r e ar e 30 atom s a ssoci at e d with th e un it ce ll o f' th e cr ph ase. Th ese at o m s ar e arrang e d with varying d e gr ees o f o rd e r o n fiv e diff e r e nt latti ce s it es in th e unit ce ll Th e sc h e mati c s h o wn in Figur e 2 7 s h o w s th e fiv e diff e rent l a tti ce sit es in th e proj ec tion of th e unit ce ll along th e c axi s o f th e cr ph ase. Th ese fiv e s i tes with th e ir e quival e nt Wy c k o ff d e signati o ns ar e A ( 2a ), B ( 4f)

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24 C (8i 1 ), D (8i~, and E (8j) (43]. Th e occupation of th ese sites by th e Nb and Al atoms was found to consist of pr e dominantly Nb atoms on th e B, C, and E sites; mostly Al atoms on the A site ; and a mixtur e o f both Nb and Al atoms on th e D site [ 42] Th e coor dination numb er does hav e an e ff ec t on this particular s it e occupancy. Th e B site has the larg est coo rdination numb er (C N = 15) th e A and D sites hav e the small est (CN = 12 ), whil e th e C and E sites have a n int e rm e diat e value (CN = 14) Th e r e f o r e, both th e atomic siz e and e lectroni c b e havior can aff ect the specific site occ upancy in the cr st ru c tur e. 2.3 Th e Gamma Phas e Th e gamma (y) phase from the binary Ti-Al system has r ece iv e d ex tensiv e research during th e last t e n y e ars This surge in r esearc h has b ee n attributed to several attractiv e pr o p e rti es of the y phase such as low d e nsity good oxidation r es i s tanc e, and high temperature strength r e t e nti on Th e d e v e lopm e nt of pot e ntial a lloy s ba se d on this phase ha s r es ult e d from st udi es on the phas e r e lations ; mi c ro struct ural formation by pr ocess ing ; and mi c rostructur e -pr operty r e lation s hip s relating to oxidation d e formation and fract1.rr e Furth er information on th ese topics for th ey phas e i s provid e d in th e r e vi e w articl es by Kim (7 44]. Th e formation of they phas e in th e binary Ti-Al system was rec e ntly modified to s h o w that it solidifi e d in the p e r itec tic L +a~ y r e action at ::: 1450 C, as shown in Figur e 2.1. Th e so lubility rang e of this pha se ex t e nds mor e to th e Al-ri c h side, rath e r than to the Ti-ri c h s id e, from th e exac t TiAl stoichiometry In studies of the t e rnary Nb-Ti-Al sys t e m th ey phase wa s found to hav e an ex t e nsiv e solubility rang e f'or Nb at temperatures in the 1200 C range [12 20-22] Th ese s tudi es indicat e d that up to 26at. %Nb wa s soluble in th e y pha se at 1200 C.

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B D c.--8 D 8 D D 25 D C z-, z0 z-0 C D D 8 D C C 8 C D B Figure 2. 7. Shows the five lattice sites in the projection of the unit cell for the cr phas e [41).

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26 Tl1 e structure of' the y phase wa s determined to be the fac e -c e nt e r e d tetragonal L1 0 structure and the P4/mmm space group [10] Th e 11nit ce ll of this structure has a lay e r e d arrangement of Ti and Al atoms on alt er nating (001) plan es, as shown in Figur e 2.8 Th e Ti atoms occupy the la and l e Wyckoff sites at ( 0, 0 0) and , 0) whil e the Al atoms occupy th e 2 e Wyckoff site at (0, and 0 ). Th e latti ce parameters for this e quiatomic TiAl composition, w e r e det e rmin e d to be a= 3.976.A. a od c = 4 049.A. [ 45] How e v e r these latti ce param ete rs also d epe nd e d o n th e Al co nt e nt. M eas uring these changes in the lattic e param e t e rs by th e c/a ratio showed that the e/a ratio increased from 1 02 to 1. 03 with in creas ing Al co nt e nt l 46]. The atomic site occ upan cy of the y phas e co ntaining t e rnary additions of Nb was studied by Kr o nitz e r et al. [47] and Jackson [48]. Kronitz e r et al. showed that Nb atoms randomly occ upi e d the la and l e Wyckoff' sites with the Ti atoms, ~ s mall additions of Nb to the y -TiAl pha se. In comparison, Ja c kson reported the existence of a n e w phas e that was bas e d on the tetragonal L6 0 structure and P4/mmm space group. Thi s structure s how e d the Ti atoms occupied th e la Wy c koff site, the Nb atoms occ upi e d the l e Wyckoff site an d the Al atoms occupied the 2e Wyckoff site. Thus the n e w structure r epo rt e d by Ja c kson showed that o rd er ing had occurred between the Nb and Ti a tom s o n the la and l e Wy c k o ff' s it es. Th e occ urr e n ce of this ordering r eac tion from the y L1 0 structure was supported by the observation of APDBs in the yL6 0 structure 2. 4 Th e B2 Phas e Th e B2 pha se ha s b ee n shown to exist in alloys of the Nb-Ti-Al system at both high temperatures and over l arge composition rang es. Th e composition range of' alloys that had been f o und to co ntain the B2 phas e includ e d compositions in the

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27 C la and le Wyckoff Sites: Ti 2e Wyckoff Site: Al Figur e 2 8 Sh o w s th e unit ce ll o f th e y -TiAl phas e

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28 central portion of the ternary phase diagram [12] compositions c los e to Ti 3 AJ. containjng Nb additions in the rang e of:= 5Nb to 30Nb (at.%) [13 24-35] th e Ti 4 Al 3 Nb composition [37] and binary Nb-Al compositions with additions of 13.5Al to 16.9Al (at.%) (49]. The 1200 C isotherm by Das et al (22] showed that th e solid solution rang e of the B2 phas e was fairly small and was locat e d c l ose to th e center of this ternary syste m a s shown in Figur e 2.5 Th e temp e ratur e of the p to B2 transition w as d ete rmin ed to increase with an increase in th e Nb co nt e nt of' alloys bas e d o n th e appro)cimate Ti 3 Al co mpositi o n. M o r e specifically, th e p to B2 transition t e mp e ratttre was f ound to incr ease from := 1100 C for 12.5at. %Nb additions to over 1400 C f o r 25at %Nb [13 24-26 28-31 ,33 -35] Lik e wis e, a high transition t e mp e ratur e o f mor e than 1400 C was suggested by B e ndersky e t al. [37] for an alloy with a Ti 4 Al 3 Nb co mpo sitio n. Th e B2 phas e was also recently o bserv e d in binary Nb-13.5Al and Nb-16.9Al (at.%) all oys aft e r b e ing heat tr ea t e d at only 800 C for 10 hours [49]. This B2 phase was d eter mined to b e a m e tastab] e phas e that had form e d inst ea d of the e quilibrium Nb 3 Al ( P 1 ) phas e, at the low aging temperature. Th e formation o f th e B2 phase, inst ea d o f th e p 1 phas e, was possibl e b ec aus e the activation e n e rgy of the seco nd or d er P to B2 transition was l o w e r than that of the first order P to p 1 transition. Th e r e f o r e, the diffusion l e ngth that wa s r e quir e d to form th e B2 pha se w as shorter and the formation o f th e B2 phas e occurred, rath e r than th e P 1 phas e. Th e B2 pha se has th e ordered BC C (CsCl) structure and th e Pm3m space group. Th e atomi c s it e occupancy of the B2 phas e in a Ti 25Al-10Nb (at.%) all oy wa s determined by Ban e rj ee e t al. (50] using c hann e lling e nhanced mi c roanalysis Th e co mpo s ition of th e B2 phas e in th e two phase a.2 + B2 microstructur e of this alloy was f o und to be Ti-24. 5A1-14Nb (at.%). Th e c hann e lling r es ults of the B2 pha se indicated

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29 that th e tw o di s tin ct s ublatti ces in th e B2 unit cell w e r e occupied mostly by Ti on th e la Wyckoff sit e at ( 0 0 0) and by a mixtur e of Ti Nb and Al on th e lb Wyck o ff site at (, ). Th e e xact atomic p e rc e nt valu es that w e re d e termin e d for the tw o sites w e r e 48Ti and 2Nb f o r the la site, and 14Ti 12Nb and 24Al for th e lb sit e Th e sc h e mati c shown in Figur e 2 9 s h o ws the B2 unit ce ll using the atomic site occupancy d ete rmin e d by Ban e rj ee e t al. A 1988 study by Stry c b o r et al. [36) s h o w e d that th e B2 phas e in Ti 3 Al + Nb alloys exhi bit ed diffuse sca tt e ring and ext r a maxima in th e selected ar e a e l ectro n diffracti o n (SAED) patterns, and tweed microstructur es in the TEM imag e s. Th e diffus e scatte ring and e xtra maxima in th e SAED patt e rns was int e rpr e t e d to s how an in s tability in th e B2 lattic e that suggested two pos s ibl e mod es of transformati o n. Th e first mod e wa s d e t e rmin e d from tl1e o bs e rvation of diffus e scattering at the l/3 < 111> positions in th e SAED patt e rns o f' th e B2 matrix This diffus e scattering wa s a ttribut e d to 2 /3< 111> longitudinal displa ce m e nt waves or phonons that l e d t o the formation of an o rd e r e d co-type phas e during low temperature aging of Ti 3 Al + Nb all oys. Th e second mod e wa s d e t e rmin e d from the o b se rvation of diffus e streaks tha t ran parall e l to the < 110> dir ectio n s and extra maxima at 1/2{110} positions in the SAED patt e rns of' the B2 matrix. Th ese observations w e r e attribut e d to 1 /2< 110>{110} type transverse lattj ce di s pla ce m e nt wav es, that c au sed lo c ali ze d strain in th e B2 latti ce, which l e d to the formation of th e tw ee d micr os tru c tur e. Th ese shear strains w e r e shown to b e consistent with the lattic e d e formation of the mart e nsitic transf'ormation of th e 2H ps e udo ortho/hex (o rth o rhombi c /h exa g o nal ) st ru ct ur e from th e B2 pha se Th e diffus e streaking and ex tra maxima observed in

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30 a a la Wyckoff Site: Ti Figure 2.9. Shows the unit cell of the B2 phase with the atomic site occupancy determined in a Ti-24.5Al-14Nb (at.%) a ll oy by Banerjee et al [50].

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31 th e SAED patt e rn s w e r e formed fr o m th e tw ee d micro s tru c tur e, and w e r e e xplain e d a s r e l-r o d s that int e rs e ct e d th e Ewald s ph e r e Thi s t w ee d mi c r o stru c tur e ha s b ee n o b se rv e d in th e B2 pha se o f many all oy s y s t e ms that e xhibit e d mart e nsiti c transformations [51] This r e lationship b e tw ee n t h e t w ee d mi c r os trt1 c ttrr e and th e mart e n s iti c tran s f o rm a tion in th ese typ es o f all oys ha s b ee n inv e stigat e d by Rob e rtson and Wayman [52-54] Tann e r e t al [55] and S c hryv e r s and Tann e r [56]. Rob e rtson and Wayman s h o w e d that th e tw ee d mi c r os tru c tur e wa s f o rm e d by < 110>{110} s tati c di s pla ce m e nt wav es, whi c h r es ult e d in th e s oft e ning o f t h e e lasti c c onstant C wh e re C' = ( C 1 1 C ~, in th e B2 ph ase o f a 63Ni3 7 Al ( at .%) all o y Tann e r e t al. and S c hryv e r s and Tann e r s h o w e d that th e t w ee d mi c r os tru c tur e in th e B2 pha se of t h e 63Ni3 7 Al (a t. %) all o y wa s c o mp ose d o f a fin e sc al e d m osaic asse mbly o f n o n-unif o rmly distort e d and mi c r o modulat e d d o main s, whi c h w e r e co in e d inh o mog e n eo u s ly s tr a in e d domain s (ISD s) Th e ISD s w e r e e xamin e d by high r eso lution e l ec tr o n mi c r osco py (HREM) Th e ISD s w e r e s h o wn t o li e p a r a ll e l to {110} plan e tra ces in th e B2 m a trix and t o hav e a si ze o f :::: 40-60 A in l e ngth with an a v e r a g e spa c ing o f :::: 13A thi c k. Th e co mput e r simulations by S c ln yve r s and Tann e r c onfirm e d tha t t h e a t o mic s tru c tur e o f th e ISD s wa s co n s i s t e nt wi t h s mall tran s v e r se shuffl es, plu s s h e ar di s t o rtion s of th e <110>{110} typ e. Th e di s pl ace m e n ts in th e ISD s w e r e co rr e lat e d with th e l o w e n e rgy tran s v e r se :E 4 (ss O ) -T ~ ph o n on m o d e t h a t w as f o und to h a v e an an o m o l o us t e mp e ratur e d e p e nd e nt in co mpl ete s oft e ning at s:::: 0.16 It wa s th eo riz e d th a t th e ISD s w e r e for m e d by th e dyn a m ic so ft e ning o f th e C co nstant whi c h r es ult e d fr o m th e TA 2 phon o n that b ec am e co upl e d t o th e s tati c s tr a .in fi e ld s of d e f ec t s in th e B2 latti ce. It

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32 was suggested that the nu c l eat ion site of the mart e nsitic phase was d e t e rmin e d by th e st r e ngth of th e s train fi e ld that was associated with th e d e fe c t. 2.5 Th e Om e ga Phas e Th e omega (co) phas e bas b ee n an intriguing phas e that was first discov e r ed in thermally aged, Psta biliz e d tit a ru11m all oys in 1954 [57] Sin ce th e n it has b ee n o b se rv e d in the group IV transiti o n e l e m e nts o f Ti Zr and I-If; in num e rous alloy s that cons i ste d o f the group IV e l e m e nts; and in the e l e m e nts imm e diat e ly t o the right o f the group IV e l e m e nts ( th e d-ri c h tran s ition e l e m e nts of Nb V and Mo) (58]. Th e co phas e has al so been observed in many alloys that w e r e not ba se d on the gr o up TV e l e m e nts s u c h as: C u-Zn [59 60] C u-Sn [59 61] Cu-Zn-Al [62] C u-Zn-Si [63] C u-Al-Ni [64] Ag-Al [65] Ag-Mg [66] F e -Al [67] and Ni-Al (68] On e of the main reasons for inv estigating th e co phase was because o f the harmful e ff ect this phase has on the m ec hanical properties of th ese alloys. Th e formation of th e co phas e ha s been s h o wn to emb rittl e t h e all oy by decreasing the du ct ility and increasing the hardn ess. Extensive s tudi es hav e b ee n co ndu cte d on the morphology ; the e ff ects on th e physical properties; the natur e of the diffus e x-ray, e l ectro n and neutron scattering; the transformation kin et i cs; and th e transformation m ec hanism o f th e co phas e. Th e r es ults of these s tudi es w e r e s11mmarized in a r e vi e w arti c l e by Sikka et al [58] Th e co phas e that forms in Ti has b ee n s ugg este d t o b e the l o w temperature high-pr ess ur e modifi ca tion of the p phase [69]. Normally th e a-H CP phas e i s th e sta bl e pha se in the unary Ti syste m below the all ot r o pi c transition of ~ 882 C a t atmospheric pressure. H o w e v er, the co pha se ha s been s l1 o wn at high pressures to be the stab l e phase, instead o f th e a phase, b ase d o n th e e quilibrium unary P T diagram that was developed for Ti [58]. Th e str u ct ur e o f the ideal co phas e in pur e Ti wa s

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33 found to b e th e H C P structure, with th e P6/mmm s pa ce group. Th e unit cell of this id e al phas e contains three atoms: one atom located on th e la Wy c koff site at (0, 0 0) and two atoms locat e d on the 2d Wyckoff site at ( 1 /s, 1/s and (1/3, 1/s ). How e v e r the ro phase has also b ee n shown to hav e th e trigonal structur e with tl1e P3ml s pa ce group in Ti-bas e d alloys. Th e dill e r e nc e in th e symm e try betwe e n th e ro-trigonal and ro-I--ICP phases was due to the small atomic displac e ment of th e two atoms on the 2d Wyck o ff' site in th e ro-I-ICP phas e. Thus th e atomic site occupancy of th e ro-trigonal phas e was shown t o hav e one atom locat e d on the la Wyckoff sit e at (0, 0 0) and two atom s lo ca t e d on th e 2d sit e at (1/s, 1/a z) and (1/ a 1 /s z). The z param e ter indicat es the magnitud e of th e atomic displa ce m e nt and that th e displac e m e nt occurs in th e dir ect ion of th e c -axis in th e trigonal unit cell. Th e p to ro transformation has b ee n d esc rib e d by a plan e collapse m ec h a ni s m that involves thr ee {lll}p plan es. In this m ec hanism one pair of plan es collapses t o g e th e r to an int e rm e diat e position whil e th e third plan e remains unalter e d (58]. Th e atomi c displac e m e nts that w e r e r e quir e d for this tra nsformati o n hav e b ee n describ e d by a soft-mode transformation m ec hanism. This mechanism involv es a l o ngitudinal sinusoidal wav e with atomic displac e m e nts U = Wsin[~x + (x)] ; a wave v ec t o r = 2/3<111 > 21t/ap ; a phas e for th e thr ee possibl e variants (x) = 0 21t/3 41t /3; and an amplitude W = ap/6 for ro-HCP or a valu e l ess than this for ro -trigonal Th e {lll}P plane collapse mechanism and th e soft mode m ec hanism of' th e p to ro transformation ar e s hown in th e schematic of Figttr e 2.10. This schematic shows th e proj ec tion of both th e ( 110) 13 plan es for th e p phas e and th e (1120)(1) p l an es for th e ro phase. If th e collapse of the two {lll}P planes is incomplet e, then th e atoms on th e B

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34 plan e ar e displa ce d. This r e sult s in a rumpl e d B plan e and th e trig o nal symm e try f o r t h e ro ph ase Th e tran s formati o n of th e J3 ph ase to th e ro pha se has b ee n shown t o f o rm tw e l ve variant s o f t h e ro pha se. Th ese v ar iant s w e r e f o rm e d from t h e or i e nta tio n r e l at i o n s hip tha t t h e ro ph ase h as with th e p ph ase: ( OOOl )ro ll{lll}J3 and < 1120 >ro 11 < 1TO > J3 Th e tw e lv e vari a nt s w e r e f o rm e d fr o m f o ur rota t i o n a l variant s t hat eac h had t hr ee tr an s l a tional variant s, a s s h o wn in th e sc h e mati c of Figur e 2 11 Th e tw o rot a ti o nal variant w e r e f o rm e d by alignin g th e ( 0001) 00 plan e parall e l t o e ith e r th e ( l l l )p o r ( 1 TI) 13 plan es o f th e p pha se. Th e thr ee tr a n s l a ti o nal vari a nts w e r e d e t e rmin e d fr o m t h e {111} 13 plan e th a t r e main e d un a lt e r e d during th e plan e collaps e. Sinc e th e stac kin g se qu e n ce o f th e {111} 13 plan es in th e B CC p pha se i s .. AB C AB C .. th e n t h e ( 0001 ) 00 plan e o f t h e A v ari a nt i s f 'o rm e d fr o m th e A plan e, th e B-variant fr o m t h e B plan e, and th o C -variant from th e C plan e Th e r e f o r e, ix o ut o f th e tw e lv e t o t a l v ar iant s ar e ill us tra te d in th e sc h e m atic o f Figur e 2 11 Th e r e hav e b ee n se v e r a l studi es in th e pa s t that hav e inv es tig a t e d th e influ e n ce o f Nb o n th e f o rmati o n o f th e ro -typ e ph ase in binary Ti-Nb all o y s [70-7 3 ] Th e s e s tudi es h a v e s h o wn that th e ro pha se c an tran s f o rm b o th ath e rmally an d i s oth e rm a lly fr o m t h e J3 phas e in th e Nb-Ti all o y s. Th e primary diff e r e n ce b et w ee n th e tw o f o rma t i o n m ec hani s m s wa s th a t th e co mp os iti o n o f th e ath e rmal ro ph ase wa s th e s am e as tha t o f t h e J3 m a tri x, whil e th e co mp os i t i o n o f th e iso th e rmal ro ph ase w as diff e r e n t fr o m t h a t of th e P m a tr ix. H o w e v e r th e i so th e rmal ro phas e wa s d ete rmin e d to fo rm fr o m th e p m a trix b y th e s am e m ec hanj s m as th e a th e rmal ro pha se, whi c h inv o lv e d th e co llap se o f alt e rnating pair s o f {111} 13 plan es

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A C B A C B A A B A B A 35 [lllJ P (lll) P trace .. ....... I~ ,.: = : =.: ,=.: :~ .. ....................... _....... ._.., .. ---------------".". ... .. .. . . ...... ....... [OOOl] co (OOOl)co trace .. ........................................ ... .. ... .. .. ......... .. .. .. f '. }::: . . ... Figur e 2. 10. Shows the {111} 13 plan e co llap se mod e l of the~ to co transformation. vi e w i s normal to th e ( 110) 13 plan es. Th e

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36 C-variant B vaiiant A-variant (111) Plane Trace '---~ (III) P l ane Trace Figure 2.11. Shows two rotational variants each one containing thre e translational variants of the co phase formed from th e J3 to co transformation. Th e vi e w i s normal to th e ( 110) 13 plan es.

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37 The most systematic studies of the ro phase in the binary Nb-Ti system w e r e by Moffat and Larbalestier [72 73]. In these studies the effect of cooling rat e and isoth e rmal aging w e r e investigated in alloys that contained 20Nb to 35Nb (at.%). Th e re s ults of thes e studies show e d that there was a competition between the formation of the ro phas e, th e metastable a" mart e nsite and the e quilibrium a phase d e pending on th e coo ling rat e from 1000 C and th e isoth e rmal aging temperatur e It wa s shown that wat e r qu e nching caused th e formation of th e a" martensit e, while air cooling and furnace cooling caused th e formation o f' th e ro pr ec ipitates. Th e aging e xp e rim e nts indicated that th e eq uilibrium a phase was formed at the higher aging temperatur es, such as 400 C to 500 C, while th e low e r aging temp e ratur e s such as 200 C and 300 C, caused th e formation of the ro phas e. Th e c omp e tition b e tw ee n th e a" mart ens it e and the m pr ec ipitat es in the alloys was d e termin e d to hav e result e d from th e low er ing of the (martensite st art) temp e ratur e of th e a" mart e nsit e by the addition of >2 0at. %Nb [72]. It was also suggested by Moffat and Larbalestier that since the ro and a" phas es w ere n e v e r observed together in th e mi c rostructure th e n which e v e r phase form e d first exc lud e d th e oth er phase from forming. Th e r e hav e been a only a f e w studies that hav e inv es tigated th e formati on of the ro-related phas es in ternary Nb-Ti-Al based alloys [36-38]. Th e study by Strychor e t al. [36] was th e first to show that ordered derivatives of th e ro phas e w e r e form e d in Ti 3 Al + 5Nb to l 7Nb (at.%) alloys. Th e ro-related phase that form e d in th ese wat e r qu e n c l1 e d samples was d e t ec t e d only in th e SAED patt e rns of the B2 matrix. Th ese diffraction patt e rn s s how e d diffu se st r e aking and e xtra diffra c tion maxima at l/3{111} 13 position s that w ere attributed t o an o rd e r e d d e rivativ e of th e ro-related phase. The effect of is o thermal aging for diff e r e nt tim es at 400 C and 500C wa s

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38 fo und t o s l o wly d egra d e th e tw ee d patt e rn o f t h e B2 m a trix and t o int e nsify th e r e fl ec ti o ns of th e o rd e r e d ro -r e lat e d phas e in appropriat e SAED patt e rns Stry c bor e t al d e t e rmin e d that t h e ord e r e d co -r e lat e d ph ase had th e B8 2 stru c tur e whi c h co nsi s t e d of th e P6 3 / mmc spa ce group Th e ro-B8 2 stru c tur e wa s found to hav e A 2 B o rd e ring wh e r e i t wa s assum e d th a t th e A r e pr e s e nt e d a mixtur e o f Ti and Nb a t o m s and th e B r e pr ese nt e d th e Al at o m Th e transf o rm a ti o n m ec l1anism of th e ro -B8 2 phas e from th e B2 phas e in a Ti 4 Al 3 Nb all oy w as inv es tigat e d by B e nd e r s ky e t al [ 3 7) Thi s s tudy s h o w e d that furn ace coo ling fr o m 1400 C r es ult e d in th e tr a nsf o rmati o n o f th e B2 matrix co mpl e t e ly t o th e t r ig o nal ro "-typ e pha se. This ro "-typ e pha se wa s d e t e rmin e d by C BED t o hav e th e P 3 ml spa ce gr o up It wa s d e du ce d fr o m mi c r o stru c tural e vid e n ce t h a t th e B2 pha se w as s tabl e at 1400 C, a nd that t h e s ub se qu e n t h ea t tr ea tm e n t, o f a furna ce co ol e d s ampl e f o r 26 day s a t 700 C r e sult e d in th e transformati o n o f th e roB8 2 pha se fr o m th e ro matrix Th e s it e occ upan c i es o f th e ro" and ro -B8 2 pha ses w e r e d e t e rmin e d b y X-ray diffr ac ti o n Thi s analysi s s h o w e d that th e Nb a to m s fl o w e d out of s it es on th e co llap se d d o ubl e lay e rs and int o sit e s on tl1 e singl e l a y e r s. 'l~ e va c a te d s it es o n t h e d o ubl e l a y e r s w e r e th e n pr e f ere ntially o cc upi e d b y Ti and Al at o m s, and th e s it es o n th e s ingl e lay e r s w e r e e nri c h e d by a mixttrr e o f Nb Ti and Al a to m s. Th e driving f o r ce for th e c h A mi c al e x c hang e b e tw ee n th e s e at o ms wa s s h o wn t o b e fr o m t h e mrutimizati o n o f th e Ti-Al b o nd s o n th e d o ubl e lay e r s Thi s co n c lusi o n wa s ba se d o n th e fa c t that th e s it es on th e d o ubl e lay e r s had a gr e at e r co ordin a ti o n numb e r o f n e ar est n e ighb o r s than th o s e o n th e s ingl e lay e rs in th e ro and ro -B8 2 phas es. Th e Ti and Al atoms w e r e s hown t o hav e a pr e f e r e n ce f o r th e s it es o n t h e

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39 double layers, since the Ti-Al bonds were the most stable and had the shortest bond length compared to the other possibl e types of bonds. In the study by Bend e rsky et al. [37] th e transformation of the par e nt B2 phase to the final co B8 2 phas e was d esc rib e d by a series of structural changes that involved s ubgroup/symm et ry r e lations in crystallography [43]. Th e transitions that d esc rib e d the B2 to co -B8 2 phas e transformation ar e shown in the schemati c of Figure 2.12 Thes e individual transitions w ere connected together by subgroup/supergroup r e lations that indicat ed how many variants of the product phas e wer e form e d from the par e nt phas e, wh et her the symmetry in the product phase was increased or decreased relative to th e parent phase, and whether there was chemical ordering or wh et h e r distortions had occurred. Th e symmetry and atomic site occupancy r es ults w e r e th e n us ed to sho w that th e transformation path from the B2 phas e to the co-B8 2 phase traversed the t rigonal co" pha se. This trigonal co" phas e had the low est symme try as compared to e ith e r th e B2 or co-B8 2 phas es. Th e transitions fr o m the B2 phase to th e co" phase occ urr e d during furnac e coo ling since th ese transition s involved only th e partial collapse of the doubl e lay e rs and the incompl ete exc hange of atoms b e tw ee n doubl e and single lay e rs How e ver the transition from the co" phas e to th e co-B8 2 phas e r eq uired thermal e n e rgy since it involv e d the full collapse of the d ou bl e lay ers, th e co mpl ete chemical e xchanges that l e d to th e disord e red single lay ers, and the fully populat e d sites o n tl1e double lay e rs by Ti and Al atoms. 2 .6 Th e Ortho/I-Iex Phas es Th e ortho/hex phases are the phases with th e closely r e lat e d HCP and o rthorhombic structures that have b ee n obs er ved in the Ti-rich alloys of the Nb-Ti-Al system. Th ese phas es are form e d from alloys with compositions that show the or

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P6/rnmm Chemically Disordered Im3m A2: bee 40 Chemical Ordering [2] Chemically Ordered Homogeneous Distortion [4] Pm3m B2: CsCl R3m '--. Homogeneous Distortion [4] Chemical Ordering [2] c'= 2c~ w-Ti tit.... w-Collapse [3] R3m Completion of Collapse [2) P3rnl Trigonal-w w-Co llaps e [ 3] Chemical Ordering Disordering between Single Layers [2] c'= 2c / (2] P3ml w' w" Figure 2.12. Shows the transformation paths from the J3 phase to the ro-related phases using subgro up and symmetry relations from Bendersky et al. [37].

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41 B2 phas e to b e stab l e a t high t e mp e ratur es. This stable~ or B2 phas e ca n th e n be qu e nch e d and r et ain e d at room t e mp e ratur e, o r it can und e rgo a vari e ty of transformations to diff e r e nt H C P and orthorhombic phas es d e p e nding on the co mpo s iti o n h ea t treatment, and coo ling rat e. 1'bu s, th e studies o f th ese diff ere nt H C P and orthorhombic phas es are divid e d int o two diff e r e nt groups : th e disord e r e d st ruci-ur es of th e a. a.' and a." pl1a ses in binary Nb-Ti alloys and th e ordered str u ct ur es of the a. 2 a. 2 and O (o rthorhombic ) phases in t e rnary Ti 3 Al + Nb alloys with co n s tant Al co nt e nt This division fa c ilitat es a comparison 01 th e influ e n ces of Nb and Al o n the str uctur e and transformation o f th ese phas es in th e Ti-rich all oys 2. 6.1 Dis o rd e r e d Stru ctt u es Th e HCP and orthorhombic phas es that w e r e observed in th e binary Nb-Ti alloy s co nsist e d o f the e quilibrium a. phase and the m etast abl e a.' and a." mart ens ii e phas es. Both e quilibrium and m e tastabl e hav e dis o rd e r e d site occupancies b et w ee n the Nb and Ti atoms. Th e structure of th e e quilibrium a. phas e wa s d e t e rmin e d t o b e the H C P st ructur e and the P6 3 / mm c space group (10] Th e m etas tabl e a.' phas e wa s d ete rmin e d to h ave f or m e d in ste ad of the a. phase wh e n a mart e nsitic transform a ti o n occ urr e d by rapid coo ling from the high temperature~ phase in Ti-bas e d alloy s co nt a ining up t o :::: 7at.%Nb [72 74]. Th e s tructur e and th e l attice param ete rs o f th e a.' phas e w e r e id e ntical t o th e a. phas e. For binary alloy co mpositions gr ea t e r than :::: 7 at. %Nb the a." phas e was d e t e rmin e d t o hav e form e d instead o f the a. phas e by a si milar mart e nsit,i c transformation from th e phase [72). Th e a." pha se was found to h ave the C -c e nt e r ed o rthorhombi c structure and C m c m space group (7 5]. Th e orthorhombic s tru c tur e o f the a." phas e r es ult e d from distorti ons b e tw een the Nb-Ti bonds. Th ese distortions b eca m e m o r e noti cea bl e in binary alloys that

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42 co ntain e d mor e tl1an 7 at. %Nb [72]. Th e rang e in compositions for the distortions in the a and b latti ce parameters, as m e a s ur e d by th e alb ratio incr ease d with the Nb co nt e nt from 0 578 for ~ 7at.%Nb to 0 654 for 20at.%Nb [74]. It was also d e t e rmin ed that the orthorhombic st ru ct ur e o f th e a." phase was produ ce d from th e a.'-HCP str u c tur e by varying the position s o f the atoms on th e Wyckoff sites in th e a." pha se [72]. Th e r e wer e f o ur lattic e sites in th e C mcm s pa ce gr o up of th e a." structure tl1at co uld b e d esc rib e d by the 4 c Wyckoff site [43]. Th e atomic positi o ns of th ese 4 c sites w e re (0, y, ); (0, y, ); (, y+ ); and( y+ ). Th e distortion in the o rthorh o mbic s tructur e was indi cated by they param ete r which was ~ o.2 for Ti-Nb alloys [76] Th ey parameter for tb e H C P structure in Ti-Nb all o y s wa s 0 1667 Thi s value f o r th e y parameter of the H C P s tr1.1 c tur e wa s bas e d only on the alb rati o, and would pr o du ce th e ideal alb rati o of 0.578 for th e undi sto rt e d HCP str u ct ur e. Th e temperature of the mart e nsiti c transformation for th e a.' and a." phases was d e t e rmin e d by J e pson e t al [74] and Moffat and Larbal es ti e r [72] to b e d epe nd e nt on the Nb co nt e nt o f th e binary Nb-Ti alloys Th e stt1dy by J e pson e t al s how e d tl1at th e ~ temperature dropp e d rapidly from ~ 850 C in Ti to ~3 00 C in th e Ti-17. a t.%Nb alloy. Th e s tudy by Moffat and Larbalesti e r showed that the Ms temperature f e ll b e low room temperature in alloy s that co ntain e d mor e than 30at %Nb. Thi s conclusion was based o n th e observation that th e retain e d f3 pha se was the only pha se that wa s observed in alloy o f tl1 ese co mpositi o ns after wat e r quenching from 1000 c. J e pson e t al also s how e d that the t e mp e ratur e was a fun c tion of coo ling r ate in the binary Nb-Ti alloys It w as f o und that rapid coo ling rates co uld suppress

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43 th e mart e nsiti c transf o rmation of the a" phas e From this r es ult it was s ugg este d that th e a" structw~e c ould form b ot h isoth e rmally and martensitically. 2.6.2 Ordered Stru c tur es Th e e quilibrium a. 2 and O pha ses and m e tastabl e a. 2 phas e form in t er nary Ti 3 Al + Nb all o ys with o rd e r e d structures. It was d e t e rmin e d that th e a. 2 and a. 2 phases formed in Ti 3 Al + Nb alloy with l ess than~ llat.%Nb [24 26 27]. Th e a. 2 phase ha s th e H C P st ru ct ur e with th e Ti 3 Al stoichiometry and P6 3 / mm c s pac e group [10] Th e ordering b e tw ee n Ti and Al in th e a. 2 structure ca us es th e a axis lattic e param e t e r to b e twice th e a-axis latti ce param e t e r o f the di so rd e r e d a-Ti s tru ct ur e. Th e a. 2 phas e has th e same HCP structure and lattic e param e t e rs as th e a. 2 phas e, but has b ee n shown to form by a mart e nsitic transformati o n during rapid cooling from th e high t e mp e ratur e f3 phas e (24 26 27]. Th e O phas e forms in Ti 3 Al + Nb alloy s that c ontain mor e than ~ 12at.%Nb [13 24-32]. Th e O phase has th e C ce nt e r e d o rth o rhombi c s tru c tur e with th e Ti 2 A1Nb stoichiometry and Cmcm s pac e group [13] In Ti 3 Al + 0 1 la t. %Nb alloy s, the a. 2 and a. 2 phase s ca n both form from the high temperature f3 o r B2 phas e during cooling. How e v e r o nly th e e quilibrium a. 2 pha se can form during i so th e rmal h ea ting of alloys co ntaining the r e tain e d f3 phas e. In fact the a. 2 and a. 2 phas es ar e so coo ling rat e d e p e nd e nt, that it is pos si ble to r eta in th e high temperature disord e r e d f3 phas e that e xists over th ese co mpo s iti o n ranges by rapid qu enc hing. Thi s disordered f3 phase th e n o rd e r s to t h e B2 ph ase at l o w e r t e mp era tur es in all oys that co ntain e d mor e than ~ 5at .% Nb (24] Th e m etas tabl e a. 2 pha se u s ually i 'o rms by mart e nsitic transformation for mo st coo ling rates. In co mp ar i son, the e quilibrium a. 2 phas e form s by nucl ea ti o n and growth processes o nly a t high temperatures and for e xt e nd e d p e riods o f time (27 ,3 2 ,35 ]. Thi s

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44 eq uilibrium a. 2 pha se can form from th e p phas e as grain boundary allotriomorphs intragranular plat es, and e quiax e d grains dep e nding on th e alloy composition and thermomechanical history Th e ~ temperature o f the mart e nsitic transformation has b een shown t o d ec r ease rapidly in the Ti 3 Al + Nb alloys with in c r eas ing Nb content (24 ,2 6 27 36]. In the binary Ti 3 Al all oy, th e mart e nsitic tra nsformation was found to b e impossibl e to s uppr ess by rapid coo ling (77] Th e a. 2 phas e that f o rm e d in this mart ens iti c t-ransformation occurred by first forming th e a' phas e and th e n ordering to the a. 2 phas e. In t er nary Ti 3 Al + Nb alloy s, th e a. 2 mart e nsiti c transformati o n co uld b e s uppr essed co mpl e t e ly pr o vid e d that the cooling rate wa s rapid e nough A v e ry high co oling rat e, such as by s plat qu e nching wa s n ecess ary to suppress th e mart e nsiti c transformation in a Ti 3 Al + 5at. %Nb alloy but a l o w e r cooling rat e co uld achieve th e sa m e r es ult in alloy s with high e r Nb co nt e nt (24]. F o r the alloy compositions and coo ling rates that s h o w e d the occurrence o f th e a. 2 mart e nsjtic r e a ctio n it w as s ugg este d that pl ates f o rm e d with the a' phas e first and th e n later ordered t o th e a. 2 pha se (27]. Th e d eve lopm e nt o f thls transformation sequence was bas e d on the obse rv a tion of midrib s and anti-phase domain b o undari es ( APDBs ) in the pl ates. TI 1ese APDB s w ere th e sa m e as th ose t hat w e re o bs erve d in the a t o a. 2 o rd er ing rea ct ion o f binary Ti 3 Al based alloys (78]. Th e formati on o f th e O pha se in z Ti 3 Al + Nb alloys with Nb co nt e nts of z 12 to 30at. %Nb r e quire s thermaJ activation since wat e r qu e nching th ese alloys from high t e mperatures r eta in s tl1e p phas e with the B2 structure at room t e mp e ratur e (13 ,24 35]. Th e r e for e, th e O pha se observed in th ese studies has b ee n found to form o nly by s l o wly coo ling from th e high temperature p or B2 phase or by isoth e rmally h eat ing

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45 the retainP.d p phas e The ma.in points concerning the O phase that w e r e d ete rmin e d from th ese studies w e r e th e O phas e was structurally relat e d to th e a 2 phas e (13], the 0 phas e could form from th e B2 phas e by a shearing mechanism that involv e d th e rmally ac t i vat ed processes (27 ,2 9] the O phas e could form from th e B2 pha se b y a co mpo s iti on invariant transformation [31] th e transformation of the O phas e co uld b e d esc rib e d c rystallographi c ally (33 34] and th e O phas e could e xist with two differ e nt atomic site occupations [31 35]. Th e d e tails of th ese points ar e covered in th e foll o wing discussion Th e O-Ti 2 AINb phas e was found to b e structurally related to th e a 2 -Ti 3 Al. phas e and to involv e additional ternary ordering that caused small orthorhombic dist o rtions [13] Th e structural r e lationship b e tw ee n thes e two phases is shown in the sc h e matic o f Figur e 2.13 Th e atomic site occupancy of th e a. 2 -Ti 3 Al. phas e s hown in Figur e 2.13a co n s i sts of Ti atoms on the 6h Wy c koff sites and Al. atoms on the 2d Wyckoff sites [43]. Kr o nitz er e t al. [47] showed that the addition of Nb to th e a. 2 phase, in am o unts that r e main e d in so lid solution pr e f e r e ntially occ upi e d th e 6h sites with Ti at o ms. H o w e v e r in Ti 3 Al. + Nb alloys that co ntain e d mor e than ~ 12at %Nb, th e O Ti 2 Al.Nb phas e was found to form inst e ad of the a 2 phase (13 24-33]. Th e 0-Ti~Nb structure shown in Figur e 2.13b involv e d further t e rnary ordering of the a. 2 phas e that ca us e d the unit cell to b e distort e d in th e dir ec tion of th e a-axis and baxis latti ce paramet e rs whi c h are co ntain e d in th e ( 001 ) planes. In th e O-Ti 2 Al.Nb st ru c tur e, the Ti Nb and Al. atoms w e r e d e t e rmin e d to pr e dominantly occupy three different latti ce sites. Th ese w e r e the 8g Wyckoff s it e by Ti th e 4cl Wy c koff site by Nb and th e 4c2 Wy c koff site by Al. [13 28]. Moz e r e t al [28] p e rform e d a s tru ct ural refinement using a sa mpl e that had the composition ofTi-25at.%Al.-25at.%Nb to

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I I I I I I I I I I (a) I ;~~f :-:. I I 46 (b) Figure 2.13. 6h Wyckoff Site: Ti 2d Wyckoff Site : Al . . . ........ . tt~rr --::'.'.-' b a 4c 1 Wyckoff Sit.e: Nb 8g Wyckoff Site: Ti 4c 2 Wyckoff Site: Al Sh o w s the r e lati onship between the crystal str u ctures of the a. 2 -Ti 3 Al ph ase and the O-Ti 2 A1Nb pha se. (a) the a. 2 -Ti 3 Al phas e ( P6 3 / mmc space group); (b) the O-Ti 2 A1Nb pha se (C m c m space gr o up ). Th e dark shaded atoms are at z 0 and the light atoms are at z c

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47 determine t h e r elative ato mi c occupancies of the thr ee diff e r e nt Wyckoff sites in th e 0 -Ti~ s tru ct ur e. Th e r es ult s o f this s tudy showed that the Nb atoms occ upi e d 18% of th e 8g sites whi c h w e r e Ti-ri c h and that th e Ti atoms occupied 18% of the 4cl s it es which w e r e Nb-ri c h Th e 4 c2 site was found to b e occ upi e d by o nly Al atoms. Th e e ff ect of t h e Nb co nt e nt o n the distortions in th e O-Ti 2 A1Nb phas e wa s stu di e d by K es tn e rW ey kamp et al [26]. Th ese distortions w e r e m eas ur ed by the alb r atio and w e r e f o und to in c r e a se from ~ 0.632 f o r th e O-Ti 2 A1Nb phas e co nt a ining 20a t.% Nb to ~ 0 .645 for th e O-Ti 2 AlNb pha se co ntaining 30at.%Nb. Th e 0-phas e has b ee n found to f o rm from th e P pha se as plat es by a latti ce shear m ec hanism in t h e s tudi es by K es tner-Weykamp [27] and B e nd e rsky e t al. [29] K est n e r-W e ykamp ex amin e d plat es that form e d during air coo ling from th e P phas e, whi c h wa s present at 1250 C, in a Ti 3 Al + 20at %Nb alloy [27]. Th ese plat es co ntain e d d e f ect str u c tur es that consis t e d o f midrib s, co lumnar APDBs and {110} twin s aligned r o ughl y parallel to the midrib From th e analy s is of th ese d e f ects, the f3 to O tran s f 'or mati o n was d e termined to b e by a latti ce invariant shear m ec hanism that ini ti ally form e d the o rtl1orl1 o mbi c a" s h ear pr o du ct with the plat e shape an d th e n la te r or d e r e d to th e O s tru ct ur e. A s ubs e qu e nt diffu s ional growth m ec hani s m wa s used to exp l a in t h e thickening o f t h e plates and the f o rmation of th e co lumnar s hap e d APDBs. Th e f o rmati on of th e O ph ase fr o m th e r e tain e d B2 phas e during i sot h e rmal h e ating o f two alloys w as inv es tigat e d by B e nd e r s ky et al [29] Th e two all oys e xamined in thi s st udy had co mp os iti o n s of Ti-12.2Al-37.2Nb and Ti-23 .9 Al 25Nb (a t.%). F o llowin g th e h e at treatment at 700 C f or 26 day s, th e O pha se wa s obse rv e d as plat es in th e Ti 12.2Al-37 2Nb all oy and as e qui axe d grains in the 1'i-23 .9 Al-25Nb

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48 alloy. The plates were determin e d by energy dispersive X-ray spectroscopy (EDX) analysis to hav e th e 1'i~Nb composition. This co mposition of the plates was diff e r e nt than th e composition of the B2 matrix which was Ti-10Al-45Nb (at.%) composition. Howev e r th e transformation of' the plates was still described using the minimization of elastic strain e nergy approach which was consistent with th e ph e nomonological theory of mart e nsitic transformations [79]. This th eo ry accurately pr e dicted th e habit plan e and 1~igid body rotation of the plates using the meas1.1red lattice parameters of the B2 and O phases in this alloy. It was determin e d for the Ti-23.9Al-25Nb alloy that th e formation of th e e quiaxed grains involved two steps: initially the r et ained B2 phase transform e d completely to a highly faulted O phas e and then later recry st alliz e d into fault fre e grains. The f o rmation of the O phas e from the B2 phas e in a Ti-24Al-15Nb (at.%) all oy was d esc rib e d by a co mposition invariant transformation m ec hanism in th e study by Mural ee dharan e t al. [31]. This alloy was heat treated for short times that Jast e d from one to sixty minut e s at t e mp e ratures of 800 C, 900C and 950 C. The shorter aging times and low er temperatures favor e d th e complete transformation of th e r eta.i n e d B2 phas e to the O phas e without a change in composition. The O phase formed as plates whi c h co ntained complex defect structures determined to be coarse APDBs with the displacement v ec tor of 1/4(110] fin e APDBs with th e displacement vector of l/2[100] and stacking fault s with th e displac e m e nt v ec tor of 1/10(025] Th e coarse APDBs o bs e rv e d in th e O plat es wer e shown to have the same size and a related displac e m e nt vector to th e APDBs that wer e observed in th e retain e d B2 matrix prior to aging Th e pr ese nc e of these APDBs wjtb no variance in co mpositi o n s upp o rt e d th e co n c lusion in this study that th e transformation from the B2 phas e to

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49 th e O pha se occ urr e d by a s h e ar typ e m ec hani s m. How e v e r this transformation m ec l1ani sm did r e quir e s hort-rang e diffusion since th e fin e APDBs observed in the 0 plat es w e r e f o rm e d fr o m an ordering r eac tion that r e quir e d atomi c exc hang es between latti ce s it es. B e nd e rsky et al. (33 34] d esc rib e d the formation o f th e O phas e from th e high te mp erat1. rr e J3 phase using a c ry s tallo gra phi c mod e l bas e d o n a se qu e n ce o f str u c tural changes r e lat e d by ub gro up and s ymm e try r e lati o ns. Thr ee all oys with co mpo si t ions o f Ti-25Al-12 5Nb Ti-25Al-25Nb and Ti-28Al-22Nb (at.%) w e r e e xamin e d after h ea ting at 1100 C for four days. Thi s h eat tr e atm e nt caused the partiti o nl ess transformation o f pl ates in these three all oys. Th e analysis s how e d differ e nt d e f ect st ru ct ur es in th e plat es o f th e Ti-25Al-12.5Nb composition as co mpar e d t o th ose in t h e two Nb-ri c h all oys. Fr o m the identification o f these d e f ect structures s uch as APDBs and stac king faults, th e transformation o f the plat es in th e Ti-25Al-12.5Nb a ll oy w as s h o wn to h ave occ urr e d from t h e di so rd e r e d J3 phase whil e t h e transformation of the plates that formed in the two Nb-ri c h alloy s occ urr e d from the B2 pha se. Th ese r es ult s w e r e th e n us e d t o s how that th e plat es pr ese nt in these a ll oys form ed al o ng t wo different transformation paths as s hown in Figur e 2.14. Th e plates in the Ti-25Al-12 .5 Nb alloy f o rm e d fr o m th e dis o rd e r e d J3 (lm3m ) phas e and tl1en follow e d th e path that pass e d through th e int e rm e diat e HCP structures b e for e r eachi ng the final O st ru c tur e. In co mpari so n the plat es observed in th e two Nb-ri c h alloys form e d from the B2 phas e a nd then follow e d th e path through th e B 19 st ru ct ur e to the final O st ru c tur e Th ere wa s n o str u c tural co nfirmation o f' th e interm e diat e transitional s tru ct ur es in eac h o f th ese t w o pa t hs. H o w e v e r e vid e n ce 01

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f 3 "" Pm~m (B2) 14 / mmm 50 ~ 3 i ~2 ~ ,. ~ P4/mmm Fmmm P6:,lmmc (A3) ~2 ~2 3-----~ ~ .... -----' Cmmm Cmcm (A20) P6:,lmmc (1)() 19 ) ~ 2 ~, + Pmma (B19) 2 t 3 Cmcm (~BC) Figur e 2 14. Shows th e transformation paths from th e f3 phase to the O phas e using subgroup and symmetry r e lations from Bend e rsky e t al. [33].

