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
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xvii, 320 leaves : ill. ; 29 cm.
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English
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Hoelzer, David Timothy, 1959-
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Materials Science and Engineering thesis, Ph. D   ( lcsh )
Dissertations, Academic -- Materials Science and Engineering -- UF   ( lcsh )
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bibliography   ( marcgt )
non-fiction   ( marcgt )

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

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University of Florida
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aleph - 026316259
oclc - 36683097
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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




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