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51 the individual transitions was obtained from th e analysis of th e d e f ec t s tru c tur es in the plat es and from the partitionl ess natur e of th e transformation in th ese all o y s. Th e O phas e has b ee n found t o e xist in tw o structural forms with diff e r e nt atomic s it e occupancies a cco rding to st udi es by Mural ee dharan e t al. [31 ,35 ]. In th e first st udy by Mural ee dharan e t al in 1992 th e O phas e was form e d at th e low e r t e mp e ratur e o f 800 C in the Ti 24Al-15Nb alloy and was d e t e rmin e d to hav e the co mmonly obse rv ed s it e occ upan cy of Ti o n the 8g Wy c k o ff site, Al on the 4cl Wy c k off s it e, and Nb on t h e 4c2 Wyck off s it e. H o w eve r th e O pha se that form e d at th e s lightly high e r temperature of 900 C showed a diff e r e nt s it e occupancy. This O pha se wa s d ete rmin e d to have a rand o m occ upan c y of Nb and Ti on th e 8g and 4c2 Wy c k off s it es, whil e Al still occ upi e d th e 4cl Wyckoff site. ln a r ece nt study by Mural ee dharan e i al. in 199 5, th e sa m e r es ult s showing two diff e r e nt s it e occ upan c i es f o r the O phas e w ere o btain e d from a series o f h ea t tr e atments that w e r e co ndu c t e d o n TI-27.5a t.% Al all oys with up to 25at .% Nb additions. In b ot h studies, the s it e occ upan c i es of the two diff e r e nt O pha ses w e r e d e t e rmin e d from int e n s ity variations between r e fl ect i o n s in the C BED patt er ns and by c hann e lling e nhanc ed mi c r oa naly s i s. Th e C BED analysis wa s co ndu cte d in thin r e gi o ns of th e O pha se t o minimi ze the dynami cal sca tt e ring e ff ec ts Th e order param ete r ( S ) wa s d e fin e d in terms of the Ti and Nb site occupation o f th e 8g and 4 c 2 Wy c koff sites. This parameter was c al c ulat e d by thermodynamic analysis and s h o w e d that a rand o m s it e occ upan cy between Ti and Nb atoms o n th ese t w o l a tti ce s it es was stab iliz e d at high e r temperatures. Thus it was s ugg este d in b o th o f these studies that Al s tabiliz e d th e disordered a." mart e n s it e str u ct ur e Thi s di so rd ered a." mart e n site str u ct ur e is a m etasta bl e phas e thai f o rm s in the binary Nb-Ti s y s t e m and then

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52 o rd e rs into th e ordered O phas e whi c h is an e quilibrium phase that forms in the ternary Nb-Ti-Al syste m.

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CHAPTER3 EXPERIMENTAL PROCEDURES 3.1 Mat e rial Th e co mp os itions of th e all o y s u se d in this inv e stigation w e r e sp ec ifi e d t o Pr a t t and Whitn e y wh o th e n manufa ct ur e d th e all o y s. Out o f a t o tal of t e n alloy c ompo s ition s originally manufa c tur e d th e r es ult s o f thr ee all o y s w e r e us e d in t lli s s tudy Th e nominal c ompositi o ns o f th e s e thr ee all o ys ( r e f e rr e d t o a s alloy s 2 3 and 4 ) ar e giv e n in Tabl e 3 .1 Th e main c rit e rion for s e l e cting th e thr ee all o y c ompositions was th a t e a c h c ontain th e B CC J3 phas e in th e mi c r os tru c tur e s at high t e mp e ratur e s. Th e r e f o r e, th e all oy co mp os ition s w e r e c h o s e n ba se d o n th e 1200 C i s oth e rm shown in Figur e 2.2 wlri c h wa s d e v e lop e d fr o m pr e vi o us s tudi e s o f th e t e rnary Nb Ti-Al pha se diagram [17-20]. Thu s i t w as e xp ec t e d that at 1200 C th e mi c ro s tru c tur es of alloy 3 s h o uld co nsi st of a s ingl e J3 ph ase and alloy s 2 and 4 s h o uld co n s i s t o f th e J3 + x pl1a ses wh e r e xi s th e y pha se ( all o y 2 ) or th e cr pha se ( alloy 4 ) On e o f th e purp oses of t hi s inv es tigati o n wa s t o th e n co nstru c t th e pha se e quilibria for th e alloys at high e r t e m pe r a tur es than 1200 Th e alloy s w e r e s uppli e d by Pratt and Whitn e y in th e form of 200 gram a r c -m e lt e d s ampl es Th e as-r ece iv e d s ampl e s had b ee n ar c m e lt e d a total o f f o ur t o s i x tim es t o e n s ur e co mpl e t e c h e mi c al mixing. Thi s wa s v e rifi e d by c omp os it io n lin e p ro fil es p e rf o rm e d o n c r oss s ect ion s o f th e a s -r ece iv e d sampl e s using th e El ec tr o n Mi c r o pr o b e Analyz e r (E MP A ) Th e s e pr o fil e s did n o t s how s ignifi c ant c h e mi c al 53

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54 Tabl e 3.1. Th e n o minal compositions of the as-received alloys and the co mpo s iti ons determined by mi c roprob e analysis o f th e r e -ar c melted alloys. Analysis Composition (at.%) Alloy M et hod Nb Ti Al Nominal 27 33 40 2 Microprob e 26 8 0 2) 33 8 0.2) 39.3 0.2 ) Nominal 50 40 10 3 Mi c roprob e 49.8 1.0) 40.6 ( 0. 7) 9.57 0.3 ) Nominal 42 28 30 4 Microprob e 41.4 ( 0.8 ) 29 .5 ( 0.3) 29 1 0.5 )

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55 inhomog e n e iti es from top to bottom or from center to outer e dg e. How e v e r, th e mi c rostru c tural characterization o f th e as-r ece iv e d alloy samp l es did show inh o m oge n eous mi cros tru c tur es that w e r e observed primarily along th e thickn ess dir ec ti o n o f th e cast sa mpl es, i. e. from top to bottom. This inhomogen e ity was most probably du e to un e v e n solidification rat es b e tw ee n th e surfac e making contact with the wat e r coo l e d Cu plates of th e ar c m e lt e r and th e untou c h e d top surface of th e samples. Du e to th e significant nont1niformiti es that w e re observed in th e as -r eceived mi c r os tru c tures a proc e dur e was adopt e d that involv e d r e -ar c m e lting fragm e nts oi' th e 200 gram samples into smaller 3 gram samp l es. Th e r e -ar c m e lting was p e rf o rm ed at the University of Fl o rida using a noncons umabl e tungst e n e l ec trod e under pressurized flowing argon gas. Th e fragm e nt e d pi eces w e r e placed in cavities on a wat e r cooled cop p e r bas e plat e and r e -ar c m e lt e d (RAM) a total of at l e ast 4 times t o e nsur e co mpl ete mixing in case th e r e w e re so m e chemical inhom o g e n ei ti es b e tw ee n th e fragm e nts Th e molt e n samp l es took approximat e ly 2 to 3 seconds to so lidify and r ea ch e d a temperature that gav e th e m a metallic lu stre. This proc e dur e r e duc e d th e microstru ct ural inhom o g e n e iti es of alloys 2 3 and 4. Th e co mpositi o n s of the r e -ar c m e lt e d samples of alloys 2 3 and 4 w e r e analyz e d by e l ec tron mi c r o prob e and ar e giv e n in Tabl e 3 1 Th e r e sults obta.ined by th e microprob e a nalysis for th e r e arc m e lt e d alloy s are c l o s e to th e n o minal compositions of th e as-r ece ived alloys. Th e int e rstitial oxygen and nitrog e n content of' th e as-r ece iv e d alloy 2 was d ete rmin e d by w et c l1 e mi ca l analysis at T e l e dyn e Wah C hang Albany (IW C A). Th e

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56 results of thls analysis showed that th e oxygen content was ~490 parts p e r million (ppm) and the nitrog e n content was ~48 ppm. 3.2 H e at Treatm e nts Th e h e at tr ea tment experiments were conducted with two types of furna ces d e p e nding on wh e th er a fast or slow cooling rat e was n ee d e d. Th e two furnac es used w ere a CM mod e l 1600 vertical tub e furnac e and a Vacuum Industri es hlgh va c uum furnac e Th e vertical tub e furnace was used wh e n fast cooling rates wer e n ee ded such as during water qu e nchlng. Thls furnac e consisted of a mullit e tub e that wa s s urround e d by MoSi 2 h ea ting elements. The mullite tub e was sealed at the top and bottom with wat e rcoo l e d r e movabl e fixtures. Wh e n the fixtur es w e re closed at both e nds th e tub e could be pr e ssuriz e d slightly with a flowing stream of argon gas that e ntered at th e top and exi ted at the bottom fixtur es. Th e top fixtur e was d esigne d to a llow botl1 a sample and a typ e B thermocouple to b e position e d within the h eati ng zo n e o f th e furna ce. Th e thermocouple was us e d to monitor the temp er ature during th e h ea t tr e atment and was also us e d to calibrate th e h ea ting zone. Th e c alibration s h o w e d that the h eat ing zone was constant to within over the tube l en gth o f cm from the center of the furnac e. The sample was placed on an alumina b oat whlch was suspended in the heating zone with molybd e num or tungst e n wir e Th e temperature diff e r ence betw ee n th e sa mpl e and th e thermocoupl e was estimated to b e l ess than since th e heating zone of th e furnac e was radially l1niform due to the cy lindri ca l d es ign. Following i s the typical e xp e rim e ntal proc e dur e: ( 1 ) Ramp the furnace at~ 10 C/ min t o th e l1 eat treatment t e mp e ratur e. ( 2) In se rt the sa mpl e into the h e ating zone and seal the top fixtur e.

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57 ( 3 ) H o ld th e s ampl e und e r flowing argon for th e duration of th e h e at tr e atm e nt ( 4 ) Op e n th e b o ttom fixtur e, c ut th e wir e, and allow th e sampl e to fall int o th e qu e n c hing m e dia. Wat e r was us e d as th e qu e n c hing m e dia in this study in o rd e r to o btain rapid coo lin g r a t e s. How e v e r th e r e w e r e a f e w h e at tr e atm e nts that w e r e p e rform e d in th e v e rtical tub e furna ce in which th e sampl e w a s simply dropp e d from th e h e ating z on e ont o th e bott o m fixtur e This p e rmitt e d th e s ampl e s to b e air cool e d which r e pr e s e nt e d a c ooling rat e that was int e rm e diat e b e tw ee n wat e r qu e nching ( fast ) and furna ce c o o ling (s l o w ). Th e va c uum furna ce wa s us e d in e xp e rim e nt s that n ee d e d a s low coo lin g r a t e. Thi s furna ce co nsi s t e d of an alumina c ru c ibl e that was surr o und e d by a tantalum r es i s tiv e h e ating c ag e. A typ e R th e rmocoupl e was us e d to m o nit o r th e t e mp e ratur e and was p o sition e d 3c m fr o m th e s ampl e in th e h e ating z on e Th e furna ce in co rp o rat e d a diffu s ion pump capp e d by a wat e r baffl e, s o that th e furnac e w as c apabl e of obtaining a pr es sur e of 1 t o 4 x 10 6 Torr d e p e nding on th e h e at tr e atm e nt te mp e ratur e. F o llowing is th e typical proc e dur e for using th e va c uum furna ce: ( 1 ) Pla ce th e s ampl e on th e alumina c ru c ibl e ( 2 ) Pla ce th e b e ll jar o v e r th e s ampl e and pump to th e b a s e vacuum pr e s s ur e. ( 4 ) Ramp th e furna ce t o th e h e at tr e atm e nt t e mp e ratur e at ~10 /min ( 5 ) H o ld at th e h e at tr e atm e nt t e mp e ratur e f o r th e duration of th e e xp e rim e nt ( 6 ) Turn th e p o w e r o ff to th e tra ns f o rm e r and allow th e sampl e t o coo l to r oo m t e mp e ratur e insid e th e furna ce ( 7 ) V e nt t o atm os ph e ri c pr e s s ur e wh e n th e t e mp e ratur e, monitor e d by th e th e rm oco upl e, i s b e low ~ 200 C.

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58 This pro ce dur e involv e d a slow h e ating rat e for th e sampl e up to th e aging te mp e ratur e s sin ce a high va c uum had to b e obtain e d prior to h e ating th e s ampl e. Th e t e mp e ratur e wa s monitor e d with th e thermocoupl e during furna c e c ooling and foll o w e d a paraboli c c urv e for th e h e at t r e atm e nt s 3. 2.1 L o ng-T e rm H eat Tr e atm e nt Exp e rim e nt s Tabl e 3.2 li s t s th e l o ng-t e rm h e at tr e atm e nt c onditions of sampl e s that w e r e u se d in thi s study Th e list in c lud es th e initial c ast c onditi o n, th e h e at tr e atm e nt param e t e rs and th e c o o ling m e thod that was e mploy e d. Mo s t of th e h e at tr e atm e nts that wer e us e d in this study w e r e conduct e d with 3 gram RAM sampl es using th e v e rtical tub e furnac e Th e s e sampl e s w e r e h e at tr ea t e d in th e asc ast RAM c ondition for 2 to 12 hours d e p e nding on th e t e mperatur e, and th e n wat e r qu e nch e d. H e at tr e atm e nts conduct e d ab o v e ~ 1200 C v a ri e d fr o m 2 t o 4 h o ur s in duration whil e th o s e th a t w e r e c ondu c t e d b e low this t e mp e ratur e w e r e e i t h e 1 12 or 16 hours. Thr ee o f th e h e at tr e atm e nts w e r e p e rf o rm e d with mat e rial that was c ut from t h e a s -r ece iv e d 200 gram arc-m e lt e d s ampl e of all o y 4 Th e thr ee sampl es u se d in th e s e h e at tr e atm e nt e xp e riment s w e r e first heat e d to 1550 C for 2 hours in th e v e rti c al tub e furnac e and th e n air c o o l e d. Furth e r h e at tr e atm e nts w e r e th e n co ndu c t e d on t w o of th e s e sAmpl e s: on e wa s h e at tr e at e d at 1515 C f o r 2 h o ur s and air coo l e d and th e oth e r was h e at tr e at e d at 1000 C for 16 hours and air co o l e d. Th e r e w e r e fiv e h e at tr e atm e nt e xp e rim e nt s that w e r e p e rform e d with th e v ac uum furnac e W it h th e e x ce ption o f th e all o y 2 sampl e which wa s h e at tr e at e d at 1200 C and furna ce c o o l e d th e s e e xp e rim e nts w e r e d e sign e d to inv e stigat e th e e ff ec t that a sl o w co oling rat e had on th e mi c rostru c tur e s of alloys 2 and 4. Th e 3 gr a m

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59 RAM sAmples were used in all of these heat treatments and all of these heat treatments last e d for a duration of 4 hours 3.2.2 Short-Term Heat Treatment Experiments Tabl e 3. 3 lists a set of 3 gram RAM samples that wer e h ea t treated for very s hort tim es lasting from 2 to 5 minut es. Thes e heat treatm e nts were conducted in the vertical tub e furnace and wat e r qu e nched. Thes e heat treatments wer e d esi gn e d to inv est igat e th e evo lution of th e high temperatur e phase equilibria. 3. 3 C haract e rization T ec hniques Th e microstru c tur es of the cast and h eat treat e d samples w e r e inv es tigat e d using optical micros copy and transmission electron microscopy (TEM) techniqu es. Op t ical microscopy was used primarily for the macroscopic characterization of the samples du e to its low magnification capabilities. A Nikon microscop e and Leica microscop e w ere us e d in this study. Th e TEM was us e d exte nsiv e ]y for the microscopic analysis of the samples du e to its combined diffraction and imag e a nalysis capabilities whil e complimented with co mposition analysis using Energy Di spe rsiv e Sp ectroscopy (EDS) o n a submicron mi c rostru ct ural scal e. Thr ee microscop es w e r e us e d over th e co urs e of this inv e stigation: a JEOL 200CX ASTEM ( AnalyticaJ S c anning Transmission Electron Mi c roscop e) and a JEOL 4000FX TEM lo c ated at the University of Florida FL and a JEOL 2000FX ASTEM locat e d at th e N e w York Stat e College of Ceramics at Alfred Univ e rsity NY. Th e analysis of mat e rials using the TEM has b ee n widely employed over th e past 30 y ears. Th e m e thods that w e r e pr e dominantly used in this inv est igation included selected area e l ectro n diffraction (SAED), co nv e rg e nt beam electron diffraction (C BED ), and amplitude co ntrast imag e formation Th e SAED analysis w as

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6 0 Tab l e 3. 2. L o n g t e rm h e at tr e atm e nts us e d fo r alloy s 2 a n d 4. I nitial T e m p. Tim e C oo ling F urna ce Alloy C o n diti o n M e t h o d Typ e 2 RAM 1 5 00 C 2 h ours W at e r Qu e nc h Tub e 2 RAM 1 4 00 C 4 h o ur s W ate r Qu e nc h Tub e 2 RAM 1 4 00 C 4 h ours Furn ac e C oo l Vacu um 2 R AM 1 3 00 C 4 h o ur s W ater Qu e n c h Tub e 2 R AM 1200 c 4 h o ur s Furn ace C oo l Vacu um 2 R AM 6 00 C 1 2 h o ur s W at e r Qu enc h Tub e 2 R AM 4 00 C 1 2 hour s W at e r Qu e n c h Tub e 4 AR 1550 C 2 h ours Air C oo l Tub e 4 AR 1 550 C 1 h our W at e r Qu ench Tub e 4 AR 1 5 1 5 C 2 h ours Air Coo l Tub e 4 RAM 1 4 00 C 4 h o ur s W at e r Q u e nch Tub e 4 RAM 1 4 00 4 h ours F ur nac e Coo l Vac u um 4 RAM 1 3 00 C 4 h ours W at e r Qu e nc h Tub e 4 R AM 1 3 00 C 4 h ours F ur nace Coo l Vacuum 4 RAM 1 2 0 0 c 4 h ours W at e r Qu e n c h Tub e 4 RAM 1 2 00 c 4 h ours F ur nace Coo l V acu t1 m 4* AR 1 00 0 c 16 hours Air Coo l Tube not e : AR As R ece iv e d 2 00 gram ar c m e l t e d samp l e. RAM R e -Ar c M e l t e d 3 gr::im sa m p l e. so l uti o niz e d at 1 550 C f o r 2 h ours a nd air coo l ed.

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61 Tabl e 3.3. Sh ort-t e rm h eat treatments u sed f' or alloys 2 and 4. Alloy Initial T emp. Tim e Coo ling Furnace Co nditi on M et h od Typ e 2 RAM 1200 c 2 min. W ater Quench Tube 2 RAM 1200 c 5 m i n Wat e r Qu e n c h Tube 4 RAM 1200 c 2 min Wat er Quench Tube 4 RAM 1200 c 5 mjn_ Wat er Quench Tube 4 RAM 1000 c 2 min Wat er Qu e n c h Tube

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62 us e d f'or routin e pha se identification and crystallographic information s uch as the analysis o f' twins orientation r e l at ionship s and s ta c king fattlts. Th e co nverg e nt b e am e l ect ron diffra c tion (CBED) technique was us e d for d e tail e d phas e analy s i s wh e r e symmetry inf o rmation wa s r e quir e d. This m e thod e nabled the crystal point group and the space group o f' th e ph ase to be d e t e rmin e d. Th ese imag es w ere ob tain e d using amplitude contrast m e thods bas e d on two-b ea m and multipl e -b ea m (i.e. n ear Lau e zone axes) conditions. This m e thod was also used to obtain e ith e r brigh t fi e ld o r dark field imag es d e p e nding on the b e am tilt co nditions Th e SAED p atte rns that w e r e used f o r phas e id e ntification w e r e m e asur e d with a Starr ett m eas uring table. Thi s instrument p e rmitt e d both lin e ar and angular m eas ur e m e nts to b e mad e with accuracies of .025mm for lin e ar m e asurem e nt s and .083 f or angular measurements. 3.4 Sampl e Pr e paration 3. 4.1 Optical Mi c r osco py Th e s ampl e pr e parati o n that was us e d for opti ca l mi c roscopy consisted o f sta ndard m e tallogr a phi c m e th o d s. Th e 3 gram RAM samples w e r e m ec hani ca lly sectio n e d into 0. 5 to 1. Omro thick sa mpl es u s ing a diam o nd e dg e d c utting wh ee l. Thes e sa mpl es w e r e then mount e d in 1in diam ete r m o ld s using ph e n o li c p o wd e r A p o li shed s urfa ce o f 0 03m wa s o bt a in e d u s ing standard grinding and poli s hing m et h ods. Th e grinding steps w ere d o n e using 240 to 600 grit Si C pap e r Th e p o li s hing steps w e r e performed with Al 2 0 3 powd e r on a ppropriat e c l o ths. Finally the sa mpl es w e r e e t c h e d u s ing Kr o ll' s etc hant

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63 3.4.2 Transmission El ec tron Microscopy Th e sample pr e paration for TEM used th e jet-polishing m e thod. This m e thod consisted of a m ec hanical pr e paration step and an e l ec tro chemical polishing step. Tl1 e mechani ca l pr e paration step involv e d cutting a wafer from th e bulk s ampl e, c utting a 3mm diam e t e r di sc from th e waf e r and finally r e ducing the thi c kn ess of th e 3mm disc to about 0.2mm (200m). Th e starting bttlk sample was ei th e r a 3 gram RAM sample or o ughly a 1cm x 1cm x 1 c m as-r ece iv e d (AR) sample. Care was taken in this proc e dur e to select TEM samples that w e r e r e pr ese nt a tiv e o f the bulk sample. In g e n e ral waf e rs that w e r e c ut n e ar th e mid-section of the bulk s ampl e w e r e selected for th e next step of obtaining th e 3mm disc. It was n ecess ary to obtain th e 3mm dis c from the center of the waf e r in order to avoid th e h e at aff ec t e d zo n e (HA V) that was observed by optical mi c r oco py in several h ea t tr ea t e d samples. Th e m e thod that wa s e mploy e d in obtaining the 3mm disc was diff e r e nt than that whi c h i s custo marily u se d s in ce the l1 e at tr ea t e d samples o ft e n fractur e d in a brittle mann e r during th e preparation wh e n using th e hol e pun c h or the ultr aso nic di sc c utt e r. Th e alt e rnativ e procedur e consisted of mounting th e waf e r using superglue o n t h e e nd of a ~ 2.8rnm diam e t e r s tainless steel r o d. Th e e dg e s of th e waf e r w e r e then smoothed down by m ec hanical grinding on 600 grit SiC paper until a circular 3mm dis c wa s obtained. Aft e r obtaining th e 3mm di sc, th e thickn ess was th e n r e du ce d by m ec hanical grinding m e thods t o a thickness of ~200 to 300m using SiC grit paper. Th e e l ec troc h e mical polishing step us e d th e j e t-polishing techniqu e in 01der t o o btain an e l ect r o n transparent r egio n n ear t h e ce nt e r of' th e 3mm di sc sample. Thi s procedure wa s co ndu c t e d with a Stru ers T e nupol j e t-p o lish e r and an e l ec trolyti c so lution Th e e l ect r o lyt e that gave the b es t p o lishing r es ults was ba se d on

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64 4 % Hf + 10%H 2 S0 4 + 86%Metba nol. Th e following polishing param e t e rs gav e c on s i s t e ntly g o od r es ult s : -50 C to -40 C t e mp e ratur e rang e, 25 volts D C, and m o d e rat e t o low fl o w r a t e. Aft e r a h o l e d e v e Jop e d in th e s ampl e, th e s ampl e w as qui c kly r e mov e d fr o m th e p o lish e r and rin se d in two successiv e methanol bath s

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C HAPTER4 EQUILIBRIUM PHASE TRANSFORMATION STUDY 4 1 Intr o du c ti o n Th e lit e ratur e r e vi e w in C hapt e r 2 indi ca t e d th e c ompl e xity o f th e ph ase e quilibria in th e Nb-Ti-Al syst e m As was pr e vi o usly s tat e d th e for e m os t r e a so n f o r th e co mpl e xity of thi s syst e m is that th e r e ar e a multitud e of e quilibrium binary pha ses and at l e a s t t w o t e rnary phas e s whi c h e xi s t in thi s t e rnary s ys te m Thi s co mbin e d with th e fa c t that th e so lidus t e mp e r a tur e o v e r m o st of th e t e rn a ry s y s t e m li es a b o v e ::: 1500 C, m e an s that most o f th e co mpr e h e nsiv e phas e e quilibri a s tudi es hav e co n ce ntr a t e d o n o n e o r tw o t e mp e ratur es f o r h e at t r e atm e nt s Th e mo s t co mm o n t e mp e r a tur e has b ee n 1200 A furth e r co mpli c ation is that a numb e r o f m e t as tabl e pha ses, s u c h a s th e mart e nsiti c a' a 2 and a" and th e m-r e lat e d ph ases, i rm in co mp e titi o n with th e e quilibrium pha ses in thi s s y s t e m Th e pr ese n ce o f th ese m e tastabl e phas e s c an compli c at e th e d e v e lopm e nt of th e e quilibrium mi c r os tru c tur es and may c au se c onfu s ion in th e e quilibrium pha se analy s i s Th e r e f o r e, th e purpo se o f c h a p te r 4 i s t w o -f o ld : t o d e t e rmin e th e e quilibrium ph ases and t o d es crib e th e ir d e v e l o pm e nt into th e e quilibrium microstru c tur es in a ll o y s 2 and 4. Th e as-c a s t mi c r os tru c tur e o f a third all o y ( alloy 3 ) with a c omp os iti o n o f 50Nb-40Ti-10Al (a t .%) wa s inv es tig a t e d in o rd e r t o co mpar e prop e rti es o f th e B CC pha se t o that o f all o y s 2 and 4 Thi s backgr o und inf o rmati o n on th e '3 ph ase will t h e n pr o vid e th e ba s i s f o r inv es ti g ati o n of th e m e ta s tabl e pha se s in c h a pt e r s 5 and 6 65

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66 The results of this chapter are divided into three sections: (1) the as-cast (2) the long-term heat treatment, and (3) the short-term heat treatment microstructures. Using these results, the discussion of this chapter is then divided into three main parts: (1) high temperature f3-phase (2) equilibrium microstructures at the aging temperatures and (3) equilibrium phase transformation mechanisms. 4.2 Results 4.2.1 As-Cast Microstructures The analysis of as-cast microstructures consisted of' re-arc melted (RAM) samples of alloys 2 3 and 4. The cooling rate associated with these samples was relatively fast due to their small size whicl1 minimized inhomogeneous microstructures in the RAM samples. However the RAM sample of alloy 2 showed a microstructure with an inhomogeneous distribution of precipitates. Therefore a RAM sample of alloy 2 was also analyzed which was electromagnetically (EM) levitated and drop quenched in order to suppress the solid-state precipitation by rapid solidification. 4. 2.1.1 Re-arc Melted The as cast microstructures that were observed in the RAM samples of alloys 2 3 and 4 are shown in Figure 4.1. These microstructures consisted of large primary grains with a coarse dendritic structure. The size of the primary grains was typically >300m (0.3mm). In alloy 2, an inhomogeneous distribution of acicular precipitates was observed near the gra.in boundaries and within the interdendritic regions (Fig11re 4. la). Alloy 4 (Figure 4. lc) occasionally showed a second phase at the grain boundaries however alloy 3 (Figure 4. lb) showed no additional phases.

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67 (a) 50m (b) 50m Figur e 4.1. Optical micrographs showing th e as-cast microstructur es. (a) all oy 2; (b) alloy 3; (continued)

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68 (c) 50m Figure 4.1. (continued) (c) alloy 4.

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69 4 2.1.1 1 Th e Primary Phas e Th e primary phas e in th e RAM samp l e s of alloys 2 3 and 4 was e xamin e d by TEM and was d e t e rmin e d to hav e th e BCC s tru c tur e, known as th e f3 pha se S e l e ct e d ar e a e lectron diffra c tion (SAED) patt e rns showing th e [001 ] 13 z on e axis fr o m th e f3 matrix in e a c h all o y ar e shown in Figur e 4 2 Th e patt e rns show diffraction spots at th e {100} p o sitions f o r alloys 2 and 4 (Figur es 4.2a and 4.2c r e sp e ctiv e ly ); how e v e r t h e on e for alloy 3 do e s not sh o w th ese spots (Figure 4 2b) Th e {100} spots d e not e th a t tl1 e f3 pha se in all o y s 2 and 4 ha s an o rd e r e d B C C stru c tur e, o r B2 (CsC l ) str u c tur e, and ar e r e f e rr e d t o a s sup er latti ce r e fl ec tion s Th e fact that n o s up e rlatti ce r e fl ec ti o ns ar e obs e rv e d in alloy 3 indicat e s that th e f3 phas e has a disord e r e d stru c tur e. Mi c r o graph s that w e r e obtain e d using a two-b e am c ondition with a ( 100 ) 13 s up e rlatti ce r e fl ec tion ar e shown in Figur e 4.3 Th ese micrographs sh o w th e pr ese n ce of l arg e anti-pha se domain boundari e s (APDBs) in th e matrix of alloys 2 and 4. Th e APDBs ar e form e d during a disord e r t o ord e r transition and indi c at e that o rd e ring occ urr e d during s olid s tat e c o oling. Th e B2 pha se that was o b se rv e d in all o y s 2 and 4 e xhibit e d s e v e ral diffus e sc att e ring anomali es Th e r e sult s showing th e s e anomali e s ar e group e d into thr ee c at e g o ri es : diffus e s tr e aking splitting of diffra c tion spots and localiz e d diffus e int e nsity maxima. A tw ee d s tru c tur e that c orr e l at e d with th e diffus e str e akin g and s p o t splitting was al s o obs e rv e d. Th e diffu se s tr e aking and th e s plitting o f diffra c ti o n sp o ts ar e b e st ob se rv e d at th e [001] 13 z on e axi s, a s sh o wn in Figur e 4.2a f o r alloy 2 and 4.2 c for all o y 4 Th e s tr e aking w a s c ontinuous in th e < 110 > dir e ctions and int e rs e ct e d both fundam e ntal

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70 (a) (b) Figure 4.2 SAED patterns showing the [001] zone axis of the f3 matrix (a) alloy 2 ; (b) alloy 3; (continued)

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71 (c) Figure 4.2 (continued) (c) alloy 4.

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72 -( a ) ii#' 0 2m (b) 0 2m Figur e 4. 3 TEM rru c rographs s h o wing th e APDBs in th e B2 matrix ( a ) all o y 2 ; (b ) all oy 4

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73 and s up e rlattic e r e fl ec tions. Th e splitting o f diffraction spots was obs e rv e d a s s at e llit e r e fl e ctions that w e r e displa ce d from th e r ec iprocal latti ce position in < 100> dir ec tions Th e s e paration of th e split r e fl ec tions incr e as e d as th e o rd e r of r e fl e ction in c r e as e d Thi s is d e m o n s trat e d in Figur e 4 4 using a SAED patt e rn that wa s tilt e d o ff th e [00 l] p z o n e axis al o ng th e g= ( l 10 ) r e fl ec ti o n Splitting i s o bs e rv e d for r e fl ec ti o n s al o ng b o th th e [100] and [010] dir ec tion s and is gr e at e r for th e ( 400 ) sp o t as co mpar e d to th e ( 200 ) sp o t. Al so, n o tic e that f o r th e ( 2"20) spot th e r e ar e f o ur s at e llit e r e fl ec tion s. Th e s e ar e co mpri se d of two pairs o f r e fl ec ti o ns that ar e split al o ng orthogonal [100] and [O 10] dir ec tions. Th e diffu se e l e ctron scatt e ring is b e st o bs e rv e d in SAED patt e rns o f th e < llO> p and < lll >p z on e ax es, a s shown in Figur e 4.5. Th e s c att e ring c onsists o f locali ze d se gm e nts o f diffus e int e nsity that e xt e nd in <112> or <110> dir e ctions and hav e a maxima l oca t e d b e tw ee n th e B C C diffraction spots at fra c tional c oordinat es. Th e diffu se int e n s i ty i s s up e rimp ose d o n th e c ontinuou s s tr e aking that is al s o o b se rv e d in th e <112 > and <110> dir ec tion s Th e fra c tiona l co o rdinat e s for int e nsity maxima in all o y 2 w e r e at 1 / s 110 >, 1 /s < 112> <112> and 1 /s p os itions in th e p zon e axis (Figur e 4 5a ) and at 1 / s and 1 / s p os iti o ns in th e <111 > 13 zo n e axi s (Figur e 4.5b ) Th e int e nsity maxima positions ob s erv e d in all o y 4 w e r e th e s am e as thos e obs e rv e d in alloy 2 with th e e xc e pti o n of th e 1/s p o si t ions that w e r e n o t o bs e rv e d at e ith e r th e < llO> p zon e axis (Figur e 4 5c) o r th e < lll >p zo n e axi s (Figur e 4.5d ) Tl1 e t w ee d s tru ct ur e co nsi s t e d o f s triation s that li e d parall e l to {110} tr a c es of th e ord e r e d J3 ph a s e, a s shown in Figur e 4.6 Th e mi c rograph was o btain e d n e ar th e [OOl] p z o n e axi s to s h o w th e tw ee d s triations along two o rth o gonal ( 110) and ( 110 )

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74 Figur e 4 4 SAED pattern from the B2 matrix showing the splitting o f di ffract ion spots. Th e specimen was tilt e d away fr o m the (001] zone axis along g= ( llO ).

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75 (a) (b) Figure 4.5. SAED patterns showing the diffuse electron scattering observ e d in th e B2 matrix of alloy 2 (a and b) and alloy 4 (c and d). (a) [110] zon e axis ; (b) (111] zone axis ; (continued)

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76 (c) (d) Figure 4.5. (continued) (c) [110] zone axis; (d) [111] zo n e axis

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77 O.lm Figur e 4 6. TEM micro graphs showing the tweed microstructure in the B2 matrix

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78 planes. The image characteristics of the tweed structure was investigated using two -b eam amplitude contrast conditions, and was found to obey the g R =O invisibility criterion (where g is the reflection vector and R is a general description of a displacement vector) and to depend on the deviation parameter s. For example, a two-beam condition using g=(lOO) caused both (1 10 ) and (ITO) striations to b e visibl e, but using g=(llO) caused only the (110) striation to be visible. The (ITO) striations are invisible using g=(llO) since gR =O, assuming R =( lTO ). Likewise the deviation parameter, s affected the tweed image causing it to hav e a coarse appearanc e for a small s magnitude and a fine appearance for a larg e s magnitude. 4.2.1.1. 2 The Pr ecipitates in Alloy 2 The inhomogeneous distribution of precipitates observed in the as-cast microstructure of alloy 2 were divided into three representative areas: the in-matrix, the grain boundary and the int er dendritic regions. The prim ary difference between the in-matrix and th e interdendritic regions was the presence of acicular shaped pr eci pitates which w e r e observed by optical microscopy as seen in Figure 4.1. Th e in-matrix region consisted of a high number density of very small pr ec ipitates as can be seen in Figure 4.3b The precipitates were homogeneously distributed within the matrix and yet were not affected by the APDBs that had formed during the J3 phase disorder/order transition. The analysis of the small precipitates identified them as being related to the class of ro -phas es, a c lo se-packed hexagonal structure [58] which will b e discussed furth e r in c hapt er 5. A TEM micrograph of th e grain boundary region i s shown in Figure 4. 7. This r e gion consisted of B2 grains witl1 the y phase based on the L1 0 tetragonal structure with TiAl stoichiometry distributed along the grain boundaries. Two morphologi es

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79 w e r e observed for th e y grains: a blocky type that formed along the B2 grain boundaries and a lath type that extended from the blocky grains into the B2 matrix. The formation of Widmanstlitten laths from grain boundary allotriomorphs [80] resembles this type of microstructur e. Th e laths wer e observed to hav e an orientation r e lationship with the B2 matrix which was determined to be as follows: 13 and {lll} 'Y II {110} 13 and is shown in th e SAED pattern in Figure 4 7b. Th e grain boundary all ot riomorphs contai n e d stac king faults that form ed on the {lll}r plan es Micrographs that ar e representative of the interdendritic regions are shown in Figur e 4.8. Th ese r egio ns co nsisted of B2 matrix with l arg e acicular shaped pr ec ipitates as well as small in-matrix m pr ec ipitat es Th e analysis of the acicular shaped pr ec ipitat es showed that th ey w e re plates and that they had an orthorhombic st ructur e (referred to from this point on as p l ates). Th e ir size was typically observed to be> lOm in length (1) and <0.5m in thickn ess (t), giving them an aspect ratio >20(1/t). Th e small m pr ec ipitat es w ere observed to b e homog e n eo usly distribut e d about the plat es. H o w e v e r occasionally larg e m pr ec ipitat es w e r e observed in co nta ct with the plat es, as sho wn in Figur e 4.8b. A d e tail e d ana l ysis of th e plates will b e cove r e d in c hapt er 6. 4.2.1.2 EM Levitat ed and Drop Qu e n c h e d Th e l e vitat ed and drop qu e nch e d sample o f alloy 2 was found to consist of an acicular mierostructur e, as shown in Figur e 4 9 The optical micrograph in Figur e 4.9a s how e d that th e microstructur e had a bask et w ea v e appearance. Th e TEM mi c r ogra ph in Figur e 4. 9b rev e al e d that a high numb e r d e nsity of l e nticular shaped plates had formed in the ordered 13 matrix. Th e ordered 13 matrix was also found t o

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80 (a) 0.2m B2 (b) Figure 4. 7 Show s the microstru cture of the as-cast sample of alloy 2. (a) TEM micrograph showing the grain boundary allotriomorphs and Widmanstatt e n laths of the y TiAl phas e; (b) SAED pattern showing the orientation r e lationship observed b et w ee n th ey laths and B2 matrix which was <110]., II <111> 13 and {111}., II {I10} 13

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81 (a) l.Om (b) 1000A Figur e 4. 8. TEM micro graphs of' the B2 matrix in the as cast sample of alloy 2. (a) the sma ll co-related precipitates and l e nticular-shaped plates ; (b) the coarse co -r e lated precipitat es adjacent to the plat e.

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\ I, \ I I I I < \ I \ I 82 (a) 20m (b) 1.0m Figur e 4.9 Shows the acicular microstructur e observed in the EM-levitated and drop quenched sample of alloy 2. (a) Optical micrograph ; (b) TEM mi crograph.

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----83 have contained APDBs. The analysis of these plates showed that they wer e the same as thos e that were observed in Figure 4.8 for the RAM sample. 4. 2. 2 Long Term Heat Tr e atments 4.2.2 1 Analysis of Alloy 2 Th e long-t erm h e at treatm e nts used for alloy 2 were shown in Table 3.2. Th e 1500 C h eat treatm en t lasted 2 hours while th e 1400C 1300 C, and 1200 C h eat treatments w e re for a duration of 4 hours. All of the heat treatments us ed a wat e r quench to cool the samples to room temperature except for the 1200C treatment which us e d a furnac e cool 4.2 2.1.1 Optical Microscopy Micrographs that are r e pr ese ntative of the microstructures observed in samples aged at 1300 C, 1400 C, and 1500 C and subsequently water quench ed are s l1own in Figure 4.10. Not e that the microstru c tures do not show any evidence of th e prior d e ndritic structure and that th e heat treatm e nts applied were sufficient to remove chemical inhomogeneities. The microstructures of the samples aged above 1300C all showed a matrix that r ese mbled tl1e acicular microstructure observed in the interdendritic regions of the as-cast sample. Th ese microstructures showed basket weave morphologi es as shown in Figures 4 .1 0a to 4.10c. Th e grain boundaries observed in th ese micr ostr uctur es indicated that th e matrix consisted of a single phase at high temperatures and had a grain size of 1 to 2mm. The e ntir e microstructur es of 1400 C and 1500 C aged samples consisted of the acicular microstructur e Howev e r the sample aged at 1300 C also showed large blocky-shaped second phas e particles that formed within th e matrix and at grain boundaries. A two-phase microstructur e was

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84 (a) 20m (b) 20m Figure 4 10 Optical micrographs showing the microstructur es of the long-term thermally aged samples of alloy 2. (a) 1500C-4hrs-WQ ; (b) 1400 -4hrs WQ ; (co ntinued)

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85 (c) lOm ( d ) lOm Figur e 4 10 ( co ntinu e d ) ( c) 1300 -4hrs-WQ ; ( d ) 1200 C4hr s -F C

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86 also o bs e rv e d in th e s ampl e ag e d at 1200 C and furnac e cooled and is shown in Figur e 4 10d 4 2 2.1.2 Mi c roprob e Analysis Th e r es ults o f th e microprob e analysis on th e samples that w e r e h e at treated at 1400 1300 and 1200 C ar e shown in Tabl e 4.1. This table showed that the cr phase whi c h h as the binary Nb~ sto i c hiom e try has a larg e solubility rang e for Ti Th e cr pha se wa s d ete rmin e d to contain 26%Ti at 1300 C and 30%Ti ( at .%) at 1200 Th e Al co nt e nt was not found to hav e changed appreciably in th e cr pha se t h at wa s pr ese nt at these t e mp e ratur es. Th e r e for e, th e stoichiometry of the cr pha se remained c los e to th e ratio of (Nb+Ti)~. Simi1arly th ey phas e, which has the binary TiAl stoichiometry, was d ete rmin e d to hav e a larg e solubility rang e f o r Nb that was found to b e ~ 21at %Nb. Th e Al content of th ey phas e was found to b e 43at. %Al which was l o w compared to th e co mp os ition of the binary TiAl pl1a se 4.2.2.1.3 Tran s mi ssio n Electron Mi c r osco py Th e general a pp e aranc es of the microstru ct ur es obs e rv e d by TEM in samples aged a t 1500 1400 C and 1300 C and then wat e r qu e nch e d ar e shown in Figur e 4.11. A co mmon observation mad e for all three h eat treatments was that the matr ix co ntain e d a larg e numb e r d e nsity of plates Th e matrix was id e ntifi e d as th e o rd e r e d j3 (B2 ) phas e APBD s w e r e obs e rv e d and w e r e consistently smaller in the aged s ampl es as co mpar e d to th e as-cast sample. Assuming that the APDBs form e d up o n coo ling the s mall er size was most probably du e to th e solid-state coo ling rate whi c h w as faster in th e aged sa mpl es b ecause of the wat e r qu e nching Th e analysi s of t h e plates d ete rmin ed that th e y w e r e si milar t o thos e that w e r e o bs er v e d in th e as-cast sa mpl e (see Figur e 4.8).

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87 Table 4.1. Microprobe results of the aged samp l es of alloy 2. T e mp e rature Phase Nb (at.%) Ti (at.%) Al (at.%) 1400 C J3 26.83 0.19) 33.85 0.18) 39.32 0.17) J3 23.06 0 06 ) 36.90 0 .07) 40.4 0.11) 1300 C er 36. 39 0 .1 8) 26. 35 0 15) 37.26 0.24) er 34.06 0.80) 29.85 0.41) 36.10 0.42) 1200 c 'Y 20. 71 2.39) 36.56 1 53) 42. 73 1.22)

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88 Figur e 4 11 TEM micrographs sh o wing th e micro s tructur es o f th e l o ng-t e rm th e rmally ag e d sampl e s o f al l oy 2 ( a ) 1400 4hr s WQ ; (b ) 1300 4hr s -WQ ( a ) lm (b ) 1.5.m

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89 Th e larg e parti c l es in th e two-phas e microstructur e of th e s ampl e ag e d at 1300 C w e r e id e ntifi e d as th e cr-(Nb 2 Af + Ti) phas e and ar e shown in Figur e 4. llb. Th e cr phas e wa s d ete rmined by SAED analysi s t o hav e a tetragonal unit c e ll with latti ce param e t e rs o f a = 9.92.A and c = 5 17 A ( 0.05.A ) Th e s e r e sults ar e con s ist e nt with th e cr pha se tha t f o rms in a vari e ty of transition m e tal alloy syst e m s [39]. Th e cr pha se ha s b ee n s h o wn to b e a co mpl e x t e trag o nal s tructur e co ntaining 30 at o ms and p ossess ing so m e d e gr ee of ord e r b e tw ee n 5 distin c t cry s tallographi c sit e s [39 41]. 1"b e f o ll o win g c har acte ri st i cs o f th e cr grain s w e r e d e t e rmin e d from this analy s i s : ( 1 ) a grain s iz e o f up t o 10-15m ; (2 ) a low numb e r d e nsity o f stacking faul ts on ( 100 ), ( 010 ), and ( 110) plan es; and ( 3) th e pr e s e n ce o f low angl e grain bound a ri es. Th e t w o -pha se microstru c tur e o f th e sampl e ag e d at 1200 C and furna ce c o o l e d is s h o wn in Figur e 4.12 Th e grains that hav e light c ontrast in th e s e mi c r o graphs w e r e id e ntifi e d from th e SAED analy s is t o b e th ey pha se whil e th e gr a in s that hav e d a rk c ontra s t w e r e id e ntifi e d t o b e th e cr phas e Th ey pha se was o b se rv e d to h a v e a l a th morphol o gy and w a s s urr o und e d by a continuously co nn ecte d m a tri x th a t w as th e cr ph ase. A l o ngitudinal vi e w o f th e y pha se lath s i s s h o wn in Figur e 4 12a and a t ransv e rs e vi e w of th e m is shown in Figur e 4.12b. Figur e 4. 12 a s h o w s that e a c h la t h co nsist e d of a co lony o f small grain s o f th e sam e ph ase. Th e gr ain si ze o f e a c h phas e was obs e rv e d to b e l ess than ~ 3m. Ann e aling twin s w e r e o b se rv e d in many o f th e r grain s and c an b e s ee n in th e normal ori e ntati o n. Th e twin s y s t e m wa s d e t e rmin e d to b e <112]{111} whi c h wa s th e s am e as that r e p o rt e d f o r m ec hani c al twins in th ey phas e [81] A diffra c ti on s tudy wa s p e rf o rm e d on t h e y phas e t o d e t e rmin e th e s it e occ upan cy o f Nb in t h e unit ce ll s in ce t h e mi c r o pr o b e r es ult s in Tabl e 4 1 s h o w e d t h a t

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90 it co ntain e d up to ~21at %Nb. Th e r es ult s from the diffraction study are shown in Figur e 4 13 Th e analysis of the tw o diffra ct i o n patt e rns showed that th e (ITO) and ( 001 ) spots w e r e observed at th e (110] zone axis (Figur e 4.13a) but that th e (TOI) and ( 010 ) spots w e r e absent at th e [101] zo n e axis (Figur e 4.13b). Th e structure factor rul e for diffracti o n in th e binary TiA1 structure pr e dicts that diffra c tion i s forbidd en and will not be observed when (h+k)=odd. Sinc e the h+k sum is odd for th e (TOI) and ( 010 ) r e fl ect ion s and is e v e n (zero is includ e d) for th e (ITO) and (001), then th ese results ar e consistent with th e structure fa c tor of th e binary y TiA1 phas e. Thus thi s indi c at es that Nb randomly occupies the Ti sites in th ey phas e of alloy 2. Tabl e 4.2 co ntains a summary of th e r es ults obtained by TEM for the h eat treated samples of alloy 2. It includ es the phas es that w e r e observed and a bri e f co mm e nt about eac h. 4 2.2.2 Analysis of Alloy 4 H ea t tr eat m e nts that last e d from 2 to 16 hours with a vari e ty of diff e r e nt coo ling m et hods w ere us e d o n alloy 4. Th e following ar e the h e at treatm e nt param e t e r s for alloy 4 samples from Table 3 2: (1) two h ea t treatments at 1550 C and 1515 C for 2 hour s with air cooling; ( 2 ) four h ea t tr ea tm e nts which includ e d one at 1550 C for 2 h o ur s and thr ee at 1400 C, 1300 C, and 1200 C for 4 hours with water quenching; (3) thr ee h ea t tr ea tm e nt s a t 1400 1300 and 1200 C f'or 4 l1our s with furnace coo ling ; (4) o n e h ea t tr e atm e nt at 1000 C for 16 hours with air coo ling ; and (5) two h eat treatments at 600 C and 400 C for 12 hour s with wat e r qu e n c hing 4 2.2 2 1 Optical Mi c roscopy Th e typi c al mi c rostru c tur es of samples that w e r e h ea t tr ea t e d at high temperatures ar e s h o wn in Figur e 4.14. Th e prior d e ndriti c structure of th e as-cast

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91 ( a ) 0.5m (b ) 1.0m Figur e 4 12. TEM mi c r o graph s s h o wing th e y lath s in th e cr + y mi c rostru c tur e o f th e th e rm a lly ag e d 1200 -4br s F C sampl e o f alloy 2 ( a ) longitudinal vi e w ; (b ) tran s v e rs e vi e w.

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92 (a) (b) Figure 4.13 SAED patterns of the y phase observed in the 1200 -4hr s-FC samp l e of alloy 2. (a) the [llO] y zone axis; (b) the [lOl] y zone axis

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93 Tabl e 4 2. Summary of th e pha ses id e ntifi e d by TEM in th e ag e d s ampl es o f all oy 2. H e at Tr e atm e nt Pha se s Pr e s e nt C omm e nt s 1500 -2hr s -WQ B2 + plat es APBDs and plat es pr e s e nt in th e B2 matrix 1400 C4hr s -WQ B2 + pl a t es APBD s and plat es pr e s e nt in th e B2 ma t ri x. 1 3 00 C -4hr s WQ B2 + cr + pla tes APBD s and plat es pr ese nt in th e B2 matrix and larg e cr grain s i ze 1200 -4hr s F C cr+ y L a th s tru c tur e a nd s mall cr and r gr a in s i zes.

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94 sample was not o bs e rv e d in any of th e ag e d sampl e s with the e xception of th e 1200 C sample. Th e optical mi c rographs showed microstructur es that evolved from pr ec ipitation of a second pha se from the primary singl e phase that was stable at high t e mp e ratur e s. Ab o v e 1550 C (Figur e 4.14a) th e microstructur e co nsist e d of single pha se grains that had a grain size of 1 to 2mm straight grain boundaries and tripl e p o int s. B e tw ee n 1515 C and 1550 C (Figure 4 14b ), a second phas e f' o rm e d as blocky particl es ( idiom o rphs) within th e matrix and at th e grain boundari es of th e matrix. This second pha se increased in volume fraction at low e r t e mperatur es and b e l o w 1300 C was observed as th e matrix i. e. th e phas e with the larg e r volume fraction(Figur es 4.14 c to 4 14t). In the samples ag e d at 1300 C and 1200 C the r e tain e d primary phas e app e ar e d as an irr eg ular circular and e longat e d morphol ogy, as ca n b e seen in Figur es 4.14d and 4.14 e The two-phas e microstructur e observed in the sa mpl e ag e d at 1000C is shown in Figur e 4.14f. Tl1 e morphology of these second pha se parti c l es from Figur e 4.14f wa s m os tly o f a c ircular s hap e and th e size wa s typi c ally l ess than lm. 4 2.2.2.2 Transmi ss ion El ec tron Mi c ros co py The microstructur e of th e sample ag e d at 1550 C and wat e r qu e nch e d wa s found t o co nsist of the B2 matrix. As shown in Figur e 4.15 the B2 mat-rix consisted of APDBs that w e r e smaller than those observed in th e RAM sample of Figur e 4. 3c. Th e p1~ ese nce of th ese APDBs indicated that th e matrix consisted of the disord e r e d J3 phas e at the 1550 C aging t e mp e ratur e. B e low 1550 C, th e cr phas e was obs e rv e d to hav e nucl e ated from th e J3 phas e to form a tw o -pha se mi c rostru c tur e. Thi s wa s d ete rmin e d from the analysis of the

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0 \ ~ a. 0 95 .. I M\11 t:) -~ .. 0 c p ~ 0 C. \ . ~di ~ c: t:7 e o .. "" ---= ~ :......._,:_ ..1 __ ~ /j ~ ,-> '1_~~~' -------' (a) 50m (b) 20m Figure 4.14. Optical micrographs showing the microstructures of the long-term thermally aged samples of alloy 4. (a) 1550C-2hrs-AC ; (b) 1515C-2hrs AC ; (continued)

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96 (c) 20m .. ( d ) 20m F ig ur e 4 14 (co ntinu e d ) (c) 1400 -4br s -W Q ; (d ) 1 300 -4hr s -W Q ; (co ntinu e d )

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. .. .. .... .. 97 ;, ( Figur e 4 14 (co n t inu e d ) (e) 1200 -4hr s -WQ ; ( f) 1000 -16hrs-A C. (e) lOm ( f) lOm

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98 0.2m Figur e 4.15 TEM mi c rograph s h o wing th e APDBs in th e B2 matrix of th e s ampl e from all o y 4 that wa s h e at tr e at e d at 1550 C and wat e r qu e n c h e d.

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99 samples heat treat ed at 1400C, 1300 C, and 1200 C and subsequently water quenched. Tl1e a grajns were betw ee n 5 and 15m in size and were typically isolated in the~ matrix at 1400 C, but formed a matrix of connected a grains below 1300C. The phas e in th ese microstructures possessed the B2 structure and contained small APDBs in all of the heat treated samples. It also e xhibited the same type of imag e and diffraction c hara cte ristics that hav e b ee n pr e viously document e d for the B2 phase. Th e irreg11lar, e l ongated, and circular morphologi es of the B2 phase which w ere recognized by optical mjcroscopy could b e id e ntified clearly in the TEM as shown in Figure 4.16. This irregularly shaped B2 phase was locat e d primarily at a grain boundaries. In general the irr e gularly shaped B2 particles were retain e d which had formed isolated grains du e to the growth and impingement of a grains with one another. This type of morphology was observed mostly at aging temperatures gr ea t er than 1300C Th e circular and elongated cross sections aros e from the rod-shaped morphology of th e B2 phas e whi c h was commonly observed at aging temperatures below 1300 C. Th e e ff ect of cooling rat e on th e microstructures was studied using th e samP heat treatments at 1550 ; 1400 C, 1300C and 1200 C, but with e ither air cooling or furnace cooling. It was found from these heat tr e atments that the stability of th e high temperature phase to the solid state phas e transformations depended on th e coo ling rat e. Th e 1550 C aged sample was air cooled and was found to consist of an inhomogeneous distribution of small precipitates in th e B2 matrix as shown in Figur e 4.17. Th e analysis of the B2 matrix revealed larg e APDBs a tweed structure, and the same general c hara cte ristics that wer e observed for th e B2 phas e found in the

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100 ( a ) 0 5m (b ) 0.3m Figur e 4.16. TEM mi c rographs sl1 o wing th e diff e r e nt morphologi e s of th e r e tain e d ph ase o bs e rv e d in th e s ampl e s of alloy 4 that w e r e h e a t tr ea t e d a t 1200 C, 1300 C, and 1400 C and wat e r qu e n c h e d ( a ) th e irr e gular and c ir c ular morphologi es; (b) th e e longat e d morphology

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101 0 .5 m Figur e 4.17. TEM mjcrograph s howing the pr ec ipitat es of the orthorhombic phase that formed on th e APDBs and in the B2 matrix of alloy 4 during air coo ling from 1550 C.

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102 as c ast mi c r o stru ct ur e Th e pr ec ipitat e s w e r e o bs e rv e d t o hav e form e d b o th o n th e APDB s and within th e domains o f th e APDB s. A d e nud e d z on e, or r e gi o n that co ntain e d n o pr ec ipitat e s wa s o bs e rv e d adja ce nt t o th e APDBs Th e analy s i s o f th e pr ec ipitat e s th a t f o rm e d in th e B2 matrix during air coo ling s h o w e d th e m t o b e c on s i s t e nt witl 1 th e orth o rh o mbi c pha se whi c h wa s ba sed o n th e Ti~Nb s t oi c hiom e try [13] 1n Figur e 4 18 th e SAED analy s is show e d tha t t h e o rth o rh o mbi c p rec ipitat es f o rm e d fr o m th e J3 phas e with th e following o ri e ntati o n r e l a ti o n s hip : [110] 13 II [001] 0 and ( 112 ) 13 II ( 110) 0 Th e a, b a nd c la tt i ce param e t e rs o f th e orthorhombi c pha se w e r e c al c ulat e d fr o m e l ec tr o n diffra c ti o n patt e rns and w e r e a 0 = 6 13A ( 0 07A ), b 0 = 9 35A ( O llA ) and c 0 = 4 57A ( 0 06A ) Th e cal c ulat e d latti ce param e t e r of th e ord e r e d J3 pha se in t h e 1550 C a g e d s ampl e was ap = 3.23A ( 0 03A ) Th e c -axis latti ce param e t e r o f th e o rth o rh o mbi c pha se was d e t e rmin e d to b e r e lat e d to th e a-axis lattic e param e t e r of th e B2 pha se by c 0 =V2 a 13 Th e s ampl es t hat w e r e ag e d at 1400 C, 1 3 00 C, and 1200 C and furna ce coo l e d s h o w e d th e sa m e tw o -phas e cr + J3 mi c rostru c tur e s that w e r e obs e rv e d in th e wat e r qu e n c h e d s ampl es. H o w e v e r th e furnac e c ool e d sample s s how e d mi c r o stru ct ur es t h a t al s o co ntain e d th e orthorhombi c pha se Th e analysi s s how e d that th e o rth o rh o mbic pha se f o rm e d from th e o rd e r e d J3 pha se with th e sam e ori e ntati o n r e lation s hip that wa s pr e vi o usly shown Th e analysis o f th e particl e s wi t h th e B2 pha se did n o t r e v e al any APDB s, but it was d e t e rmin e d that th e APDBs had nu c l e at e d and grown t o th e siz e of th e parti c l es during furnac e cooling Th e analy sis a ls o s h o w e d that th e o rth o rh o mbi c phas e l1 a d a plat e m o rph o logy and form e d a t th e

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103 00211 11211 @ 2000 @ .. 1100 1100 e @ 110 11 0200 Figure 4 18. SAED pattern showing tl1e orientation relationship b et w een th e orthorhombic phase and the B2 phase in the 1550C-2brs AC sample of alloy 4. Th e OR was consistent with [001] 0 II <110> 13 and (110) 0 II {112} 13

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104 er/p interface as shown in Figure 4. 19. The plates were observed to have contained a low number density of stacking faults that were parallel to the (0 01) 0 planes An attempt was made to observe the domain structure in the orthorhombic plates using dark field imaging techniques, however, this study showed that there were no APDBs pres e nt in these plates. The habit plan e of the plat es was determined at the (001] 0 zo n e axis and was f o und to lie ~9 from th e {2 ll}p planes of the p matrix. There was no apparent dislocation structure at the interface between the plates and the P phase which was exam in e d using conventional d.iffraction contrast imaging techniques. The s~mple that was aged at 1000 C and air coo l ed showed a two-phas e microstructure that cons ist ed of the er and orthorhombic phases. The matrix was composed of larg e er grains with small particl es of the orthorhombic phase at the er grain boundaries and is shown in Figure 4.20. The SAED pattern in Figure 4.20b shows the (110] 0 zone axis of the orthorhombic phase. The inv esti gation of the er matrix in the 1000 C and 1200C aged samples revealed large colonies that consisted of multiple er grains with similar orientations. Evid e nc e of this is shown in Figure 4.21 by tilting a single er grain to the [OOI]cr zone axis. It was also observed that when this zone axis was reach e d many of the surrounding er grains were also oriented close to subsequent [OOl]cr zone axes. This is shown in the micrograph of Figure 4. 2 la as a collection of er grains that all hav e a dark contrast associated with them. The dark contrast is generated by amplitude contTast conditions and signifies that all of the dark grains are oriented for strong diffraction i. e. oriented c l ose to a laue zone axis. This condition was also verified by using a large se l ected ar ea apertur e that sampled several er grain s and is shown in

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105 O lm Figure 4.19. TEM micrograph showing the orthorhombic plates that formed in the retained B2 phase of alloy 4 during furnace cooling from the heat treatment temperatures of 1200 C, 1300 C, and 1400 C.

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106 (a) O.lm .. @@11' t (b) 1l ll lQJ .. .. Figure 4 20 Shows the microstructure observed in the 1000C 16hrs AC sample of alloy 4. (a) TEM micrograph showing particles of the orthorhombic phase observed at the grain boundaries of the cr phase; (b) SAED pattern showing the (110 ] 0 zone axis of the orthorhombic phase

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107 (a) 0.5m (b) Figur e 4.21. Show s the microstructur e observed in the samples of alloy 4 that w e r e h eat treated at 1000 C and 1200 (a) TEM micrograph showing the colony of cr grains; (b) SAED pattern showing multipl e [00 l] a zone axes from the cr grains present in the co l ony.

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108 Figure 4.21b. This diffraction pattern shows several [OOlla zone axes from eac h cr grain. Not e that th ese ar e both slight ly tilt e d and rotated from one anoth e r A summary o f t h e TEM r es ul ts f 'or th e l ong t e rm h eat treat e d samp l es of alloy 4 is giv e n in Tabl e 4.3. It in c lud e s a l i t of th e phas e s that w e r e pr ese nt in th e mi c r ostr u ct ur es and a bri e f co mm e nt about eac h at th e aging t e mp e ratur es. 4 2 3 Sh ort -T e rm H ea t Tr eat m e nts Th ese h eat tr ea tments in c lud e d e xp e rim e nt s that last e d a duration of 2 or 5 minutes and th e n th e sa mpl es w e r e imm e diately wat e r qu e nch ed. Th e t e mperatur e, tim e, and coo ling m e thod e mpl oyed for th e h ea t tr ea tm e nts w ere shown in Tabl e 3.3. TI1 e purpos e of th ese e xp e rim e nt s wa s to und e rstand the m ec hanisms of these pha se transformations that l e d to th e f ormat i o n o f th e two-phase mi c rostru ctures in alloy 2 and 4 4 2 3.1 All o y 2 Th e sa mpl es aged at 1200 C for 2 and 5 minut es co nsi s t e d of two-phase cr + y mi c r ostr u c tures. Ev e n in 2 minut e s th e p phas e of the as cast sample had compl e t e ly transform e d as s h own in Figur e 4. 22. Th e transform e d mi c r ost ructur e n e ar a prior P grain boundary is sho wn in Figur e 4.22a. This figur e shows that th ey grains ca n ge n e rally b e distingui s h e d fr o m th e cr grains du e to a light e r co ntrast from a low e r ma ss cont r as t. Th e r e f o r e, th e lin e of y grains o b serve d with light contrast indi ca t e d t l1 at th ey e i t h e r form e d along th e pri or p grain boundaries during the h e at treatm e nt o r w e r e pr ese nt pri o r to th e h eat tr e atm e nt. Th e latt e r may b e possibl e s in ce y grains w e r e observed at primary p grain boundari es in th e as-cast mi c r os tru ct ur e as wa, s h o wn in Figur e 4 7 Figur e 4.22b s h o w s a cr grain that was purposely tilt e d to th e l l lO] (J ori e ntation in orde r to e xhibit st r o ng diffraction co ntrast. This pr o duced a

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109 T a bl e 4.3 Summ ar y of th e pha ses id e ntifi e d by TEM in th e ag e d sa mpl es o f all oy 4 H e a t Tr eat m e n t Ph ases Pr ese n t Co mm e n ts 1550 -2hr s -WQ B2 APDB s pr e s e nt in th e B2 matrix. 1550 -2hr sA C B2+0 Orth o rh o mbi c plat es f o rm e d o n t h e APDB s in th e B2 m a trix I 400 -4hr sWQ er+ B2 Larg e er grain siz e a nd APBDs pr ese nt in th e B2 parti c l es 1400 -4hr s -F C er+ B2 + 0 Orth o rhombi c pla tes f o rm e d at th e int e rfa ce b e tw ee n th e B2 and er ph ases. 1300 -4hr s -WQ er+ B2 J ., am e llar s tru c tur e and APDB s pr ese nt in th e B2 p ar ti c l es. 1300 -4hr s -F C er+ B2 + 0 Orth o rh o mbi c plat es f o rm e d at th e in te rf ace b e tw ee n th e B2 and er ph ases. 1200 C4hr sWQ er+ B2 LRm e llar s tru ct ur e a nd APDB s pr ese nt in th e B2 parti c l es. 1200 -4hr s-FC er+ B2 + 0 Orth o rh o mbi c plat es f o rm e d a t th e in te rfa ce b e tw ee n th e B2 and er pha se s. 1000 -16hr s -A C er+ O Larg e er grain s i ze with small par t i c l es o f th e o rthorh o mbi c ph ase.

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110 dark contrast for the cr grain which demonstrated that the grain size of the cr phase was larger than ~5m. This was larger than the grain size of the y phase wlri ch was dete rmined to be in the 0.2 to 1.0m size range. A significant difference in the morphology of the cr + y microstructures was observed between the short-term and long-term heat treatments. The short-term heat treatment produced essentia lly isotropic grains of both cr and y phases as shown in Figure 4.22. This morphology was different from that of the long-t erm heat treatment which produced laths consisting of rows of cr and y grain8. This long-term heat treatment morphology was previously shown by optica l microscopy in Figure 4.10d and by TEM in Figure 4.12. The view that was normal to the cr + y lath microstructure in Figure 4.12b resembled that which was observed in the short-term heat treatment. The different lath morphologies were most probably related to different heating rates, since the long-t erm heat treatment e mploy ed the vacuum furnace and the short-term heat treatments were conducted in the vertical tube furnace. The samples heated in the vacuum furnace w ere pumped to the operating vacuum pressure before the furnace was ramped up to the aging temperature. The samples heated in the tube furnace were simply dropped into the heating zone. Thus the samples heated in the vacuum furnace experienced a slower heating rate, while the samples heated in the tube furnace experienced a faster heating rate. 4.2.3.2 Alloy 4 The analysis of the samples heat treated at 1200C for 2 and 5 minutes revealed that the as-cast microstructure transformed completely to a two-phase cr + J3 microstructure. Figure 4.23 shows the microstructure as observed by optical microscopy using bright field and dark field conditions. TEM micrographs shown in

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111 (a) 2.0m (b) 0.5m Figure 4. 22. TEM micro graphs showing the microstructure observed in the 1200 C2min-WQ sample of alloy 2 (a) the cr (dark) and the y Q.ight) phases at a low magnification; (b) the cr grain size.

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112 Fi g ur e 4 24 d e m o ns trate th at th e int e r co nn ec t e d pha se i s cr and th e e l o ngat e d l a ths hap e ph ase is di s ord e r e d p ph ase. Th e s ampl e s hown in th e mi c r o gr a ph s w as tilt e d so that a s ingl e f3 grain was o ri e nt e d on th e (001] zon e axis. This s ub se qu e ntl y c aus e d all o f th e P gr a ins t o b e o ri e nt e d o n th e [00 I] zo n e axis and is ob se rv e d in Figur e 4.24a by th e f ac t that all o f th e f3 grains hav e th e s am e dark co ntra s t. Th e s ampl e wa s th e n til te d su c h tha t th e two a grains w e r e ori e nt e d o n th e (0011 zo n e axis. Thi s diff e r e n ce in tiltin g am o unt e d t o a mis o ri e ntation of about 1 -16 fr o m th e [00 l] z on e axi s o f th e f3 grain s Thi s pr oce dur e c au se d e a c h o f th e a gr a in s t o s how a d a rk co ntra s t and d e m ons trat e d th a t e a c h gr a in wa s gr e at e r than 10-15m in s i ze a nd co nt a in e d a numb e r o f small f3 grain s Ii s hould al so b e m e nti o n e d that th e f3 gr a ins w e r e obs e rv e d t o hav e an o rd e r e d stru c ture and t o c ontain small APBD s. Th e s ampl e that was h e at tr e at e d at 1000 f o r 2 minut es w a s o bs e rv e d to hav e a partially transform e d mi c r os tru c tur e a s s hown in th e opti c al mi c rogr a ph o f Figur e 4 25 Thi s mi c rograph indi ca t es that th e transform e d produ c t f o rm s in co l o ni es and th a t nu c l e ati o n occ ur s a t th e primary J3 grain b o undari es. Th e TEM anal ysis indi c a te d th at th e mi c r os tru c tu.r e e volv e d by t h e coo p e rativ e f o rmati o n o f th e tw o pha ses A mi c r o graph illustr a ting this pr oce ss at th e r e a c tion fron t o f a co lony i s sh o wn in Figur e 4.26 Thi s ori e ntati o n s hows that th e co l o ny co n s ist e d o f a r o ds h a p e seco ndary f3 phas e l oc at e d b e tw ee n th e a pha se. It was d e t e rmin e d by SAED analy s i s that th e P phas e in th e tw o -phas e s tru c tur e had a disord e r e d BB C cry s tal s in ce n o s up e rlatti ce r e fl ec ti o n s w e r e o bs e rv e d a t th e [100] z o n e axi s, as s hown in Figur e 4.26b. Th e r o d -s hap e di so rd e r e d f3 grain s w e r e o b se rv e d t o b e se parat e d fr o m th e r eta in e d B2 grain s b y an int e rfa ce.

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113 ( a) 20m (b) Figure 4.23 Optical micrographs showing the mjcrostructure observed in the 1200 -2mjn-WQ sample of alloy 4. (a) bright fi e ld micrograph ; (b) dark fi e ld mi c rograph.

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114 (a) 4IJ J )f -c; 1 0m (b) Figure 4.24. TEM mi c rographs showing the microstructur e observed in the 1200 C 2min-WQ sample of alloy 4. (a) the thin foil specimen was tilted to th e (001] 13 zo n e axis of the f3 partic l es; (b) the thin foil spec im e n was tilted to the [OOl] ozone axis o f th e cr matrix

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115 20m F ig ur e 4.25 Opti cal mi c r o graph s h o wing th e mi c ro s tru c tur e o b se rv e d in th e 1000 C2min-WQ s ampl e o f all oy 4

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116 (1 (a) B2 300A (b) Figure 4.26. Shows the microstructure observed in the 1000C-2min-WQ sample of alloy 4. (a) TEM micrograph showing the reaction front of a colony that partially transformed from the B2 matrix; (b) SAED pattern showing the [100] 13 zone axis of the disordered p phase lo cated between the cr grains in the colony.

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117 The transverse view of the colony structure near the interface between the colony and the retained B2 matrix is shown in Figure 4.27. The analysis at this orientation showed that this structure consisted of a continuous cr matrix. The following orientation relationship between the cr and the B2 phases was observed at this direction: <00l>B2 II 13 II
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118 (a) 500 A (b ) Fi g ur e 4.27 Sh o w s t h e mi c r ost ru c tur e o b se rv e d in th e 1000 -2min WQ s ampl e o f all o y 4 (a) TEM mi c r o gr a ph s howing th e tran s v e r se vi e w o f th e co l o n y s tru c tur e n ea r th e int e rfa ce b e tw ee n th e co lony and th e B2 phas e; (b ) SAED p at t e rn s h o wing th e o ri e nt a ti o n r e lati o nship b et w ee n th e B2 and cr ph ases, whi c h w as < 100 >B 2 II [OOl] cr and {110} 8 2 II {llO )cr.

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119 (a) 500 A (b ) Fi g ur e 4 28 Sh o w s t h e mi c ro s tru ct ur e o b se rv e d in th e 1000 -2min-WQ s ampl e of all o y 4 (a) TEM mj c r o gr a ph s h o win g th e tr a n s v e r se vi e w o f th e co l o n y n e ar th e ce nt e r of th e c ol o ny ; (b) C BED patt e rn sh o wing th e ori e ntati o n r e l a ti o nship b e tw ee n th e cr and pl1a ses, whi c h w as < 100 > 13 II [lOO] (f an d {110} 13 II {1 lO ) a.

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120 Th e analysis rev e aled that there w e re other possible orientation r e lationships between th e cr and primary p phas es. Thes e adclitional orientation r e lationships w ere observed for th e cr phas e that precipitat e d from the primary p phas e at temperatures above ~ 1200 C and for the cr phas e that form e d in co l onies from the primary J3 phas e at 1000 C. It wa s d e termin e d that th ese orientation r e lationships involv e d th e [OOl] (J zone axis which was th e 4-fold symm e try axis in the tetragonal structure of th e cr phas e Figure 4.29 s hows two of th e additional orientatio n r e lati onships that wer e co mmonly observed The first orientation r e lationship is shown in Figtu e 4. 29a: < 1 TO>a 2 II a 2 II <001J a and {110} 8 2 II {110)(1 since both of th ese orientation r e l ationships ar e p e rpendicular to e ach other. The second orientation r e lationship is s hown in Figure 4.29b: < 113 > 82 II
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121 (a) 'u ~!IDcb ~'D@m3 I(b) Figure 4.29. SAED patterns s l1owing two additional orientation relationships that wer e observed b e tw ee n the cr and phases in the h eat treated samples of alloy 4. (a) < 1TO >B2 II B 2 II [OOl]cr and {110} B 2 II {llO)cr.

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122 In this chapter the discussion will fo c us on th e e quilibrium phas es and th e ir transformation me c hanisms. Th e d e tailed analysis of the metastable phases observ e d in alloy 2 will be covered in chapters 5 and 6. This discussion is divid e d into thr ee sections: (1) the l3-phas e 1 (2) phase equilibrium, and (3) phase transformation m ec hanisms. 4.3.1 Th e 13 Phas e Th e results showed that th e BCC 13 phase solidifi e d as the primary phas e in alloys 2, 3 and 4. This is in g e neral agre e m e nt with recent data concer ning th e liquidus projection of the 13 phase in the Nb-Ti-Al ternary syste m [12). The heat tr ea tments prov e d that the 13 phase existed at elevated temp e ratur es as a single phase from the 13-s o lidus temperature to approximately 1400C in alloy 2 and to betw ee n 1550 C and 15 l 5 C in alloy 4. Th e refor e, tl1ese results combined with th e as-cast results d e monstrat e th e broad rang e in temperature and composition that th e 13 phase has at e l e vated temperatures in the Nb-Ti-Al system and add to the results of previous studies [12 21 22,24,37). Th e impli cation of th ese results i s that basic physical m e tallurgy principles involving heat treatments and quenching proc e dures can b e us e d for microstructural development [82 83]. It was shown in the as-cast and aged samples that the 13 phase exhibited a disord e r to order transition in alloys 2 and 4. This was assessed by the presen ce of APDBs which form e d during so lidstate coo ling from th e high temperature disordered 13 s tru c tur e to th e l o w t e mp e ratur e ordered B2 structure. A numb e r o f studies hav e s h own that th e B2 phase exists over an extens iv e composition rang e in th e Nb-Ti-Al system [12 22,24 37]. The results of' this study seem to support th ese findings since they showed the 13 phase to be ordered in alloys

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12 3 2 and 4 Th e fa c t th at o rd e ring wa s n o t ob se rv e d in all o y 3 indi ca t es that it d e p e nd s o n co mp os ition Th e mo s t s ignifi c ant diff e r e n ce b e tw ee n th e co mp os iti o n o f all oy 3 ( 10 % Al ) and all o y s 2 ( 40 % Al ) and 4 ( 30 % Al ) wa s in th e Al c ont e nt o f th e m Th e r e f o r e it will b e argu e d th at Al is r e sp o n s ibl e f o r ord e ring in th e phas e. Th e s e r e sult s diff e r fr o m t hat g e n er ally a cce pt e d in t h e lit e r a tur e whi c h co rr e lat e o rd e rin g with t h e Ti co nt e nt (50] A numb e r o f st udi e s h a v e s h o wn a g e n e ral tr e nd indi c ating t hat Al alw a y s occ upi es a se par a t e s ublatti ce in vari o u s pha ses in th e Nb-Ti-Al sy s t e m whil e Nb a nd Ti a r e f o und t o occ up y a sublatti ce e ith e r c ompl e t e ly o r at l e ast limit e dly. F o r in s tan ce, K o nit ze r et al. (47] d e t e rmin e d that t h e 611 (Wy c k o ff d es igna tio n ) s it es in a. 2 Ti 3 Al w e r e o cc upi e d by Ti and Nb with n e ith e r substituting for Al on th e 2d s it es. In thi s sa m e s tudy by K o nitz e r e t al it wa s s h o wn that Nb rand o mly s ub s titut e d with Ti o n l a and l e s it es whil e Al occ upi e d th e 2 e s it es in a dilut e Nb co ntaining y -TiAl ph ase. Th e s am e co n c lusi o n wa s obtain e d f o r th e y pha se that w as pr ese nt in all o y 2 a g e d a t 1200 C. Int e r e stingly th e y phas e in alloy 2 co ntain e d up t o ~ 22at. %Nb and y e t th e diffr ac ti o n r es ults w e r e co n s i s t e nt with a rand o m occ upan c y o f Ti and Nb o n t h e l a and l e s it es (see sec ti o n 4 2.2 1 ) Ev e n in th e sto i c hiom e tr ic TizAINh o rth o rhombic ph ase e xamin e d by Mo ze r e t al. [28] in which Ti oc c upied the 8g s it es, Al occ upi e d th e 4 c l si t es, and Nb th e 4 c 2 s it es [13] it wa s d e t e rmin e d that Ti and Nb s how e d s om e d e gr ee o f co -o cc upan c y o n th e 8g and 4 c 2 sit e s but that n e ith e r s ub s titut e d for Al o n th e 4 c l s it es. Th e r es ult s o f th e ir s tudy also s h o w e d that up t o 18 % o f th e 8g sit es w e r e o cc upi e d by Nb with th e sam e fra c tion o f 4 c 2 sit e s occ upi e d by Ti. Fin a lly th e s tudy by K o hm o to e t al [49] s how e d th a t Al wa s r es pon s ibl e f o r

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124 th e J3 to B2 ord e ring s in ce th e m e tastabl e B2 phas e was obs e rv e d in two binary Nb-Al a ll o y s that c ontain e d only 13 5at %Al and 16 9at.%Al This s am e s it e occ upan c y tr e nd may al so e xi s t f o r th e B2 phas e. Ev e n t hough th e r e sult s by Ban e rj ee e t al [50] show e d that th e la s ublattic e of th e B2 pha se wa s occ upi e d pr e f e r e ntially by Ti and th e lb sublatti ce by both Nb and Al this study still indi c at e d that Al pr e f e r s to oc c upy sit e s on a s e parat e sublatti ce Th e co occ upan cy b e tw ee n Al and Nb o n on e sublattic e may simply b e du e to an insuffici e nt amount o f Al s in ce th e alloy B a n e rj ee e t al. e xamin e d and e v e n thos e e xamin e d by Kaufman e t al. (25] had z 25at. % Al Th e refore o nly 50% of th e sites on th e sublattic e c ould b e fill e d by Al whil e th e r e st had to b e o c cupi e d by e ith e r Nb and/or Ti. It is also s ignifi c ant to n o t e th a t Ban e rj ee e t al d e t e rmined that Nb and Ti (admitt e dly in limit e d quanti t i es) co occ upi e d s it es o n b o th sublatti ce s but that Al only o cc upi e d s it es on on e s ublatti ce This r e asoning is c onsist e nt with th e o bs e rvations of alloy 3 in which ord e ring did not occur s in ce it co ntain e d only ~ lOat %Al. Thi s r e pr ese nt e d only a 20% sit e o c cupancy of th e sublattice and was too small of a fra c tion for o rd e ring to o cc ur Th e r e s till m a y b e a pr e f e r e n ce b e tw e en Ti and Al in th e ord e ring b e havi o r o f th e J3 phas e This may b e understood b e tter by c onsidering th e bonding b e tw ee n Nb Ti and Al N e ar e st n e ighbor (NN) L e nnard Jon e s pot e ntials (84] indicat e that th e Ti-Al b o nd i s th e m os t s t abl e a t ~ -56 kJ/m o l with th e Nb-Al bond just slightly l e ss s tabl e at ~ -51 kJ / m o l Th e s e valu es ar e signifi c antly diff e r e nt from that of th e Ti-Nb bond whi c h wa s d e t e rmin e d t o b e ~ -5 kJ/mol by B CC int e raction param e t e r (69]. Th e latt e r agr ee s with th e binary Ti-Nb phas e diagram which shows a broad solubility rang e b e tw ee n Nb and Ti and no B2 ord e ring [10 11 72, 73] Th e slightly strong e r

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125 Ti-Al bond is stabilized by positioning Ti on the la Wyckoff sites and Al on the lb sites Th ese sites represent th e NN bonds, in terms of bond lengths, and are the s hort est in the BC C str ucture This atomic arrang e m e nt is shown in Figur e 4 .3 0 for ( 110 ) lattic e plan es of two unit ce lls of the ordered p (B2 ) phas e. Tl1 e NN bond is represented by the a/2<111> v ec tor in this figure The slightly l ess stable bond between Nb-Al results in a co rr esponding ly slightly long er b on d l e ngth Th e r efo r e, it seems reasonable that Nb will preferentially occupy the lb sites along with Al since the next-nearest-neighbor (NNN) bond in th e B CC structure defines th e latti ce translation b et w ee n two lb sites, which is the a v ec tor shown in Figur e 4.30. Thus the slight diff e r e nce b et w ee n th e Ti-Al and Nb Al bond strength could account for the site occupancy r es ults of Ban e rj ee et al. [50] whil e still maintaining that Al is responsible for ordering in th e p phas e. In summary the gen er al pi c tur e that develops from these considerations is that or d er ing in the P phas e is du e to Al adopting one o f two sublattices. Th e other sub l attice is pr e f e r ent ially fill e d by Ti du e to the st rong e r b o nd b e tw ee n Ti-Al. D eficie nt s i tes o n the Al sublattice are pr e f e r e ntially fill e d by Nb since its bond st r e ngth with Al is s lightly weak e r H o w eve r in r eality, an alloy with a t e rnary co mposition will d eviate from th ese g e neral rules and ultimat e ly th e ordering will b e dictated by th e rmodynamic paramet e rs. Th ere w e r e also s e v e ral anomali es that w e r e co nsist e ntly observed in the ordered P phase, regardless of wh e th er it e xist e d as th e matrix or as r e tain e d parti c l es in the mi crostr uctur e. Th ese anomali e s w e r e observed in the SAED patterns as the diffuse e l ec tron scattering, stre aking, and spot splitting (Figures 4 .2, 4.4 a nd 4.5). Th e matrix wa s also found to hav e a tweed microstructur e whi c h

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126 .. .. ----<{:{ ? };~ :i-~ :;:;:;:;::,~ .... .-.~: . ~--{ J f '.; }J ,-......a[OOl] a/2[111] . . ........ ~--;: {;t ~ Jt~: :J,,, -[001] (110) B 2 Lattice Titanium/Niobium Aluminum Figure 4.30 Shows the translational vector for t h e n earest n e ighb or (NN) and n ext n e arest neighbor (NNN) s it es in the unit ce ll of the B2 phase. Th e atomic site oc c u pancy sho w s Nb and Ti atoms randomly occupying th e la Wy ckoff site and Al atoms occ upying the lb W yckoff s it e.

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127 consisted of striations that were parallel to the {lTO}p planes in the p directions. The analysis of the tweed striations showed them to be visible or invisible depending on the g reflection. This indicated that strain was associated with the striations. '"fhis type of image characteristic was similar to that of dislocations which can be made visibl e or invisible depending on the g b condition, where b is the Burgers vector of the dislocation and g is the reflection vector [85]. Thus, these observations were consistent with those from previous studies and indicated that the anomalies were caused by incommensurate shear strains on every other {110} 13 plane in the <110> 13 direction of the B2 phase [51-56]. The origin of the shear strains was attributed to phonons which had the wave vector k = and the polarization vector <110> [36,51]. The atomic displacements that result from these lattice shears was shown to be consistent with the heterogeneous deformation that occurs in the mart e nsitic transformation of' the 2H ortho -h ex agonal phase from the p phase [36]. Th e results of this study seem to support tllis possibility for alloy 2 since a martensitic transformation of plates formed from the B2 phase was observed in this alloy during water quenching and during rapid solidification. This martensitic transformation will be covered in chapter 6. 4.3.2 Phase Equilibrium The phase equilibria that was determined for al]oys 2 and 4 is s11mmarized using the flow diagrams shown in Figure 4.31. The results showed that the single disord e red P phase was stable from th e p solidus temperature down to the cr transus t e mperature in both alloys. The cr phase nucleated from the p phase at the cr transus temperature, which was slightly below 1400 C for alloy 2 and between 1550C and 1515 C for alloy 4. These alloys had equilibrium two-phase cr + p microstructures

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128 b e l o w th e er transus te mp e ratures Th e er and p phas es wer e th e equilibrium phas es in the microstructur es down to t e mp er atures that w e r e 1300C in alloy 2 and below 1200 C in alloy 4. Th e phase equilibria that was observed b e low th ese temperatures was found to b e diff ere nt for alloys 2 and 4 Th e microstru c tur e in alloy 4 at 1000C was d ete rmined to consist of th e er phas e and O phas e, whi c h was bas e d o n th e O-Ti 2 A1Nb phas e [13], in th e long t erm h eat treated sample but consisted of the er and disorder e d p phas es in th e short-term heat treated sample. A possible reason for the diff e r e nt two-phase micro st ru ct ur es may hav e been du e to an insuffi c i e nt amount of tim e f o r the O phas e to form during the short term h ea t treated sample. This idea was bas e d on pr e vious studies that had shown the transformation of th e O phas e from th e high temperature p / B2 phase to be tim e and temp e ratur e d e p e nd e nt [ 13 24 35]. Th ese studies hav e shown that this transformation occ urs e ith e r during slow coo ling from th e p/B2 phas e, or during isothermal h eat ing of the qu e nch e d and r e tain e d p/B2 phas e Th e co mposition of the retained P parti c l es which w e r e pr ese n t in th e tw opha se er+ p microstructur es of a ll oy 4 wa s d eter min ed pr e vi ous ly by Gomez [86]. In this study, the composition of t l1 e P part ic l es was found to b e ~3 6Nb 43Ti-21Al ( at.%) which indicat e d that the Ti and Al co nt e nt w e r e in c r e as e d whil e the Nb co nt e nt was decr e as e d compared to the com po s ition of alloy 4 which wa s 42Nb-28Ti-30Al (at.%) Thus the c hang e in the com position of th e J3 particl es would hav e favor e d the formation of th e O phas e, provided that suffi cie nt time was allow e d for this phas e to form. This suggests that th e disord ere d P phase whi c h form e d in 2 minut es during th e short t e rm h e at treatment, was a n on-e quilibrium ph ase that would hav e transform e d to th e O phas e with a long e r aging ti m e.

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129 p phase Solidus to 1400 C p + cr phases 1300C (a) p + cr + y phases 1300C > T > 1200C cr + y phases 1200 c p phase Solid11s to 1550C p + cr phases 1515c p + cr phases 1400 c (b) p + cr phases 1300 C p + cr phases 1200 c p + cr phases 1000c Figur e 4 31. Shows the equilibrium phases that formed at the aging temperatures in alloys 2 and 4. (a) alloy 2; (b) alloy 4.

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130 Another con sideration for explaining the formation of different two-phase microstructures in the 1000C aging experiments of alloy 4 was the use of different cooling rates in the short-term and long-term aging experiments. In these experiments, the sample used in the short term aging exper im ent was water quenched while that in the long-term aging experiment was air cooled It was thought that if the disordered f3 phase was in equilibrium at 1000C then the slow cooling rate used in the long-term aging experiments would have allowed the O phase to form at lower temperatures However, the results by Kaufman et al. [25] showed that the O phase was stab l e in a 45Ti-25Al-30Nb (at.%) alloy at 1000 The composition of this alloy was c l ose to the composition of the retained f3 particles in alloy 4 [86] Therefore, the cooling rate eff ect was disregarded as a plausible explanation, since the long term heat treatment of alloy 4 showed the presence of the 0 phase at 1000C. Thus, it was concluded that the equilibr ium phases of alloy 4 were the cr and O phases at 1000C. The different two phase microstructures observed in alloy 2 at 1300C and 1200 C indicated that a bivariant three-phase cr + y + f3 tie triang l e crossed over the composition of 27Nb-33Ti-40Al (at %) of alloy 2. It was determined that between these two temperatures, the cr + f3 microstructure which was present at 1 300C changed to the two-phase cr + y microstructure whicl1 was present at 1200 C. The passage of the bivariant three-phase cr + f3 + y field in alloy 2 is consistent with the reported liquidus projection of Perepezko et al. [12] and the 1200C i sotherm of Das et al. [22]. The liquid us projection, which was shown in Figure 2. 3, showed that a class II four-phase reaction between L + cr + f3 and L + cr + y occurred that produced the L + f3 + y and f3 + cr + y tie-triangles. The former tie-triangle continues along the

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131 liquidus projection while the latt er persists to low er temperatures and wa s s h own in the 1200 C isotherm in Figur e 2.4. B eca us e th e rmodynamics allows for two degrees of fr eedo m in ternary three-phase e quilibria the co mp os itions of the three-phases in the p + cr + y ti e -triangl e ca n change as a fun c tion of temperature ( with pr essure fixed). Th e r e for e, it can b e asserted that at 1300 C the co mposition of alloy 2 was within the two-phase cr + p fi e ld that bordered on one side of the cr + f3 + r tie-triangle. As the temperature decreased, the po s ition o f the ti e -triangl e shifted until it star t ed to c ross the co mp osition of this alloy At this point the f3 phas e started to di sap p ear at the expe ns e of the d e v e loping cr + y phases Upon passag e of th e ti e -triangl e, th e co mp os iti o n o f the alloy would then be within the two-phase cr + y fi e ld. 4 3. 3 Pha se Transformation M ec hanisms 4.3.3 1 Alloy 4 At high aging temperattrres, th e transformation of the cr phas e from the P matrix occ urs by nu c l ea tion and growth proc esses consisting of bulk diffusion. The o rd e r e d f3 phas e r etaine d upon qu e nching at room t e mp e ratur e becomes disord e r ed upon h eat ing to aging temperatures > 1200 C. This was assess e d from Figur es 4 15 and 4 16 s howing APDBs in the B2 grains that wer e formed during quenching fr o m the aging temperatU1es. Th e r e for e, these e l e m e nt s p ossesse d high diffusivity rat es at high aging temperatures, since the f3 phas e has a disord e r e d struct,ure. This e nabl ed long-rang e diffusion to occur through the f3 matrix, which would account for the nucleation of individual cr grains along the p grain boundari es and within the p matrix as seen in Figures 4.14ae. Th e fully d eve loped microstructur es at the aging temperatures may consist of is o lat ed cr grains in the f3 matrix or of several cr grains that hav e grown and imping ed

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132 on o n e an o th e r d e p e nding on th e aging t e mp e ratur e Th e cr grains ar e typi c ally i so lat e d fr o m o n e an o th e r at t e mp e ratur e s gr e at e r than 1400 As th e aging t e mp e ratur e i s d ec r ease d th e v o lum e fra c ti o n o f th e cr pha se in c r e a ses r e lativ e t o t h e f3 ph ase. At som e p o int th e v o lum e fra c tion in c r e as e s t o an e xt e nt that th e cr grains s t ar t to imping e o n o n e an o th e r. Thi s pr ocess r es ult s in th e c hang e o f th e m a trix fr o m th e f3 phas e t o t h e cr pha se b e tw ee n 1400 C and 1300 C. This occ urr e n ce c an b e see n in Figur e 4 14 which sh o w s th e J3 matrix a t 1400 C (Figur e 4 14d ) and th e cr m a trix a t 1300 C (Figur e 4 14 e) Th e impin ge m e nt o f th e cr grain s t o f o rm th e matrix c au se d th e parti c l es o f th e r e tain e d f3 phas e t o b e com e isolat e d in th e micr os tru c tur es th a t d e v e lop e d b e l o w th e 1300 C agin g t e mp e ratur es Th e diff e r e nt morphologi es of th e r e tain e d f3 particl e s w e r e c ontroll e d by th e f o rmati o n m ec hani s m o f th e cr g r a in s. Th e irr e gular s hap e d f3 parti c l es form e d at t h e hi g h e r a ging t e mp e r a tUI es. At th ese t e mp e ratur e s th e long rang e bulk diffusi o n pr ocesses ca u se d t h e nucl e ation of i so l a t e d cr grains, which th e n gr e w with a co nv e x c urv a tur e a t th e cr / f3 int e rfa ce. Th e imping e m e nt o f th ese cr grain s l e d t o th e irr e gular s h a p e o f th e r e tain e d f3 parti c l es At low e ring a ging t e mp e ratur es, th e cr grain s w e r e o b se rv e d to form a s a co l o ny with th e r e tain e d f3 parti c l e s c ontain e d as inclusi o n s. Th e co l o ny structur e was shown in Figur e 4 21 for a numb e r of clos e ly ori e nt e d cr gr a in s sh o wing co mm o n [0011 z on e ax e s. H o w e v e r t h e cr grains o f a c olony mu st ha ve nu c l e at e d s e parat e ly sinc e som e of th e [001] z on e ax es in th e c o lony w e r e r o ta te d by as mu c h a s 45 fr o m o n e an o th e r. Thi s w o uld imply tha t high angl e grain boundari es exis t e d b e tw ee n th e cr grain s o f th e co l o ny At th e l o w e r a ging t e mp e ratur es, th e transf o rmati o n r e s e mbl e s that o f a dis co ntinuous pr ec ipitati o n pr ocess [87 88] Dis co ntinu o us tran s f o rmations inv o lv e a

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133 supersaturated matrix phase that r eac ts to form a lam e lla e made up of the depl e ted solid solution phas e and the precipitated phase. The transformation is based on diffusional proc esses that occur at th e mi grating boundary of the reaction front Th e diffusional requirem e nt of this typ e of transformation is that the diffusivity of th e boundary between the phases of the lamellae must b e lower than th e diffusivity of the r e action front. In this case, th e lamellar spacing and the driving force are high e nough for th e surfa ce e n e rgy of the r e action front to allow its migration. This typ e of tra nsformation s ee ms to hav e occurred at th e temperature of 1000C wh e r e th e m e tastabl e f3 phas e r e mains supersaturated and ordered at this aging temperature. Th e ordered structure would low e r the solute diffusion and mak e it mor e likely for diffusion to occur along grain boundaries rath e r than through the bulk lattic e. The r e action starts at th e matrix grain boundary by precipitation of th e secondary phas e, in this case the cr phase which advances into the matrix by solut e diffusion al o ng th e advancing reaction grain boundary front. Th e r e action in this particular case c an b e writt e n: f f3 ~f3+cr wh e r e th e f3 pha se is th e ordered (B2 ) matrix th e f3 phas e has the disord e r e d structure and is th e depleted solid solution, and the cr phase is the precipitated phase. Th e f3 and f3 phas es form an int e rfac e that acts as th e r e action front. Confirmation of this typ e of r e action front was obtained in th e 2 min11te heat treatment of alloy 4 in whi c h th e transformation was "frozen in" by water qu e nching. The diffraction analysis clearly s how e d ther e t o b e f3 grains in conjunction with th e cr grains in th e transforming two phase structure. The analysis also confirmed that th e coloni es form e d at the J3 grain b o undari es and gr e w into the matrix It is suggested that th e

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134 same type of transformation occurred in the aged 1200C and 1300C samples, since colonies were observed by optical microscopy and verified by TEM to consist of large cr grains that contained a multitude of f3 inclusions. There were several interesting aspects of the discontinuous transformation that were observed in this study. For instance the precipitated cr phase was the major constituent of' the two-phase microstructure. Therefore, the relationship of the cr phase with the f3 matrix appeared to control the formation of the colony. This was confirmed by showing several orientation relationships between the cr grains of the colony and the f3 matrix and by showing facets at the cr/f3 interface. Faceting was observed along the sides and at the reaction front of the transforming co l onies. The analysis of' the different orientation relationships consistently showed the {110)(1 planes to be parallel to {llO}p planes. These planes were also commonly observed as faceted planes at the cr/f3 interface However other facets were also observed, such as the one between {100} 13 and {lOO}a planes. The facets imply that they have low interfacial energies and, because of this, they influence the crystallographic directions that the transformation front will proceed with in the matrix. Together these results have shown that crystallography plays a role in the nature of the discontinuous transformation observed in this study. 4.3.3 2 Alloy 2 The formation of the cr phase from the f3 phase below:::: 1400 C in alloy 2 occurs by nucleation and growth that requires long range diffusion through the f3 matrix. This type of transformation is consistent with the formation of large blocky-shaped cr grains along p rim ary f3 grain boundaries and within the f3 matrix of the 1300C aged sample as shown in Figure 4.10c. The development of the cr + y microstructure at

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135 1200 r ese mbl es that o f an e ut ec toid transformation This typ e of transformation occurs by the J3 phas e, which was th e matrix in th e RAM sampl e, transformjng to th e cr + r phas es at the aging t e mp e ratur e. Th e e utectojd transformation is als o bas e d o n nucl eat ion and growth that r e quir e s diffusional proc esses. Th e diff e r e nt morphologi es o f th e cr + r phas es that w e r e observed in the mi c r ost ructur es o f th e short-term and l ong-term 1200 aged samp l es indicat es that tl1e d e v e l opment o f these mi c rostructur es are affect ed by the heat rat e. Th e l ath m or ph o logy whi c h was observed in Figur e 4 12 in tl1 e long-t e rm aged sa mpl e wa s form e d by slowly h ea ting th e RAM sampl e from room t e mperatur e to 1200 In this h ea t t r ea tm e nt th e nu c leation and gr o wth of th e cr and y phas es from t h e J3 matrix occurs b e low 1200 Therefor e, it is e xp ec t e d that the diffusion path through ih e J3 matrix will be smaller at t e mp e ratur es l ower than at 1200 C, rath e r than at 1200 C. Th e transformation o f the J3 pl1as e to the cr + y phas es at the l ower temperatur es i s t h o u g ht L o hav e occ urr e d as a co l o ny of these two pha ses. Th e diffusi o n occurs n ear the transformation front and ha s a path l engt h that is in th e order of the grain size observed for th e cr a nd y phas es in th e lath str u ctures. Th e small diffusion path in the J3 matrix also ca u ses th e nu c l eat i o n o f n e w grains o f th e cr and y pl1as es al1ead of the advancing transformation front. Th e e quiaxed grain morphology which w as observed in Figur e 4.22 for th e s hort-t e rm aged sample, was form e d by rapidly h ea ting the RAM sa mpl e to 1200 C. The nu c l ea tion of th e cr and y phas es from th e J3 phas e occurs by site sa turation since the samp l e quickly r e ach es the aging te mp e ratur e o f 1200 C. Th e high numb er density o f nu c l e i and th e rapid diffu sio n rate at thl s t e mp e r at ur e r es ults in th e f o rmati o n of small e quiax e d y grains disp e rs e d in the cr matrix

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136 Th e analysis of' the short term aged samples revealed that the r e tain ed f3 phase transformed to the eq uilibrium microstructur es very quickly at 1200C and higher. This was demonstrated by aging a RAM sample for just tw o minut es at 1200 C and observing the equilibrium cr + y phases in the microstructure as shown in Figure 4.22. This r es ult also indicat es that th e ro-related pr ec ipitat es and plates that formed initially in the f3 matrix of th e RAM sample dissolv e d rapidly to form the e quilibrium cr + y phases in this short amount of tim e. Howev e r this transformation rate pertains to only phase e quilibria and d oes not indi ca t e how stable the mi crostructure is to grain coarsening processes. It is unlik e l y that the initially co mpl ex microstru c tur es observed in the as-cast samples l1ad any affect on th e d eve lopm e nt of the cr + f3 microstructur e at 1300 C. This is based on the results whi c h will be co v ere d in c hapt er 5 and 6 that show the ro-phase and the plates are m e tastabl e phas es that form e d from th e f3 phase with o ut the need for l o ng r ange diffusion. Th ere for e, the dissolution of these m etas tabl e phases s h o uld occur very fast at the high aging temperatures and that the ordered B2 ph ase s h o uld rapidly r e v e rt ba c k to the di sor d e r e d f3 phas e.

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C HAPTERS THE OMEGA-RELATED (w -D ) PHASE 5. 1 Introdu ctio n In c hapt er 4 o n phas e e quilibria it was shown that an ro-type phas e h a d pr ec ipitat e d in th e p matrix of th e as c ast RAM sample of alloy 2 How e v e r these precipitates w e r e n ot observed in th e h ea t tr ea t e d and wat e r qu e n c h e d samp]es of all oy 2 or w e r e they o bs e rv e d in alloy 4 which includ e d th e as-cast RAM and h eat treated sa mpl es. It was al so p re viously shown that the ro -typ e pr ec ipitat es that form e d in alloy 2 h a d two size distributions whi c h d epe nd e d o n their l ocatio n in the micr ost ru ct ur e. A hi g h numb e r density of small pr ec ipitat es with a size rang e of ~ 0 05m w e r e di s tribut e d homog e n eo usly in th e matrix as was shown in Figur e 4 .3 b. Th e r e wa s n o apparent int e ra ction b et w ee n them and th e anti-phase domain boundari es that f o rm e d in the p matrix during th e disord e r to o rd e r transition Th e larg e ro-type precipitates w e r e usually connected t o th e l e nti c ular precipitates (plates) as w as s h o wn in Figur e 4.8b. Th e size range of these ro -typ e pr ec ipitat es wa s b e tw een 0 5 and l Om. Th e purp ose of t hi s c hapt er is to exa min e th e transformation m ec hani sm of the m etasta bl e ro type phas e fr o m th e p phas e in gr ea t e r d e tail. Results co n cer ning the st ru c tur e, the o ri e ntati o n r e l at ionship with r es p ect t o the p ph ase, and th e t e mp e r a tur e d e p e nd e nt param ete rs that affect the f o rmati o n o f tllis pha se will b e pr ese nt ed. Thi s will b e foll o w e d by a di sc ussion of th e transformation m ec hanism of 137

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138 th e ro pha se and will b e tr e at e d in te rm s o f a gr o up / subgr o up r e pr ese ntati o n It will b e s h o wn that th e ropha se that wa s obs e rv e d in alloy 2 i s a n e w pha se that was di sco v e r e d in thi s st ud y Th e r e f o r e, th e di sc us s i o n will al so f oc u s o n it s r e lati o n t o th e "family" o f ro -ph ases and will in c lud e a pr o po se d sit e oc cupan c y for it ba sed o n r es ult s o b t ain e d in t hi s s tudy 5. 2 R es ult s Th e r es ult s prese nt e d in t hi s c hapt e r co v e r th e s tru c tural anal ys i s o f th e ro -typ e pha se and th e e ff ec t o f h ea t tr e atm e nt and coo ling rat e o n it s f o rmati o n Th e s tru c tural analy s i s sect ion d esc rib es th e d e t e rminati o n o f th e c ry s tal st ru c tur e, th e p o in t group th e s p ace gr o up th e l a t t i ce param e t e r s, and th e ori e n t ati o n r e lati o n s hip wi t h th e f3-pha se and th e plat es Th e o th e r two s ec ti o ns d e al with th e influ e n ce o f coo lin g r a t e and l o w h e at tr ea tm e n t te mp e ratur es o n th e f o rma t i o n o f th e ro -iyp e pha se. 5 2 1 Stru ct ural An al y s i s 5. 2 1 1 C ry s tal P o in t G r o up a nd Sp ace G r o up D e t e rminati o n Th e c ry st al str u c tur e o f th e ro t yp e phas e wa s d e t e rmin e d u s ing C BED tec hniqu es o n th e l ar g e pr ec ipitat es th a t f o rm e d in th e as cas t RAM sa mpl e as w as s h o wn in Figur e 4.8b o f chapt e r 4 Figur e 5 1 s how s C BED patt e rns that w e r e o b ta in e d fr o m o n e of t h e s e pr ec ipit a t es. Th e C BED wh o l e patt e rns o bs e rv e d in Fi g ur es 5 l a and 5 lb w e r e o btain e d with a larg e and a s mall ca m e ra co n s tan t, r es p ec tiv e ly and s h o w 6 fold s ymm e try with two o rth o g o nal mirr o r lin es It w as f o und th at t h e high e r o rd e r lau e zo n e (H O LZ ) lin es in th e dir ec t tran s mitt e d o r bri g h t fi e ld di sc w e r e n e v er o b se rv e d and th a t v e r y f a int fir st o rd e r lau e zo n e (FO LZ ) r ing s w ere occas i o nally o b se rv e d but o nly af t e r v e ry l o ng e xp os ur e tim e s H o w e v er,

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139 (a) (b) Figur e 5 1. CBED whol e patt e rns showing the 6mm s ymmetry observed in th e [0001] zo n e axis of the ro-related phas e in alloy 2 (a) long camera l e ngth s howing th e zero order lau e patt e rn; (b) short ca m e ra l e ngth showing the faint FOLZ ring s.

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140 wh e n th e FOLZ rings wer e observed they always show e d th e 6mm symm e try in th e C BED whol e patt e rn This r e sult indi c at es that the structur e of th e ro-typ e phas e b e l o ngs t o th e 6mm lR proj ec ti o n diffraction group [89 90] as shown in Tabl e 5 1 Fr o m thi s tabl e, th e two p o ssibl e diffraction gr o up s that b e l o ng to the 6mm I R pr o j ec ti o n diffra c tion gr o up and show 6mm whol e patt e rn symm e try ar e e ith e r th e 6m m o r 6mm IR. B ot h o f th ese s ymm e tri e s ar e c o nsist e nt with th e h e xagonal c los e pa c k e d (H C P ) Brav a is lattic e. Th e r e for e, th e diffra c ti o n patt e rns shown in Figur e 5 1 ar e c onsist e nt with t h e [0001] 00 zon e axis Th e point gr o up of th e ro-typ e phas e was d e t e rmin e d by CBED using a sec ond zo n e axis o f high s ymm e try or th e [lTOO] co zon e axis in this s tudy and is s hown in Figur e 5 2. Th e s e C BED whol e patt e rns show 2mm symm e try for this z on e a xi s u s i ng a larg e (Figur e 5 2a ) and a small (Figur e 5 2b ) c am e ra constant This r e sult i s co nsi ste nt with th e 2mm I R pr o j ec ti o n diffr ac ti o n group whi c h has tw o po s sibl e diffra c tion s ubgr o up s a sso c iat e d with it a s s h o wn in Tabl e 5. 2 [89 90]. A cc ording t o thi s tabl e th e two p o ssibl e diffra c ti o n gr o up s ar e e ith e r 2mm or 2mm lR Finally by co mbining t h e r e sult s of th e [0001] 00 and [lTOO] co zon e ax es t h e p o int group of th e ro -typ e phas e wa s id e ntifi e d A compil e d list from Buxton e t al. [91] of all th e p o ssibl e point groups f o r e a c h of th e four diffraction groups is sh o wn in Tabl e 5 3 Th e e xamination of th e tabl e shows that th e r e is only on e p o int gr o 11p t h a t occ urs for th e diffraction group s of b o th z on e ax e s. Thi s is d e t e rmin e d to b e th e 6 / mmm p o int gr o up and s in ce th e r e ar e n o o th e r p o int group s that m ee t thi s c rit e r io n th e n thi s i s th e point group o f th e ro -typ e pha se Th e s pa ce gr o up o f th e co -typ e phas e wa s d e t e rmin e d by noting th e pr ese n ce o f additi o nal tran s la t i o nal symm e try e l e m e nts i .e. a s c r e w axis and/or a glid e plan e

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Table 5.1. 141 Shows the relation between the possible diffraction groups and the symmetries observed in the Convergent Beam Electron Diffraction (CBED) patterns at the [0001) 00 zone axis. [89]. Observ e d Projection Possible Whol e Pattern Symmetry in Diffra ctio n Diffra ction Symmetry Whol e Pattern Group Groups 6mRmR 6 6mm 6mm 6mm 6mm1R 6RmmR 3m 6mmlR 6mm

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142 (a) (b) Figur e 5 2. C BED wh o l e patt e rns showing th e 2mm s ymm et ry o bs e rv e d in tl1e [1100] zo n e axis of th e ro -r e lat e d pha se in all o y 2 (a) long ca m era l e ngth showing th e zero order l au e patt e rn ; (b) short c am e ra l e ngth showing the FOLZ ring s.

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143 Tabl e 5.2 S h o w s th e r e lati o n b e tw ee n th e p o ssib l e diffra c ti o n gr o ups and th e symm e tri e s obs e rv e d in th e Co nv e rg e nt B e am E l e ctron Diffra c ti o n (CB E D ) patt e rns at th e [ lT O O ] w zo n e axi s. [ 89 ] Obs e rv e d Pr o j ec ti o n Possib l e Wh o l e Patt e rn Symm e tr y in Diffra c ti o n Diffra c tion S y m m e try Who l e P atte rn Gr o up Gr o up s 2m.RmR 2 2mm 2mm 2mm 2mml R 2Rm m R m 2 m m l R 2mm

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144 Tabl e 5.3. Sh o w s t h e c rystal point groups that ar e co nsist e nt wi t h the diffr act i on groups o b ser v e d in th e C BED wh o l e patt er n s [89] Diffra ctio n Z one Axis C ry sta l Point Gro up s Gro up s 6mm 6mm [0001] 6mm1 R 6 / mmm 2mm mm2 6m2 [1100] 2mm1 R mmm 4 / mmm 6 / mmm m3 m 3m

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145 Figure 5.3 shows a CBE D patt e rn of the [1100) 00 zone axis that wa s obse rv ed to s how evi d e nc e of both a scre w axis and a glide p l ane. This patt e rn was obtained by tilting the speci m e n in order to position th e kinematically forbidden (0001) r e fl ection at the Bragg con dition This co nditi o n s h o w s dynami c absences as a black cross, kn own as Gjonnes -M oo di e lines or G M lines, for the ( 0001 ) r e fl ect i o ns [89] and indi cates the sim 1tltan eo us pr esence of both a screw axis and a glide plan e in tl1 e structure. Th e presence of a screw axis at this zone axis is consistent with a 6 3 screw axis that li es parall e l to the ( 0001 ) plan e normal. Th e glid e plan e d ete rmin e d from this orientation li es in the (0001) plane and is parall e l to the [1120] dir ec tion or is d e fined as a c-g lid e plan e. T o e liminat e P6 3 / m cc as a po ssib l e space group, the pr ese n ce of a glide plan e was checked at the (1120] zone axes but was not observed Th ese observations indicate that the space group o f t h e ro -typ e phase that was o bs e rv ed in all oy 2 is cons i ste nt with P6 3 /mc m [ 43 92]. 5.2.1.2 Orientation Relationship with the (3 Phas e This inv estigatio n s how ed that th e ro-type precipitates had an o ri e nt atio n relationship with the p matrix in the as cast RAM sample. Figur e 5.4 shows this orientation r e lati ons hip: [0001) 00 II [111] 13 and (1100)(0 II (110) 13 R e fl ections at l/3{110} p position s co rr es pond to {1100} 00 r e flections and indi cate that both latti ce plan es are parall e l to e ach ot h e r. Th e o b se rv ed orie ntati on r e lati ons hip indicat es that tw e lv e variants w ere formed in the transformation. Four rotational variants w e r e produced by alignin g the [0001]0) direction o f eac h variant parallel to eac h of the four <111> 13 dir ectio ns Th e r e w e r e then three trans l at ional variants that w ere formed by aligning th e ( 1100 ) 00

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146 Figur e 5 3 CBED whol e patt e rn o f th e [1100] zone axis with th e b e am tilt e d slightly to show th e bla c k c r oss in th e kin e matically forbidden (0001) r e fl ectio n at th e Br agg co ncliti o n.

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147 12lp 1100 Ollp . 10~0 0110(4 2 110 (4 1210(1) . llO p 1120"' 101 p . 1121} 2 11 11 Figure 5.4. SAED pattern showing the orientation r e l ationship that was obs e rved for th e co re l ated and B2 phas e s. The OR was d etermined to b e [ OOOl] (J) II [ lll ] p and (lTOO)(J) II ( l TO)p.

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148 plan es parallel t o eac h of th e thr ee {110} 13 plan es of the four rotational variant s. Th e r e for e, the SAED patt e rn shown in Figur e 5 4 c an co nsist of diffraction spots fr om a ll twelve variant s, which w o uld b e superimposed o n the [111] 13 zo n e axis Th e twelve variants would b e accounted for in this diffra c tion patt e rn by the [00011 w, [4223] 0)1 [2423] 00 and [22 43 ] (1) zo n e axes. The or i e ntati o n r e lati o n s hip b et w ee n the ro-type and J3 pha ses is al so s h o wn in the s t e r eogra phi c projection o f Figur e 5.5 Th e proj ec ti o n shows tl1e [OOOl] w and [1 1 1) 13 p o l es and r eprese nt s o n e o f th e four rotational variants. Th e int e r sectio n of th e three <4223> 00 p o l es with th e thr ee other <111> 13 p o les is seen in this proj ect ion. Thi s accounts f or the possibility of three a dditional rotation variants that c an b e s up e rimpos e d on th e [000 l] w variant at the [11 l]p zone axis Figur e 5.6 shows th e SAED pattern o f th e ro-type phas e at the [110] 13 zone axis. It co n s i sts o f two zo n e axes ob tain e d from two grains of the ro-type phas e, th e [1100] (1) and [1010] 0)1 t ha t are s up e rimp ose d o n th e [110) 13 zo n e axis. Th e patt e rn shows that th e ( OOOl )(j) reflections from both grains are parall e l to the (111) 13 and ( 111 ) 13 r e fl ect i ons and th at tl1e diffra ct i o n int e nsity is gr ea t e r for one of the grains. Th e v e ry faint diffra c tion s p ots observed at l /3 {110} 13 positi ons co rr es pond t o the {1100} (1) r e fl ec tions of ro-type pr ec ipitat es oriented at th e <1116> (j) zone axes. Th ese originate fr o m the s mall ro-type pr ec ipitat es that w e r e within the selected ar ea diffra ct i o n aperture wh e n the pattern wa s recorded. 5.2.1.3 Orientation Relationship with the J3 Phas e and Plat es In addition to forming an o ri entat i o n relationship with the J3 pha se t h e ro -typ e precipitates also f orme d o n e with the plates that w e r e pr ese nt in th e as-cast mi c r ost ru ct ur e. Figure 5. 7 s h o w s an area o f the mi c ro s tructur e co ntaining all three

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149 ph ases a nd th e SAED pa tte rn o f th e o ri e ntati o n r e lati o nship Th e analy s i s o f th e diffr ac ti o n patt e rn with th e plat es ind e x e d a cco rding to th e ( 0) orthorhombi c s tru c tur e s h o w ed t h e f o ll o wing o ri e n ta ti o n r e la t i o nship : [0001] 00 II [110] 0 II [111] 13 a nd ( 1100 ) 00 II (001 ) 0 II ( 110 ) 13 Th e r e l a tion s hip o b se rv e d b e tw ee n th e co and f3 pha ses i s th e sam e as pr e viously s h o wn in Figur e 5.4 Th e r e lati o nship o b se rv e d b e tw ee n th e plat es a nd th e f3 pha se will b e di sc uss e d f urth e r in c hapt e r 6 Th e mi c rograph in Figur e 5. 7a shows thr ee r ota ti o nal variant s o f th e pl a t es Th ese ar e mark e d 1 2 and 3 and ar e r ot at e d 60 r e l a tiv e t o e ach o th e r Th e larg e co gr a in i s m a rk e d with th e co symb o l and i s o b se rv e d to b e co nn ecte d to t h e plat es. An int e r est ing o b se rva t i o n mad e in th e mi c r o gr a ph wa s th e V shap e o f th e co grain in co nta c t with th e tw o plat e s This ob se rvati o n will b e co nsid e r e d lat e r in t h e di sc us s i o n s in ce i t may provid e inf o rm a ti o n co n ce rnin g th e o rd e r in whi c h th e co phas e and th e plat es f o rm e d from th e f3 pha se. Th e diffr ac ti o n patt e rn in Figur e 5 7b i s c omp ose d of s e v e ral patt e rn s that ar e all s up e rimpo se d on th e [111] 13 z on e axi s Th e d e tail e d a naly s i s o f this c omp os it e patt e rn is shown in t h e cal c ulat e d diffra c tion pa tt e rns o f Figur e 5.8 Figur e 5.8a s h o w s th e co mp os i te o f fiv e diffr act i o n p a tt e rn s fr o m th e f3 matrix co pr ec ipi tate, and thr ee plat e s Figur es 5 8b t o 5.8f show th e individual patt e rns of th e s e phas es. Th e [111] 13 zo n e axi s o f th e f3 matrix i s s h o wn in Figur e 5.8b Th e [0001] 00 z on e axi s o f th e co -typ e pha se i s s h o wn in Figur e 5 8 c and co rr es p o nd s t o th e grain mark e d with th e co s ymb o l in th e mi c r o graph (Figur e 5 7a ) Tl1 e n e xt thr ee diffracti o n patt e rns s h o wn in Figur es 5 8d to 5 8f ar e from th e thr ee plat e s mark e d 1 2 and 3 in th e mi r ..r o graph. Th e t hr ee diffra c ti o n patt e rns s how th e sam e [110] 0 zon e a xis of th e plat es and ar e r o t ate d r e l a tiv e t o eac h oth e r by 60 This r o tation r es ult s from th e o ri e ntati o n

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150 [101J J3 [21l] f3 [OllO] ro [1120 ] ro :i------c k. [112J J3 D [111J f3 C [2243] ro a 00 1] f3 D D [011J f3 C [llOOJro C C [l 26] ro [010] D ] [2113]ro [101J f3 . [ 0 0 0) [lll] f3 C [2116] ro C [ 1 21] f3 [4223Jro [121 [2110]ro C C [110J f3 [1010 ] ro a ' [ 2 11J J3 -[10 ] f3 [1 23 ) ro [1120] ro 0) a [111J f3 [ 2 4 2 3) a C [101] f3 [OllO] ro [110] J3 [1010J ro .. D [112] f3 [1210J ro [121J f3 [2110 ] ro [011J f3 [llOO ] ro Figur e 5 5 Sh o w s th e st e r eo graphi c pr o j ec tion of th e OR r e lati o nship that w as obs e rv e d for th e ro-r e lat e d and J3 phas es in alloy 2. Th e proj ec ti o n s how s th e [lll] p and [0001) 00 p o l e s

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G 0 0 e 151 002 11 112 13 O 001 13 O G 111 11 . 1~10 2 1120 ~ 0002 ~ e') OI O e 0 0 0 O 0 0 Figur e 5 6 SAED patt e rn s h o wing th e OR b e tw ee n th e ro-r e lat e d and~ pha ses a t th e [110) 13 z on e axi s Two r o tati o nal variant s with th e [lTOO] (l) and [lOTOJ (I) zo n e ax es ar e s up e rimp ose d o n thi s diffra c ti o n patt e rn

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152 (a) 0.2m I i ii (b) ... 9 ,. Figure 5. 7. Shows the microstructure observed in the as-cast sample of alloy 2. (a) TEM micrograph showing three plates and a coarse co-related grain observed in the B2 matrix; (b) SAED pattern s howing the orientation relationship observed for the three plates, the co-related grain, and the B2 phase at the [111] 13 zone axis

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011 110 101 . 1100 . ~1210 (a) (b) (c) 110 001 110 110 001 001 (d) (e) (f) Figure 5.8. Shows the calculated diffra c tion patt e rns of the three plat es and the coarse co-re l ated grain at the (111] zone ruris of the B2 matrix. (a) the co mpo site patt e rn ; (b) the (111] zone axis of the B2 matrix; (c) the (0001] 00 zone axis of the co-related grain; ( d ) th e (110] zone axi s o f th e o rthorl1ombic plate l ; (e) the (110] zone axis of the orthorhombic plat e 2 ; (f) t h e (110] zone axis of the orthorhombic plat e 3. Ol C,l,j

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154 r e la t i o nship that th e plat es hav e with th e P phas e ( di sc us se d furth e r in c hap te r 6 ) whi c h c aus e s th e ( 001 ) plan es o f th e pl a t es to li e parall e l to e ach of th e thr ee {lTO} p plan es o f th e p ph ase. 5 2 1 4. L a tti ce Param e t e r D e t e rminati o n Th e lat tice p ar am e t e r s of th e ro pha se w e r e d e t e rmin e d fr o m SAED p at t e rns Th e a a xi s param ete r wa s c al c ttlat e d fr o m th e ds p ac ing m e asur e m e nts o f th e {1100} r e fl ec ti o n s at th e (0001] 00 z on e axi s. Th e c -axi s param e t e r wa s d e t e rmin e d dir ec tly fr o m t h e d -s pa c in g m eas ur e m e n ts o f t h e ( 0001 ) r e fl ect i o n a t b o th th e [1100] 00 a n d [11 2 0l n, zo n e a xes Th ese m eas ur e m e nt s indi c at e d th a t th e latti ce param e t e r s of th e a axis w as 7 96 A and th e c -axi s wa s 5 63 A. It was d e t e rmin e d fr o m th e diffr ac ti o n analysi s that t h e latti ce param e t e r s o f t h e ro ph ase w e r e r e l a t e d t o th e lat t i ce param e t e r of t h e P pha se in th e f o ll o win g m a nn e r : aw = V6 ap a nd c 00 = V3 ap Th ese r e lationship s co rr e sp o nd e d t o th e ds p ac ing s b et w ee n th e {ITOO} CiJ a nd {112} p plan es f o r a 00 a nd b e tw ee n th e ( 0001 ) 00 a nd th e {lll} p pl ane f o r cQ) 5 2 2 Eff ec t o f C o o ling Rat e Th e inv es ti gatio n o f th e effect o f coo l i n g r a t e o n t h e f o rmati o n o f th e ro -t ype ph ase inv o lv e d h eat ing a RAM s ampl e o f all o y 2 t o 1400 C f o r 4 hour s and furna ce coo lin g it in th e v a c uum furn ace. Th e mi c r os tru c tur e that d e v e l o p e d in thi s sa mpl e w as co mpar e d t o t h at o f th e wa te r qu e n c h e d s ampl e und e r th e sa m e h e at tr ea tm e nt co ndi t i o ns (see Figur es 4 10 and 4.11 ) An o pti c al mi c r o graph s h o wing th e mi c r o stru c tur e o bs e rv e d in t h e furn ace coo l e d s ampl e i s s h o wn in Figur e 5.9 It co n s i s t e d of s traight grain boundari es with t ripl e p o in ts and a h o m o g e n eo u s di s tribution o f a c i c ular shap e d pr ec ipitat es in th e

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155 matrix region. How e v e r, the acicular precipitates were not observed along some r eg ions of the grain boundaries. Th e microstructure in these regions resembl ed that of a discontinuous type transformation that was in progr ess. Th e s e observations are consistent with the results obtai o e d in th e phas e e quilibria study (chapter 4) which indicated that only the f3 phase existe d at th e aging t e mperatur e of 1400 Th e additional phases that are observed in th e optical micrograph of Figur e 5. 9 occ urr ed during solid state cooling. Th e TEM analysis of the furnace cooled sample identified two regions in the m ic rostructur e that wer e comprised of diff e r e nt phases. Figur e 5.10 shows the m ic rostructures of these two regions. The microstructure of the in-matrix region (Figure 5.10a) was complex, but was d e t e rmin e d to consist of just two phases. Th e matrix in this region co nsist e d o f an int e rc o nn ec t ed netw o rk of small grains or domains, of the co-type phas e, which showed th e same crystal structure as thos e of the as-cast RAM sample. Th e second phase in this r e gion consisted of plates that w e r e distributed in the matrix. Th ese plates w e r e the same as thos e identified in the h eat treated and wat e r qu e nched samples of this alloy. Th e grain boundary r e gion (Figure 5.10b) was l ess complicated than the in-matrix region and also consisted of two phas es. The matrix was composed of large cr grains with a number of small y grains distributed at the cr grain boundaries. The microstructure of this region was sim ilar in appearance to that of th e eq uilibrium cr + y microstructure that was observed in the 1200 aged sample (compare to Figur e 4 12). Th e results of this investigation indicat e that the co-type phase is aff ecte d by the coo ling rat e from e l e vat e d temperatures and that a s l ow cooling rat e r es ult s in the formation of th e ro-type phas e, instead of plates. How e v e r if the cooling rate is

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156 Figur e 5.9 Opti c a1 micr o graph s h o win g th e mi c r os tru c tur e ob se rv e d in th e 1400 -4hrs F C sampl e o f all o y 2. IOOm

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157 (a) 0.2m (b) 0.2m Figure 5.10. TEM micrographs of the 1400 -4hrs-FC sample of alloy 2. (a) the in matrix r e gion consisting of th e ro-related phase and plat es; (b) the prior grain boundary region consisting of the cr and y phas es.

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158 v e ry s low th e n th e cr a nd y pha se s will form in s t ea d o f e ith e r th e ro-typ e pha se o r pla tes. 5.2 3 Eff ec t of L o w T e mp e ratur e H e at Tr ea tm e nts Th e inv es tig at i o n o f th e e ff ect o f l o w t e mp e r a tur e o n th e form a ti o n o f th e ro t, yp e ph ase w as co ndu cte d wi t h RAM s ampl es th a t w e r e h e at tr e at e d at 400 C a nd a t 600 C f o r 12 h o ur s. Th e TEM an a ly s i s o f th e s a mpl e ag e d at 400 C did not sh o w any s ignifi c ant c hang es in t h e mi c r os tru c tur e as c ompar e d t o t h e as-c a s t co nditi o n. Tl1 e APDB s th a t f o rm e d in th e J3 matri x o f th e a s c ast s ampl e w e r e s till pr e s e nt aft e r aging H o w e v e r it wa s d e t e rmin e d that th e ro-typ e pr ec ipitat e s that f o rm e d in th e J3 matrix w e r e larg e r in s iz e than in th e a s c a s t mi c ro s tru c tur e. B o th o f th e s e ob se rvation s ar e s h o wn in Figur e 5 11 Th e in c r e a se in th e siz e of th e ro pr ec ipitat es c an b e o b se rv e d b y co mparing Figur e 5 .1 la o f th e 400 C a ge d sa mpl e wi t h Figur e 4. 3 o f th e a s-c a s t sa mpl e. Th e r e al so a pp e ar e d t o b e m o r e ro pr ec ipit a t es that w e r e atta c h e d t o ih e APDB s in th e J3 m a trix foll o wing th e 400 C aging tr e atm e nt This i s o b se rv e d in Figur e 5.1 lb whi c h c an al s o b e c ompar e d with Figur e 4 3 o f th e as c ast s ampl e. Th e e ff ec t th at t e mp e ratur e had on th e mi c r o stru c tur e o f all o y 2 wa s mor e pr o n o un ce d in th e 600 C ag e d s ampl e than in th e 400 C ag e d sampl e It wa s f o und tha t th e e ntir e B2 m a trix whi c h was pr e s e nt in th e as c ast mi c r o stru c tur e, h a d co mpl e t e ly transform e d to th e ro -typ e pha se in th e 600 ag e d sa mpl e as shown in Figur e 5 12. H o w e v e r th e analysi s s how e d that thi s aging t e mp e ratur e did n o t a ff ec t th e g r ai n b o undary y allotriom o rphs and in-matrix plat es that w e r e pr e s e nt in th e as cas t s ampl e s in ce n o furth e r nu c l e ati o n gr o wth or diss o luti o n o f th ese ph a s es w as o b se rv e d. Th e a n a l ys i s o f th e tran s f o rm e d ro -typ e pha se s h o w e d th a i th e r e w e r e

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159 (a) 0 4m !, "ft .. p "' ., .., ., .., (b ) 0 2m Figur e 5 11. TEM mi c rograph s o f th e 400 C -12hrs-WQ s ampl e o f all o y 2 (a) s h o w s th e pr ec ipitat es of th e ro -r e l a t e d phas e ob se rv e d in th e B2 matrix ; (b) s h o w s t h e APDB s o b se rv e d in t h e B2 m a trix

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160 (a) O.lm (b) 0.2m Figur e 5.12 TEM micrographs of th e 600 -12hrs-WQ sample of alloy 2. (a) shows the fin e ro-related domains that formed from the B2 phase ; (b) shows the coarse ro -r e lat e d domains that formed at the prior B2 grain boundari es.

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161 tw o s i ze distributi o ns that d e p e nd e d o n th e l oc ation in th e micr os tru c tur e. Th e s mall d o main s w e r e ~200A in siz e and w e r e obs e rv e d in th e prior B2 matrix r e gion (Figur e 5 12 a), whil e th e co ars e d o main s w e r e ~ 0 1 t o 0 5m in s i ze and w e r e o b se rv e d n ear th e pri o r B2 grain b o undari e s and prior plat e s (Figur e 5.12b) Th e c oars e domain s th a t w e r e pr e s e nt in th e 600 C age d s ampl e w e r e s imilar in siz e t o tho se th a t w e r e o b ser v e d in th e as cas t RAM sa mpl e (co mpar e Figur e 5.12b t o Fi g ur e 4 8b ). Th e SAED p at t e rn s tha t ar e s h o wn in Figur e 5.13 w e r e fr o m four diff e r e nt zo n e axes o f th e pri o r B2 ph ase that tran s f o rm e d t o th e co -typ e ph ase. Th e f o ur z on e a x es w e r e id e ntifi e d acc ording t o th e B2 phas e as t h e [lll] p (Figur e 5 13a ), th e [llO] p (Figur e 5 13b ), th e [lOO] p (Figur e 5 13 c), and th e [112] p (Figur e 5.13d ) Th e [11 l] p and [l lO] p zo n e ax e s w e r e s imilar t o th o s e pr e viou s ly sh o wn in Figur e 5. 9 of th e structural a nal ys is From th e pr e vious analy s i s o f th e [lll] p and [llO] p zon e ax es, th e f o ur SAED p a tt e rns s h o wn in Figur e 5 13 indi c at e d that th e mi c rostru c tur e s h o wn in Figur e 5 12 co nsi ste d o f thr ee tr a n s lati o nal variant s f o r e a c h of th e f o ur r ota ti o nal v ariant s, o r tw e lv e to tal vru ~ i a nt s, o f t h e co -typ e ph ase. Th e analy s i s o f th e c oars e ro -typ e domains in th e 600 C ag e d sampl e r e v e al e d t h e pr ese n ce o f int e rnal int e rfa ces that r ese mbl e d APDB s. Th e typ ic al a pp eara n ce o f thi s typ e o f int e rfa ce i s sh o wn in Figur e 5.14 This int e rf a c e wa s obs e rv e d in a singl e r ota ti o nal variant of th e ro-typ e phas e at th e [lTOO] co z on e axis using th e g= ( 1120 )U) r e fl ect i o n. Th e s e in te rfac es w e r e wavy in shap e and w e r e n e v e r s ee n t o int e r sec t with o th e r s imilar typ e s o f int e rfa ce s Thi s o bs e rvation indi c at e d that th ese w e r e tr a n s l a ti o nal int e rf aces and wa s co nfirm e d by th e SAED analysi s that s h o w e d n o extra r e fl ec ti o n s th a t w o uld hav e indi ca t e d tw o r o t a tional variant s

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162 (a) SJ/fj ~ ( (b) Figure 5.13. SAED patterns fr o m the 600 C12hrs-WQ s ampl e o f all oy 2 that s h o w th e diffraction patt e rns o f the form e r B2 matrix (a) [lll] B 2 zone axis; (b) [l 10] 82 zone axis; (co n t inu e d )

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163 I (c) (d) Figure 5.13 (continued) (c) [100] B 2 zo n e axis; ( d ) [112]B 2 zone axis

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164 0.lm Figure 5.14. TEM micrograph of the 600C-12hrs-WQ sample of alloy 2 that shows the APDBs observed in a coarse rotational domain of the co-related phase. The g = (1120) reflection was used to show the APDBs in the coarse domain.

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165 5 3 Discussion In s ummary th e r es ult s o f thi s inv e stigati o n sh o w e d that an ro-typ e pha se (re f e rr e d t o a s roD from this p o int o n ) transform e d from th e high t e mp e ratur e f3 phas e in all o y 2 It wa s d e t e rmin e d fr o m th e as -c a s t R AM and th e furna ce co ol e dsampl e s that a s l o w cooling rat e fav o r e d th e ro -D transformati o n r e lativ e t o that o f th e plat e s whi c h form e d during rapid c o o ling rat e s Th e structural analy s is o f th e ro -D pha se indi c at e d that it was sim i lar in many way s to th e "fam i ly" o f ro ph ases, h o w e v e r th e r e w e r e al so s e v e ral diff e r e nc e s Th e signifi c anc e of th e s e diff e r e n ces w as that th e y indi c at e d that th e ro -D phas e w a s a n e w phase in th e t e rnary Nb Ti-Al s y s t e m that wa s id e ntifi e d in thi s s tudy Th e di sc u ssio n that f o l l o w s will co v e r many asp ec ts of th e B2 pha se t o ro -D ph ase transf o rm a ti o n th a t occ urr e d in all o y 2 Th e obs e rv e d o ri e ntation r e lation s hip h exa g o nal stru ct ur e, and diffus e s cat te ring all signify that th e ro D phas e tran s f o rm e d from th e B2 pha se a nal o g o us to th e {111} 13 p l an e c ollap se m e chanism ( 48]. H o w e v e r t h ere w e r e s e v e ral diff e r e nc e s whi c h se parat e d th e ro -D phas e fr o m th e pr e viously re port e d ro r e lat e d phas es It will b e s hown that th e s e diff e r e n ces ar e c au se d by t h e o rd e ring in th e ro D phas e which c hang e d th e symm e try la t ti ce param e t e r s, and m o difi e d th e o ri e nta t i o n r e l ationship a s co mpar e d to th e ro -r e l at e d pha se s o f thi s t e rnary s y s t e m Th e s it e o c cupan c y o f th e ro-D pha se wil l th e n b e s hown t o b e co n s i s t e n t with th e e xp e r i m e ntal o b se rv a tion s Finally th e c rystallogr a phi c a s p ec t s o f th e f3 t o ro -D pha se t ran s f o rmation will b e di sc us se d bas e d o n co nsist e n c i es with th e e xp e r i m e ntal o b se rvati o ns and with pr e vious w o rk by B e nd e r s ky e t al [ 3 7]

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166 5.3.1 Microstructural Aspects of the ro-D Phase The results of this study indicated that the composition of the ro-D phase must be close to the composition of alloy 2 which is 27Nb-33Ti-40Al (at.%). The J3 phase in the as cast microstructure t1ansformed complete ly to the roD phase upon slow cooling from 1400C and upon holding at 600C. The small amount of y phase that was present in the as cast microstructure before the heat treatments was not expected to have much effect on the composition of the J3 phase matrix and neither was the plates since they will be shown in chapter 6 to form by a displacive transformation. Results similar to these were observed by Bendersky et al. [37] when they performed heat treatments at 700C and higher followed by furnace cooling in an alloy with the Ti 4 Al 3 Nb composition. The significance of this observation is that when it is used with the results from the structural analysis then the site occupancy of the ro-D phase may be determined. The results of the microstructural characterization showed that the precipitation of the ro D phase occurred homogeneously and with limited diffusion from the ordered J3 (B2) matrix. This was determined from Figure 4. 3, which showed a very high number density of small ro-D precipitates that nucleated within the B2 matrix of the as-cast sample. The fact that roD precipitates did not nucleate on the APDBs of the B2 matrix is consistent with this assessment. It is expected that the APDBs sho uld have acted as a heterogeneous nucleation source for the ro -D precipitates since this was observed by Strycor et al. (36] in their investigation. The results obtained from the 400C aged samp l e suggested that the ro precipitates can nucleat e without long-range diffusion, since the APDBs in the B2 matrix were unaffected by this aging temperature as shown in Figure 5. llb. This observation

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167 indicates that diffusion in th e B2 matrix did not occur or this would hav e d est roy e d the a/2<111> 13 APDB v ec tor Th e lack of long rang e diffusion is consistent with r es ults that show a decrease in diffu s ion rates in ordered BCC phases aft e r the disorder to order transition (93]. Th e r e for e, if ro precipitates formed at 400 C as postulated then they did so without th e n ee d for long-range diffusion Thus these observations indicat e d that the ro-D phas e nucleat e d from the ordered 13 (B2 ) matrix of alloy 2 aft er the disord e r to order reaction, that formed th e APDBs l1 ad occurred. Thes e observations also suggest that the transformation occurred with limit e d diffusion in the B2 matrix. Th e coarse roD grains that w e re observed adjac e nt to the plat es suggests a h e t e rog e n eo us nu c l eat ion m ec hanism. It was shown in Figur e 5. 7 that an ro -D grain had a V-shap e morphology that resulted from the impingem e nt of two plates. This o bs er vation may imply that the plat es form e d from the B2 phas e b e for e the formation of the ro -D grain s If this scenario w e r e t o hav e occurred th e n th e coarse ro-D grains co uld hav e then nu c l ea t e d at the interface of the plates at a higher temperature than the fine in-matrix ro -D precipitates. This co uld hav e occurred since heterog eneo us nu clea tion eve nt s occ ur at high er temperatures o r l o w e r und e rco o lings during quenching than h omoge n eous nu c l eat ion eve nts Th e high e r nu c l e ation t e mp erature w o uld fa c ilitat e fast er diffusion rat es, whi c h would ca us e mor e rapid grain growth. This appears to be cons ist e nt with the obs e rvations in this study since low er temperatures would result in limit e d diffusion and in th e formation of smaller in-matrix ro-D pr ec ipitat es. How e v e r an appar e nt contradiction to this was d ete rmin e d from the analysi s o f the plates, which showed that the se plat es formed fr om th e B2 matrix by a displaciv e transformation that requir e d no long-rang e

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168 diffusion (see chapter 6). This r es ult supports the notion that long-rang e diffusion is not required for the precipitation of both fine and coarse co-D grains. Thus there is so m e un ce rtainty as to why the ro -D grains that nucleated heterogeneously on th e plates were larger than the in-matrix ro -D pr ec ipitat es. Th e r e i s a possibility that the larg e co-D grains have a different c rystal str u c tur e than the small in-matrix ro-D pr ec ipitates. This possibility arises from th e results of B e nd ers ky et al. [8] in whi c h co ars e and fin e ro grains w ere observed intermix e d together in th e transformed B2 matrix. Their analysis showed that the fine grains bad trigonal p--g-ml (ro") symmetry which r es ulted from the incompl ete co llap se of' the doubl e lay e rs during the transformation. In comparison, the coarse ro grains w e r e fo11nd to hav e h exago nal P6 3 / mmc (ro-BB;, symmetry, which r e sult e d from che mi cal ordering and co mpl e t e co llaps e of th e doubl e lay e rs. Th e r es ult s of this st udy also showed that the larg e ro-D grains had hexagonal symmetry and P6 3 / mcm space group. Unfortunately, this symmetry could not b e verified using CBED analysis for the in-matrix ro pr ec ipitat es du e t o th e small size. How e v e r it wa s int e r est ing that the SAED patterns of th e small in matrix ro precipitat es of th e as cast microstru cture showed the pr esence of 1/3{110}p r e fl ectio ns. Furth e rmor e, these r e fl ec tions w e r e also observed in SAED patterns of the ro precipitat es pr ese nt in the fully transformed microstructure of the 600 C heat treated sample. Thes e r eflectio ns w e r e dete rrnin e d to co rr es pond to the {1100} (1) lattic e planes of th e ro-D phase bas e d o n the hexagonal st ru cture. Tht1s if the small in matrix ro pr eci pitat es hav e trigonal symmetry, then they must also share a similar chemical ordering to that of th e h exagonal st ru c tur e. Th e diff e r e nc e between the two st ructur es may then be du e only to th e incomplete collapse of the doubl e lay e r and partial ordering on th e

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169 single layers. This seems lik e ly since the in matrix co precipitates were previously determined to have formed at low temperatures where atomic movement is limited. 5.3.2 Comparison of the co-Related Phases in the Nb-Ti Al System A comparison will be made in this section between the co-D phase discovered in this investigation and the co-related phases from the literature This comparison will only include the co-Ti (disordered) [48] and the co-B8 2 [36 37] phases that have hexagonal cry stal structures. The trigonal co structure will not be included since it is considered an intermediate structure that results from the incomplete transformation of the J3 phase to the hexagonal co phase [ 48 37]. Table 5. 4 shows the results of this comparison The results indicate that the co -D phase has several characteristics that are sjmilar to those of the co-related phases. It possesses a hexagonal crystal structure that has 6/mmm point group symmetry. It forms from the J3-phase with the (0001) 00 planes parallel to the {lll}p planes, or simi larly with the [OOOl] w plane normal parallel to the [lll]p plane normal Furthermore in terms of the parent J3 phase, diffuse e l ectron scatteri ng with intensity maxima's near 1/3{111}p reflections are present in the diffraction patterns. These scattering anomalies are commonly observed in co forming systems and are associated with atomic displacements along the p direction. These similarities suggest that the transformation mechanism of the co -D phase is analogous to that whi c h has been determined for tl1e co-related phases [ 48]. However, there were some differences that also existed between the co-D phase and the co -r elated phases. These differences included changes in the space group, the orientation relationship with the J3 phase and the latti ce parameters.

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Tabl e 5.4. Com paris on of c haract e ri.<3tics betw ee n th e co -D the disordered co-Ti, and the o rd e red ro-B8 2 phas es fr om the Nb-Ti-Al t e rnary system. Lattic e Param e t e rs Ori e ntati o n Relati o nships Point Spa ce ro phase G r o up Group a(A) c(A) Parall e l to Parall e l to a ( 1120 )0) co B8 2 P / mmm P6 3 / mmc 4.60 aro= V2a13 5.68 Cw = V3cp [000 l] m ( 1120 )(;) co-D P / mmm P6 3 / mcm 8.03 aco=V6ap 5.68 C 00 = V3c 13 [0001] (1) ( 1100 )(1) the values li s t e d for the l attice param e t e rs w e r e o btain e d using th e r e lations shown in the col11mn to th e right of them with a p = 3.28A

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171 In terms of the space group, the ro-D phase was determjned by CBED analysis to have th e P6 3 / mcm space group symmetry. This differs from the P6/mmm space group of the disord ere d ro -Ti phase [48] and the P6 3 / mmc space group of the ro-B8 2 phas e [36 37]. B e nd e rsky et al [37] proved that th e 6 3 screw axis a nd th e c-glide plane in ro -B8 2 resulted from chemical o rd e ring. Since both of these translati on al symmetry e lem e nt s wer e also observed in th e roD phas e, then this indicates that chemical ordering also occurred in the structure. Th e chemical ordering has to be different how e v er, since the c-glide plane changed from the {1100} 00 plan e in the ro-B8 2 phas e to the {1120} 00 plan e in the ro -D phase. This represents a 90 rotation around the [OOOI] (j) screw axis and r es ults in a symmetry change from the P6 3 / mmc, space group of the ro-B8 2 phas e to the P6 3 / mcm space group of th e ro-D phas e. It was also o bs e rved in th e orientation r e lationships that the {1100}0) planes of the ro-D phase w e r e parallel to the {lTO}p plan es at the [0001]0) II [lll]p zone axes. For the ro-Ti and ro -B8 2 phases th e {1120} 00 plan es li e d parallel to the {110} 13 planes. Tiris difference amounted to a 30 rotation between the unit cells of the ro D phas e and the ro-Ti and ro -B8 2 phas es with ref ere n ce to the lattic e of th e f3 phas e. This result was consistent with the 90 rotation that was observed in the c -glid e plane betw ee n the ro -B82 and ro-D phas es. Finally the dimensions of the unit cell for the ro -D phase wer e co nsid erab ly larg er than that of the ro-related phases. The comparison shown in Table 5.4 shows that the am param eter for the ro-D phas e is 8 03A which is larger than 4.60A for both the ro-Ti (48] and the ro-B8 2 phases [36 37]. Th ese values w ere calculated from the latti ce correspondence b et ween the f3 and ro phases as seen in th e orientation relationslrips. As an exa mple for th e ro-D phase the <112> 13 direction was observed to

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172 lie parallel to the <1120> ro dir ectio n which th e n d efines the a 00 lattic e para m eter of the h exagona l unit ce ll Th e unit l engt h o f' the <112> 13 direction in the BCC unit cell is eq ual to V6ab whi c h gives 8.03.A. for ap = 3.28A Th ere for e, th e relationship sho wing a 00 = V6ab is then obta in ed for th e co-D phas e. In comparison, the orientation r e la t ionship that has been observed for the co -Ti and co-B8 2 phas es indi cates that th e < 110> 13 dir ec tion li es parallel to the <1120> 00 direction Th e unit l e ngth o f th e p direction in the B CC unit cell i s equal to V2a 13 whi c h gives 4 60A for a 13 = 3.25A. Thu s, the relationship s howing a'=V2a 13 i s then obtained for the co-Ti and co-B8 2 phases. A co mparis on of the c 00 latti ce param e t er o f the co-D phas e shows that it is the sa m e as that of the co-B8 2 phas e, how e v e r it is twice that of th e co-Ti pha se Thi s determination wa s based on the r e lationship b et w ee n c 00 and c 13 lattic e parameters whi ch showed it to be c 00 = V3ap for the co D and co B8 2 phases and c 00 = (V3a~/2 for the ro -Ti phase. Th ese relationships w e r e d e riv ed in a s imilar mann e r as described before except that they r e lat e th e [0001] 00 dir ect ion to the <111> 13 dir ectio n Th e /3 and /312 terms are d e riv e d from the body diagonal o f the B CC unit cell which has a l e ngth of V3ap, or 5.68A assum ing a 13 = 3.28.A.. Thi s l engt h is the same as th e c 00 lattice parameter o f the co -D a nd coB8 2 phases, but it i s twice that of th e ro -Ti phas e which is 2.84A. or 1/2 the p diagonal. 5. 3. 3 Transformati o n and Sit e Occupancies of the ro Pha ses Th e transformation of th e ro phas e from the c ubi c f3 phas e has b ee n studied extens iv e ly in the past and wa s exa min e d bri e fly in the lit e ratur e review o f cha pt er 2 Th e purpose of trus sectio n will b e t o exa min e the co phase transformations that hav e been studied in the Nb-Ti-Al system. In parti c ular the relationship b et w ee n the s it e

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173 occupancies of the parent J3 phase and the product ro phase will be examined. Finally the site occupancy of the roD phase will be proposed based on its relationship with the site occupancies of the B2 phase in alloy 2 and the ro B8 2 phase analyzed by Bendersky et al. (37]. 5. 3. 3.1 The ro-Ti and roB82 Phases The transformation of the J3-Ti phase to the ro-Ti phase involves disord e r e d structures of both phases and is shown in the schematic of Figure 5.15. The BCC unit cell of J3 Ti has two lattice sites at (0,0,0) and(,,) at 2a Wyckoff positions and that are occupied by Ti atoms (or a random solid solution). The atoms at these sites are shifted during the transformation such that atoms on every third {111} 13 plane (layers O and 3) remain undisturbed but atoms on the other two successive planes (layers 1 and 2) ar e displaced towards each other. The latter process describ e s th e collapse that forms the double layer in the ro structure. If the collapse is incomplete, i. e the double layer is non planar or commonly referred to as rumpled, then the ro phase has P]"ml trigonal syn-,metry [48]. However if the collapse is complete, then the hexagonal ro structure with P6/mmm symmetry results and the double layer is planar. The structure in this case has three lattice sites at la (0,0,0) and 2d (1/3,1/s,) Wyckoff positions and that are occupied by Ti atoms. The dimension of the hexagonal ro-Ti unit cell along the c-axis is half the length of the body diagonal in the BCC J3 Ti unit cell, or a 13 /2<111> 13 The same treatment applies to the transformation of the B2 phase to the w-B8 2 phase except that both phases in this case have ordered structures. This transformation was examined in detail by Bendersky et al. and related the site occupancies of the B2 phase to the ro phases (37] A schematic of the site occupancies

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[111] ~ (11 l) p plan es Disordered B C C P-Ti phase G 2a Wyckoff Site: Ti 174 Layer Number 3 2 1.5 1 0 [0001] (1) (0001) 00 planes Hexagonal ro-Ti phas e la Wyckoff Site : Ti 2d Wyckoff Site : Ti Figur e 5 1 5. Sh o w s tl1 e at o mi c s it e occ upan c i es o f th e dis o rd e r e d p phas e and th e di so rd e r e d ro pha se f o r Ti

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175 for these two phases from their study is shown in Figure 5.16. The site occupancy information that they used for the B2 phase was from previous w ork by Banerjee et al. [50] which showed that Ti atoms occupied the la Wy ckoff position at (0,0,0) and Al plus Nb atoms co-occupied the other lb Wy ckoff position at(,,). The BCC unit cell is shown in this schematic along the [111] 13 body diagonal and, therefore, the lb Wyckoff position is located at half of this l ength along this direction. The structure of' the co-B8 2 phase is produced in the same manner that was previously described, in which two out of every three {111} 13 planes collapse to form double layers while the third plane remains stationary. However, as observed in Figure 5.16, two major differences distinguish the co-B8 2 phase from the disordered co -Ti phase: an ordered site occupancy in the unit ce ll and a c 00 lattice parameter that has doubled in length. The ordered structure shows Ti and Al atoms occupying two distinct sublattices at the 2d and 2c Wy ckoff positions on the collapsed double lay ers and Ti, Nb, and Al randomly occ upying the lattice sites on the stationary single layer at the 2a Wyckoff position. The ordered s it e occupancy results in the doubling of the cro-axis in the co -B8 2 unit cell. Also shown in Figure 5.16 is that the l ength of the c(l)-axis corresponds to the full body diagonal of the BCC B2 unit cell, or a 13 <111> 13 A significant observation that can be made from Figure 5.16 is that the latti ce sites in the co-B8 2 structure are not fully inherited from the B2 phase, as would be the case if it were truly a displacive transformation. This indicates that some atomic: movements are necessary in order to form the co -B8 2 pl1ase from the B2 phase. The transformation was examined by Bendersky et al. [37] a nd was s h own to occur in steps that involved both displacive and replacive (atom i c exchange between lattice sites) components. A purely displacive transformation from the B2 phase was shown

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[1 ll] p (lll) p planes Ordered BCC p (B2) phase ft la Wyckoff Site: Ti G lb Wyckoff Site: Al and Nb 176 Layer Number 6 5 4.5 4 3 2 1.5 1 0 (0001) 00 planes 2a Wyckoff Site: Random Ti, Nb Al 2c Wyckoff Site: Al 2d Wyckoff Site: Ti Figur e 5.16. Sh o w s th e atomi c sit e o c cupanci e s of the order e d J3 (B2) phas e and th e ro-B8 2 phase from th e r e sults of Bendersky e t al. [37].

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177 to produce an co phase, referred to as co', with trigonal p--g-ml symmetry. It wa s shown that the c h e mi cal order in this structure could occur by atomic mov e m e nts that had no effect on the trigonal symmetry, which was th e n referred to as co". Th e c h e mi cal ordering that produced the co" structure was shown to hav e the prop e r site occupancy on the rumpled double lay e r as wa s determined on the planar doubl e lay er of t h e h exago nal co-B8 2 structure. This p er tains to th e 2c ( Al ) and 2d (Ti ) Wy c k off po sitions of the co-B8 2 unit ce ll shown in Figur e 5.16. Furth e r atomic exc hang e b et w ee n the sing l e and double l ayers in tb e co" struc tur e resulted in an in c r ease in symmetry to the h exago nal coB8 2 structure. This occ urr e d by co mbining th e la and lb Wy ckoff sites of' co" to f o rm the 2a Wy c koff site of co-B8 2 which wer e det e rmin ed to hav e a random occupancy of Ti Nb, and Al. Thus the transformation of the B2 phas e to the co -B8 2 phase can b e viewed as trav ersing through an intermediate trigonal co" str u c tur e with both displacive and r ep laciv e components. 5.3.3.2 Th e Pr o pos ed co-D Phas e Th e co-D pha se was shown to hav e se v e ral similarities to the co-related phases that indicated it transformed from the p pha se by the sa m e {111} 13 plane co llap se mechanism. This wa s detern1ined from the orientation r e lationship in the form of diffuse e l ectron scatte ring at l/3{111} 13 reflections whi c l1 indicat e d that the ( 0001 )"' planes of the ro -D phase w e r e parallel to the {111} 13 plan es of th e p phas e. Furthermore, the c 00 lattice param ete r o f th e co-D phas e was d e t e rmin ed from SAED patte1ns to be ~5. 68 A, or e quival e nt to the a 13 <111> 13 b o dy diagonal o f the BB C unit ce ll Thi s is the same as the c 00 latti ce parameter that was o bs e rv e d by B e nd ers ky et al. l37] and Stryc or et al. [36] for the coB8 2 phase and is twice that o bs erve d 1' or the coTi phase. F1.1rth ermo r e, it was pr evio usly s h o wn that the co -B8 2 phas e has an

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178 ordered structure with Ti and Al occupying lattice sites on the collapsed double lay e r and a random site occupancy on the single lay e r. Thus, th e s e observations indicat e that the ro-D phas e also has an ordered structure. In order to d e t e rmin e th e site occ upancy of th e ordered ro-D phas e it is n ecess ary to first es tablish th e total numb e r of atoms within the h exa gonal unit cell. 'Ibis ca n b e don e by superimposing the unit ce ll of the ro D phas e onto that of the B2 phas e using the orientation relationship that was previously det e rmined. Th e sc h e matic of th e 11nit ce ll correspondence is shown in Figur e 5 17 for the six layers o f the {111} 13 plan es that c orr e lat e to the (0001) 00 plan es. This figure also shows the {1100} 00 plan es, whi c h d e lin ea t e the unit cell of the ro D structure (seen as lin e traces), to b e parall el t o the {110} 13 plan es. Th e site occ upan c y of the B2 structure is ba se d o n the co n c lusion r e a c h e d in chapter 4 and shows Al occupying th e la Wy c koff sites at ( 0 0 0 ) and Nb and Ti occupying th e lb Wyckoff sites in th e B CC uni t cel l. Th e assessment of the two unit ce lls shows that th e r e ar e 18 lattic e sites, o r ato m s in th e B2 structure th at convert to sites in a single unit cell of th e ro D phas e Th e next ste p in th e site occupancy d e termination is to inv e stigat e th e co rr es pond e nc e b e tw ee n the lattic e sites in th e B2 phas e and th e ro-D phas e Figur e 5.17 demonstrat es what happens to the B2 sites wh e n th e {111} 13 plan e s co llaps e occ urs during a di sp la c iv e transformation that involv es no chemical ord e ring If this occ ur s, then th e r es ulting doubl e lay ers of the ro D structure will b e occupied by Al atoms and Nb and Ti atoms on two se parat e sites a nd singl e lay e rs that alt ern at e b etwee n lay e rs of Al atoms and Nb and Ti atom s on two separate s:ites as w e ll Th e probl e m with this arrangement is that the alternating single lay e rs d est roy t h e h ex agonal s ymm etry and, thus produ ce the trigonal symmetry.

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Figur e 5.17 179 :;: ::: ;:, :i. -------i:. ::{t~:;. ::: ~lt : .. . ... .. . . .... . . . ' .. ... :::::: ::;:: O ( a) T itanium Niobium A l uminum & (b) Th e ( 111 ) pr o j ec tion o f t h e B2 ph ase with th e ato mi c s it e occ up a n cy s h o win g Nb and Ti a t o m s occ up y ing th e la Wy c k o ff si te an d Al ato m s occ up y in g th e lb Wy c k o ff s it e (a) s h o w s t h e plan es a t z O O A ( dark ) 0 095 A (light ), and 0 189 A (light ); (b ) sh o w s th e plan es at z 0 284A. ( dark ) 0.379A (light ) and 0 473 A (light ). Th e dash e d lin es d e n ote th e unit ce ll o f th e co -D pha se

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180 In order to obtain hexagonal symmetry there must be atomic shuffling between atoms on the single and double layers during the B2 to c.o -D transformation. The most likely scenario based on the results of Bendersky et al. [37] would be for the sites on the double layers to be occupied solely by Ti and Al atoms as in the c.oB8 2 structure This would have to occur by the diffusion of Nb atoms from the double layers to the single layers and for Ti atoms from the single to the double layers. Assuming that this occurred then it is possible to determine the Wy ckoff positions of the Ti and Al atoms on the double layers since it was determined that the c.o-D phase has P6 3 /mcm symmetry. Using the tables of crystallography [64] it is found that the double layers consist of two 6g Wyckoff sites for each of the Ti and Al atoms. This means that 6 Ti atoms and 6 Al atoms are required to comp l ete the ordering on the double layers or in atomic percent 33%Ti and 33%Al (# atoms on 6g site/# atoms in unit cell) in the hexagonal c.o-D unit cel l. The interesting correlation made from this assessment is that the composition of alloy 2 was determined to be 27Nb-33Ti-40Al (at.%) and, therefore, virtually the entire 33at. %Ti associated with the alloy composition goes to occupying the 6g sites of the c.o-D phase This is based on the observation that the entire p matrix transformed to the c.o -D phase at low aging tempera tures. Furthermore the single l ayers must then be occupied mostly by Nb atoms and excess AL atoms (left over from the 6g site occupancy). Thus, the single layers, which contain 6 atomic sites must be occupied by 27 at. %Nb and ~7 at. %Al based on the alloy composition The site occupancies determined up to this point are still those that have been determined for the c.o -B8 2 phase which indicates that further ordering must have occurred in the ro D phase. The most likely occurrence for the further ordering would

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181 be on the single layers that contain Nb and Al atoms. This is probable since the single lay ers of the ro B8 2 phase were determined to have a random occupancy of 50%Ti, 12.5%Al, and 37.5%Nb (37]. This equates to ~79%Nb and ~2 1%Al (after normalizing) for the sing l e lay ers of the ro D phase. Therefore, if ordering occurred on the single layers as proposed then it only involved Nb and Al atoms. Using this information and applying it to the crystallography tables of the P6 3 /mcm space group, it can be shown that there are two Wyckoff positions, the 2b and 4d sites, that closely match tl1e chemical composition of the single lay ers. This match is obtained if the 2b sites are occupied by two Al atoms and the 4d sites are occupied by four Nb atoms which give 11%Al and 22%Nb occupancy (in relation to the unit cell) on the single layers. These values are c lo se to those which w ere determined from the alloy composition, 7at %Al and 27at.%Nb. A summary of the site occupancies previously determined is shown in Table 5.5. It shows that the single layers are comprised of 2b sites occ upi ed by two Al atoms and 4d sites occupied by four Nb atoms The double lay ers consist of two 6g sites that are occupied by six Ti atoms on one site and six Al atoms on the other site. The exact sto i chiometry of the ro-D phase determined from this site occupancy is 22Nb 33Ti-44Al (at.%), or Nb 2 Ti 3 Al 4 This agrees reasonably w ell with the composition of alloy 2 which was 27Nb-33Ti-40Al (at.%) Finally, in order for the site occ upan cy to be correct, it must also be consistent with the observed symmetry restrictions, i e. the 6 3 screw axis and the c axis glide plane parallel to the {1120}(1) plane. However, both of these symmetry elements are accounted for by the site occupancy as shown in the schematic of Figure 5.18. The six -f o ld symmetry req11irement of the hexagonal crysta l point group is observed for

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182 Table 5. 5. Proposed site occupancy for the ro D phase with Ti 3 Al 4 Nb 2 stoichiometry and P6 3 /mcm (193) space group Site Site Position Number X 1 0 2 0 3 0.333 4 0.667 5 0.667 6 0.333 7 0.333 8 0 9 0.667 10 0.667 11 0 12 0.333 13 0.667 14 0 15 0.333 16 0.333 17 0 18 0.667 note: 6g 1 x = 0. 333 6g 2 X = 0.667 y 0 0 0.667 0.333 0.333 0 667 0 0.333 0.667 0 0 0.333 0 0.667 0.333 0 0.333 0.667 Element Wyckoff z Position 0 Al 2b 0 5 Al 2b 0 Nb 4d 0 5 Nb 4d 0 Nb 4d 0.5 Nb 4d 0.25 Ti 6g1 0 25 Ti 6g1 0.25 Ti 6g1 0.75 Ti 6g1 0.75 Ti 6g1 0 75 Ti 6g1 0 25 Al 6g 2 0 .25 Al 6g 2 0.25 Al 6g 2 0.75 Al 6g 2 0. 75 Al 6g2 0 75 Al 6g2

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183 the (000 l)(i) planes of the single layers at z = 0 and z = cm in Figures 5.18a and 5.18b respectively. Both of these lay ers show Al atoms at the unit cell corners surrounded by 6 Nb atoms. Th e double layers consisting of Ti and Al atoms at z=c(i) and z=cw provide evidence of the 6 3 screw axis. This is observed by focusing on the Ti atoms at z = cru (Figure 5.17 a) and then applying a translation of z = cw along the [000 l] (i) dir ec tion with a 120 rotation that brings them into coincidence with the Ti atoms in th e next layer at z = c(I) (Figure 5.17b). Likewise the c -glid e plane is also observed in Figur e 5.18 and is obtained by the symmetry operation that reflects specific atoms across th e {11"2}0) plane and translates it by a distance of c/2 along the c -axi s. 1 ,hus the site occupancy of the co-D phase is consistent with the observed symmetry r e quir e ments. It should b e mention e d that this argument was fortuitous since th e Wyckoff positions w e re obtained for th e P6 3 / mcm space group and therefore, must incorporat e th e nec ess ary symmetry r es trictions. It is now possible to show in Figure 5.19 the transformation of the B2 phas e to the co -D phase using the site occupancy information previously determined. This figur e demonstrates the displacive character of the transformation in which two out of every three {111} 13 planes (i e layers 1 and 2) co llap se together and the third {111} 13 plane (layer 3) r e mains stationary to form the planar (000 1 )6) plane of the co-D s tru c tur e. If it is assumed that both phases possess their perspective stoichiometries (i.e AB for B2 and A 4 B 3 C 2 for co-D phase) th e n it is possible that th e atomic sites on the collapsed {111} 13 planes can be converted directly into sites on th e planar doubl e layers of the co-D structure without atomic diffusion. However, symmetry dictates that atoms occupying sites on the stationary {111} 13 planes must exchange positions

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Figur e 5.18. 184 (a) .: ftl: t : :::: Niobium Titanium Aluminum (b) [1010] [1120] Th e (0001) proj ec tion o f tl1 e ro-D pha se whi ch is based on th e P6 3 /mcm s pa ce gr o up and Al 4 Ti 3 Nb 2 stoic hi omet ry ( a ) sho w s th e single la yer at z O OA ( dark) and doubl e lay e rs at z = l /4c (light) ; (b) sho ws th e sing l e lay e r at z l/2 c (dark) and doubl e lay ers at z 3/4c (light).

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185 with their s urr o unding s. Th erefo r e, th e rmal activation must als o b e in corpo r ate d in the transformation. Th e ternary co mp os iti o n of all oy 2 i s another reason why atomic shuffling must hav e occurred during th e B2 to co-D transformation. The stoichiometry of the B2 phas e with the ternary composition of' alloy 2 would be A 4 (B 3 3 /C 2 7 ) wh e r e A= Al B = Ti C = Nb and the par e nth es i s indicat es random occupancy on one sublatiice of the B2 structure. There is only 40at. %Al in tl1 e alloy which m ea ns that th e d efic i e n cy in the occupancy of the la Wy c koff sites must b e fill e d by Nb or Ti atoms. This isn't critical si n ce th e real probl e m s occur for th e random occ upan cy o f Nb and Ti on th e lb Wy c k off site shown in Figur e 5.19. Th e co mplications arise from the fact that Nb atoms must b e r e j ecte d from sites on the collapsed {111} 13 planes at layers 1 and 5 and exc hang ed with Al atoms on the stationary plan es of lay e rs O and 6 Th e va cated sites then must be filled by Ti atoms from th e stationary plan e at lay e r 3 Furthermore Al atoms must also exc hang e with Nb or Ti atoms on layer 3 of the B2 s tru c tur e in order to satisfy th e 2b Wy c koff site occupancy requir e m e nt of th e co -D structure. Thus, the ordered ro -D structure n ecess itat es that atomic exc hang es b et w ee n the {111} 13 planes hav e to occur during its transformation from th e B2 phase. Th e atomic exc hanges h o w e v er, do not imply that long rang e diffusion is necessary for the transformation to occur. In fact the diffusion path l e ngth must be on the order of the unit cell dim ens ion for th e ro-D phas e, s inc e the sto ichiom etry of the unit cell is n ear ly the same as that of th e alloy 2 composition. This partly exp lains why the ro-D phase can transform fr o m th e B2 phase at low temperatur es e v e n though long range diffusion does not occur. This was verifi e d in Figur e 5.11 whi c h showed that APDBs wer e still pr ese nt in the B2 matrix aft e r the h e at

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[111] ~ (111) ~ planes Ordered BCC p (B2) phase f) la Wyckoff Site: Al lb Wyckoff Site: Ti and Nb 186 Layer Number 6 5 4.5 4 3 2 1.5 1 0 [0001]0) (0001) 0) planes Ordered ro-D phase 2b, 6g Wyckoff Sit,es: Al 6g Wyckoff Sit,e: Ti 4d Wyckoff Sit,e: Nb Figur e 5 19. Shows th e atomic sit e occupancies of the B2 phase and the ro-D phase

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187 tr ea tm e nt at 400 C f o r 12 hours. In this case th e APDBs had form e d in th e a s c ast microstructur e prior to th e h e at tr e atm e nt. Thus th e high numb e r d e nsity of' fin e ro -D pr ec ipitat e s that transform e d in th e as-ca s t microstructur e simply coars e n e d during th e 400 C h e at tr e atm e nt. 5. 3. 4 Crystallographic Aspects of the ro -D Pha se Transformation Th e r e sults o f this study indicat e d that th e transformation of th e ro-D phas e f ro m th e B2 phas e co uld b e d e s c rib e d c rystallographi c ally as a s e ri e s of transitions t h a t inv o lv e d symm e try Thi s typ e o f d e s c ripti o n was pr e viously c ov e r e d in c hapt e r 2 a nd wa s b ase d o n th e s tudy by B e nd e rsky e t al [37] It was shown in Figur e 2 11 of c h a pt e r 2 that th e transformation of th e ro-B8 2 phas e from th e disord e r e d f3 phas e occ urr e d by a s e qu e n ce of transitions that cross e d a stat e of minimum symm e try whi c h was th e ro"-P 3 ml spac e group. It was also shown that the individual tran si tion s o f th e transformation w e r e conn e ct e d t o g e th e r by subgr o up/sup e rgr o up r e l a tion s in c rystallography Thu s, th e purpos e of tlris s e ction is to show that th e tran s f o rmati o n of th e ro -D phas e o cc urr e d by a similar se ri e s of transitions ex ce pt th a t t h e la s t t ran s iti o n s o f th e ro -D pha se w e r e diff e r e nt fr o m th e ro-B8 2 pha se Th e tran s f o rmation of th e roD phas e was previously shown to hav e o cc urr e d aft e r th e f3 phas e ord e r e d t o th e B2 phas e Th e r e f o r e, th e first s e ri e s of transitions th a t oc c urr e d for th e ro-D phas e w e r e s imilar t o thos e for th e ro-B8 2 phas e b a s e d on subgr o up/supergroup r e lations (64]. Th e s e initial transitions ar e d e scribed using th e s p ace gr o up anrl pha se notations : Im3m ( f3) Pm3m(B2 ) R 3 m P3ml(ro") wh e r e ro i s th e low s ymm e try trig o nal phas e tl1at ha s und e rgon e som e d e gr ee o f' c h e mi c al ord e ring o n th e co llap se d d o ubl e lay e rs [37] Th e physical d e scripti o n giv e n

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188 t o th ese transition s i s a s foll o w s: th e Im3m t o Pm3m transition involv e d ch e mi c al or d e ring in th e J3 m a trix that pr o du ce d tw o translati o nal variants of th e B2 pha se fr o m th e di s ord e r e d J3 phas e; th e Pm 3m t o R3m tran s ition d e t e rmin e d whi c h o n e of t h e f o ur p oss ibl e b o dy diagonal s in th e B2 unit cell b ec ame th e 3 axis of th e trig o nal ro phas e; and th e R 3 m t o P3ml tran s ition d e t e rmin e d whi c h on e of th e t hr ee p os sibl e {lll}J3 plan e s r e main e d stationary whil e th e two oth e r plan es c o llap se d t o g e th e r 1."bus th e s tructural c hang es that o cc urr e d in thes e thr ee transitions w ere th ose that produ ce d th e ro c hara c t e ristics in th e ro-D phas e On ce t hat th e ro "-P3ml pha se was form e d th e r e majnjng transformati o n th a t f o rm e d th e ro -D pha se co uld hav e occ urr e d by t w o diff e r e nt transiti o n path s Th e tw o paths ar e s h o wn in t h e sc h e mati c of Figur e 5.20. In thi s sc h e mati c, th e o rd e r param e t e r whi c h i s see n as th e bra c k e t e d numb e rs indicat e s how many vari a nts o f t h e pr o du c t pha se will b e form e d and th e dir ec tions of th e arr o ws show wh e th e r th e r e will b e an in c r e as e o r a d ec r e as e in th e symm e try of th e produ c t pha se Th e fir s t tran s ition path wa s bas e d on th e r es ults publish e d by B e nd e rsky e t al [37] that s h o w e d thr ee tran s lat i onal variants o f an ro-r e lated phas e that had pr ec ipitat e d fr o m th e co ars e ro B8 2 domain s in a 700 C ag e d sampl e of th e Ti 4 Al 3 Nb alloy Th e ro-r e lat e d phas e that wa s o b se rv e d in th e s tudy by B e nd e r s ky e t al was similar to th e ro -D ph ase that w as o b se rv e d in thi s s tudy s in ce both phas e s bad similar latti ce p ara m e t e r s. Th e roB8 2 phas e wa s s h o wn in th e e arli e r study by B e nd e rsky e t al t o hav e form e d fr o m a matrix co n s isting o f th e ro "-P3ml phas e in th e Ti 4 Al 3 Nb alloy Th e sec ond transiti o n path shown in Figur e 5 20 i s propos e d bas e d on th e r es ult s o f thi s s tudy. This path s hows that an int e rm e diat e trigonal P3lm ( ro"' ) pha se f o rm e d th at p o s se ss e d a minimum in symm e try co mpar e d to th e ro"-P3ml phas e and t h e ro-D

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189 Im3m (A2) [ 2] Pm3m (B2) [ 4] R3m [ 2] [ 3] P3m l (0 '' [ 3] [ 2] P31m ro'" Figure 5.20. Show s the transformation paths from the f3 phas e to the roD phase d esc ribed by s ubgr o up /symmetry relation.'3.

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190 phas e. Thus it was po ss ibl e that th e ro -D phas e co uld hav e transformed along both of th e paths shown in Figur e 5.20. Th e result showing th e 1/3{011} 13 refl ec tions in th e SAED patt e rns of th e B2 matrix in th e ascast RAM sampl e was found t o b e consistent with th e second path in Figur e 5 20 This r es ult indicat e d that th e chemical ordering that form e d th e ro-D phas e had occurred in the very e arly stages of the ro nu c l ea tion proc ess How e v e r in th e Ti 4 Al 3 Nb alloy studied by B e nd e rsky e t al [37] it was shown that a 400 C/ min coo ling rat e c aus e d the transformation of the high t e mp e ratur e f3 phase complete l y t o the co phase. Th e 400 C/ min, coo ling r a t e was much slower than that of the as-cast RAM sample of alloy 2 which co ntain e d a high numb e r d e nsity of s mall ro-related pr ec ipitat e s distribut e d l1omog e n eo usly in the B2 matrix as shown in Figur e 4 3. Thus it was e xp ecte d that th e ro phas e should hav e form e d in th e as cas t RAM sa mpl e, since th e ro" phas e involv es the l ea st amount of c h e mi c al ordering and the partial co llaps e o f th e doubl e l ayers compared to the ro -D phas e. How e v er, th e small co-re lat e d pr ec ipitat es in th e B2 matrix of the as cast RAM sample showed refl ect ions l oca t e d at th e l/3{011} 13 p os iti ons in the SAED pattern o f Figur e 4.5. Th ese w e r e {lT00} 11 r e fle ctions o f th e ro-D pha se whi c h indi c at e d that furth e r ordering over that o f the co '' had alr ea dy occ urr e d in the small ror e lat e d pr ec ipitat es. This was s urprising s in ce tl1 e limit e d t im e that w as availabl e for th ese precipitates to form should hav e pr e v e nt e d the co mpl ete co llaps e of the doubl e lay e rs whi c h was n ecess ary to f o rm the H C P s ymm e try o f the co D pha se Th e r es ult showing no tripl e p o int junctions for th e interfac es obs e rved in th e coarse ro -D d o main s was consistent with th e second path shown in Figur e 5 20 Thi s co n c lu sio n wa s bas e d o n th e numb e r of variants that would hav e formed and,

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191 d e p e nd e d on the transformation path that was tak e n. If th e first path had occurred, then three translati on al variants of the coD phas e would have form e d from th e intermediate co-B8 2 domain, since th e o rd e r param e t er o f th e last transition was three. Th e thr ee translational domains ca n b e visualiz e d using th e schematic shown in Fi.gur e 5.18 by translating the Al atom at the origin o f th e unit cell by 1/3<1 TOO> into the two positi ons occupied by Nb atoms, which ar e on the 4d Wy c koff sites in the zero lay e r o f th e ( OOOl )CI) plan e. If the second path had occurred, th e n by th e sa m e argument it w o uld hav e b ee n expecte d that two translational variants of th e co D phas e w o uld hav e formed from the int e rm e diat e trigonal co"'-P31m phas e. This path agrees with th e observations o f th e 600 C aged sample shown in Figur e 5.14, whi c h s how e d that the single rotational variants of th e coarse co-D domains contained wavy in te rf aces t hat did not int e r sect eac h o th er. Thi s o bs e rvation was co nsist e nt with the presence of tw o translational variants separated by the int er fac e. Th e tab l es o f c rystallography indi cate d that th e co"' was form e d by chemical ordering that occ urr ed in the co" phas e. Th e P3lm space group of the co '" pha se is list e d as a typ e IIb maximal non-isomorphic subgroup of th e co"-P3ml space group. Th e o rd e ring ca us ed an incr ease in th e latti ce param ete rs of th e co" phas e whi c h m a d e th e unit ce ll the same size as that of th e co-D phas e, and rotat e d the 2/m mirr o r sy r nm et ry axis by 90 from th e <1120> dir ec ti o n in the co" phase as compared t o th e < 1100> dir ectio n in t h e co'" pha se. Thi s rotation caused th e 2/m syrnmetry axis in the co "'-P '3" lm phas e to become parall e l to th e 2/m symmetry in the co-D pha se, whi c h had the P6 3 / m c m space group. Tlti s c h e mi c al o rd e ring wa s associated with th e single lay e:r s, since it was previously shown that ordering occurred b e tw ee n Al and Nb atoms o n th e single lay e rs o f th e co -D phase. Th e Wy c k o ff sites on the single lay e r s of

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192 the co"' phase that were consistent with this ordering scheme w e r e the la site at ( 0 0 O) ; the lb site at ( 0 0 ); th e 2c site at ( 1/3 1/s 0) and (1/s, 1/3 O) ; and the 2d site at ( 1 /3 1/s and (1/s, 1/3 ). Thus the ordering occurred betw ee n atoms on the la and 2c sites in th e z = 0 lay e r and atoms on th e lb and 2d sites in the z =layer. This would hav e in c r ease d th e lattic e param e t e rs of the co'" phas e to th e same s iz e as those of the co -D phase. Th e l o w symmetry of the co'" phas e indi ca t e d that the atomic occ upan cy of these si t es was general. Th e o nly r est riction o n the atomic site OCCLl pan cy w o uld hav e been that dissimi1ar atoms had t o occ upy the two types of' Wy c koff sites o n eac h s ingl e lay e r Th e Wy c k o ff sites that w e r e consistent with th e d o ubl e ]ay e rs of the co"' phas e w e r e the 6k 1 and 6k 2 sites. Th e 6k sites accounted for the partial co llaps e o f th e doubl e lay e rs in th e co"' phas e, since th e z-parameter in th e Wy c koff coo rdinat es of th ese sites was variabl e. Th e advantag e to forming th e trigonal co"' phas e as opposed to th e co-B8 2 phase, as th e int e rm e diat e transitional phas e was that th e formation of this phas e inv o lv e d very littl e modification to the s it e occ upan c y and th e in co mpl e t e co llaps e o f th e double lay e r s as co mpar e d to th e co phas e. In comparison the transition path that form e d th e co-B8 2 phas e involved an incr e ase in symmetry which would hav e r e quir e d co n s id e rably mor e atomi c e x c hang e s b e tw ee n th e single and d o ubl e l a y e r s and the co mpl e t e collapse of the doubl e lay e rs. Th e tran s ition fr o m the co" phas e t o t h e co -B8 2 phas e w o uld hav e also r e quir e d th e rmal e n e rgy since Bend e rsky et al. [37] showed that thi s transition occurred after h ea t tr e ating the Ti 4 A1. 3 Nb alloy f o r 26 d ays at 700 C or coo ling s l o wly from l 100 C at ~400C/minute. This result indicat e d that th e coo ling rate associated with the as cas t RAM sample of alloy 2 would hav e b ee n too fa st for the transition o f th e coB8 2 phas e to hav e occ urr e d. Thus it wa s m ore

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193 lik e ly that th e s mall pr ec ipitat es th at f o rm e d in th e B2 matrix o f th e a s c a s t RAM sa mpl e co nsi ste d of t h e ro "' pha se. Th e final tran s iti o n from t h e ro "' phas e t o th e ro-D phas e was e xp ecte d t o r e quir e th e rmal e n e rgy s in ce e x c hang es b e tw ee n Nb and Al atoms on th e s ingl e la ye r s w e r e n ecess ary Th e sit e occ upan c y of th e B2 phas e that was initi a lly inh e rit e d in th e P 3 1m s tru c tur e o f t h e ro "' pha se would hav e pr o du ce d Al ri c h s ingl e la ye r s with th e l a and 2 c sit es and Nb and Ti ri c h s ingl e l a y e r s with th e lb and 2d s it e s. Thus th e Al and Nb atoms would hav e had to e x c hang e sit es b e tw ee n th e two s ingl e l a y e r s H o w e v e r this r e pr ese nt e d th e long e st path b e tw ee n sit e s in th e unit ce ll a nd th e e x c han ge w o uld hav e had t o p ass thr o ugh th e doubl e lay e r s Th ese e x c h a ng es w e r e n ecessa ry t o f o rm th e ro -D s tru c tur e s in ce th e singl e lay e r s in thi s s tru ct ur e hav e t h e sa m e Al and Nb co nt e nt Th e r e for e, it was e xp ec t e d that th e ro -D ph a s e wa s abl e t o o nly f o rm during furna ce c o o ling from th e J3 pha se at 1400 C o r during i so th e rm a l h e ating at 600 C f o r 12 h o urs.

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CHAPTER6 THE ORTH(HEX) PLATES 6.1 Introduction In c hapt e r 4, it was shown (see Figur es 4 8 4.9, and 4 11) that l e nticular precipitates (referred to as plat es) form e d in the BCC 13 phas e of 27Nb-33Ti-40Al (alloy 2) and that composition and cooling rat e affe c ted their formation. Th e effect of composition was observed by the pres e nce of plates in the interdendritic regions of th e as-cast sample which indicat e d that th e composition in thes e regions must hav e b ee n conducive for the nucl ea tion of the plat es to occur. The e ffect of cooling rate was o b se rv e d in th e l evitate d and drop qu e n c h e d sample by th e pr ese nce of a high number density of plates throughout the 13 matrix Thi type of microstructur e wa s also o b served in the h eat treated samples of alloy 2 that w e r e wat er qu e nch e d. Thus these observations indicated that a fast cooling rat e favor e d the formation of the plates from the 13 phas e, while a slow coo ling rate favor e d the formation of e ith e r th e ro -D pha se (intermediate cooling rate) or th e e quilibrium a and y phas es (very slow cooling rat e). Th e purpos e o f this chapter is to examine the formation of plat es from the 13 phase in greater d et ail. Results co nc e rning the structure, the shape morph o logy and the orientation relationship of the plat es with resp ect to th e 13 phas e will be pr ese nted This will b e followed by a discussion of the transformation mechanism and will be treated in t e rms of a displacive transformation using a group/subgroup representation and the invariant lin e concept. 194

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195 6 2 Results Th e inv e stigation of plat e s that form e d in sampl e s during rapid cooling from high t e mp e ratur e s r e veal e d two c l o s e ly r e lat e d phas e s which w e r e bas e d on th e orthorhombi c (ORTH) and the h e xagonal clos e packed (HCP) Bravais structur e s Th e r e sults indicat e d that th e ORTH stru c tur e was th e most pr e val e nt phas e. 1 h e H C P stru c tur e wa s f o t1nd o nly in thi c k pla te s of th e 1300 C h e at tr e at e d s ampl e, h o w e v e r pl a t es with th e ORTI-I s tru ct ur e w e r e al so o b se rv e d in thi s sampl e. Th e r es ults s h o w e d that se v e r a l variant s o f th e OR1"1! stru c tur e e xi s t e d and w e r e d e p e nd e nt o n plat e thi c kn e s s. In addition th e latti ce param e t e rs for th e s e ORTH phas e s w e r e alr ec t e d by th e h e at tr e atm e nt t e mp e ratur e s. Th e r e for e, a total of thr ee diff e r e nt ORTH phas e s and o n e H C P phas e ( som e c ontaining a third phas e) w e r e id e ntifi e d fr o m th e co rr e lation o f plat e thi c kn e ss with h e at tr e atm e nt t e mp e ratur e in thi s s tudy Thi s s tudy al so r e v e al e d that th e plat es contain e d a vari e ty of planar d e f ec t s tru c tur es in c luding anti-phas e domain boundari e s ( APDBs). How e v e r th e r e w e r e s ignifi c ant diff e r e n ces b e tw ee n th e APDBs of th e ORTH and HCP plat e s. Thus th e r es ult s ar e divid e d int o four sec ti o ns that d e al with th e structural analy s is o f th e plat es, det e rmjnati o n o f th e plat e morph o l o gy zig-zag plat e m o rph o logy and int e rnal d e f e ct s tru c tural analysis of th e plat es Th e o ri e ntati o n r e lati o nship is also in c lud e d sin ce it c o ntains information r e garding th e rigid b o dy r o tati o n b e tw ee n th e plat e s and th e f3 phas e. This information was us e d to di s tingui s h b e tw ee n th e H C P and th e ORTH phas e s

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196 6 2 1 Stru c tural Analy s i s Th e s tru c tural analy s i s o f th e plat es i s divid e d int o thr ee sec t io ns : thi c k pl ates with th e H C P ph ase, m e dium t hi c k plat es wiLh th e ORTI 11 ph ase, and thin plat es with th e O RTH2 and ORTI-13 pha se s In e a c h of th e s e s ec ti o n s, th e r es ults p e rtin e nt to th e d e t e rmin a ti o n o f th e c ry s tal s tru ct ur e, latti ce param ete r s, ori e n ta ti o n r e l a tion s hip with th e J3 phas e, and h a bit plan e of th e plat es is d es crib e d 6.2 1.1 Thi c k Pl a t es with H C P Stru c tur e It wa s d ete rmin e d that plat es with c ro ss sec tions typi c ally ~0.15m p o s sesse d t h e H C P c rystal stru c tur e Som e o f th e s e plat e s w e r e found to hav e a c ompli c at e d s tru c tur e du e t o th e pr e s e n ce of an o th e r pha se in addition to th e H C P pha se H o w e v e r b o th typ es o f thi c k plat es w e r e o nly o b se rv e d in th e 1300 C h e at tr ea t e d a nd wa te r qu e n c h ed s ampl e As wa s d esc rib e d in c hapt e r 4 thi s h e at tr ea tm e nt r es ult e d in th e f o rm a ti o n o f th e cr ph ase at 1300 C and th e s ub se qu e nt f o rmati o n o f th e plat es fr o m th e J3 pha se during rapid coo ling Th e mi c r o stru c tur e di s playing th e di s tributi o n o f th ese thr ee pha ses wa s s hown pr e vi o usly ju Figur e 4. llb In th e s ub se qu e nt analysi s of thi s sampl e, it was found that th e thick H C P plat es pr e f e r e ntially f o rm e d n e xt t o cr grain s and oft e n w e r e conn ec t e d to cr grain s Plat es th a t had th e ORTH phas e w e r e als o o b se rv e d in this h e at tr e at e d sampl e, h o w e v e r t h ese pl a t es w e r e l oca t e d a way fr o m t h e cr gr a in s a nd had c r oss se cti o n thi c kn e s ses o f ~0 15m S e v e r a l c har ac t e ri s ti cs o f th e thi c k plat e s ar e s h o wn in th e TEM mi c r o graph a nd SAED patt e rn o f Figur e 6.1 Th e plat e ob se rv e d in th e mi c r o graph of F i gur e 6 la wa s d e t e rmin e d to h a v e a thickn e ss of ~ 0.16m. Th e SAED patt e rn of Figur e 6. lb s h o w e d that th e plat e fo1m e d fr o m th e J3 phas e with th e Burg e r s ori e ntati o n

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197 (a) 1000A (b) Figure 6.1. Sh o ws a thick plat e with the HCP structure observed in th e 1300 C aged sa mpl e. (a) TEM mi c rograph o i' the H C P plat e; (b) SAED pattern of the o ri e ntation r e lati o nship observed b et w ee n th e HCP plat e and th e p phase whi c h wa s [OOOl]H II (011] 13 and ( lTOO )H II (211) 13

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198 relationship (93]: [OOOl]H II [011] 13 and (1120)g II (TIT ) 13 This orientation relationship can be observed in Figure 6.2 with the stereographic projection of the [OOOl]g II [011] 13 poles. From this orientation relationship, twelve variants of the HCP plates were formed from the f3 phase. Six variants were produced by aligning the [OOOl] H direction parallel to the six possible <011> 13 directions and the other six additional variants by aligning the (1120)u planes parallel to the two possible {IIT} 13 planes present at each <011> 13 direction. The diffraction pattern also showed the (lTOO)a and (211) 13 reflections to be parallel and the (011) 13 and (TOlO)H reflections to be separated by an angle of 5.50 (.20). The HCP structure of the thick plates was confirmed by CBED analysis. This is demonstrated in Figure 6.3 using CBED patterns to show the 6-fold symmetry of the [OOOl]H zone axis. Two orthogonal mirror lines are also observed in these CBED whole patterns using a large camera length (Figure 6. 3a) and a small camera length (Figure 6.3b). The presence of the mirror lines and the 6-fold axis indicate that these CBED whole patterns have 6mm symmetry. The symmetry of the whole pattern shown in Figure 6.3b was determined by faint HOLZ lines, since outer HOLZ rings were never observed even after long exposure times. The 6mm symmetry indicates that the HCP phase belongs to the 6mm lR projection diffraction group, as shown in Table 6.1. There are two possible diffraction groups showing the 6mm symmetry in the CBED whole patterns: 6mm or 6mm lR. In eit h er case, both of these diffraction groups are consistent with the HCP crystal structure The point group of the HCP plates was determined by using a second zone axis of higher symmetry, or the [1126] R zone axis in this study, and is shown in

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[211J f3 [llOO]H [lll] J3 [2110]H [II2J f3 Oil D 199 [1210]H (100] (3 [OllO]H l21IJ J3 D D [lll] J3 0 [211] J3 C [101J f3 [l ; OJ J3 [1120]H [lll] J3 D [12l] A [112] J3 JJ [121J J3 [1010]H [011] (3 [001] J3 D D D D [011J J3 [1010]H D D [112J J3 [111) (3 [1120]H [2ll] J3 [OllO]H l \ 12J J3 [121] (3 0 [12lJ J3 [111J J3 [1Ql ] J3 D [211J J3 D D D D [100) (3 [1210]H [110] J3 [2110 ]H [111J J3 [ llOO]H Figur e 6 2 Th e ste r eo graphi c proj ec ti o n o f th e o ri e ntati o n r e l ati o n s hip b e tw ee n th e H C P pha se and th e j3 phas e whi c h show s th e [O O Ol]n II l Oll] p p o l es.

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200 (a) (b) Figur e 6.3. CBED patterns showing the whole pattern symmetry of the [OOOl]H zone axis observed for the thick HCP plates. (a) a large camera constant; (b) a small camera constant.

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Tabl e 6 1 201 Show s the r e l ation b e tw ee n th e possibl e diffraction groups and the symmetr i es observed in the Co nv e rg e nt B e am Electron Diffra c tion (CBED) patterns of the HCP phase at th e [OOOl]H zon e axis. [89] Observed Pr o j ectio n Po ss ibl e Wh o l e P attern Symm etry in Diffra ct i o n Diffraction Symm etry Wh o l e Patt e rn Group Gro up s 6mRmR 6 6mm 6mm 6mm 6mm1 R 6RmmR 3m 6mm1 R 6mm

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202 Figur e 6 4. Th ese C BED wh o l e patt e rns showed a 2-fold symmetry with two o rthogonal mirr o r lin es, o r 2mm s ymm e try using a larg e c am e ra l e ngth (Figur e 6 4a ) and a s mall ca m e ra l e ngth (Figur e 6.4b) Th e 2mm symmetry of' th ese C BED whol e p atte rn s are cons i ste nt with th e 2mm1 R pr o j ect ion diffra c tion group as s h o wn in T a bl e 6 2 Th e 2m R m R and 2 R mm n diffra c tion groups are e liminat e d since th ey d o not s h o w a 2mm symmetry in th e whol e p atte rn. Thi s l ea v es the 2mm or 2mml R sy mm etries as th e possible diffraction groups. Th e final s t e p in th e d e t e rmination of th e point group c onsist e d of co mbining th e r es ult s o f the [OOOl] H and [11] H zone ax es A compiled list from Buxton et al. [91] o f all the p oss ibl e point group s for eac h of th e four pos s ibl e diffraction groups i s s h o wn in Tabl e 6 3. Th e e xamination o f th e tabl e showed that there was only o n e point gr o up that occ urr e d from the diffra c tion groups of both zone axes. This was determined to be th e 6 / mmm point gr o up Sin ce th e r e w e r e no other point gr o ups th at m et t hi s c rit er i o n th e n thi s mus t be the p o int gr o up o f the IICP phas e. Th e s p ace group of th e I-ICP pha se was d e t e rmin e d by th e pr ese n ce of additional translational s ymm e try e l eme nts i .e. a scr e w axis and/or glid e plan e. Thi s st udy e xamin e d the [lTOO] H and [11"2"0] H zone axes and was co ndu cte d o n a numb er o f pl ates. Figur e 6.5 shows a CBED wh o l e pattern of th e [11"] 8 zone axis. This patt e rn shows that the ( 000 l)H r e fl ec tion which is kin e matically forbidd e n do es hav e limit e d int ens ity du e t o dynami c al doubl e diffra ctio n e ff ects. V e ry diffus e dark lin es rtmning n o rmal and parallel t o th e g=(OOO 1) 8 r e fl ec tion w e r e also disc e rnabl e in th ese disc's. Th ese lin es were interpreted as bla c k c r osses kn o wn as G jonn es -M oo di e lin es or GM lin es [89] Th ese bla c k c r osses indicat e d that a sc r e w axis li e d parallel to the c -axi s and that a glide plane li e d parall e l t o th e zone axis and was co ntain e d in the

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203 (a) (b) Figur e 6 4. CBED patterns showing the whol e pattern symmetry of the [1126]H zone axis observed for the thick HCP plat e s. (a) a large camera constant; (b) a small camera constant.

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Table 6.2. 204 Shows the relation between the possible diffraction groups and the symmetries observed in the Convergent Beam Electron Diffraction (CBED) patterns of the HCP phase at the [1126]H zone axis. [89]. Observed Projection Possible Whole Pattern Symmetry in Diffraction Diffraction Symmetry Whole Pattern Group Groups 2mRmR 2 2mm 2mm 2mm 2mm1R 2RmmR m 2mm1R 2mm

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205 Tabl e 6 3. Show s the relation b e tw ee n diffraction groups and crystal point gro up s for th e CBED patt e rns of th e HCP plates. [89] Diffr action Zon e Axis Crystal Point Groups Groups 6mm 6mm [0001] 6mm1R 6/mmm 2mm mm2 6m2 [11 26 ] 2mml R mmm 4/mmm 6/mmrn m3 m3m

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206 Figure 6.5. C BED pa tt e rn showing t h e diffus e bla c k c ross in th e ( 0001) dis c observed in the [1120] H zone axis of th e HCP plat es

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207 c-axis The only possible space group that was consistent with these observations was P6 3 / mmc [89,91] Occasionally boundaries wer e observed to run across the HCP plates. Figure 6. 6a s hows a plat e that contains such a boundary. Figur e 6. 6b shows the SAED pattern from a r egion co ntaining this boundary and it do es not indicate the pr ese n ce of two diff e r e ntly oriented crystals. Furtl1er tilting of the specimen consistently co nfirm ed the existence of' one single pattern that was consistent with the H CP phase How e v e r analysis of the CBED patt e rns from the two sides (A and B shown in Figur e 6. 7 a) of the boundary rev ea l e d the existence of possibly another phase in addition to the HCP phase Th e CBED patt e rn shown in Figure 6. 7a confirms that the A side had 6mm symmetry, which was co nsistent with the [OOOl]H zone axis of the HCP phas e. How ever, the CBED patt e rn shown in Figure 6. 7b reveals that the B side bad 3m s ymm etry. From Buxton e t al. [91] the po ss ibl e diffraction groups s h o wing 3m s ymm etry in the CBED whol e patt e rn ar e 6RmmR, 3mn, and 3m. Non e of' these diffraction groups ar e possible for th e HCP phas e. Thus this n e w phas e must hav e e ith e r the cubic or trigonal Bravais lattice since only th e y are consistent with the 6 R mm R, 3mR, and 3m diffraction groups. Sinc e the tilting of the specimen did n ot r e v eal any other diffra c tion patt e rns beside the one for th e HCP then it was co ncluded that th e phase e xisting on the B side had a trigonal lattic e which can r esul t in an id e ntical SAED patt e rn as one belonging to an HCP phase. The lattic e param e t e r s of only the HCP phas e w e re d ete rmin e d in this study and w e r e ob tain e d from SAED patt erns of th e [OOOl]H and [1120] H zone axes: aH = 5.792A ( 0.026A ) and cH = 4 .6 25A 0 016A) Th e a 8 and cH lattic e param ete rs

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208 Figur e 6 6 Sh o w s th e boundary obs e rv e d in a thi c k H C P plat e ( a ) TEM mi c r o graph ; (b) SAED patt e rn of th e two r e gions s e parat e d by th e boundary whi c h was co nsist e nt with th e ori e ntation r e lati o nship o bs e rv e d b e tw ee n th e H C P plat es and th e f3 pha se ( a ) 1000A (b )

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209 (a) (b) Figure 6. 7. CBED patterns of the two sides (A and Bin Figure 6.6). (a) shows the 6mm symmetry that was cons ist ent with the [OOOl]H zone axis of the HCP phase; (b) shows the 3m sy mm etry of possibly a different phase.

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210 w e r e c al c ulat e d fr o m m e a s ur e m e nt s o f th e ds pacing s f o r {lTOO}u and ( 0001 ) 1 1 r e fl ec tions r e sp ec tiv e ly Th e habit plan e of th ese thi c k pl a t e s was o bt a in e d at th e (0001] 8 II (011] 13 o ri e ntati o n It wa s o b se rv e d that plat e s w e r e ori e nt e d e dge-on wh e n vi e wed in th e [0001] 8 II (011] 13 dir ec ti o n but w e r e in c lin e d slightly wh e n vi e w e d in th e [1120]H II [IlTJ 13 dir ec ti o n Thi s i s d e m o n s tr a t e d by co mparing th e plat e m o rphol o gy s h o wn p1 1 e vi o u s l y in Figur e 6 la at th e (011] 13 zo n e axis to that s hown in Figur e 6.8a a t th e [IlTJ 13 zo n e axis Th e int e rfa ce b e tw ee n th e plat e and th e f3 pha se at th e [Oll] p zo n e axi s s how e d no ov e rlap in th e pr o j ec t e d imag e, whi c h indi c at e d that th e plat e wa s o r ie nt e d e dg e o n r e l at iv e t o th e e l ec tr o n b e am dir ec ti o n. Th e n o rmal t o this int e rfa ce, o r habit normal show e d that it li e d ~10 ( fr o m th e ( 211) 13 plan e normal t o ward s th e ( lll ) f3 pl a n e n o rmal dir ec ti o n. Wh e n th e plat e s w e r e o b se rv e d at th e [I1T] 13 zo n e axi s, as i s th e plat e s hown in Figur e 6.8 a, th e plat e /f3 pha se int e rf ace w as alw a y s parall e l t o th e ( 211 ) 13 plan es and in c lin e d r e lativ e t o th e b e am dir ec ti o n Thi s i s o b se rv e d in th e mi c r o graph o f Figur e 6.8a a s tw o parall e l tr aces o f th e int e rf ace whi c h int e rs ec t th e t o p and bottom o f th e thin foil al o ng b o th sid e s o f th e plat e Th e am o un t o f in c lin a ti o n wa s d e t e 1 1 min e d by tilting th e pl a t e ab o ut th e g = ( 000 l )H r e fl ec tion until th e int e rfa ce was ori e nt e d e dg e on and th e n m e asuring th e tilt fr o m th e g o niom e t e r. Th e r e sult s of thi s e xp e rim e nt ar e s h o wn in tl1 e mi c rograph o f Figur e 6.8b and indi c at e that th e am o unt o f tilt m e asur e d fr o m th e [IITJ 13 zon e axi s co rr e lat es with th e angl e m e a s ur e d b e tw ee n th e habit n o rmal and th e (211 ) 13 plan e n o rmal fr o m Figur e 6 la Th e r e f o r e, applying this m e th o d to a numb e r of thi c k H C P pl a t es s h o w e d th a t th e habit plan e normal was 10 ( fr o m th e ( 211 ) 13 plan e n o rmal

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211 (a) 1000A (b) Figure 6.8 TEM micrographs of a thick HCP plate (a) shows the plate inclined relative to the beam, at the [T1T] 13 zone axes; (b) shows the plate edge-on after tilting ~11 along the g = (011) 13 II (OOOl)H reflections.

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212 6.2.1.2 Medium Thick Plates with the ORTHl Structure The results indicated that plates with thicknesses of 0.05m t 0.15m possessed the orthorhombic (ORTH) str uctur e. The lattice parameter of this phase depended on the heat treatment as will be shown in the following section. Since the diffraction pattern of this phase was different than the ORTH phase observed in thin plates the thicker plates were referred to as ORTHl. In addition, many of the characteristics observed for the ORTHl plates were simi lar to the HCP plates, including the number of plate variants derived from the same type of orientation relationship with the p phase. The plates with the ORTHl phase were observed in the RAM samp l es heated at 1300C, 1400C, and 1500 C and the EM l evitated samples. The SAED patterns of tl1ese plates were indistinguishable from those of the thick HCP plates. However, the CBED analysis showed that they possessed orthorhomb i c and not hexagonal symmetry. Figure 6. 9 shows micro graphs of a medium thick plate that had a thickness of ~0.065 m (650A) and the diffraction characteristics of the ORTHl phase. The SAED pattern of Figure 6.9b was indexed with the [001] 01 direction parallel to the [011] 13 direction and the (110) 01 planes parallel to the (211) 13 planes and sho w ed the orientation relationship: [001] 01 II [011] 13 and ( 110 ) 0 1 II (211) 13 This orientation relationship indicates that twelve variants were formed during the transformation. Six variants were formed by aligning the [001] 01 direction parallel to six possible <011> 13 directions and six more variants were formed by aligning the ( 110 ) 01 planes parallel to either of two (211) 13 planes at each <01 1 > 13 direction.

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213 (a) 1000A (b) Figur e 6 .9 Show s a m e dium thick plate with th e ORTHl s tructur e. ( a ) TEM mi c r ogra ph ; (b) SAED patt e rn of the o ri e ntation r e lationship b et w ee n the ORTHl plat e and the~ phase, whi c h wa s [001] 01 II [Oll] p and ( 110 ) 0 1 II ( 211) p.

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214 The stereographic projection of Figure 6.10 demonstrates that the above crystallographic relationship also includes coincidences of two other directions between the ORTHl and~ phase. It shows the [001] 01 II [011] 13 poles and the additional orientation relationship: (110] 01 II [111] 13 and (001) 01 II (011) 13 The analysis of the orientation relationship observed in the SAED pattern of Figure 6.9b showed that the angular relationship between the (130) 01 plane of the ORTH! phase and the (111) 13 plane of the B2 phase was affected by the heat treatment temperature. It was found that the plates which formed during water quenching from 1300C showed that the (130) 01 planes were parallel to the (T1T) 13 planes. Thus 1 the orientation relationship observed for these plates was exactly the same in appearance as that observed for the HCP plate shown in Figure 6. lb. This result was different than that which was observed for the ORTHI plates that formed during water quenching from z 1400C. The orientation relationship observed for these ORTHl plates showed that the (130) 01 planes were rotated slightly from the (T1T) 13 planes. This rotation was due to the distortions in the a-axis and b-axis lattice parameters of the orthorhombic structure compared to those of the HCP structure. The relationship between the orthorhombic and HCP structures will be covered in the discussion section. The CBED analysis showed that crystal structure of the ORTHl plates was consistent with the orthorhombic symmetry. Figure 6.11 shows a CBED pattern obtained from the same plate previously shown in Figure 6 9. This CBED whole pa t tern showed that the intensity of the (020) 01 reflection was greater than that of the (110) 01 and (110) 01 reflections This result indicated that the symmetry of this

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[lllJ P [ 110)0 [Oll] P [010]0 [110]0 [ 211] p (110)0. C c [Ii21 p a C D a C 215 [100]0 [100J f3 D [2ll] J3 [10l] J3 c [lll] P D [110] f3 [112] J3 [ 21 J P (110)0 [21iJ P [121] f3 D C [111J f3 [110)0 [OOlJ P D D a [010]0 [011 J f3 C C (110) 0 D [010J f3 C D [121J f3 a D [112] J3 ii [10l] J3 [lll] P [2ll] J3 D [100J f3 [010]0 D [110] J3 a o D [ 21IJ P (110)0 [110 ] 0 [lll] f3 Figur e 6.10 Th e stereo graphi c pr o j ect i o n o f the o ri e ntation r e lationship betw ee n th e ORTHl phas e and the J3 pha se which shows th e (001] 0 1 II [Oll]p pol es.

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216 Figur e 6.11 C BED patt e rn showing t h e wh o l e p a tt e rn s ymm e try o f th e [001] 01 zo n e a xi s o b se rv e d for th e ORTHl plat es

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217 CBED pattern was 2mm sinc e it showed a 2-fold syrnmetry with two orthogonal mirror lin es. Thus it was determined from Table 6 4 that the 2mm I R proj ec tion diffraction group was consistent with th e 2mm symmetry of th e whole pattern. From this pr o j ectio n group, there w e r e two possibl e diffraction groups that w e r e consistent with the 2mm symmetry: th e 2mm and th e 2mm IR Th e ORTHl phas e was al so s tudi e d at th e [110] 01 zone axis in order to d ete rmin e th e point group. Figur e 6.12a shows th e CBED whol e pattern of th e [110] 01 zone axis which was found to hav e 2mm symmetry These observations and the us e of Tabl e 6.4 indicat e that the two possibl e diffraction groups w e r e th e sam e as d e t e rmin e d for th e [001] 0 1 zon e axis. This information was compiled in Tabl e 6.5 t o fa c ilitat e th e d e t e rmination of th e c rystal point groups for the ORTHl phase. All p oss ibl e p o int groups ba se d o n h e xagonal c ubi c, and t e tragonal Bravais s tru ct ur es w e r e e lim jnated from this tabl e, since no orientations w e r e found that sho w ed 6-fold 3-fold o r 4-fold axes for the ORTHl phas e. Thus, ther e w e re only thr ee possibl e point groups r e maining: mm 2 222 and mmm [91]. In o rd e r t o d e t e rmin e wh e th er a mirror plan e e xist e d p e rp e ndi c ular to th e c-axis, as in th e mmm point group an ORTH l plat e was tilted from th e [110] 0 1 zone axis along th e ( 001 ) 01 reflection to form a systematic row of reflections as shown in Figur e 6.12b. Tiris procedur e Pliminat ed the problems du e doubl e diffraction at tl1e [110] 0 1 zon e axis whi ch may c ause int ens ity to b e see n for the forbidden r e flections, for e xample th e doubl e diffra ctio n f o r th e ( 001 ) 01 r e fl ect i o n ca n occ ur between the ( 111 ) 0 1 and (TI0) 0 1 r e fl ect ions 1,h e r es ult s o f this tilting procedure indi cate d that r e fl ec tions at odd numb e e d ( 001 ) po s itions disapp eare d and that o nly e v e n numb e r e d (00 1 ) r e fl ec tions r e mained visibl e. 1..,his meant that the ( 001 ) 0 1 r e fl ec tions w e r e forbidd e n and indicat e d that a

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218 Tabl e 6.4 Sh o w s th e r e l ation b et w ee n the possibl e diffraction groups and the symmetries observed in th e Convergent B e am E l ectron Diffracti o n (CBED) patterns of the ORTH! phase at the (001 ] 0 and the (110] 0 zone axes. (89]. Obs e rv e d Projection Possibl e Wh ole Pattern Symm etry in Diffracti o n Diffra c tion Symm e try Whol e Patt e rn Group Groups 2m R mR 2 2mm 2mm 2mm 2mm1R 2nmmR m 2mm1R 2mm

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219 ( a ) (b ) Figur e 6.12. Th e s tru c tural analy s i s o f th e m e dium thi c k ORTHl plat e s ( a ) C BED patt e rn s howing th e 2mm whol e patt e rn symm e try of th e [110] 01 z o n e axis ; (b) SAED pattern obtain e d by tilting th e thin foil sp e cim e n from th e [110] 01 zon e axis along th e g = ( 001 ) r e fl ec tion.

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220 Tabl e 6.5. Sh o w s th e r e l ati o n b e tw ee n th e diffra c ti o n gr o up s and c rystal p o int group s for th e C BED patt e rns o f th e ORTHl plat es [ 89]. Z o n e Ax e s Diffr ac ti o n C ry s tal Point Groups Gr o ups 2mm mm2 [001] 2m R m R 222 and [110] 2mmlR mmm ml R mm2

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221 mirr o r plan e existed p e rpendicular to th e c -axis which was consistent with th e mmm p o int group. Thus this result combined with those showing that the (010) and (100 ) r e fl ec tions wer e al so forbidd e n indicat e d that th e ORTHl phas e has a bas e ce nt e r e d orthorhombic Bravais lattice with th e mmm point group. This suggested that th e ORTHl phas e bad th e Cmmm point space group if no other symmetry elements, such as a sc r e w axis and/or glid e plan e, w e r e pr ese nt. Th e [100] 01 zone axis was e xamin e d by CBED analysis in order to v e rify wh e th e r th e ORTHl phas e had th e same C m c m space group as th e O-Ti 2 A1Nb phas e [13). Th e [100) zone axis of thi s space group would contain evidence of th e 2 1 screw axis in th e [001] dir ect ion and th e c -glid e plan e parall e l to th e (010) plan e s as a black c ross in kinemaiically forbidd e n (001) 01 r e fl ec tions [89]. Th e r e sults of this CBED analysis did not d e t ec t th e black c r oss in th e ORTHl plat e s. How e v e r this analysi s wa s co mplicat e d a s a r e sult of th e thin cross-section of th e ORTHl plat e s and th e proj ec tion probl e ms that th ese plat es had with th e Stu-rounding B2 matrix at th e [100] 01 zone axis. Th ese probl e ms wer e e xaggerated by the fact that dynami ca l diffraction conditions must b e m e t to e nsur e the proper c onditions for forming tl1 e black c oss [89] whi c h m ea nt that the plat es pr ese nt in tlrick ar e a s of th e thin f'oil had t o b e analyz e d Th e latti ce paramet e r s of th e ORTHl phas e wer e det e rmined from SAED patt e rns at th e [001) 0 1 and [110] 01 zone ax es The a 01 b 01 and c 01 lattic e param e t ers w e r e c alculat e d from the d-spa c ing m e asur e m e nts of th e (200) 01 (020) 0 1 and (002) 01 r e fl ec tions r es p ec tiv e ly. Tabl e 6.6 shows that th e latti ce param e t e rs d e pend e d on th e h e at tr e atm e nt It can be seen that th e c paramet e r is almost ind e pend e nt of th e h eat tr e ating co ndition how e v e r th e a and b lattic e parameters in c reas e with th e

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222 qu e n c hing t e mp e ra t ur e Not e that th e l e vitat e d sp e cim e n has th e largest a and b latti ce param e t e r s In this sp ec i m e n b ec aus e of th e v e ry fast cooling rat e, no cr phas e pr ec ipitat e d up o n coo ling. H e n ce, th e co mp os iti o n of th e f3 phas e is e xpe c t e d t o b e s imilar to a s ampl e qu e nch e d from abov e th e f3 tran s u s, i .e. abov e 1400 Th e fa c t t h a t t h e ORTHl ph ase wa s ob se rv e d in th e r e gion s of th e f3 matrix away from th e cr grain s in th e 1300 C ag e d s ampl e s ugg es t s that c h e mi c al co mposition o f th e f3 pha se may play a rol e in th e variation o f th e l atti ce param e t e rs. Th e habit pl a n e of plat es that p o ss esse d th e ORTI-Il phas e was d e t e rmin e d at th e [001] 0 1 II [ Oll ] p o ri e ntation using th e sam e m e thod pr e vious l y d e s c rib e d f o r th e H C P phas e Th e r es ults consi s t e ntly s how e d that th e hab i t p l an e normal of th ese pl a t e s w as 11 ( from th e ( 21 l ) p plan e normal in th e rur ec tion towards th e (ITI ) p pl a n e n o rmal Figur e 6 9 s h o w s th e habit plan e n o rmal r e lativ e t o th e (211 ) 13 p l a n e n o rmal f o r th e plat e. 6 2 l.3 Thin Pl a t es with th e ORTH Stru c tur e Th e thin pl ates with tlti c kn e s ses of ~500.A s h o w e d th e ORTH s tru c tur e, alth o ugh th e ir diffra c tion patt e rn sho w e d distinct diff e r e nc e s from th e ORTHl s tru c tur e. D e p e nding o n th e djffra c ti o n patt e rns th e ORTH stru c tur es obs e rv e d in th e thin plat es ar e c at e goriz e d a s e ith e r ORTH2 o r ORTI-I3 Th e ORTH2 plat es show e d diffu se diffr ac ti o n in th e form of s mall s tr e aks inst e ad o f distinct s p o t s f o r so m e r e fl ec ti o ns in a ppropriat e SAED pat te rns Th e ORTH3 plat es s h o w e d n o diffr ac ti o n f or th e sa m e r e fl ec ti o ns Th e ORTH2 plat es had a thi c kn ess of 100.A and w e r e f o und in a ll o f t h e h ea t tr ea t e d and c a s t s ampl es. Th e ORTH3 plat es w e r e v e r y thin with a thickn ess of 100.A and wer e only obs e rv e d in the tw o sampl e s that w e r e h ea t tr ea t e d a t 1300 C and 1500

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223 T a b l e 6.6 S h o w s t h e l attic e param e t e rs of pl at e s wi t h t h e O RTHl structur e, wh e r e ~ 1 i s t h e ang ] e b e t w e en th e ( O l l )p an d ( 020 ) 02 re fl e ctio n s S a m p l e L atti ce para m et e rs ( A) ~ l Con d ition a ( 0 2 1 ) b ( .059 ) C ( 0 53) L e vitat e d 5.944 9.9 1 4 4.6 2 2 4 67 1 5 00 C 5 922 9 836 4.625 4 67 1 4 00 C 5 808 9 796 4. 62 4 5 33 1 3 00 C 5 77 4 9.95 6 4.625 5 50

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224 Figur e 6 1 3 s h o ws a thin plat e that has a thi c kn e ss o f ~0 020m ( 200A ) and t h e diffr actio n c har ac t e ri s ti c s o f th e OR1..,H2 phas e. 1."b. e SAED patt e rn o b se rv e d in Fi g ur e 6 13b w as ob tain e d fr o m th e [011]~ zo n e axi s and s hows th e sam e pl ate o ri e ntation that w as pr e vi o u s ly s hown f o r th e H C P (Figur e 6.1) and ORTHI (Figur e 6 9 ) pha se s. A magnifi e d SAED patt e rn o f Figur e 6.13b is sh o wn in Figur e 6 14 in o rd e r t o m o r e e a s ily s ee th e diffus e diffra c ti o n o f th e ORTH2 pha se In this patt e rn ve r y f ai nt s p o t s, lab e ll e d a s Rl and R2 w e r e o b se rv e d at positi o ns that corr e sp o nd e d t o t h e ( 020 ) and ( 200 ) r e fl ec tions o f th e ORTHl phas e. How e v e r thos e p o sition s tha t co rr e s p o nd e d t o th e ( I 10 ) a nd ( ITO ) r e fl ec ti o ns o f th e ORTHl pha se lab e ll e d as R 3 and R4 s h o w e d s h o rt diffus e s tr e ak s for th e ORTH2 pha se Thu s, th e e ntir e patt e rn s h o w e d that it co n sis t e d o f al te rnatin g r o w s o f diffra c tion s p o t s and diffus e s t re ak s. Th e diffu se str e ak s vari e d in l e ngth but w e r e typi c ally l es s than ~ 0.045 / A and w e r e e l o ngat e d in a dir ect ion slightly r o tat e d fr o m th e ( lOO )p r e fl ec ti o n It wa s al so o b se rv e d in thi s p a tt e rn that th e r e l at iv e int e nsity o f th e R5 and R6 spots wa s grea t e r than that o f t h e Rl and R2 s pot s. Th e m e a s ur e m e n t s fr o m t h e SAED patt e rn o f Figur e 6.13 w e r e c ondu cte d with o nly th e di s tin c t diffr ac ti o n s pot s Thi s in c lud e d only th e s p o ts that ar e ind e x e d in t his patt e rn a s RI R2 and R 3. Th e c al c ulat e d ds pa c ing s of th e s e r e fl ect i o n s w e r e a s f o ll o w s: Rl = 4 941 0.018A ) R2 = 2.985 ( 0 007 A ) R 3 = 2 550 ( 0 005A ). Th e angl e b e tw ee n th e RI and ( 01 1) 13 r e fl ec tions was m e asur e d t o b e 4.67 ( 20 ). Th e p a tt e rn also s h o w e d th e R3 s pot t o b e r o tat e d s lightly from th e {I11} p r e fl ec ti o n

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225 (a) 1000A (b) Figur e 6.13. Show s a thin plat e with the ORTH2 structure. (a) TEM microgr ap h ; (b) SAED pattern showing the orientation r e lationship b e twe e n th e ORTH2 plat e a nd the 13 phas e at the [Oll] p zone axis.

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226 Figure 6.14. Enlarged SAED pattern showing the diffuse streaks and diffraction spots observed for the ORTH2 plates. The diffraction pattern shows the [O 11] 13 zone axis for the J3 phase.

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227 Th e habit plan e of the plat es with th e ORTH2 phas e was d e termin e d at the l001] 02 II [Oll]p orientation using th e same m et hod d esc rib e d e arli e r. Th e r es ult s co n s i s t e ntly showed that the habit plan e normal of these plat es was 11 ( from the ( 2Tl )p plan e n o rmal in th e dir ect i o n towards the ( lll) p plane normal. Figur e 6.13 a s h o w s the habit plan e normal r e lativ e t o th e (211) 13 plan e normal for th e plat e. Figure 6 15 shows a v e ry thin plat e that was observed t o have a thi c kn ess o f ~ O.Ollm ( 110A) and th e diffra ctio n c haract e ri s ti cs of the ORTH4 phas e. Th e SAED patt e rn observed in Figur e 6.15b was obtained from th e [011] 13 zone axis. A magnifi e d SAED patt e rn of Figur e 6.15b is shown in Figur e 6.16 in o rd e r to b e tt e r o bs e rv e the diffus e diffra c tion of th e ORTH4 phas e. This pattern r ese mbl e d that of th e ORTH2 phas e (Figur e 6.13 ) exce pt that th e r e was n o diffraction int e nsity obse 1~v ed for the po sitio ns that corres pond e d t o diffus e st r e aks in the ORTH2 phase. In this patt e rn v e ry faint diffra c tion int e nsity was observed for spots, lab e ll e d as Rl and R2 which co rr es p o nd e d to th e (020) and (200) r e fl ec tions of the ORTHl phas e (Figures 6.9). How e v e r n o diffraction int e nsity was observed for thos e p os iti ons that co rr es pond e d to th e (110) and (110) r e fl ec tion s o f th ese phas es It was observed in t his patt e rn that th e r e lativ e int e nsity was gr ea t e r for th e R3 R4 and R5 spots than for the Rl and R2 spots. Tl1 e m e asur e m e nt of tl1 e SAED patt e rn shown in Figur e 6.15 was conducted in a si milar mann e r as w as d o n e f o r the previous ORTH2 phase. F o ll o win g are the c alculat e d cts p ac ings of thre e r e fl ec tions whi c h are shown in th e patt e rn mark e d as Rl R2 and R3: Rl = 4 997 ( 0 018A ) R2 = 2.925 ( 0 007 A) R3 = 2.518 0.005A)

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228 (a) 1000A (b) Figur e 6.15. Shows a thin plat e with th e ORTH3 structure (a) TEM micrograph ; (b) SAED patt e rn showing th e o ri e ntation r e lationship b et w ee n th e ORTII3 plat e and th e ~ phas e at th e [Oll]p zone axis.

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229 Figur e 6.16. Enlarg e d SAED patt e rn showing th e missing diffraction spots for the ORTI-13 plat es. Th e diffraction patt e rn shows th e (011] 13 zone axis for th e phase.

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230 Th e angl e b e tw ee n th e Rl and ( 011) 13 r e fl ec tions was measur e d to b e 5.33 (.20). Th e patt e rn also sh o w e d that th e spot mark e d as R3 was again rotat e d slightly from th e {111} 13 r e fl e ction. Th e habit plan e o f th e plat e sh o wn i11 Figur e 6.15a was d e t e rmin e d at t h e [001] 02 II [011] 13 o ri en t a tion u s ing th e s am e m e thod d esc rib e d e arli e r. Thi s analy s is s h o w e d th a t th e l1 a bit plan e normal o f this plat e was 11 from th e (2Tl ) f3 plan e n o rmal in th e dir ec tion toward s th e ( lll)f3 plan e normal. Th e habit plan e n o rmal r e lativ e to th e ( 2Tl ) f3 plan e normal is shown for th e plat e in Figur e 6.15a 6 2.2 Plat e Morphol o gy Th e d e t e rmination of th e plat e morphology involv e d a c omparison of th e imag es of similar plat e s at thr ee diff e r e nt ori e ntations. Figur e 6 17 shows plat e s at th e following thr ee o ri e ntation s: [001] 0 RTH o r [OOOl] HcP [110] 0 R TH o r [11"2JH c P f 1T0] 0 Rm o r [lTOO]H cP (F igur e 6 17 a) (Figur e 6 17b ) (Figur e 6 17 c ) A t t l1 e ori e ntation s s h o wn in Figur es 6 17 a and 6.17b th e plat es hav e w e ll d e fin e d int e rfa ces with tl1 e B2 phas e and high asp ec t ratios of l e ngth (1 ) to thickn e ss ( t ), o r (l i t ). Th e int e rfac e is in c lin e d r e lativ e to th e b e am dir e ction for th e plat e shown in Figur e 6.17b but i t is edge on for th e plat e sh o wn in Figur e 6 17 a Th e s e o bs e rvations ar e co nsist e nt with th e s id es o f a plat e shap e sin ce th e angl e b e tw ee n tl1 e [001] 0 and lll0] 0 dir e ctions i s 90 and this r e pr e s e nts a rotation around a dir ec tion p e rp e ndi c ular t o th e pl a t e fa ce Th e a c tual fa ce o f th e plat e wa s confirm e d by c h oos ing a third o ri e ntati o n s h o wn in Figur e 6 17 c, t o sh o w that th e plat e wa s o ri e nt e d with it s fa ce n o rmal c los e to th e b e am dir ec tion In thi s figur e, th e plat e d oes n o t hav e a w e ll d e fin e d int e rfa ce and its thi c kn e ss wa s c onsid e rably larg e r than

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231 (a) 1000 .A (b) Figure 6 .17. TEM micrographs showing the plate m orp hol ogy that wa s d eterm in ed from the imag es of plat es at th e three o rthogonal dir ections. (a) the [001] 0 zone axis ; (b) the [110] 0 zone axis; (contin u e d )

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232 Figure 6.17 (continued) (c) the (110] 0 zone axis (c) 1000A

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233 a t the o th e r tw o o ri e nt a tion s. Thus this o bs e rvation was also consistent with a plat e shape since th e [110) 0 dir ec tion li es 90 from th e [001] 0 and ~ 60 from th e (110] 0 dir ec tions. Figur e 6 18 shows a plate that was obs e rv e d to be almost e ntir e ly e nclos e d within th e thickness o f the thin foil s p ec im e n This orientation e nabl e d a mor e co mpl e t e assess m e nt o f the plat e morphology. Th e dark and light co ntrast oscillations indi cate that th e thickness of th e plat e incr e as es from it,s out e r e dg es towards it s ce nt e r Th e r e ar e several ot h e r plat es observed in this figur e that are o ri e nt e d n e arly e dg e o n within th e s p ec im e n and hav e thin c rosssec tions. 6.2 .3 Zig-Zag Plat e Morphology It was shown in chapter 4 that the microstru c tures of samples co ntaining plates r ese mbl e d a basket weave app e aranc e Thes e microstructur es w e r e pr e viously shown in Figur e 4.8 for the l e vitat e d and drop qu e nch e d sample and in Figur es 4 9a to 4 9 c f o 1 t h e th e rmally aged and wat e r quenched sa mpl es. Furth e r analysi s o f these mi crostr u c tur es revealed that the b as k et w e av e appearance was du e to plat es tha t had f o rm e d with a z ig-zag morphology from the p phas e, as shown in Figur e 6.19. Th ese mi c rographs wer e obtained by orienting th e sample in such a mann e r so as t o s how the two pr e dominant variants of the zig-zag morphology. Th e g e n e ral o bs e rvations that w e r e mad e from th ese BF and DF micrographs w e r e that pairs of plates w e r e alway s co nn ec t e d at an angle of ~ 60 to one anoth e r Figur e 6.20 shows two plat es that w e r e connected to e ach other along a twin boundary and r o tat e d ~6 0 from e a c h ot h e r Th e twin r e lationship b et w ee n th e two p lat es i s d e mon st rat e d jn th e SAED pattern o f' Figur e 6.20b. This patt e rn shows that

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234 0.2m Figure 6.18. TEM micrograph showing plates that were partially enclosed within the thin foil specimen of the as-cast RAM sample of alloy 2.

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235 (a) 0.5m (b) Figur e 6.19. TEM micrographs showing the zig-zag morphology of the plat es. (a) bright field micrograph ; (b) dark field micrograph

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236 both ORTH plates formed an orientation relationship with the f3 phase as follows: [110]P 1 II [IIT]p and ( 001) P 1 II ( lOT) p [110] p 2 II [IlT] p and (00l)p 2 II ( llO)p For HCP plates, this relationship was as follows : [1120] P 1 II [IIT]p and (0001)P 1 II (lOT)p [1120]p 2 II [T11] 13 and ( 0001 )p 2 II (llO)p (plate 1) (plate 2). (plate 1 ) (plate 2). The diffraction pattern indicated that the (221) 0 or (2"2 0T) H, planes corresponded to the twin boundary since these reflections were superimposed on the (0 l l )p reflection of the f3 phase. The diffraction pattern also indicated that the ( 001) 0 or (OOOl)H, planes o f each of these twinned plates were aligned parallel to the (TOl)p (plate 1 ) and the ( 110 ) 13 (plate 2) planes. Thu s, the int erp lanar angle which separates the {110} 13 planes accounts for the 60 angle and is a contributing factor in the development of the zig-zag morphology of the twinned plates. 6.2.4 Defect Structural Analysis of the Plates The defect structures present in the ORTH and HCP plates consisted of planar faults interfacial dislocations, and anti phase domain boundaries (APDBs). In genera l the thick HCP medium t hick ORTHl, and thin ORTH2 plates s how ed all of these defects. The thin O RTl-13 plates did not appear to show any of these defects. However, these plates co uld not be identified at all of the diff erent orientations that were necessary to conduct the comp l ete defect analysis due to their thin size. The planar faults of the plates w ere always found to have the same crystallograp hi c relationship with the f3 phase, irregardless of the plate phase. The interfacial dislocations w ere always observed for the HCP plates, only occasionally for the

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237 (a) 1000A (b) Figure 6.20. Shows two plat es co nn ec t e d together along a twin b o undary (a) TEM mi c rograph ; (b) SAED pattern at the [IIT] 13 zone axis of the p phas e showing the twin r e lationship b e tw ee n plat e 1 and plat e 2

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238 ORTHl plat e, and n e v e r for th e ORTH2 plat e s The APDBs wer e distin c tly diff e r e nt in th e tI C P plat e s a s c ompar e d to th e ORTH plat e s. Th e planar fault s w e r e ob se rv e d in th e H C P, ORTHl and ORTH2 plat es. Th e s e faults w e r e always obs e rv e d in plat e s that w e r e ori e nted edg e -on at th e [f IT) p z on e ax.i s. From t h e o ri e ntation r e lati o nship s o bs e rv e d f o r th e s e plat es, th e e dg e o n o ri e nt a ti o n o f th e pl a t es p e rtain e d t o th e [110] 0 zo n e axi s f o r th e OR1.1Il plat e and th e [1120] H z o n e axi s f o r th e H C P plat es Th e s e faults w e r e lo c at e d parall e l t o th e ( 001 ) 0 o r ( OOOl ) H plan e s and w e r e f o und t o o b e y th e invisibility c rit e ri o n as sh o wn in Figur e 6.21 Figur e 6 21a indicat e s th e planar fault s w e r e invisibl e using th e (002 ) 0 o r ( 0002 ) u r e fl ec ti o n but w e r e vi s ibl e using th e ( 211 ) 0 or ( 2201 ) H r e fl ec ti o ns as s l1 o wn in Figur e 6.21b. Thes e r e sults ar e co n s i s tent with s tacking fault s and indi c at e that th e y ar e runnin g parall e l t o th e ( 001 ) 0 o r ( 000 l ) H plan e s Di s l oc ati o n s w e r e obs e rv e d at th e int e rfa ce b e tw ee n th e plat e s and f3 matrix Th e int e rf ac ial di s l oca ti o ns o b se rv e d f o r th e pl a t e sh o wn in Figur e 6.2l a app e ar e d t o b e al so assoc iat e d with t h e s ta c king fault s. How e v e r it wa s f o und that th e imaging c onditi o n s that m a d e th e sta c king faults invisibl e did not always mak e th e di s l oca ti o ns invisibl e a s w e ll. Unf o rtunat e ly th e Burg e rs v ec t o r analy s i s o f th ese int e rf ac ial dislo c ation s was in co n c lusiv e N e v e rth e l e ss th e s e int e rfacial disl oca tion s w e r e b e li e v e d t o b e mj s fit di s l oc ati o ns that f o rm e d during plat e thi c k e ning Th e APDB s w e r e obs e rv e d in th e H C P ORTHl and ORTH2 plat e s Th e H C P plat es s how e d APDB s t hat w e r e diff e r e nt from th e APDB s in th e ORTH plat es. Thi s o b se rv a ti o n wa s b ase d o n a co mpari so n o f DF im a g es usin g e quival e nt r e fl ec ti o n s f o r plat es o ri e nt e d e dg e o n a t th e [01 l] p zo n e axi s. Th ese r e fl ec tions w e r e pr e s e nt in th e SAED patt e rns o f t h e [001] 0 z on e axi s f o r th e ORTHl plat e and th e [000 l] H z o n e axis

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239 ( a ) 500A. (b ) Figur e 6.21. TEM mi c rographs showing th e stacking faults obs e rv e d in th e plat es a t th e [110] 0 or [ 1120 ] 8 II [ T1T] B 2 zone ax e s (a ) shows that th e st ac king faults w e r e invisibl e using g = (002) 0 or (0002) 8 ; (b) shows that th e stacking faults wer e visib l e using g = (211) 0 or (220 l) g

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240 for th e H C P plat e s Following ar e th e e quival e nt r e fl ec tions that w e r e us e d t o form th e DF imag e s: g = (I 2TO )H = ( 200 ) 0, g = ( 2110 )a = ( 1 ~ 0 ) 0 g = ( 1120 )H = ( 130) 0 Additional DF imag es w e r e als o obtain e d as n ece ssary in ord e r t o d e t e rmin e th e pr o babl e APDB v ec t o rs in th e appropriat e ORTH and H C P phas e s. 6.2 4 1 APDBs in H C P Plat e s Typi c al APDBs that w e re obs e rv e d in th e H C P plat e s using r e flections a t th e [OOOl]H II [011] 13 zon e axes ar e sh o wn in Figur e 6 22 Th e micrographs indicat e that s mall domrun s f o rm e d in th e plat e s and w e r e visibl e in all thr ee DF imag es. Th e DF im a g e o f Figur e 6 22 a wa s obtain e d using th e (I2TO ) H and ( 100) 13 r e fl ec tions in o rd er to s how th e APDB s o f th e H C P and B2 pha ses r e sp ec tiv e ly. Th e r e was s om e i ndi c ation in thi s mi cr ograph that th e APDB of th e B2 phas e pass e d through th e plat e/ p int e rfa ce and int o th e plat e. How e v e r th e s e APDBs w e r e n e v e r o bs e rv e d t o pas s e ntir e ly acro ss t h e plat e, but inst e ad c urv e d into th e plat e and join e d th e s mall e r domain stru c tur e. It was also obs e rv e d that th e APDBs s ee n wi t h th e (I2TO )H r e fl e cti o n w e r e e longat e d in a dir e ction parall e l to th e plat e /p int e rfa ce. Th e DF imag es of Figur e 6 22b and 6.22 c w e r e form e d with th e (2110) n and (1120) a r e fl ec tion s r e sp ect iv e l y a nd sh o w e d a c olumnar morph o logy that e xt e nd e d fr o m t h e midsec tion o f th e pl a t e o utward t o th e plat e /p int e 1fa ce Th ese c olumnar s hap e d APDB s e xt e nd e d in a normal dir ec ti o n from th e midrib whi c h was not always w e ll d e fin e d t o th e int e rf ace.

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241 (a) 1000.A (b) Figur e 6.22. TEM dark fi e ld micrographs showing the APDBs observed in the HCP plat es. ( a ) g = (I2TO)g and ( lOO )p; (b ) g = (2110) H; (continued)

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242 Figure 6 22 (continued) (c) g = (1120)H (c) 1000A

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243 Typi c al APDB s that w e r e obs e rv e d in th e H C P plat e s also show e d a columnar m o rph o logy wh e n vi e w e d at th e [1120] H II [111 ] 13 z o n e ax e s as shown in Figur e 6 2 3 Th ese APDBs w e r e o bs e rv e d in DF imag e s tha t w e r e form e d with th e (lTOO) H and ( llOl )H r e fl e ction s. Figur e 6 23b s h o w s th e APDB s using th e ( lTOl)H r e fl e ction and indi c at es that th e co lumnar APDB s in th e plat e w e r e s e parat e d by a midrib F i gur e 6.2 3a s h o w s a BF im age o f th e plat e f o rm e d with th e ( 2200 )H r e fl ec tion in o rd e r t o s h o w that th e APDB s and midrib w e r e invisibl e at this imaging c o ndition. Th e APDBs w e r e found t o e xhibit an e quia xe d morph o l o gy for plat es ori e n te d a t th e [21 lO]H zon e axi s as shown in Figur e 6 24 The APDBs at this o ri e ntati o n w e r e b e st o bs e rved u s ing th e DF imaging conditions of Figur e 6.24b which w e r e form e d using th e ( Ol l l ) H r e fl ec tion. H o w e v e r th e APDBs still e xhibit e d som e r es idual c ontrast using th e BF imaging co nditions of Figur e 6 24a whi c h was f o rm e d with th e ( OlTO )H r e fl ec ti o n In additi o n s ta c kin g fault s w e r e also o b se rv e d in th ese figur es t o b e parall e l to t h e ( 000 l )H pl a n es H o w e v e r th e APDBs app e ar e d to b e unaff ec t e d by t h e s t ac king faults s in ce th e y did not t e rminat e at th e faults. 6 2 4.2 APDBs in ORTHI Plat es Typical APDB s that w e r e obs e rv e d in th e ORTHl plat e using r e fl e cti o n s at th e [001] 0 II [011] 13 z o n e ax es ar e shown in Figur e 6.25. Th e DF imag e of Figur e 6 25a wa s o btain e d with tl1 e (200 ) 0 and ( 100 ) 13 r e fl ec tions in ord e r to show th e APDBs o f th e ORTH and B2 phas es, r es p ec tiv e ly. This micrograph shows that th e APDB o b se rv e d in th e B2 matrix maintain e d c ontinuity up o n passing through th e plat e /f3 int e rfac e s o f t h e p l a t e. Th e r e w e r e n o o th e r d o main s o b se rv e d in th e ORTH plat es wh e n imag e d with th e ( 200 ) 0 r e fl ectio n. Th e r e f o r e, th e spa c ing b e tw ee n th e s e gm e nts of th e co ar se d o main s ob se rv e d in t h e ORTH plat e s w as c omparabl e t o th e siz e o f th e APDB s

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244 (a) O.lm (b) Figure 6.23. TEM micrographs showing th e columnar shaped APDBs observed in the H C P plat es at the [1120]H II [111) 13 zone ax es. (a) bright fi e ld micrograph formed with g = (2200)H; (b) dark fi e ld micrograph form e d with g = (lTOl)g.

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245 (a) O.lm (b) O.lm Figur e 6.24. TEM micrographs showing the equiaxed morphology of the APDB s observed in the I-ICP plat es at the [21 lOJn zone axis. (a) dark field micrograph formed with g = (OlTO)n; (b) dark field micrograph formed with g = (0111) 8

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246 o bs e rv e d in the B2 matrix. Th ese coarse APDBs of th e ORTH plat es co ntr as t ed gr eat ly with th e APDBs observed in the HCP plates of Figur e 6 21a as both DF imag es w e r e obtained at comparable r e fl ec tions. Th e DF images obtained with the ( 130 ) 0 and ( 130) 0 r e fl ect ions in Figur es 6 25b and 6 25 c, r es p ec tiv e ly showed small APDBs with a co lumnar morphology that e xt e nded from a mid-rib near th e ce nt e r of the plate o utward towards th e plat e/ ~ interface. Th e appearance of the APDB s wa s s jmilar in both DF im ages. Th e growth direction of th ese co lumnar shaped APDBs w as ~3 1 from th e int e rf ace normal whi c h was n e arly parall e l to the (310) 0 plan es. Th e width o f th e d o mains wa s generally l ess than 200A. Th e fin e APDBs observed in the ORTH plates also showed a columnar morph o logy wh e n vi e w e d at th e [110] 0 II [011] 13 zone ax es as shown in Figur e 6.26. Th ese fin e APDB s w ere observed in DF imag es f o rm e d with ( 110) 0 and ( 111 ) 0 r e fl ect ions. Figur e 6.26 was obtained with the ( 111 ) 0 reflection to show th e APDB s a nd the midrib that separated the fine co llimnar APDBs in this plat e. Th e fine APDBs exhi bit ed an e quiax ed m o rph o l o gy f o r plat es o ri e nt e d at the [110] 0 zo n e axis as shown in Figur e 6.27. Th e DF mi c rograph in this figure wa s o btain e d with the ( 111 ) 0 r e fl ec tion. Sta c king fault s w e r e also observed in Figur e 6 .2 7 to be parall e l to the ( 001 ) 0 planes How e v e r th e fin e APDBs appeared to b e unaff ec t e d b y th e stacking faults since they did not terminate at th e faults 6.2.4.3 APDBs in ORTI-12 Plat es Th e analysis of the APDB s o f the ORTH2 plat es showed that th e r e w ere many s imilar i ti es with th e APDBs that w e r e observed in th e ORTHl plat es. Co ar se APDB s w ere o b se rv e d that pa sse d through th e int e rfa ce from the B2 phas e into th e ORTH2 p lat es. Lik e wis e, a high numb e r d e n s ity of v e ry fin e domains w e r e also pr ese nt in

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247 (a) 1000A Figure 6.25. TEM dark fi e ld mi c rographs showing th e APDBs observed in the ORTHI plates. (a) g = (200) 0 and (IOO)p; (b ) g = (130) 0 ; (co ntinu e d ) (b)

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248 Figur e 6.25. (continued) (c) g = (130) 0 (c) 1000A

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.. 249 1000A Figur e 6 .2 6. TEM dark fi e ld micrograph s h owing the co lumnar shaped APDBs observed in t h e ORTHl plat es at th e [110] 0 II [011] 13 zone axes using g = ( 111 ) 0

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250 1000A Figur e 6.27. TEM dark field micrograph showing the e quiaxed morphology of the APDBs observed in th e ORTHl plat es at the [l 10] 0 zone axis using g=(lll) 0

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251 th e ORTH2 plat e s and ar e sh o wn in Figur e 6 28 r rh e s e fin e domain s w e r e d e t ec t e d in DF imag e s form e d with th e small diffus e str e aks obs e rv e d for thos e plat es o ri e nt e d a t th e (011) 13 zon e axi s, similar to Figur e 6 15 for th e ORTH2 structur e. Th e position o f thi s str e ak was n e ar th e (130 ) 0 r e fl ec tion in r e f e ren ce to th e ORTHl plat es Th e d o main s w e r e f o und t o hav e a columnar morph o logy that e xt e nded from a midrib in th e pl a t e t o ward s t h e int e rfa ce in a dir ec ti o n n o rmal t o t h e s tr e aks Th e width o f th e d o m a in s w as g e n e rally l e ss than 100.A a nd th e r e ciprocal length of th e diffus e str e ak s in th e SAED patt e rn s wa typi c ally l ess than 0.04 / .A In summary th e analysi s of th e APDB s indicat e d that th e APDBs ob se rv e d in th e H C P plat e s w e r e diff e r e nt from th ose obs e rv e d in th e ORTI-Il and ORTH2 plat es. Th e thr ee diff e r e nt ORTH phas es sh o w e d th e sam e typ e of APDBs in th e s e plat e s 11i e APDB s pr ese nt in th e H C P plat e s w e r e n o t found t o b e d e p e nd e nt ttp o n th e diff e r e nt r e fl ec tion s that w e r e u se d t o f o rm th e DF imag e s. This was opposit e of that o b se rv e d f o r th e ORTI-I plat es whi c h s h o w e d b o th fin e and co ars e APDB s that d e p e nd e d o n th e s p ec ifi c DF imaging co ndition s Tabl e 6. 7 lists all o f th e r e fl ec ti o ns t h a t w e r e u se d in t h e APDB study and th e visibility/invisibility co nditions ob se rv e d i 'o r th e APDB s in th e H C P ORTHl and ORTH2 plat e s. 6. 3 Discussion of th e Plat e Transformation 6 3 1 Mart e n s itic Transformati o n Th e r e sults of this study sugg e st that th e plat e s form from th e J3 phas e t hr o ugh a mart e n s iti c transf o rm a tion R a pid c o o ling rat es o btain e d by wat e r qu e n c hing and s pl a t qu e n c hing r es ult e d in th e f o rmati o n of plat es H o w e v e r s l o w e r coo ling rat es w e r e f o und t o r e sult in th e formati o n of th e co -pha se and th e cr + y

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252 Figure 6.28. TEM dark fi e ld mi c rograph s h owing the co lumnar shaped APDBs observed in the thin ORTH2 plates using the diffuse streak. Th e spec im en was tilted fr o m the [011] 13 zone axis of the~ phas e. 500A

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253 Tabl e 6. 7 Th e im a ging co ndition s of th e APDBs o b se rv e d in th e p l at es. R ec i procal Plat e Str u c tur e Latti ce g V e ct o r H C P O RTHl O R TH2 O R TH3 {1 1 2 0 }H + -------{ l lOO }n + -------{22 00 }H 0 ------{1 l O l }n + ---------{0002} H 0 ------{2202} H 0 -------( 200 )0 ---A A A ( 020) 0 --A A A ( 110 ) 0 ---B B B ( 220 )0 ---0 0 0 ( 130) 0 -B B B ( 111 ) 0 ---B B B ( 0 0 2) 0 ---0 0 0 Th e O m e an s th e APDB s w e r e n o t o b se rv e d For t h e H C P phas e, + m e an s th e APDBS w e r e o bs e rv e d but with n o diff e r e n ce s b e tw ee n th e m. For th e ORTH p h as es, A p e rtain s t o t h e co ar se APDB s whil e B p e rt a in s t o th e fin e APDBs. Th e fin e APDB s in Lh e O RT H p l at es co uld po ss ib l y hav e o b sc ur e d t h e d e t ec ti o n o f th e coars e APDBs du e to th e s mall s i ze and high numb e r d e nsity

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254 microstructures. Th e lenticular plat e morphology and th e associat e habit plan e ar e al s o consist e nt with th e mart e nsit e morphology Th e following discussion d e scrib e s th e basic conc e pt of th e invariant lin e th eo ry [94] so as to s how th e proc e dur e for cal c ulating th e habit plan e Th e pr e dict e d habit plan e is compar e d to thos e obs e rved by TEM. Th e agr ee ment betw ee n th e c al c ulat e d and ob se rv e d habit plan e s c onfirm that th e ORTH and H C P plat e s ar e f o rm e d fr o m th e ph a s e by a mart e n s itic tran s f o rmation. Th e invariant lin e th e ory is c onsist e nt with th e c lassical or ph e nom e n o logical th eo ry of mart e nsiti c transformati o ns [79]. Both th e ori e s ar e bas e d on a diffusionl e ss tr a n s f o rmati o n that inv o lv e s th e minimization o f s train e n e rgy whi c h occ ur s fr o m a s h e ar-typ e latti ce d e formation A latti c e corr e spond e n ce is first chos e n b e tw ee n th e par e nt and produ c t phas e s. Th e habit plan e is th e n calculat e d by applying to th e l a tti ce c o rr e spond e n ce : a homog e n e ou s strain that conv e rts on e Bravais lattic e into an o th e r ; an inhomog e n eo us strain that pr o vid e s additional displac e m e nts but c aus es n o c hang e in s hap e to f o rm th e final s tru c tur e; and a rigid body rotation that f o rm s t h e invariant habit plan e [79]. Th e s trajn e n e rgy fr o m th e mart e nsiti c tr a n s formation aff ects tl1 e kin e tics and th e morph o logy o f th e product phas e. Th e pr o du c t phas e typi c ally has a plat e morph o logy and glis s il e int e rfa ce s tl1at all o w th e co mp o sition including atomi c ord e r and latti ce d e f ec ts s u c h as APDBs of th e par e nt pha se t o b e inherit e d by th e produ c t phas e [95]. Th e classical th e ory and th e invariant lin e tl1 e ory both stat e that an invarian t lin e e xists in th e habit plane [79 94] Th e invariant lin e is a v e ctor c ontained in both t h e par e nt a nd pr o du c t phas e s that i s un e xt e nd e d and unrotat e d by th e mart e n s iti c t ran s f o rm a ti o n Alth o ugh both th eo ri es c an b e us e d t o c al c ulat e th e invariant lin e,

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255 they will predict different values since the invariant line is treated differently in these theories In the classical theory the invariant line is considered a strain-free direction while, in the invariant line theory it is considered a stress free direction. Th e invariant line theory considers that a thin plate with principal strains of m:ixed signs are in a two-dimensional hydrostatic stress state, or plane stress condition [94]. The plate changes from a two-dimensional to a uniaxial stress state by rotating into a position that contains the uniaxial stress axis and the invariant line in the habit pla ne. The uniaxial stress axis is normal to the invariant lin e in the habit plane and is the axis that the rigid body rotation occurs around. This co ndition requires that two of the three principal strains 8 1 and -8 2 must have opposite signs and that the third one, 8 3 is nearly eq ual to zero [9 4 ] If this condition is met then the e la stic strain 8 3 that is present in the direction of tl1e uniaxial stress axis causes a Poisson contraction that is superimposed on the two other principal strains by an amount equal to v8 3 Thus the calc ulat ed position of the invariant line using the invariant line theory will be modified from that using the classical theory when the Poisson contraction on these two principal strains are also considered. The procedure that was us ed in this study for ca l c ulating the invariant line of the ORTH and HCP plates is covered in the following discussion. This procedure is based on the invariant line theory and is described for the ORTHl plate observed in the 1400 C aged sample shown in Figure 6.9 and for the HCP plate observed in the 1300 C aged samp l e shown in Figure 6.1. The calculations were done using th e specific latti ce parameters of the ORTHl and HCP plates observed in these figures. This was necessary since the calc ulati ons sho w e d that the habit plane was sensitive to the lattice parameters of the J3 phase and the different phases of the plates. In

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256 o rd e r t o t es t th e a cc ura c y of th e invariant lin e th e ory th e habit plan e m e asur e d f o r th e ORTHl plat e in Figur e 6 9 and th e H C P plat e in Figur e 6 1 had to b e c ompar e d to th e habi t plan e that wa s c al c ulat e d u s ing thi s th eo ry Th e first st e p in th e pro ce dur e for c al c ulating th e habit plan e of th e ORTIIl and H C P pl a t es wa s to c al c ula te t h e h o m o g e n eo us distorti o n (B ) fr o m th e prin c ipal axes o f s tr a in ( A ) th at f o rm e d during th e m a rt e nsiti c tran s formati o n s o f th ese p lat es. Th e r e ar e thr ee prin c ipal ax es o f s train whi c h ar e th e e ig e nvalu es o f th e lin e ar tran s f o rm a tions that d e fin e thr ee o rthog o nal dir ec ti o ns in th e par e nt J3 phas e th a t tran s f o rm into thr ee o rth o gonal dir ec tion s in th e produ c t phas e [ 79] Th e s e o r t h o g o n a l dir ec ti o ns al s o und e r go a rigid b o dy r o tati o n (R ) as part o f th e mar te n s iti c transformation, but thi s rotati o n do es n o t c aus e any dim e nsional c hang e s to th ese dir ec t io n s. Th e h o m o g e n e ous di s torti o n (B ) is d e s c rib e d by a 3 x 3 matrix with th e t hr ee prin c ipal ax es o f s train as f o ll o w s: r "'1 B = 0 0 0 1 0 + 0 83 0 0 0 1 0 0 1 wh e r e th e ma t rix to t h e l e ft o f th e e qual s ign s h o ws t h e p rin c ipal ax es o f s train or e ig e nvalu e s whil e th e first matrix f o ll o wing th e e qual sign s hows th e prin c ipal s tr a in s and th e la st ma t rix i s an id e ntity m a trix. Th e s ub sc ript s 1 2 and 3 r e pr ese n t dir ec ti o ns al o ng th e (100] 13 (011] 13 and (011] 13 axes, r e sp ec tiv e ly o f th e J3 pha se. B y d e finiti o n th e prin c ipal s train s o f th e h o m oge n eo u s tr a n s l 'o rmation ar e d ete rmin e d f r o m th e e ig e nvalu es by th e r e lati o nship ei = "-i 1 wh e r e i d e fin es eac h o f th e thr ee o rth o g o nal ax es [79]

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257 The principal axes of strain w e r e calculated from the lattice par::imeters and the lattice correspondences b e tween the p phase and each of the ORTHl and HCP phas es The lattice correspondences were first determined from the orientation relationships as follows for the ORTHl phase (Figure 6.9): (100] 13 II [100] 0 [011] 13 II [010] 0 [011] 13 II [001] 0 and for the HCP phase (Figure 6 1): [lOO] p II [1120]H [011) 13 II [IIOO]H [O 11] 13 II [000 l]H. The principal axes of strain were then calculated from the lattice correspondences with the lattice parameters of the p, ORTHl, and HCP phases. Thus, following are the matrices of the homogeneous distortions determin e d from these calculations for the ORTHl phas e: 2ap/a 0 B = 0 0 for the HCP phas e : 0 .J2ap /aHcos(30) 0 1 083 0 0 1119 0 0 0 0.927 0 0 0.914 0 0 0 0.994 0 0 0.996 From the definition of the principal strains, it was found that the magnitudes of e 1

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258 and e 2 were the largest but of opposite signs, and that e 3 was the c l osest to zero in the homogeneous distortions of the ORTHl and HCP plates. This result indicated that the uniaxial stress axis for each of these different plates was parallel to the [011) 13 direction of th e p phase. Ther e fore the [001] 0 direction of the ORTHl and th e [OOOl]H dir ect ion of' the HCP plat es were the uniaxial stress axes that th ese plat es rotated around to form the habit plane with the p phase. These directions wer e also confirmed by the calculations that showed that both e 2 and e 3 had negative valu es, but that th e value of e 3 was an order of magnitude l ess than the value for e 2 Therefore according to the invariant line theory this co ndition indicated that the plate was uniaxially stressed along the smaller of the two simi lar strain values which was e 3 Th e next procedure in the invariant line theory was to add the effect of th e elastic strain along the uniaxial stress axis to th e homog e n eo us distortions of the ORTH l and HCP plates. Sinc e the invariant line in the habit plane is stress fr ee according to this theory, then the elastic strain in th e uniaxial stress axis causes an ex t e nsion of -e 3 in this direction. This extension then causes a Poisson contraction by an amount of ve 3 on the other two principal strains. Thus following is the additional e ff ect of this extension on the principal axes of strain in the homogeneous distortion (B ) : B' = 0 A 2 0 0 7 o I+ Ve3 0 0 0 0 0 0 0 0 0 7 0 wher e v is the Poisson ratio and is assumed to be 0 3. The e ffect of the modified homogeneous distortion B' is to minimize the elastic stra.in energy of the plat es and to leave them in an orientation containing the habit plane. The sc h ematic in Figure

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259 6.29 shows the plate in this uniaxially stressed state and the three principal axes of strain. The habit plane of the plate shown in this figure contains both the uniaxial stress axis along A 3 which is the direction that the plate rotated around, and the invariant line which is normal to the uniaxial stress axis and is unextended. The final procedure is to calculate the invariant line in the habit plane of the ORTHl and HCP plates using the modified homogeneous distortion B' and the rigid body rotation R These calculations were done following the method described in the classical theory of martensitic transformations [79] From this theory, the total transformation T of the ORTHl and HCP plates in matrix notation is as follows: 0 A' 2 0 0 0 cos e sine 0 sine cos e 0 01 O 1 The rotation axis in R is the [O 11] 13 direction of the J3 phase, which is parallel to the [001] 02 direction of the ORTHl plate and the [OOOI]n direction of the HCP plate. These directions define tl1e uniaxial stress axis of th e transformation T. A vector u will be the invariant line in the l1abit plane of the plates when the following is true: Tu=u which indicates that u is undistorted by the transformation T The invariant line u was calculated from the eigenvalues A of the transformation T and is described by the matrix equation: Tu= AU, or IT Al I u = 0. The non-trivial solutions of this equation were then determined by setting the

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Figure 6.29 260 Uniaxial Stress Axis (Rotation Axis) Invar i ant Line (Stress Free Axis) e Shows the uniaxial stress state of the plate and the three principal axes of strain which are 1v 1 1v 2 and 1v 3

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261 d e t e rminant of th e e quation t o z e ro: d e t I T Al I = 0 whi c h has th e thr ee e ig e nvalu es A 1 A 2 and A 3 Sin ce th e invariant lin e is unr o tat e d and undi s t o rt e d th e n t h e v ec t o r u wa s f o und by se tting on e o f th e e jg e nvalu es t o 1.1ni ty a nd so lving t h e d ete rminant Th e rigid b o dy r o tati o n 4> b e tw ee n th e 13 ph ase and th e ORTHl and H C P pha ses i s s olv e d from th e pr e vious e quation for A = 1 and is as f o ll o w s: Thi s r o tati o n bring s th e undist o rt e d v ec tor s in th e 13 phas e and th e ORTHl or H C P pha ses into c oin c id e n ce with on e anoth e r to form th e invariant lin e u in th e habit plan e of th ese plat es. Th e c al c ul a ti o n s f o r th e ORTHl plat e s hown in Figur e 6.9 indi c at e d tha t th e i nv ar iant lin e i s 42 71 ( 0 23 fr o m th e ( 100 ) 0 t plan e Co nv e rting thi s an g l e from th e O RTHl phas e t o t h e 13 pha se s h o w e d that this lin e i s 47.15 ( 0.23 fr o m th e ( 100 ) 13 plan e Th e diff e r e nc e b e tw e en th e se two angl e s i s e qual to th e rigid b o dy r o tati o n whi c h i s 4 44 ( 0 23). Th ese c al c ulations s h o w e d go o d agr ee m e nt with th e habit plane and rigid body rotati o n o bs e rv e d for th e ORTHl plat e in Figur e 6 9. Th e h a bit plan e of this plat e was m e asur e d a s 12.41 ( 0 08) from th e ( 211 ) 13 plan e o r 47 67 ( 0.08 fr o m th e (100 ) 13 plan e, sinc e th e angl e b e tw ee n th e ( 211 ) 13 and ( 100 ) 13 plan es is 35.26. Th e r e for e, th e c a l c ulat e d angl e for th e invariant lin e of th e habit plan e i s o nly ~ 0.5 fr o m t l1 e m e a s ur e d a ng l e f o r th e habit plan e o f thi s ORTlll pl a t e. Th e rigid b o dy r o ta tio n w as d e t e rmin e d fr o m th e o ri e ntation r e lati o nship s h o wn in t h e SAED patt e rn of Figur e 6 9 by m e asuring th e angl e b e tw ee n th e ( 010 ) 01 plan e of

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262 the ORTHl phas e and th e (011) 13 plan e of th e f3 phase. Th e measur e d angl e was 4.67 ( 0.08) which i s 0.23 larger than th e calculated angle of the rigid body rotation Th e HCP plate shown in Figur e 6.1 was also found to show good agre e m e nt b et w ee n th e calculated and the experimentally m eas ur e d habit plan e. Th e calculations showed that th e invariant line in the habit plane is 39.27 0 23 from the (lTOO)H plan e of this HCP plat e, which is also 45.05 0.23) from the (100) 13 plan e of th e f3 phas e. Th e diff e r e nc e b et w ee n th e two angles of the HCP and f3 phases is e qual to th e rigid body rotation whi c h was calculated to b e 5. 78 ( 0 23). Th e co mparison of the calculated angle to the m e asured angl e of the habit plan e shows that these two angl es ar e diff e r e nt from on e another by only 0.13 This r es ult was d e t e rmined from the m e asur e m e nt s that showed that the habit plan e of th e H C P plat e shown in Figur e 6.1 was 44.92 0.08) from th e (100) 13 plan e. Likewis e, there i s good agr ee m e nt b e tw ee n th e calculated and th e measured angles of th e rigid body rotation. From the SAED patt e rn of the orientation r e lationship shown in Figur e 6 1 the m ea sur e d angl e between th e (100) 13 and (lTOO)H reflections was 5.58 0 08 ), whi c h is 0.20 l ess than the c alculat e d angl e In summary, it was shown in this discus s ion that the invariant lin e th eory is able t o ac c urat e ly pr e dict th e invariant line in th e habit plan e of th e ORTHl and H C P plat es. The s p ec ifi c lattic e param e t e rs and the habit plan e m e asured for the ORTHl and HCP plat es had to b e used so that the ability of this th eor y to pr edict th e h a bit plan e co uld b e tested. It was found in this study that wh e n th e sp ec ific valu es of the lattic e param eters w e r e us e d f o r a parti c ular ORTIIl, ORTH2 or H C P plate then the calculations bas e d on th e invariant lin e consistently showed good agreement with th e m eas ur e d habit plan e of th ese plat es. Th e agreement b e tw ee n th e

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263 calculated angles and the measured angles was within acceptable ranges based on the measurement errors associated with the 1"'EM. Thus it is concluded that the ORTH and HCP plates formed from the f3 phase by a martensitic transformation based on the results of this study. 6.3.2 Structural Analysis of the Plates The structural analysis of the ORTHl and HCP plates using SAED and CBED techniques showed them to be c los ely related to each other. This was observed in the orientation relationships with the f3 phase and in the interplanar d-spacings and angles measured from reflections in SAED patterns. The schematic in Figure 6 30 shows the relationship between the ORTHl and HCP phases. The unit ce lls of these phases are superimposed upon each other to illustrate how the distortions of the orthorhombic phase are related to the hexagonal phase The dotted lines outline the ORTHl unit cell with the ao and b 0 lattice parameters normal to the [001] 0 direction. The HCP unit cell is shown normal to the [OOOl]B direction. The HCP unit cell is shown with both the usual h exagonal representation using the aH latti ce parameter and also with the orthorhombic representation using the aH,o and bH,o lattice parameters. The hexagonal structure requires that the two aH axes be separated by 120 This makes the alb ratio a constant of V3 or 0.577. From this schematic, it can be seen that distortions in the aH, o and bH, o directions can destroy the hexagonal structure and form the orthorhombic structure with an alb ratio that would normally not be equal to 0.577. The alb ratios of the ORTHl structure determined from Table 6.6 are ~0.580 to 0.602 which are larger than 0.577. These values are obtained by a decrease in the a 8 10

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..... 0 ,.Q [lOlO] H [010] 0 264 3o 1 and a H/o [i2IO] H [100] 0 Figure 6.30. Sh ows the unit ce ll s of the ORTHl and H C P phases superimposed up on each other to illustrat e how th e distortion s of the orthorhombic phase are related to the h exago nal phas e. Th e drawing shows the (001] 0 dir ection of the ORTH! unit cell and the [OOOl]R direction of the HCP unit ce ll

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265 latti ce param e t e r and an incr e as e in th e ba ,o lattic e param e t e r which ar e due t o th e orth o rhombic dist o rti o ns A differ e n ce in th e atomi c sit e o ccupan c y i s sugg e sted to b e r e sponsibl e for th e di s tortions of th e H C P phas e to an orthorhombic phas e How e ver b e for e discussing th e at o mic sit e o c cupan c y of th e ORTHl and H C P phas e s it is important to dis c uss th e APDBs that w e r e pr e sent in th e s e phases and th e e ff e ct that th e ~ co mpo s iti o n had on th e m. Information obtain e d in e a c h of th e s e two following s e ctions will b e u se d in th e two s ec ti o ns that d e al with th e atomi c sit e o c cupancy of the ORTH! and H C P phas es 6.3.2 1 Th e APDBs Th e pr e s e n ce o f APDBs in th e plat e s indicat e d that disord e r to ord e r transitions had occurr e d in th e plat es as th e y d e v e lop e d. Thus th e presen ce o f diff e r e nt typ e s of APDB s impli e d that th e tI C P and ORTHl plat e s follow e d diff e r e nt stru c tural d e v e l o pm e nt paths. Th e m o st appar e nt diff e r e n ce s in th e APDBs w e r e obs e rv e d in th e DF imag e s of th e H C P plat e (Figur e s 6 22a) and th e ORTHl plat e (Figur e 6 25a ) that w e r e form e d with th e (I2TO) a and ( 200) 0 r e fl e ctions r e sp ec tiv e ly By in co rp o rating th e ( 100) 13 r e fl ec ti o n a long wi t h e a c h o f th e s e r e fl e ctions in th e SAED a p e rtur e; th e APDB s pr e s e nt in th e B2 matrix th e num e rous small APDBs in th e H C P plat es, and th e f e w coars e APDB s in th e ORTHl plat e s w e r e simultan e ou s ly o b se rv e d. This t ec hniqu e show e d that th e APDBs in th e B2 matrix w e r e c onn P.c t e d t o th e coars e APDBs in th e ORTHl plat e s which w e r e oft e n observed to pass acr o ss th e plat e s This observation was not mad e for th e APDBs pr e s e nt in th e HCP plat e s. Th e r e w e r e additional diff e r e n ce s that w e r e observ e d in th e APDB s pr ese nt in th e H C P and ORTHl plat e s wh e n oth e r r e fl e ctions w e r e us e d to form th e DF imag es.

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266 Th e coars e APDB s in th e ORTHl plat e s w e r e not obs e rv e d wh e n th e (130) 0 and ( 130 ) 0 r e fl e ctions w e r e us e d, but inst e ad a high numb e r density of small APDBs w e r e ob se rv e d (co mpar e Figur e 6 25a t o Figur es 6 25b and 6 25 c). Thi s r e sttlt was diff e r e nt f o r th e APDB s pr ese nt in th e H C P plat e s whi c h s h o w e d th e sam e g e n e ral a pp e aranc e in all thr ee imaging conditions ( c ompar e Figur e 6 22a to Figur e s 6 22b and 6 22c ) Ev e n th o ugh a columnar morphol o gy was o bs e rv e d for th e APDB s in th e H C P plat e s and for th e fin e APDB s in th e ORTHI plat e s th e scal e and th e ori e ntation of th ese c olumnar APDBs w e r e diff e r e nt Th e r e sult s indicat e d that th e c olumnar APDBs in th e HCP plat es w e r e ori e nt e d normal to th e habit plan e, whil e th e fin e APDBs in th e ORTHI plat e s w e r e ori e nt e d by up to ~30 from th e habit n o rmal. How e v e r t h e c olumnar APDB s w e r e o bs e rv e d normal to th e plat e int e rfa ce f o r th e H C P pl a t es a t th e [1120J n zo n e axi s and f o r t h e ORTHI plat e s a t th e (110] 0 zo n e axi s Th e r e f o r e, t h e [1120J n and (110] 0 zo n e ax e s s h o w e d th e APDB s in th e H C P and ORTHl pla tes, r es p ec tiv e ly at an o ri e ntati o n that was 90 fr o m th e [OOOl] H and [001] 0 z on e ax es. Th e analy s i s o f th e plat e s indi c at e d th e s e tw o ori e nta tio n s re pr e s e nt e d a 90 rotation about th e habit normal of th e H C P and ORTHI pla te s Thu s, th e v ec t o r r e pr e s e nting th e ~30 angl e b e tw ee n th e habit n o rmal and th e dir e cti o n of th e c ol1-1mnar APDBs in t h e ORTHl plat es was parall e l to th e e l ec tron b e am at th e [110] 0 zo n e axis. This e xplains why th e s e APDBs w e r e obs e rv e d n o rmal t o t h e int e rfa ce of th e ORTI-Il plat es a t thi s o ri e ntat i on Th e di s pla ce m e n t v ec t o r o f th e APDBs w as important to d e t e rmjn e, s in ce it will b e us e d l a t e r in t h e dis c ussi o n t o s how th e at o mi c s it e o cc upan c y of th e ORTHl and H C P phas e s. Th e displa ce m e nt v ec t o r wa s c al c ulat e d for possibl e v ec tor s fr o m th e lit e r a tur e using th e imaging prin c ipl es o f a-boundari e s [96]. Th e s e boundari es, o r

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267 int e rfa ces, impart a phase shift describ e d by th e formula a = 21t(g R ) on th e normal scatte ring phas e = 21t(g s) wh e r e R is th e displac e m e nt v ec tor associated with th e APDB s is the d e viation param e t er, and g is th e r e fl ec tion v ec tor Th e APDB s will b e invisibl e wh en (g R ) = n1t wh e r e n is an e v e n int e g e r since they will th e n hav e the s am e phas e as the normal scattering phas e. Lik e wis e, the APDBs will b e visibl e wh e n n is an odd int eger, or out of phase with the normal scattering phas e. Thus these vi s ibl e/ invi s ibl e imaging c hara cter isti cs of th e APDBs are responsible for th e imaging condi tions s h o wn in Tabl e 6. 7. Th e most pr oba bl e displa ce m e nt v ec tor out of all of th e possibl e v ec tor s, for the APDBs in th e H CP plat es was l/6<1120>. Previous studies had shown that this vector was associated with th e ordering reaction of th e a-Ti phas e to th e a 2 -Ti 3 Al phase [78]. Th e calcula t e d phas e factors using a= 21t(g R) for this displa ce m e nt vector and for several g r e fl ect ions ar e shown in Tabl e 6 8. A Co mpari so n of these calc ulat e d valu es in this table to the o bs e rv e d imaging conditions in Tabl e 6. 7 shows good agree m e nt between th e visible and invisible co ndit io ns. Thus the l/6<1120> displacement vector is d esc riptiv e o f the APDB s in the H C P phas e Analy s i s of the APDBs in th e ORTHl and ORTH2 plat es rev e al e d that the di sp la ce m e nt vector for the fm e APDBs i s V2[100] whil e the displa ce m e nt v ec tor f or the coarse APDBs is 1/4[110]. Tabl e 6.9 lists th e calculated phas e factors of th ese APDB v ecto r s for several g r e fl ect ion vectors. Co mparing th e calculated valu es in Table 6 9 to the o b serve d imaging co nditi o ns in Tabl e 6. 7 showed that 1/2(100] is t h e correct v ecto r for the fine APDBs and that 1/4(110] is the v ec tor for th e coa rs e APDB s in these ORTHl and ORTH2 plat es. Th ese two specific displa ce m e nt vectors for the

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268 Tabl e 6.8. Calculated phase factor s, a= 2n(g R ), for th e APDB vectors in the HCP phase. Refl ec tion g APDB Vectors V ecto r 1/6[1120] 1/6[1210] 1 /6(2 110] (1120) 0 -7t -7t ( 1210 ) -7t 0 -7t (2110) -7t -7t 0 (1100) 0 7t 7t (1010) 7t 0 7t (0110) 7t 7t 0 (1101) 0 7t 7t ( 1011 ) 7t 0 7t (0111) 7t 7t 0 (2200) 0 0 0 (2020) 0 0 0 (0220) 0 0 0 (0002) 0 0 0

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269 Table 6.9. Calculated phase factors, a= 2n(g R ), for the APDB vectors in the ORTHl and ORTH2 phases. Reflection g APDB Vectors Vector l/2[100] l/4[110] (200) 0 7t (020) 0 7t (002) 0 0 (110) 7t 7t (220) 0 0 (130) 7t 0 (260) 0 0 (111) 7t 7t

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270 fine and coarse APDBs correlate with results previously published in the literature for the 0-Ti~Nb phase (31 ,34]. 6.3.2.2 Effect of the p Composition Assuming that the plates have formed by a martensitic reaction then they must have inherited the composition of the f3 matrix. From chapter 4 it was found that the microstructures of the aged samples consisted of the single f3 phase at temperatures of 1400 C and the cr + f3 phases at 1300C. The composition values shown in Table 4.1 indicated that the composition of the f3 matrix at 1400 C was increased by 3%Ti (to 37%Ti) and decreased by ~4%Nb (to 23%Nb) upon formation of the cr grains at 1300 The Al content in the f3 matrix was observed to stay nearly constant between 1400 C and 1300C. This composition change explains the change in lattice parameters observed upon a variation in quenching temperature. Results by Kestner-Weykamp e t al. [26] showed a similar dependency on the composition of th e f3 phase in which a change in the composition of the f3 phase from 20Nb to 30Nb (at.%) in Ti-25Al-XNb (at.%) based alloys increased th e alb ratio of the O-Ti 2 A1Nb phase from 0. 632 to 0. 645. There were two ways that the f3 composition may have affected the structure of the plates: one is related to the content of the Al and the other to the content of the Ti. The primary effect of the Al content i s to lower the magnitude of the orthorhombic distortion. This effect is determined by comparing the alb ratios of the ORTHl plates which w e re 0.580 and ~0.602, to the equilibrium O-Ti 2 A1Nb phase which was >0.630 (28] The low er alb values were attributed to the higher Al content, which is ~40Al (at.%) in alloy 2 as compared to ~25Al (at %) in the stoichiometric compos ition of th e 0-Ti~Nb phas e. This comparison shows that Al must have

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271 sub s tituted for Ti sinc e th e ~27Nb (at %) cont e nt of alloy 2 was e ss e ntially th e sam e as that in th e 0-Ti~Nb phas e Th e low e ring of th e alb ratio by an incr e a se in th e Al c ont e nt is th e opposit e of what has b ee n obs e rv e d in pr e vious work by an increas e in t h e Nb cont e nt [26] Thus it is d e du ce d from this study that th e Al c ontent must l1av e r e du ce d th e orthorhombic distortions which caus e s th e alb ratio t o appr o a c h th e alb ratio of th e H C P phas e in all o y 2 If th e Ti cont e nt of th e J3 matrix w e r e also higl 1, in a ddition t o th e ~ 40 % Al c ont e nt th e n tog e th e r th e y might hav e s t abiliz e d th e H C P pha se in t h e pla tes Th e r e for e, th e e ff ec t o f Ti i s to act with Al to stabiliz e th e H C P phas e, inst e ad o f th e ORTH phas e in t h e 1300 C ag e d sampl e of alloy 2. 6.3 2 3 Atomi c Sit e Oc c upan c y 6 3.2.3.1 Th e ORTHl and ORTH2 Phas e s It i s ass e rt e d fr o m this study that th e ORTHl phas e is r e lat e d to th e a'' mart e nsit e and th e O-Ti 2 A1Nb phas e, e xc e pt that th e atomi c sit e o ccupancy o f th e O RTHl phas e is modifi e d by th e gr e at e r amount of Al c ont e nt in alloy 2. Th e r e w e r e many s imilariti es b et w ee n th e ORTI-11 pl a t es that f o rm e d in all o y 2 th e a" ma r t e nsit e that f o rm e d in Ti-Nb all o y s [62] and th e 0-Ti~Nb pha se that form e d in t h e Ti 3 Al + Nb all o y s [13]. Th e ORTHl plat e s form by a mart e nsitic transformat io n s imilar t o th e a" m ar t e nsit e, but th e ORTHl plat e s ar e s tru c turally s imilar to t h e O-Ti 2 A1Nb phas e ba se d on th e latti ce param e t e rs ori e ntation r e lationship with th e J3 phas e, bas e ce nt e r e d orthorhombi c Bravais lattic e with mmm point group and APDB s It will b e s h o wn in this sec ti o n that all thr ee pha se s ar e s tru c turally link e d t hr o ugh th e sam e C m c m spac e group Th e s tructural diff e r e n ce b e tw ee n th e s e pha ses d e p e nds prim a rily on wh e th e r th e s e phas es hav e th e disord e r e d C mcm s tru c tur e ( a" mart e nsit e) o r th e o rd e r e d C m c m s tru c tur e ( 0-Ti~Nb ph a s e and

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272 ORTHl plates). Unfortunately, direct proof of the structural relationship between the ORTHl phas e and the 0-Ti~ phase was unobtainable, since the CBED analysis of the ORTHl plates was comp li cated by the sma ll thicknesses and internal defects. Internal def ec ts s uch as stacking faults and APDBs are lattice imperfections that destroy the conclus iv e evidence of glide-plane and screw-axis symmetry elements in CBED patterns [89]. Indirect proof of the structural relationship between the ORTHl phase and the O-Ti 2 A1Nb phase is establis h ed by the determination of a similar ordered Cmcm space group for the ORTHl phas e. Determination of the ordered Cmcm space group for the ORTHl pl1ase is based on the orientat i on relationship; mmm point group symmetry; absent (100) (0 10 ), and (00 1 ) reflections; and APDBs. The results of this study establis h ed that the ORTHl phase has the mmm point group symmetry by showing a 2mm symmetry for tl1e [001] and [110] zone axes and a mirror plane para lie! to the (001) planes. The SAED patterns of the [001] zone axis showed that the (10 0 ) and (0 10) reflections w ere forbidden and that the (110) reflection wa s allowed. These were consistent with the structure factor of a base centered Bravais la ttice i.e. (h + k) = n where odd n is forbidden and even n is allowed. Finally, the APDBs that were observed in the ORTHl plates s h o w ed indirect proof that the ORTIIl phase had the Cmcm space group. This proof was based on the APDB analysis that showed that the displacement vector for the fine col umnar APDBs was 1/2(100] and for the coarse APDBs was 1/4(110]. These two specific displacement v ec tors in the ORTHl phase are the same as those observed in the O-Ti 2 Al.Nb phas e (31 34] which suggests that the three separate s ublatti ces in the ORTHl phase are similar to those in the O-Ti 2 Al.Nb phase. These sublattices co uld have different

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273 atomic arrangements but still show the necessary 2 1 screw-axis and c -glid e plane of the Cmcm space group. Thus, all of these observations are consistent with the Cmcm space group and indicate that the ORTHl phase has a simi lar space group to the 0-Ti~ phase. The proposed atomic site occupancy of the ORTHl phase is based on the three sublattices and the higher Al content of alloy 2, compared to that in the 0-Ti~ phase. This proposed model is shown in Figure 6. 31. The difference between this model and the 0-Ti~Nb phase is that the Ti and Al positions are exchanged between the 8g and 4cl sites with the Al occupying the 8g Ti the 4cl, and Nb the 4c2 Wyckoff sites. The atomic site occupancy of the ORTHl phase is listed in Table 6.10. In this model the unit cell of the ORTI-11 phase consists of alternating rows of Ti-Nb an d Al Al bonds in the [100] directions, and alternating rows of Ti Al and Nb-Al bonds in the [110] directions. The strong bonds between Ti-Al and Nb Al act as rigid chains in the lattice along the [110] direction. Thes e chains then keep the weak bonds between Ti-Nb strictly in the [100] direction. This suggests that the orthorhombic distortions in the ORTHl phase are due to the Ti-Nb and Al Al bonds in the [100] directions. Thus the coordinates of the atoms listed in Table 6.10 and shown in Figure 6.31 are assumed to be (0, 0 1667, 0.25) for Ti on the 4cl site; (0, 0.1667, 0.25) for Nb on the 4c2 site; and ( 0.25 0.0833, 0.25) for Al on the 8g site. This model for the site occupancy of the ORTHl phase is derived from the O-Ti 2 Al.Nb phase by considering that Al substitutes for Ti Therefore an increase in the Al composition of th e 0-Ti~Nb phase with constant Nb must be accomplished by occupying either the 8g sites with Ti or the 4c2 sites with Nb. The most likely scenario would be for the Al to occupy the 4c2 sites This maximizes the number of'

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274 Table 6. 1 0. The proposed atomic site occupancy of the ORTH l phase with th e Al 2 NbTi stoichiometry and the Cmcm (63) space group. Site Site Position W yckoff E l ement Number Position X y z 1 0.25 0.4167 0 .25 Al 8g 2 0.25 0 9167 0.25 Al 8g 3 0.75 0.4167 0.25 Al 8g 4 0.75 0.9167 0 .25 Al 8g 5 0.25 0.0833 0.75 Al 8g 6 0.25 0.5833 0 75 Al 8g 7 0.75 0.0833 0.75 Al 8g 8 0.75 0.5833 0 .75 Al 8g 9 0 0.6667 0.25 Ti 4C 1 10 0 5 0.1667 0.25 Ti 4C 1 11 0 0.3733 0.75 Ti 4C 1 1 2 0.5 0.8333 0.75 Ti 4C 1 13 0 0.1667 0. 25 Nb 4c 2 14 0.5 0.6667 0.25 Nb 4c 2 15 0 0 8333 0 75 Nb 4c 2 1 6 0 5 0.3733 0.75 Nb 4c 2 Th e fo ll owing Wyckoff values were us e d: Al (8g) (x, y z) = (0 25 0.0833, 0 25) Ti (4cJ (x, y z) = (0, 0 .1 667 0.25) Nb (4cJ (x, y z) = (0, 0.3333, 0.25)

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Figur e 6 31. 275 ... t~J~J~: '. .... . . . . .. . .. 4c2 Wyckoff Site: Nb 4cl Wyckoff Site: Ti 8g Wyckoff Site: Al b a Shows the proposed atomic site occupancy of the ORTHl phase bas e d on th e Cmcm space group with Al atoms occupying th e 8g Ti atoms occupying the 4c 1 and Nb atoms occupying th e 4c2 Wyckoff positions.

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276 Ti-Al bonds which have tl1e strongest bonds and the shortest bond length [84]. Wh e n a sufficient numb e r of 4c2 sites are filled by Al and in addition to the 4cl sites already occupied by Al then these form the basis of the 8g sites in the propos e d ORTHl pl1ase. As a consequence of the in crease in Al and subsequent occupancy of the 8g sites, the Ti is forced to occupy the 4cl sites and Nb is reassigned to th e 4c2 sites in th e new ORTHl phase. Evidence that Al occupies th e 8g Wyckoff sites in the proposed ORTHl phase is obtained in this study by comparing th e relative intensities of the (020) 0 and (1 10 ) 0 reflections with th e calculated structure factors. The intensity of these reflections were shown in the CBED pattern of Figure 6.11. The structure factors of' the (020) 0 (110) 0 and other reflections at the [001] orientation were then calculated bas e d on th e atomic occupancies of the ORTHl and 0 Ti~Nb phases. Figure 6.32 shows th e relative intensities of these two phas es in the form of ca l culated CBED patterns for comparison with th e actual CBED pattern in Figure 6 11 The ca lculat ed patterns showed th e best match between the ORTH phase (Figure 6.32a) and the CBED patt e rns of the ORTHI plates (Figur es 6.11) since both patterns indicated that the intensity of th e (020) r e fl ec tion was greater than the (1 10 ) reflection. The opposite was true of the 0-Ti~Nb phase (Figure 6.32b) wlrich showed that the intensity of the ( 110 ) reflection was greater than the (02 0 ) reflection. Thus, this comparison shows that Al and not Ti, must occupy the 8g sites in the ORTHl phases Further proof of the propos e d site occupancy for the ORTHl phas e in alloy 2 was obtained by conducting the CBED analysis on plates that w ere present in a Ti-rich alloy (designated as alloy 1 ). The composition of alloy 1 was 40Ti 32Nb-28Al (at.%). This composition is close to the 45Ti-30Nb-25Al ( at.%) composition of an all oy

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. .. .. ... ......... 277 (a) (b) ... . .. .. .. .. .. .. .. . . .. . . . 000000 0@ 0@~@ 000000 Reflection Al 2 NbTi Ti 2 NbAl 020 17 12 7.25 110 6.44 12.36 200 23.74 1.52 130 7.79 9.12 040 35.49 60.23 220 36.73 24.07 Figure 6. 32. Shows the calculated CBED patt e rns of the orthorhombic structures based on two possible atom i c site occupancies. (a) the proposed ORTHl phas e with the stoichiometry of Al 2 TiNb; (b) the O-Ti 2 A1Nb phase [13]

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278 that c ontain e d a singl e O-Ti 2 A1Nb phas e microstructur e aft e r aging at 800 C [26]. Th e r e for e, alloy 1 w a s h e at tr e at e d at 1400 C for 4 hours to form th e singl e P phas e and t h e n furnac e co ol e d which r e sult e d in th e formation of a high numb e r d e nsity o f O-Ti 2 A1Nb plat es fr o m th e B2 matrix. Th e structur e of th e 0-Ti~ plat e s was co nfirm e d by th e S AE D and C BED analy s i s Th e SAED analysi s show e d that t h e o ri e ntati o n r e lati o nship b e tw ee n th e O-Ti 2 A1Nb and B2 phas e s was [001] 0 11 [011] 13 zo n e ax es with th e ( 110 ) 0 plan e s parall e l to th e (211 ) 13 plan e s Th e C BED whol e patt e rn is shown in Figur e 6. 33 and indicat e s that th e int e nsity of th e ( 020) r e fl ec ti o n wa s l es s than th e ( 110) r e fl e ction. This CBED patt e rn mat c h e s b e tt e r with th e c al c ulat e d C BED patt e rn sh o wn in Figur e 6. 32 for th e O-Ti 2 A1Nb phas e, wh e r e Ti occ upi e s th e 8g sit es. Th e r e for e, this mat c h c onfirms th e pr e vious c onclusion that Al o cc t1pi e d th e 8g sit es, rath e r than Ti in th e propos e d ORTHl phas e. Anoth e r r es ult that c o nfirm e d that Al oc cupi e d th e 8g Wyck o ff sit e s in th e O RTHl phas e i s th e o bs e rvation th a t th e co ar se APDBs pas se d a c r oss th e int e rfa ce b e tw ee n th e ORTHl plat es and B2 m a trix as see n in Figur e 6.25a From c hapt e r 4 1 i t wa s shown th a t Al o c cupi e d th e lb s it es in th e B2 phas e, and this l e ft th e occ upancy of th e la s it e s and unfill e d lb sit e s (~ 20% ) for Nb and Ti Th e atoms occ upying th e s e sit es must hav e r e mainP.d in pla ce during th e transformation from th e B2 pha se to th e ORTHl pha se, sin ce th e APDBs that f o rm e d in th e P to B2 o rd e ring transition w e r e inh e rit e d by th e plat es. This obs e rvation can b e s ee n in Figur e 6. 34 whi c h s h o w s th e APDB v ec tor s o f th e B2 phas e in th e ( 011) 13 plan e s and t h e O RTHl pha se in t h e ( 001 ) 0 plan es. Th e s e plan es ar e lying parall e l to e a c h o th e r as pr e vi o u s ly d e t e rmin e d fr o m th e o ri e ntation r e lationship From Figure 6.34 it i s co n c lud e d that th e 1 / 2 < 111 > v ec t o r pr e s e nt in th e B2 phas e is transform e d int o th e

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279 Figure 6. 33. CBED pattern showing the whole pattern symmetry of the 0 -Ti ~Nb phas e observed in alloy 1.

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280 ... , ... : :: :::::: :f :: ~:: ::::::::~::::::::: ... ' .-:-:-:-:,:. <{'.'.;;:}/ -----~ ll}?: ' ' ' a/4[110] a/2(100] a/2(111] [001] [100] Titanium Niobium Niobium Figure 6.34. Aluminum Titanium Aluminum Sh o w s the APDB vectors in th e ( 011 ) 13 planes of the B2 phas e and in th e ( 001 ) 0 planes of th e ORTHl phas e

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281 1/4(110] v e ctor in th e OR1Hl phas e. Th e continuity of th e APDB v e ctors acro s s th e B2/plat e int e rfac e indicat e s that th e o c cupan c y of' th e atomic sit e s in th e B2 and ORTHl phas e s is unalt e r e d during this part of th e transformation. Thus th e la ( Nbtri) and lb (Al) s it es in th e B2 phas e ar e transform e d into th e two 4c ( 4 c l and 4c2 ) and th e 8g sit es, r e sp ec tiv e ly in the ORTHl phas e Th e r e for e this r e sult pr e di c ts th a t Al will e nd up on th e 8g s it es in th e pr o pos e d ORTHl phas e. Th e r es ults o f t hi s study indi ca t e d that an OR1H3 phas e, diff e r e nt from th e propos e d ORTHl pha se shown in Figur e 6.31 e xist e d during th e d e velopm e nt of th e ORTHl plates. Th e ORTH3 phas e i s b e li e v e d to e xist imm e diat e ly aft e r th e plat e s f o rm e d wh e n th e sit e occupan c y of th e B2 pha se i s lock e d in and a disord e r e d o cc upan c y b e tw ee n Nb and Ti on th e two 4 c s it e s in th e OR1Hl phase had oc c urr e d Pr o of o f this disord e r e d o c c upan c y is obtain e d by SAED analysis of v e ry thin OR1H3 plat e s and also by th e obs e rvation of t h e c olumnar shap e d APDBs in th e ORTIIl pla tes. Th e SAED p a tt e rn of a v e ry thin plat e that had th e ORTH3 phas e was s h o wn in Figur e 6 15. In thi s patt e rn th e (hkO ) r e fl ect i o ns wh e r e h and k ar e odd int e g e r s a r e a b se nt f o r th e ORTII3 phas e. Th e s e sam e r e fl e ctions initially appear e d a s s tr e ak s in th e SAED patt e rn o f th e ORTH2 phas e (Figur e 6 13) and finally a s diffraction spots in th e SAED patt e rns of th e ORTHl phas e (Figur es 6 9 ). Th e co lumnar shap e d APDBs obs e rv e d in th e ORTHl plat e w e r e shown in Figur e 6.25 and in th e ORTH2 plat e in Figur e 6 28. Th e pr e s e nc e of th e s e APDBs indi c at e d th at a disord e r t o ord e r transiti o n o c c urr e d and form e d tw o s ublatti ces in th e ORTHl phas e. It wa s pr e viously d e t e rmin e d from th e s tru c tural analysis that th e displa ce m e nt v ec t o r o f th ese APDB s was 1 / 2(100]. 'fhi s displa ce m e nt v e ct o r as s hown in Figur e 6.34 c onn ec t s th e 4cl (Ti ) with th e 4 c 2 ( Nb ) sit es in th e ORTHl pha se and

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282 is th e s am e displa ce m e nt v ec tor that has b ee n ob se r v e d for tl1 e APDBs in th e 0 -Ti~Nb pha se (31 34] Thus th e pr e s e n ce of thes e fin e APDBs indicat e s that Nb and Ti ar e initially di s ord e r e d on th e 4cl and 4c2 s it e s during th e formation o f th e ORTHl plat e s. Th e implication of thi s r e sult i s that a dis o rd e r e d sit e o cc upan c y o n th e 4 c l and 4c2 s it es pr e dicts a diff e r e nt ORTHl structur e bas e d on symm e try. Th e atomi c s it e occupancy of th e ORTH3 phas e is det e rmin e d by c onsid e ring th e inh e rit e d sit e occ upancy of th e B2 phas e and th e o c curr e nc e of two kinds of latti ce di s t o rti o ns. A mod e l for th e sit e oc c upancy of th e ORTH3 phas e is c reat e d by ass uming that th e {011} plan e s of th e B2 unit c e ll is c o mpr e ss e d in th e <100> dir ect i o n and e xpand e d in tl1 e <011 > dir ec ti o n. Thi s i s th e n follow e d by at o mi c s huffl es of alt e rna t ing {011} pl a n e s to obtain th e unit ce ll shown in Figur e 6.35. Thi s figur e al so s hows th e unit ce ll of th e ORTHl phas e but with disord e r e d o cc upan c y b et w ee n Nb and Ti o n th e 4 c l and 4 c 2 Wy c koff sit e s. Th e disord e r b e twe e n Nb and Ti indi c at e s that th e ORTH3 phas e ha s a diff e r e nt structur e and a high e r symm e try than th e C m.cm symm e try of th e ORTHl phas e. Th e results of this study ar e c onsi s t e nt with th e Pmma spa ce group as th e stru c tur e for th e ORTH3 pha se Thi s co n c lusion i s bas e d o n th e subgroup/ s up e rgr o up r e lations in crystallography that s h o w e d that th e Pmm a s pa ce gr o up e xist s a s th e minimal non isomorphi c s up e rgroup o f t h e C m c m spa ce group with its latti ce param e t e rs r e duc e d in half along th e a and b dir ec tions o f th e ORTHl pha se Th e sc r e w axis and th e glide plan e that e xist in th e C m c m stru c tur e a r e r e tain e d in th e Pmma stru c tur e, but with diff e r e nt or ie nt a ti o n s This diff e r e n ce in ori e ntation aris es from th e rotation in th e orthogonal axis s y s t e m whi c h mak e s th e [001] dir e ction in th e Cmcm structur e th e [100] dir e ction in th e Pmma s tru c tur e Thus th e 2 1 s c r e w axis is parall e l to th e [100] dir ec ti o n and th e

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283 a-glide p l ane is parallel to the (001) plane in the Pmma structure Th ere are two Wyck o ff sites in the Pmma structure that e volv e from th e site occupancy of the B2 phase, and these w ere the 2 e sites at(, 0 z) and (, 0 z) by the Nb and Ti and the 2f sites at (, z) and (, z) by Al Th e magnitud e of the atomic displa ceme nts from the {O 1 l}J3 plane sh uffl es d etermined t h e z coordi nat es of tl1ese two Wyck off sites and w e r e assumed to b e 0 167 for tl1e 2e site and 0.667 (0.167 + 0.500) for the 2f s it e. Table 6.11 list s the coordinates of th e atoms in the proposed ORTH3 phase shown in Figure 6. 35. Figure 6.36 shows a calculated SAED patt e rn for the ORTH3 phas e using the atomic site occupancy given in the mod e l in Figur e 6.35 This diffraction pattern was ca l c ulat e d at the [011] 13 zone axis to show that there was consistency with the orientation relationship observed b et w een the ORTH3 and J3 phases in Figure 6 15. According to the ca l c ulat e d patt ern, the [100] 04 anrl [011] 13 zone axes w e r e parall e l and the ( 011 ) 04 and (211) 13 plan es w ere parall e l. This patt e rn does not show the syste mati c rows of (hkO) r e fl ec tions wh e r e h and k are odd int ege rs that are observed in the ORTHl phas e Furth e rmor e, there is good agreement between the structure factors of several r e flections shown in Figur e 6.36 and th e int ensity of these reflections observed in the SAED patt e rn of Figure 6.15. Thes e structural factors and reflection intensities show that the (002) and (011) reflections hav e similar int ens iti es and that these r e fl ectio ns w e r e greater in int e nsity than the (001), (010), and ( 012 ) reflections. Thus th ese r es ults showed that there wa s good agreement between the c al c ulat e d and the exper im e ntally obtained SAED patterns of the v e ry thin ORTH3 plates and prove that they hav e th e Pmma structure.

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Figur e 6. 35. I I 284 I I I I I . '' .. {!l!~!tl!ltltlt!~!~t I ., 2e Wyckoff Site: 11 and Nb 2f Wyckoff Site: Al C b Shows the propos e d site occ upancy of th e ORTH3 phase based on the Pmma space group with Nb and Ti atoms occupying the 2 e Wyckoff sites and Al atoms oc c upying the 2f Wyckoff sites.

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285 1,able 6.11. The proposed atomic sit e occupancy of the ORTH3 p h ase with th e Al(NbTi) stoichiometry and t h e Pmma (51) space group. Site Site Posit i on W yckoff E l ement Number Postiion X y z 1 0.25 0 5 0.1667 Al 2f 2 0 75 0.5 0.8333 Al 2f 3 0 25 0 0.3333 Nbtri 2e 4 0.75 0 0.6667 Nb/I'i 2e Th e fo ll owing Wyckoff values w ere used: Al (2f) (x, y z) = (0.25, 0.5 0 1667) Nbtri (2e) (x, y z) = (0.25, 0, 0.3333).

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(211) 13 (011) e (111) 13 G (011) 13 (001) @ 286 G (200) 13 (211) 13 (010) (011) (012) 8 (111) 13 (001) (002) @ G (000) (011) 13 G .. (hkl) Reflection Structure Fact-Or (001) 4.0 (002) 8.8 (010) 5 6 (011) 9 1 (012) 2.1 Figur e 6 36 Show s th e c al c ul a t e d SAED patt e rn of t h e o ri e ntation r e lation s hip b e tw ee n th e ORTH3 plat e and th e B2 phas e. Th e OR s h o w e d tl1 e [100] 0 4 and [Oll] p z on e axes t o b e pa1'ail e l and th e (011 ) 04 and ( 21 l )p plan es t o b e parall e l

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287 6.3.2.3.2 Th e H C P Phase The results of this study showed that the HCP plates are structurally similar to the ex 2 -Ti 3 AJ. (D0 19 ) phase [10]. Th e similarities b et w ee n these two phas es includ ed the same Burgers orientation r ela tionship with th e p phas e, lattic e param eters, l/6<1120>H displac emen t vector for the APDBs and P6 3 / mmc space group Th ese similarities indicat e that the atomic s it e occupancy o f' the HCP phase is the same as that of th e ex 2 -Ti 3 AJ. phase, which was previously shown in Figure 2.13. Thus th e atomic site occupancy of the HCP phas e is such that the 2d Wyckoff sites are occupied by the Al co nt e nt and th e 6h sites are co-occ upied by th e Ti and Nb contents of the p m at rix in the 1300 C aged sample of alloy 2. Th e r es ults showed that th e transformation mechanism of the phas e to the HCP plates is similar to that of the ex 2 mart e n s ite that has b ee n reported to form in Ti-25Al + Nb (at.%) based alloys [72 74]. It was shown in chapter 2 that the disordered ex' mart ens it e form e d initially upon qu e nching in the Ti-25AJ. + Nb alloy s and subsequently ordered to form th e ex 2 mart e nsit e Th e r e for e, tl1 e ex 2 mart ensi t e bas th e same structure and APDBs as do es th e ex 2 -Ti 3 Al phase except that it forms by a mart ens iti c transformation. A similar type of transformation to that of the ex 2 mart e n s it e was observed f or the H C P plates, since the APDBs observed in these plates (Figure 6 24) were previously found to hav e a displac e m e nt v ec tor of 1/6<1120>. Thi s is the same displa ce m e nt v ecto r that forms during th e ordering reaction from the ex-Ti phase to ex 2 -Ti 3 Al phas e and produces four domains from the l / 6<1120> translational v ec tor s [78]. This r es ult indicat e s that the HCP plate s form initially with th e di so rd e r e d ex phas e, and th e n they o rd e r to the ex 2 phase during wat er quenching. Thus, th e rapid transformation and the ordering of the ex to the ex 2

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288 phas e s obs e rv e d for th e HCP plat e s of alloy 2 indicat e s that th e y f o rm via a similar m ec hanism as do e s th e a 2 mart e nsit e. Th e s imilari ty o f th e stru c tur e o f th e H C P plat e s to th e a 2 phas e is un e xp ecte d d u e t o th e high Nb co n te nt and th e l o w Ti co nt e nt o f th e J3 matrix in th e 1300 C ag e d s ampl e. Th e c ompo s i t ion of th e J3 matrix was giv e n in Tabl e 4.1 t o b e 37Ti 23Nb 40Al ( at %). If it is a ss um e d that th e H C P plat es form fr o m th e J3 matrix with thi s co mp os ition th e n th e 23%Nb co nt e nt is gr e at e r than th e solubility limit o f Nb in th e a 2 phas e. Th e s o lubility limit has b ee n shown t o b e ~ 12%Nb in th e a 2 phas e of Ti 3 Al + Nb alloy s [32] This solubility limit in c lud e s th e a 2 martensit e as w e ll sin ce th e ~ 12%Nb c ont e n t in th e s e all o ys s uppr es s e s th e formati o n of th e a 2 mart e nsit e upon qu e n c hing [24 26 32]. Th e s toi c hiom e try o f th e a 2 phas e with Ti and Nb s haring t h e s am e s ublatti ce a nd Al o cc upying t h e ot h e r s ubla t ti ce shows that th e rati o bet w ee n Ti+Nb a nd Al t o b e 3 : 1. Th e rati o th a t wa s d e t e rmin e d for th e J3 matrix in t h e 1300 C ag e d sampl e wa s 3:2 and i s diff e r e nt than th e stoichiom e tri c rati o Thi s s ugg e sts that th e in c r ease d Al co nt e nt in th e J3 matrix aff e cts th e solubility limi t o f Nb and r a is es this limit in th e H C P plat e s Th e r e ar e seve r a l possibl e r e a s ons why th e H C P plat e s form inst e ad o f th e ORTHl plat es, in th e J3 matrix n e ar th e cr grains at 1300 C. On e of th e p os sibiliti es i s t h a t s train wa s pr ese nt in th e J3 matrix in th e r e gions surrounding th e cr grain s This s train wa s e xp ec t e d to hav e ari se n from th e volumetri c diff e r e n ce s b e tw ee n th e s e tw o ph ases Thi s s tr a in is al s o e xp ec t e d to b e larg e, s in ce th e cr pha se i s a top o l o gi c all y c l ose -p ac k e d pha se [ 3 9] whil e th e B CC J3 pha se i s an open s tru c tur e with po o r pac king e ffi c i e n c y Str a in indu ce d by pla s ti c d e f o rmati o n in f e rrou s mat e rial s ha s b ee n s hown to rais e t h e t e mp e ratur e of th e mart e nsit e transformation [95] Th e r e for e, th e s train that i s e xp ec t e d to hav e f o rm e d in th e J3 matrix from volum e tr ic

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289 differences with the cr grains could also have raised th e i temperatur e of the plat es that formed near the cr grains. It is possible that th ese plates f'ormed before the J3 matrix ordered to the B2 phase. This is supported by th e result that showed that the APDBs in the B2 matrix did not pass through the HCP plates. Thus these plates that formed from the disordered J3 matrix may have then follow e d a differ e nt transformation path compared to th e ORTHl plates which d e v e loped aft e r the J3 phase ordered to th e B2 phas e It is als o possible that composition gradients, d e v e lop e d in the J3 matrix during the d e v e lopm e nt of the cr grains favor e d the formation of the HCP phase over the ORTHl phas e. Th ese composition gradients could hav e arisen from both the nu c leation and the growth stages of th e cr grains since both of thes e stages r eq uir ed l o ng-rang e diffusion through the J3 matrix. The composition gradients are expected to form by the partitioning of elements across th e moving cr/J3 int e rfac e. From Tabl e 4.1 the most probabl e elements implicat e d in th e partitionjng ar e Ti and Nb. This is because Nb diffus es to the cr grains, which is a Nb-rich phase and Ti is partition e d into the J3 matrix during the dev e lopm e nt of the cr grains at 1300 C. Likewis e, the J3 matrix was e nri c h ed only slightly in Al at 1300C. This indicates that th e solubility 1imit of' Al in th e cr phase is similar to the J3 phase and is ~40%Al The increas e in Ti and the decreas e in the Nb content in the J3 matrix near th e cr grains could hav e favored the formation of plat es with the HCP phases. An incr eased Ti content in th e J3 matrix n ear th e cr grains could also hav e low e r e d the temperature of the disorder to order transition in the J3 phas e This id ea is consistent with pr evio us studies that hav e shown that an incr e ase in Nb content raises the temperature o f th e J3 to B2 reaction in th e Ti 3 Al. + Nb alloys [24 25]. Th ere for e, th e opposite b e havior is also possibl e, in which an incr e ase in the Ti

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290 content in the f3 matrix lowers the f3 to B2 reaction temperature. This idea supports the scenario that shows that the formation of the plates with the HCP phase occurred from the disordered f3 phase. Thus, a combination of strain small differences in composition due to compositional gradients, and a lower temperature for the f3 to B2 ordering reaction may all have contributed to the formation of plates with the HCP phase. It is suggested that the observation of a third phase in some of the HCP plates indicates an instability in tl1e H C P phase at low temperatures. This idea is based on the concept that when the plates form initially at high temperatures, there is a greater amount of configurational entropy that stabilizes the disordered site occupancy between Nb, Ti, and Al on the 2d Wy ckoff site of the a-HCP phase. Upon lowering the temperature, these three different atoms then order into two different sub l attices which are the 2d site occupied by Al and the 6h site occupied by Ti and Nb to form the a 2 -HCP phase. This ordering reaction involves a decrease in the configurational entropy and a decrease in the symmetry of the a-HCP phase. Likewise as the configurat ional entropy is decreased further at lower temperatures, then the composition of the HCP phase can cause an instabilit -y in these plates. It js suggested that the instability occurs on the 6h sublattice of the a 2 -HCP phase, since the ~23%Nb content of these plates i s larg er t han the ~ 11 % to 12%Nb solubility limit that has been reported f'or Nb in the a 2 -Ti 3 Al phase of ternary Ti 3 Al + Nb alloys. The excess ~40%Al content in the a 2 -HCP phase i s not expected to cause this instability since Al occupies the 2d site and ha s been shown in the binary Ti-Al alloys to have a solubility 1imit in the a 2 -Ti 3 Al phase by as much as ~4 0%Al [10]. Thus, as the temperature is continuous ly lowered during w ater quenching, the excess Nb content on the 6h sites shared with Ti causes t h e instability in the a 2 -HCP phase.

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291 A structure for this third phase which will b e referred to as a. 3 is postulated based on the r es ults o f the CBED analysis and the subgroup/supergroup relationship between the a. 2 -HCP and a. 3 phas es. From the preceding discussion low er ing the t em p erature causes a symmAtry reduction in the disord e r e d a HCP to ordered a. 2 -HCP transition. This transition involv es chemical ordering and has to occur at intermediate temp e ratures sinc e the order ing requires thermal activation energy. It i s s ugg este d that a similar r e duction in symmetry accompanied th e transition from the a. 2 -HCP phas e to tl1e a 3 phas e as the temperature is furth e r low e r e d Th e s mall amount of th e rmal activation energy at th ese low t e mp e ratur es is expected to prevent furth e r chemical exc hang es. Thus, this transition must have involv e d only distortions in the a. 2 -HCP lattic e and r es ult e d in the a. 3 phas e. Th e stru cture of th e a. 3 phas e was shown by the CBED patt e rn in Figur e 6. 7b to hav e a 3m s ymm etry axis. Th e 3m symmetry can only exist in structures based on the c ubi c and trigonal Bravais latti ces. Furth e rmor e, it was found that the SAED patterns of the a 3 phase at diff e r e nt orientations w ere indistinguishabl e from those of th e a. 2 -HCP phas e. Thi s observation rules out th e cubic Bravais lattic e as a po ssib l e structure for th e a. 3 phase for th e following r easo n Th e SAED patt er n of a cubic a. 3 phase at the [011] 13 zo n e axis of the f3 phas e shown in Figur e 6.6b would hav e been indexed as the [lll] a. 3 zo n e axis Thus th e SAED pattern of the a. 3 phase at the [111) 13 zone axis of the f3 phas e shown in Figure 6. 3a would then hav e b ee n ind exe d as the [Oll] a 3 zone axis since the 35 26 angl e that separated the [011) 13 from the [011] 13 zone axes of the f3 phas e would hav e also se parat e d the [lll] a 3 from the [110] a 3 zo n e axes of the a. 3 phas e du e to the orientation r e lationship Th e ratio m et hod wa s used to co mpar e the calculated ratios based on diff e r e nt lattic e param e t e rs that wer e determined from th e [lll] a. 3 zone axis to the m eas ur e d ratios whi c h w ere d eter min ed

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292 from the reflections in the SAED pattern of the [110]a 3 zone axis. This analysis showed no consistency between the calculated and measured ratios based on the cubic a. 3 structure. In comparison, the same type of ratio method showed that the trigonal Bravais lattice is possible, since the trigonal unit cell can be indexed using the same (hki1) Miller indices as the HCP unit cell. Furthermore, there are two possible trigonal structures, tl1e P3ml and P31c space groups, that are maximal non-isomorphic subgroups of the P6 3 /mmc space group, which is the space group of the a. 2 -HCP phase (43]. Both of these trigonal space groups contain the necessary 3-fold axis that was observed by CBED analysis. However, the P31c space group showed that a c-glide plane is present in the same direction as the c-glide plane in the a. 2 -HCP phase, and that it could have formed by small displacements of Nb and Ti atoms from the 6h Wyckoff sites in the a. 2 -HCP phase. These displacements would have destroyed the 6-fold symmetry in the a. 2 -HCP phase and would have subsequently created the 3 symmetry axis and yet retained the c glide plane in the a. 3 phase. Thus it is suggested that the P31c space group is the most probable structure for the a. 3 phase. However, further work is necessary to confirm this as the possible structure. 6. 3. 3 Crystallographic Treatment of' tl1e Plate Transformation In this section, it will be shown that the transformations of the ORTI-Il and HCP plates occur in a series of transitions that start from the disordered p phase and end with either the ORTHl or HCP phases. Each of these transitions involves a chan ge in symmetry between the parent and product phases that conforms to the subgroup/supergroup relationships in crystallography [ 43]. This formalism assigns a set of order parameters to the individual transitions and indicates how many variants of the product phas e are formed from the parent phase (97]. In this regard, the

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293 r e sults of the stru c tural analysis of th e ORTHl and HCP plat e s c an b e us e d t o sh o w tl1 e s e individual tran s itions which may th e n b e us e d to show the e ntir e transformation of th ese phas e s a s a se qu e n ce of transitions Th e r e sult s o f thi s s tudy s h o w e d that th e r e ar e two transformation p a th s f o r th e plat es : o n e tha t l e ads t o th e ORTHl phas e and on e that l e ads t o th e HCP phas e. Th e tw o p a ths ar e b ase d on th e s tud y by B e nd e rsky e t al. [33] and ar e s h o wn in Figur e 6.37 Th e dir e ction of th e arr o w s in th e two paths of this figur e s how wh e th e r an in c r e as e (up ) o r d ec r e as e (down ) in symm e try oc c urr e d in th e transition of th e produ c t phas e from t h e par e nt pha se. Th e symm e try c hang e s shown by th e v e rti c al arr o ws involv e ch e mi c al ord e ring whil e thos e of th e in c lin e d arrows involv e s tru c tural dist o rtion s Th e bra c k e t e d numb e r s ar e t h e o rd e r param e t e rs and indi c at e h o w many variant s of t h e produ c t phas e ar e f o rm e d fr o m th e par e nt phas e in t h e tran s i t i o n Th e discus s i o n that follows will c on s id e r th e s e two diff e r e nt transf o rm a tion pa t hs s e parat e ly und e r th e s ec tions path 1 and path 2. Th e r es ults from th e s tru c tural analysis will b e us e d in support o f th e s e paths. 6.3 3 1 Path 1 of th e ORTHl Plat e s Th e r e sttlts o f thi s study indi c at e d that th e ORTH plat e s d e v e lop e d along path 1 in Figur e 6 37. Th e fir s t transiti o n o f this path is c onsist e nt with th e obs e rvations t hat s h o w e d that th e APDBs pr e s e nt in th e B2 matri x pass through th e ORTHl plat es (Figur e 6 27a ). This r e sult indi c at es t hat th e ORTHl plat e s form from an o rd e r e d f3 (B2 ) matri x whi c h impli e s that th e f3 pha se first ord e rs to th e B2 ph ase and th e n th e forma tio n of th e ORTHl plat e s occ ur s. Th e r e f o r e, th e d e v e l o pm e nt o f t h e ORTH plates b e gin s with th e sit e oc c upan c y of th e B2 phas e.

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Im3m(l3) [2] Pm3m(B2) 14/mmm [3] P4/mmm Fmmm Cmmm Path 1 Path2 [3] Cmcm(OD) Pmma(ORTH3) [2] Cmcm(ORTHl) P6 3 /mmc(a) [4] P6 J mmc(aJ Figur e 6 .3 7. Shows the two transformation paths that l e d to the formation of the ORTHl plates (path 1 ) and th e HCP plat es (path 2) from th e J3 phas e.

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295 Th e first transition from the~ (Im3m) to the B2 (Pm3m) phas es involves che mi cal o rd e ring From the tables of c ry s tallography th e Pm3m space group is a type Ila maximal n onisomorpru c subgroup of the Im3m space group that forms tw o translational variants [ 43]. Th e two variants f' o rm the APDBs in the B2 matrix with the a/2<111> displa ce m e nt v ec tor as was shown in chapter 4. Th e impli ca tion of th e c h e mi c al ordering in the~ phas e is that two sublattices form with Al on the lb Wy c koff s it e and a disord e r e d mixtur e o f Ti and Nb on the la Wy c koff site. Thus th e a t o mi c occ upancy of these two sublattices in this first transition r e main s through o ut the r est of the phas es along path 1. Th e n ext thr ee transitions whi c h occ ur after the B2 phas e forms in path 1 are c rystallographic d esc ripti o ns o f' th e pla te f o rmation that are consistent with the ph e n omeno l ogica l th eo ry of mart ensitic transformations [79] Th ese three steps ar e eq uival e nt to the h o m oge n eous and h e t eroge n ous distortions o f the classical approach to treat ing mart ensitic transformations. Th e r e fore th e three transitions, whi c h ar e cons is te nt with th e h omoge n eous and h e t e rog e n eo us dist o rtions and the rigid body r ot ati o n (see section 6.3 1), ar e responsible for th e formation of th e twelv e plate variants the plat e morphology (Figure 6 19 ), and th e habit plan e (Figure 6.1 ) o f th e O RTI-Il plates. Th e first distortion in th e plat e d e v e lopm e nt al o ng path 1 of Figur e 6. 37 i s th e l1 o m ogeneous di stortion. Thi s dist o r tio n i s a lin e ar transformation [79] that distort s the unit ce ll of the B2 phase and forms the unit ce ll of the C mmm (bas e ce nt e r ed orthorhombic) pha se. Th e unit ce ll s of these two phas es and ot h e rs that follow are s h o wn in Figur e 6. 38. F o ll o wing i s the crys tall ogra phi c description o f this distortion in s p ace gr o up n o tation: Pm3m P4 / mmm + Cmmm.

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296 The first step formed three variants of the P4/mmm (primitive tetragonal) phase since the P4/mmm space group is a type I maximal non-isomorphic subgroup of the Pm3m space group (43]. These three variants are formed by a compressive distortion in each of the three possible <100>B 2 face normals. The second step forms two variants of the Cmmm phase from each of the three tetragonally distorted variants of the P4/mmm phase. This is a tensile distortion that affects all of the atomic positions cooperatively and is accounted for by the relationship that shows the Cmmm space group as a type I maxjmal non-isomorphic subgroup of the P4/mmm space group [43]. The combination of the compressive and tensile distortions decreases the [100]B 2 and increases the [Oll]B 2 axes in the (0Tl)B 2 planes of the Pm3m phase (Figure 6.38a) to f'orm the (010) planes of the Cmmm phase (Figure 6.38b). The site occupancy of the Cmmm phase is determined from that of the B2 phas e, since no chemical exchanges occur between the atomic sites during the homogeneous distortion Taking into account the small distortions this indicates that the la (Nbtri) and lb (Al) sites in the B2 phase are converted into the 2a (Nb/I'i) and 2c (Al) sites in the Cmmm phase [43]. The 2a and 2c Wyckoff sites have fixed coordinates and are at (0, 0, 0) and( 0 ), respectively in the Cmmm phase. The unit cell of the Cmmm phase in Figure 6 38b shows the occupancy of the 2a and 2c sites normal to the [010] axis since this direction was derived from the [0Tl]B 2 axis in tl1e B2 phase (Figure 6.38a) Thus the base-centered symmetry on the (001) planes of the Cmmm phase are shown to be parallel to the (100) planes of the B2 phas e in this figure. Th e second distortion or heterogeneous distortion, involves the transition from the Crnrnrn phase to the Pmma phase in path 1 of Figure 6.37. The ORTH3 phase was shown previously in the structural analysis discussion to have the Pmma

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Ti N b (a) Al ., (100] Ti Nb (b) Al 0 [001] (c) T i Nb [010] :;() : :: .. .... .. .. .. (d) Nb (100] Figur e 6.38 Sh o w s th e uni t ce ll s o f t h e diff e r e n t stru ct ur es that l e d to th e f o rm at i o n o f t h e ORTHl pha se fr o m th e B2 ph ase. ( a ) th e B2 (Pm:f m ) s tru c tur e; (b ) th e o rth or h o mbi c (C mmm ) s t ru c tur e; (c) th e ORTH2 (Pmma ) st ru c tur e ; ( d ) t h e o r d e re d O RTHl (C m c m ) s tru cture.

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298 (primitive orthorhombic) structure Thus tl1e ORTH3 phase is the Pmma phase that is shown in path 1 of Figure 6. 37. This transition consists of the displacement of atoms on alternating (010) plan es, or shuffles, in th e Cmmm phase. These shuffles ca n b e described using Wyckoff positions and the subgroup/supergroup relation b et ween the Cmmm and Pmma space groups. The Pmma space group is a type Ila maximal non-isomorphic subgroup of the Cmmm space group. This indicates that there are two variants of the lower symmetry Pmma phase that form from the Cm mm phase. The occurrence of the heterogeneous distortion, which displaces atoms on every other (001) plane transforms the 2a and 2c Wyckoff sites of the Cmmm phase (Figure 6.38b) into the 2e and 2f Wyckoff sites of the Pmma phase. The unit ce ll of the Pmma phase normal to the [100] direction is shown in Figure 6.38c The site correspondence between th e two phases indicates that Nb and Ti share the 2e sites at (, 0 z) and (, 0, z) and Al occupies the 2f sites at (, z) and (, z). Thus the site occupancy of the Pmma phase which is derived from the B2 phase by changes in symmetry is the same that was experimentally determjned for the ORTI-I3 phase (Figure 6.36) in the structura1 analysis discussion. The two variants of the Pmma phase, that were mentioned previously to have formed from the Cmmm phase, are responsible for the stacking faults present in the ORTHl plates. These stacking faults were shown in Figure 6 23 to lie parallel to the ( 001) planes of the ORTH 1 plates. The connection between the two variants and th e stacking faults was determined by considering the shuffles on alternating (010) plan es during the Cmmm to Pmma transition. The direction of the sh1.tffles can occur in opposite [100] directions in the Cmmm phase. This will result in either a positiv e (+z) or negative (-z) displacement in reference to the unit cell of the Pmma phase. Th e unit cell of the Pmma phase in Figure 6.38c was constructed with a positive z

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299 displacement and shows an ( ... ABABA ... ) typ e of stacking sequence for the (100) planes. Howev e r, a negative z displacem e nt would hav e created a diff e r e nt stacking sequence and a stacking fault betw ee n th e two variants. The defect analysis showed that the stacking faults w e re parall e l to (001) 0 planes and w e re visibl e with all of the g r e flections exce pt for th e (002) 0 r e fl ec tion. Thes e observations are consistent with the 1/10<025> displa ce m e nt v ector for th e s ta c king faults and i s similar to the vector that has pr evio usly been observed in the O-Ti 2 A1Nb phase [31]. Th e disruption in the stacking sequence of the (001) 0 plan es du e to these faults is Aimilar to that caused by stacking faults in the HCP structures. Th ese changed the stacking sequence of the ( OOOl) H plan es from ( .. ABABAB .. ) to ( .. ABABA I C B C BC .. ), where I is the stacking fault [98]. Th ese stacking faults originate in the Cmmm to Pmma transition since the (001) 0 plan es of the ORTH! phase ar e derived from th e (100) planes of th e Pmma phase. Thus the stac king faults provided proof of th e Cmmm to Pmma tran si ti o n in path 1 Th e last transition in path 1 of Figur e 6.37 i s r es ponsibl e for forming the small co lumnar APDB s observed in the ORTHl (Figur e 6.27b and 6.27c) and ORTH2 (Figure 6. 30) plat es. From th e st ru c tural analysis section, it was found that the formation of th ese APDBs involv ed chemical o rd e ring b et w ee n Ti and Nb on two separa t e 4c Wyckoff sites in th e ORTH! phas es, whi c h ar e connected by the 1 / 2(100] displacement v ec tor This same conclusion can b e obtained from consideration of the su bgroup/sup e rgroup r e lations and th e Wyckoff positions in the Pmma and C mcm space groups. Thes e two space groups are present in the last transition of path 1 Th e C m c m space group i s a type IIb maximal non-isomorphi c subgroup of the Pmma space group [ 43]. This subgroup r e lation indicates that c h e mical ordering occurs in the Pmma space group that doubl es the b and c lattic e param ete rs and forms two

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300 variants of th e C m c m space group Th e two C m c m variants ar e observed as th e fin e APDBs in th e ORTHl phas es The C BED analysis showing Al occupying the 8g sit e in th e ORTHl phas e indicates that th e chemical ordering involves the 2 e Wyckoff site which is occupied by Ti and Nb and not th e 2f site which is occupied by Al in the Pmma phas e. Figur e 6.38 demonstrates this r e sult by showing th e unit cells and th e atomic site occupancies of th e Pmma (Figur e 6. 38c) and Cmcm (Figure 6. 38d) phases. Wh e n ordering occurs betwe e n Ti and Nb, two 4c sub latti ces ar e form e d in th e Cmcm pha se This causes th e b and c lattic e paramet e rs of th e Pmma phas e to doubl e in l e ngth to form th e a and b lattic e paramAters of th e Cmcm phase. This ordering transition r es ults in the alt e rnating occ upan c i es of Ti and Nb along th e rows pointing in t h e (100] dir ect i o n Th e translation v ecto r c onnP-cting Ti to Nb along thes e row s is 1/2(100] which i s the displac e m e nt vector that is r es ponsibl e for forming the fin e APDBs. Finally th e 2f sit e in th e Pmma phas e transforms into th e 8g sit e in the C mcm phas e and is occupied by Al Thus this last transition se e n in path 1 of Figur e 6.37 co nclud es th e d eve lopment of th e ORTHl plates pr e sent in alloy 2. 6 .3 .3.2 Path 2 of th e H C P Plat es Th e transformation of th e HCP phas e along path 2 (Figure 6.37) is diff e rent from that of th e ORTH phas e al o ng path 1 since it was found tl1at th e HCP plates form from a disord ere d j3 matrix. This transformation path was demonstrat e d by the o b se rvation of APDBs in th e B2 matrix that did not pass through th e HCP plat es ( Figur e 6.24a) Th e implication of this result is that th e disord e r e d sublattic e of th e j3 pha se is inherit e d by th e plates that form along path 2. This disorder e d sublattic e p e r s ists in th e subsequent transitional phas es of path 2 until the transition from th e ex. to cx. 2 phases occurs, which involv es c h e mi c al ord e ring. Th e occurrence of this transition was d e t e rmin e d from th e structural analysis of the APDBs that w e r e

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301 present in the HCP plates. This analysis showed that the displac ement v ector of these APDBs is 1/6<1120>, which is the same vector that forms during the disord er to order transition from the cx.-Ti to cx. 2 -Ti 3 Al phases [78] The first three transitions in path 2 of Figure 6.37 d esc rib e the homogenous and heterogenous distortions which are consistent with martensitic transformations using symmetry and subgroup/supergroup relations. These three transitions have the same purpose as the three transitions in path 1 that w ere previously described for the Pm3m to Pmma phases, except that these transitions involve phases that have a disordered sublattice. Thus, these three transitions which were consistent with the homogeneous and h eterogeneo us distortions and rigid body rotation (see section 6.3 1 ) w ere responsible for the formation of the twelve plate variants the plate morphology (Figure 6.19), and the habit plane (Figure 6.1) of the HCP plates. The homogeneous distortion consisted of the first two transitions in path 2 of Figure 6.37. This distortion c hang es the unit ce ll of the f3 phase to form the unit cell of the Fromm (face-centered ort h orho mbi c) phas e and is s hown in Figure 6. 39 along with the other phases that follow in path 2. Following i s the crystallographic description of this distortion in space group notation: Im3m 14/mmm + Fmmm. The first step of the homogeneous distortion form s three variants of the I 4/mmm (body-centered tetragonal) phase, whi c h is a type I maximal non-isomorphic subgroup of the Im 3m space group [43] These three variants are formed by a compressive distortion in each of the three possible < lOO >p face normals. A tensile distortion in either of the two < 110] directions in the 14/mmm phase forms two variants of the Fmmm phase. The distortions in this tranBition are uniformly applied to all of the atoms in the sublattice, since the Fromm space group is a type I maxim.al

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302 non-isomorphic subgroup of th e 14/mmm space group [43] These two transitions formed six plat e variants and two mor e variants formed for e ach of th ese six by the rigid body rotation Thus this accounts for th e tw e lv e plat e variants that w ere observed. Th e r e lationship betwe e n th e unit ce lls of th e Im3m and Fmmm phas es aft er the homogen eo us distortion has occurred is shown in Figur e 6.39 Th e uniqu e feature of th e site occupancy of th e Fromm phas e is differ e ntiated by th e 4a Wyckoff sites, s in ce these sites are d e riv e d from 2 a Wyckoff sites in th e p phas e The 4a sites hav e a di so rd e r e d site occupancy and are lo c at e d at th e (0, 0 0) origin and th e three fa ce ce nt e rs Figur e 6.39 indicates that th e (001) plan es of th e Fromm phas e (Figur e 6.39b ) ar e parall e l to the (011) plan es of th e p phas e (Figure 6.39a ). Th e third transition in path 2 of th e H C P plates is consistent with th e heterogeneous distortion for this mart e nsitic transformation [79] This transition inv o lv es displac e m e nts of atoms on alternating (001) plan es, or shuffles, in the Fromm phas e. Th ese shuffles can b e d e scrib e d with Wyckoff positions and by th e s ubgroup/sup e rgroup r e lati o nship b et w ee n the Fmmm and Cmcm space groups. Th e C m c m s pa ce group is a type Ila maximal non-i so morphic subgroup of th e Fromm s pa ce group and indi cates that the atomic displa ce m e nts ar e not uniformly appli e d to all th e atoms in tl1e unit cell of the Fromm phas e Th e r e ar e also two variant s of the l o w er symmetry C m c m phas e that f or m from th e Fromm phase in this transition Th e h e t e rog e n eo us distortion that causes the displa ce ment of atoms on every other (001) plan e transf o rm s th e 4a Wyckoff sites in th e Fromm phas e into th e 4c Wyckoff sites in the C mcm phase. Th e coordinates of the 4c sit es ar e (0, y and ( 0 y, and tw o mor e with the( , 0 ) latti ce translation. Th ese sites hav e a disord e r e d site occ upan cy. Th e r ef 'or e, th e shuffles are in the [010] dir ec tion of the Fromm pha se

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(a) (b) Nb Ti Al I Nb ) Ti Al I (100] (100] (c) Nb Ti Al ,......, 0 8 [100] . :-::::::: : (d) [1210] [2110] Figure 6.39. Sl1ows the unit cells of the different structures that led to the formation of the HCP phase from the p phase. (a) the disordered J3 (Im 3m) structure; (b) the disordered orthorhombic (Fmmm) structure; (c) the disordered orthorl1ombic (Cmcm) and th e disorder e d HCP (P6 3 /mme) structures; (d) the ordered HCP (P6 3 /mmc) structure.

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304 and are displaced by an amount equal to y in the [O 10] direction of the Cmcm phase The resulting unit cell of the Cmcm phase after these displacements have occurred is shown in Figure 6. 38c for the (001) planes. There are two special conditions that could have arisen in the transition from the Fromm to the Cmcm phase. The two conditions depend on the magnitude and the direction of the atomic displacements. These two parameters are indicated by the y parameter of the 4c Wyckoff site in the Cmcm phase. There are two scenarios that can occur for these special values of the y parameter. If y is equal in magnitude but opposite in direction for two variants of this transition, then a stacking fault will form between them. However, if y is equal to 0.333 then the Cmcm phase will degenerate into the higher symmetry P6 3 /mmc phase, provided that the lattice parameters of this phase have the critical ratio of 0.577, which it does for the hexagonal phase. The results of this study indicated that both of these scenarios occurred during the development of the HCP plates. Th e first scenario was indicated by the formation of stacking faults and formed during the Fromm to Cmcm transition. This was determined by the observation of stacking faults in the HCP plates. The analysis of the stacking faults seen in the plate of Figure 6.23 indicated that they were parallel to the (OOOl)H planes and were visible when imaged with all of the g reflections, except the (0002)H reflection These observations were consistent with the 1/6<2023> displacement v ec tor for the stacking faults and is common for stacking faults in the HCP phases l99]. The stacking faults signified that a disruption in the stacking sequence occurred for the (OOOl)H plan es from one showing ( .. ABABAB .. ) to one showing ( .. ABABA I CBCBC .. ), where I is th e stacking fault [98]. Thus these stacking faults

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305 originated in the Fmmm to Cmcm transition, since the (OOOl)H planes of the HCP phase are deriv e d from th e (001) plan es of th e Cmcm phas e. The occurrence of the second scenario was indicated by the degeneration of the Cmcm phase into the higher symm etr y P6 3 / mmc (a) phase. This was inferred since the plates that tran s formed along path 2 consisted of the a. 2 -HCP phase. As m e ntioned already the observation of APDBs with the 1/6<1120> displacement vector indicates that the a 2 -H CP plat es poss esse s a disord e r e d a-HCP structure prior to this ordering rea ct ion. Furthermore the r es ult that showed that the lattice param e t e rs of th e ORTHl plates formed by qu e nching from 1300 C had the same alb ratio of 0.577 as the HCP phas e impli e d that the condition for h exa gonal symm etry was met. Th e ORTH! plat es d e v e lop along path 1 and involve the inheritanc e of tw o sublattices from the B2 phase. Th e two sublattices prevent the Pmma phas e from d egener ating into a higher symmAtry HCP phase since the HCP phase do es not occur as a possibl e supergroup of the Pmma phase. How e ver in the case of the HCP plat es that d e v e lop along path 1 there is th e possibility of degeneration into th e lligher symmetry HCP structure. This is possibl e b ec ause the Cmcm phase of this path possesses a disord ered sublattice and does show the P6 3 / mmc as a minimal non isomorphic supergroup [43]. 6.3.4 Formation of Plate Martensit e from the 13 Phase Th e structural analysis and crystallographic treatment of the plat e martensit e in th e ce ntral portion of the Nb-Ti-Al system suggests that th e growth of the plates r eq uires some th e rmal activation The thermal activation is required during the formation of both the HCP and ORTI-Il plates since the disorder to order transitions that occur in thes e plates involv e the exc hang e betwe e n diff e r e nt types of atoms into separate sublattices in these structures. The thermal dep en dency of these

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306 mart e nsit e plates i s c onsistent with th e formation of mart e nsit e s in binary Ti-Al and 1:i-Nb alloys and in ternary Ti 3 Al + Nb alloys which was cov e r e d in chapt e r 2. Th e c on ce pt o f mart e nsit e start CMs) and finish (Mr) t e mperatur e s is us e ful t o co rr e lat e th e diff e r e nt o rth o rhombi c ( ORTH) s tru c tur e s with th e d e v e lopm e nt s tag es. It wa s pr e vi o usly s h o wn in chapt e r 2 that th e t e mp e ratur e was d e cr e a se d by th e addi t ion o f Nb to binary Ti-Nb alloys but was in c r e as e d by th e addition of Al t o binary Ti Al all o ys (73 74]. In t e rnary Ti 3 Al + Nb alloy s, th e studi e s hav e sh o wn th a t th e t e mp e ratur e o f th e a. 2 mart e nsit e was suppr e ss e d to b e low room t e mp e ratur e for Nb additions gr e at e r than ~12at .% (24 26 36]. Thus mart e nsit e plat e s did not form during rapid co oling in th e s e alloys with high Nb additions Th e c omposition o f a lloy 2 is 33Ti-27Nb-40Al (at.% ) and indicat es that th e amount of Al is suffi c i e n t t o c omp e nsat e f o r th e amount o f Nb in raising th e t e mp e ratur e abov e room t e mp e ratur e s o tha t t h e pl a t es co uld form during rapid co oling. This p e rmits pl a t es that f o rm c l o s e t o th e t e mp e ratur e to hav e mor e tim e to thick e n than tho se that f o rm c lo se t o th e M r t e mp e ratur e, whi c h i s b e low r o om t e mp e ratur e Th e d e v e l o pm e nt of diff e r e n t plat e s tru c tur es for diff e r e nt thickn e ss es in thi s st udy c an b e int e rpr e t e d to r e pr e s e nt various stag e s of plat e d e velopm e nt. This r e asoning c an b e appli e d t o th e d e v e lopm e nt of th e ORTHl structur e. Sin ce th e sam e habit plan e was obs e rv e d for all of th e s e plat e s ind e p e nd e nt of thickn es s th e n thi s impli es tl1at th ese plat e s all form e d fr o m th e f3 matrix by th e sam e m e chanism. Th e habit plan e o b se rv e d f o r th e s e plat es wa s co ns i st e ntly f o und t o li e ~ 11 fr o m th e ( 211) 13 plan es with t h e plat e rota t i o n about th e (011] 13 dir ec ti o n. Th e r e sult s of t hi s s tudy s ugg e st that t h e ORTH3 plat e s form e d initially fi~om th e B2 phas e, sin ce th ese pl a t e s w e r e th e thinn es t Assuming that th e ORTH3 plat e form e d clos e t o th e t e mp e ratur e, th e n this structur e can e volv e int o th e ORTH2 phas e as th e plat e

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307 thickens. The obs erv ation of small diffuse streaks in the SAED pattern of Figur e 6.14 showed proof of this transition. Th ese streaks in the ORTH2 phase wer e at positions cor r es ponding to a no scattering int e nsity for th e ORTH3 plates (see Figure 6.16) but to diffraction spots for the m e dium thick ORTHl plat e s (see Figures 6.9) If the disorder to order tran s ition from th e ORTH3 to ORTHl structure is complete then t h e final d e v e lopm ent o f the ORTHl structure occurs. Th e r e for e, th e ORTHl structure r e pres e nt s the last stage of d e v e lopm e nt and also correlates with these plates b e ing th e thickest o f th e diff e r e nt ORTH plat es. Th e ORTH3 plates showing that the Pmma symmetry is considered to b e a m etasta bl e structure that forms at room temperatur e Previous studies hav e s ugg es t e d that th e Pmma structure played a part in the d e v e lopment of the 0-Ti~Nb phase, but it was not observed [33 ,3 4]. Th e r e for e, the composition of th e phase in alloy 2, whi c h is inherit e d by th e ORTH3 plat es, must have stabilized the Pmma structure. Th e stability of the Pmma phase was expected to occur at low temperatures wh e r e the th e rmal activation process es w e re minimal to pr e v e nt the o rd e ring between Ti and Nb atoms o nto separate 4cl and 4c2 sites. The driving for ce b e hind this ordering was also e xp ec t e d to hav e b ee n high because of the diff e r e nt bond strengths b e tw ee n Ti Al and Nb-Al. However thermal energy was n ecessar y to driv e this ord e ring reaction, since Nb and Ti had to exc hang e atomic sites. Thi s reasoning then accounts for the ORTH3 plat e s b e ing v e ry thin (~ lOOA) sin ce they formed at l o w t empe ratur es witl1 insuffici e nt thermal e nergy for ordering to occur. 1"b. e ORTH2 plat es, that w ere slightly thicker than th e ORTH3 plates showed st r e aks d e v e loping in the SAED patt e rns at the (hkO) refl ec tions wher e h and k w ere o dd int egers Th ese streaks w e r e ca u se d by the fin e APDBs that w e r e form ed during the or d e ring from the Pmma to the C m c m (O RTHl ) structures. Th ey w ere caused by

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308 small displacements from the correct atomic positions in the Cmcm (ORTHI) structure and indicated that the ORTH2 structure was incommensurate. Therefore when the Ti and Nb atoms were locked into the correct atomic positions on the 4cl and 4c2 Wyckoff sites, then the commensurate Cmcm (ORTHI) phase was formed. This commensurate Cmcm phase caused the diffraction spots that were observed in the SAED patterns of the ORTHl phase (Figure 6.9). In this study, the HCP plates were found to be the thickest of the different plate structures. This observation is believed to be due to the enrichment of Ti in the composition of the J3 phase that the HCP plates formed from. The increase in the Ti content of the J3 phase lowers the disorder to order transition and increases the temperature. Thus the 1-ICP plates can form from a disordered J3 matrix and also at a higher temperature than the ORTH plates. This implies a longer development time for the HCP plates than for the ORTH plates and agrees with their larger thickness.

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In summary th e purpos e of this inv e stigation was to pr o vid e basic r ese ar c h o n t h e pha se e quilibria and pha se transf o rmations in th e c e ntral portion of th e t e rnar y Nb-Ti-Al syst e m as r e pr e s e nt e d by two alloys with compositions o f 27Nb-33Ti-40Al ( alloy 2) and 42Nb-28Ti-30Al (alloy 4 ). Following ar e th e main findings from this s tudy Th e B CC f3 pha se solidifi e d as th e pr i mary pha se in both alloy s 2 and 4 S o lid sta t e coo ling r es ul te d in th e o rd e ring o f th e primary f3 pha se t o th e B2 phas e, whi c h w as m e ta s tabl e at r o om t e mp e ratur e in th e s e alloy s In alloy 2 solid stat e c o oling a lso r e sult e d in th e formation of th e y -Ti.Al phas e as allotriomorphs and lath s along t h e primary f3 grain boundari e s plat e s in th e int e rd e ndritic r e gions and ro-r e lat e d pr ec ipitat es in th e f3 matrix Th e pr e cipitation of th e cr-NbzA]. phas e o c curr e d from th e f3 phas e at high t e mp e ratur e s in b ot h alloys 2 and 4 Th e cr transus was d e t e rmjn e d to b e slightly b e l o w 1400 C in all o y 2 and b e low 1550 C in alloy 4 Th e cr phase injtially form e d a s i s ola te d grain s in th e f3 matrix and al o ng f3 grain boundari e s in both alloys S e v e ral diff e r e nt ori e ntation r e lati o nship s w e r e o bs e rv e d b e tw ee n th e cr and f3 phas es, h o w e v e r th e {llO} (f plan es w e r e alw a ys parall e l to th e {110] 13 p l an e s. At l o w e r aging t e mp e ratur e s th e cr pha se form e d fr o m th e f3 phas e as c oloni e s of gra.ins that w e r e s imilarly ori e nted al o ng th e (00 l ] a dir e cti o ns Th e r e tain e d f3 parti c l e s in th ese 309

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310 mj c rostru c tur es s h o w e d irr e gular c ir c ular and e longated morphologies which dep e nd e d on th e formation and growth process of th e cr grains A eutectoid transformation of th e J3 phas e to th e cr + y phas e s occtrrred at 1200 C in alloy 2 This transformation was fa s t and was completed in th e RAM sa mpl e in l ess than two minut es Th e morphology of th e cr + y tw ophas e m i crost ru c tur e d e p e nd e d on th e h ea ting rat e Slow h e ating in the va c uum furna ce l e d to th e lath morphology whil e fa st h e ating in th e tub e furna ce l e d to e quiax e d y grains in th e cr ma trix. A discontinuous transformation of th e B2 phas e to the cr + J3 (disordered) occ urr e d at 1000C in alloy 4. This transformation was still in progr ess after two minut es of aging wh e n the sample was subsequently water qu e nch e d Th e lam e lla e structure co nsisted most l y of cr grain s and some disordered p particles An int e rfa ce wa s observed b e tw ee n the B2 and J3 phas es at tl1 e r eac tion front It was assumed that the P pha se w as m etasta bl e and w o uld hav e transf'orm e d with long e r aging tim es to the 0 phas e, which was ba se d on th e O-Ti 2 A1Nb phas e. This assumption wa s supported by the furnac e coo ling e xp e rim e nts and showed that plat es of the 0 phas e form e d from th e r e tain e d p particl es Th e J3 phas e transformed c omp l e t e l y to th e co D phas e in alloy 2 by e ith e r slow coo ling from high temperatures or aging at l ow t e mp e ratur es The structural analysjs showed that the ro D phas e had th e H C P structure and th e P6 3 / mcm space group. Th e stoichiometry of the co-D phas e was Al 4 Ti 3 Nb 2 and was based on the co mp os ition of the p phase. Th e SAED analy s i s s h o w e d that th e a latti ce param ete r wa s in crease d and t h e orientation r e lationship with th e J3 pha se was rotat e d by 90 as co mpar e d t o o th er ro related pha ses r e port e d in this system. I t was shown that t h e

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311 f o rmation of th e ro-D phas e from th e B2 phas e occurr e d by the {lll}J3 plan e co llaps e m ec hanism and inv o lv e d additional c h e mical ord e ring in th e atomic sit e o cc upan c y t hat was inh e rit e d by th e B2 phas e Th e pr o pos e d sit e occupancy of th e ro-D pha se indi c at e d that th e s ingl e lay e rs consist e d of Al atoms on th e 2b Wyckoff sit e and Nb a t o ms o n th e 4d Wy c koff sit e, whil e th e doubl e lay e rs c onsist e d of Ti atoms o n th e 6gl Wyckoff s it e and Al atoms on th e 6g2 Wy c koff sit e. Th e transformation path was d e scrib e d using subgroup and syrnm e try r e lations as Pm3m(B2) P31m(ro"' ) P6 3 / m c m (ro -D ) F ast coo lin g ra tes fr o m high t e mp e ratur es r es ult e d in th e mart e n s iti c tr ansformation o f th e J3 pha se to plat e s in a ll o y 2. Rapid s o lidification by e l e ctro m a gn e ti c (EM) l e vitati o n and drop qu e n c hing r es ult e d in a bask e t w e av e app e aran ce in th e mi c rostru c tur e o f alloy 2 that r e s e mbl e d a c i c ular mart e nsit e. Th e ag e d s ampl es o f alloy 2 th a t w e r e wat e r qu e n c h e d also show e d th e aci c ular mart e n s it e in r e gi o ns o f th e micr os tructur e that co n t ain e d th e J3 phas e. Th e obs e rv e d ha bit plan e o f t h e plat es agr ee d with that c alculat e d using th e invariant lin e th e ory Th e f o rmati o n o f th e mart e nsit e plat e s was th e rmall y a c tivat e d and involv e d ch e mical ord e ring during qu e n c hing Th e J3 composition aff ec t e d th e stru c tur e of th e plat e s a nd co n s ist e d of e ith e r th e orthorh o mbi c s tructur es with th e Pmma and Cmcm spa ce gr o up s o r th e H C P s tru c tur e with th e P6 3 /mm c spa ce group Th e stru c tur e o f th ese plat es was d e p e nd e nt on tlri c kn ess Th e prop o s e d s it e o c cupancy o f th e ORTHl s tru c tur e was ba se d on th e Al 2 TiNb s t o ichiom e try and Al atom s o c cupi e d th e 8g Wy c koff sit e Ti at o ms o c cupi e d th e 4 c l Wyck o ff sit e, and Nb at o ms o cc upi e d th e 4 c 2 Wy c koff sit e Analy s is of d o main s tru c tur es, s ta c king faults and e l e ctron diffra c ti o n s ugg e st e d tw o pos s ibl e transf o rmati o n paths: Im3m ( J3 ) Cmcm(disord e r e d )

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312 P6 3 / mmc(disord e r ed) P6 3 / mmc(D0 1 9 ) for HCP plates and Pm3m(B2) Pmma C mcm(ordered) for orthorhombic plat e s.

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CHAPTERS FUTURE WORK In this study, as in any study, more questions and id eas aris e as one finds the answers and solutions to pr e vious qu estio ns Following are some ideas of futur e w or k that could b e pursu e d : 1) D e t e rmin e th e transus t e mperatur e of th e high temp e rature f3 phase for other compositions in this central area of th e Nb-Ti-Al system. This knowl e dg e is n ee ded for furth e r alloy d e v e lopm e nt since h e at treatments could then b e d es ign e d that w o uld p e rmit control over th e microstructur e 2) Do a systematic st udy of th e competition b e tw ee n th e ro-D phas e and th e mart ensi t e plat e formation in the 27Nb 3311 40Al alloy. Thi s study would involve s tudying the diff e r e nt phonon mod es, which w e r e observed as diffus e e l ect ron scattering in th e diffra ctio n patt e rns, and how th ey might affect the transformation of th e f3 phas e to eac h of th ese structures. 3) D o channel e nhanc e d microanalysis to confirm the atomic site occupancies of the B2 phas e, the ro-D phas e, and the mart e nsit e plat es that formed in th e 27Nb 33Ti 40Al alloy. This would h e lp clarify the role that aluminum may hav e had on th e o rd e red structures of these phases. 313

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318 69. D L. Moffat and U.R. Kattn e r Met. Trans. A 19A ( 1988 ), 2389. 70. A T Bal ce rzak and S.L Sass M e t Trans. A a (1972) 1601. 71 0. Lyon J L es s Co mm o n M et., 81 ( 1981 ), 103 72. D.L Moffat and D. C Larbal estie r M e t Tran8. A 19A (1988), 1677. 7 3. D .L Moffat and D C Larbal est ier M et Trans. A 19A (1988), 1687 74 K.S. J e ps o n A.R.G Br o wn and J A Gray Th e S c i e n ce T ec hnology and Application of Titanium R J Jaff ee and N E. Promis e e ds P e rgam on Pr ess, London (1970), 677. 75 Y A. Bagariatskii G.l. Nos o va and T.V Tagunova Sov. Phys Dok.lady 3, (1958), 1014 76 A R.G. Brown D C lark J. Eastabrook and K.S. Jepson, Natur e, 210, ( 1964 ), 914. 77. S.M.L. Sastry and H A Lipsitt, M et all Trans A SA (1977), 1543 78. M.J. Bla c kburn Th e S c i e n ce, Technology and Appli ca tion of Titanium R.J J affee and N.E. Promis e e d s., P e rgamon Pr ess, London ( 1970 ), 633 79 C .M. Wayman Introdu ct ion to the Crystallography of Mart e nsitic Transformation Macmillan N e w York (1964). 80 T. Furuhara H.J. Le e, E.S.K. M e non and H I Aaronson M eta ll. Trans. A 21A (1990) 1627. 81. S Far e n c, A Co ujou and A Co ur et, Phil. Mag. A 67 1 ( 1993) 127 82. R .E. Smallman Mod e rn Phys ci al Metallurgy Butt e rworth London ( 1985 ). 83. R E. R ee d Hill, Physical M e tallurgy Principl es, Brooks/Co l e Engin ee ring Division M on t erey, CA, (1973). 84. M. Enomoto a nd H Harada M e tall Tran s., 20A ( 1989 ), 649 85. P. Hir sc h A H o wi e, R B Ni c holson D W Pashl ey, and M J. Wh e lan Electron Mi c ro sco py of Thin C rys tals, Rob e rt E. Kri e g e r Publishing Co., Malabar FL ( 1977 ). 86. J R .C. Gomez MS Th esis, UF ( 1990 ). 87 D. Turnbull Acta Metall ., 3 ( 1955 ), 55

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319 88. A Pawlowski and W. Truszkowski, Acta Metall. Mater 38, 1 (1990), 135. 89 J C.H. Spenc e and J M Zuo Electron Microdiffraction, Plenum Press, New York, (1992). 90. M H. Loretto, in Electron Beam Analysis of Materials, Chapman and Hall London, (1984). 91 B F. Buxton J.A. Eades J.W Steeds and G.M Rackham Phil. Trans. R. Soc. 281 (1976), 181. 92. M. Tanaka and M. Terauchi Convergent Beam Electron Diffraction JEOL LTD. Tokyo Japan (1985). 93. H. Ba kker, Mat. Res. Soc. Symp Proc 21 (1984), 319. 94. U. Dahm e n and K.H. Westmacott Acta Metall 34, 3, (1986), 475. 95. G.B. Olson, Frontiers in Materials Techno l ogies E l sevier, Amsterdam (1985), 43. 96. J. W. Edington Practical Electron Microscopy in Materials Science N V. Philips' Gloeilampenfabrieken, Eindhoven (1974). 97. L D Landau and E.M Lifshitz Statistical Physics Pergamon Press London (1980). 98. G.E. Di ete r Mechanical M eta llurgy McGraw-Hill New York (1976). 99. R.E. Smallman, Modern Physcial Metallurgy, Butt e rworth, London, (1985).

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BIO G RAPHI C AL SKET C H D a vid Tim o thy H oe l ze r wa s b o rn Jun e 17 19 5 9 at Fort Dix NJ H e wa s f o r t un a t e t o b e b o rn in a f'amily t hat was e nlist e d in th e Unit e d Stat e s Air For ce This all o w e d him t o gro w up in f o r e ign c ountri e s s u c h as th e Philippin es and G e rmany and in many s t a t e s in Am e ri c a. H e gradua te d fr o m T a bb High Sch o ol VA in th e s pring o f 1977 H e th e n e nr o ll e d at th e Univ e rsity of Florida Gaine s vill e, FL in th e Fall of 1977 and dis co v e r e d th e world of m a t e rial s sc i e n ce in 1978 as h e work e d part tim e in t h e M a t er i a l s S c i e n ce and Engin ee ring D e partm e n t. Wh e n financ es b eca m e e xhau s t e d h e w o rk e d for Rus s s Sh eet M et al in G ain es vill e, FL from 1980 t o 1982 and gain e d a valua b l e s kill installing h e ating and air co nditi o ning s yst e m s. H e r e turn e d to th e Univ e rsity of Fl o rida in 1982 and c ompl e t e d his Bach e l o r s D e gr ee in Mat e rials S c i e n ce a nd Engin ee ring in 1985. H e was admitt e d to th e Graduat e S c h oo l a t t h e Univ e rsity o f Florida in 1985 Soon aft e r b e ing admitt e d into graduat e s ch oo l h e m e t th e form e r Amy Ann Tuttl e whom h e marri e d in 1987 In 1991 h e r ece iv e d his M as t e rs D e gr ee f o r r ese ar c h w o rk that wa s d o n e b e tw ee n 1985 and 1988 H e work e d o n hi s do c t o ral r esearc h fr o m 1988 until 1991 H e wa s bl e ss e d with th e birth o f his d aught e r Ra c h e l Lynn in D e c e mb e r 1991. In 1992 h e b ec am e e mploy e d at th e N e w Y o rk Sta te Co ll e g e o f Ce rami c s Alfr e d NY wh e r e h e is still pr e s e ntly e mpl oye d 320

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I certify that I have read this study and that in my opinion it conforn1s to acceptable standards of scholarly presentation and is fully adequate, in scope and quality, as a dissertation for the degree of Doctor of Philosophy. Fereshteh Ebrahimi, Chairperson Associate Professor of Materials Science and Engineering I certify that I have read this study and that in my opinion it conforms to acceptable standards of scholarly presentation and is fully adequate, in scope and quality, as a dissertation for the degree of Doctor of Philosophy. Robet T DeHoff Professor of Materials cience and Engineering I certify that I have read thi s s tudy and that in my opinion it conforrns to acceptable standards of scholarly pre se ntation and i s fully adequate, in scope and quality, as a dissertation for the degree of Doctor of Philo sop hy. Michael J. Kaufm Prof essor of Ma rials Sci enc nd Engineering I certify that I have read this s tudy and that in my opinion it confor111s to acceptable standards of scholarly presentation and is fully adequate, in scope and quality, as a dissertation for the degree of Doctor of Philosophy. Elli s D Verink, Jr. Distinguished Service Professor of Materials Science and Engineering I certify that I have read this study and that in my opinion it confor111s to acceptable standards of scholarly pre s entation and is fully adequate, in scope and quality, as a dissertation for the degree of Doctor of Philosophy. (' Anna J. Brajter-T th A ss ociate Professor of Chemistry

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This dissertation was submitted to the Graduate Faculty of the College of Engineering and to the Graduate School and was accepted as partial fulfillment of the requirements for the degree of Doctor of Philosophy. December 1996 W--~ Winfred M. Phillip s Dean, College of Engineering Karen A. Holbrook Dean, Graduate School

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LO 1780 1990 UNIVERSITY OF FLORIDA II I I l I I I I I 11 I II I 3 1262 08554 9425


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