Acta metall, mater. Vol. 40, No. 11, pp. 3035-3049, 1992
Printed in Great Britain. All rights reserved
0956-7151/92 $5.00 + 0.00
Copyright © 1992Pergamon Press Ltd
MECHANISMS OF ISOTHERMAL PHASE TRANSFORMATIONS
IN RAPIDLY SOLIDIFIED Ti-24Al-11Nb
L. M. HSIUNG, W. CAI and H. N. G. WADLEY
Department of Materials Science and Engineering, University of Virginia, Charlottesville,
VA 22903-2442, U.S.A.
(Received 6 April 1992)
Abstract--Isothermal phase transformations of a rapidly solidified Ti-24AI-I 1Nb alloy at temperatures
below 900°C have been studied using electron microscopy (SEM and TEM) and X-ray diffraction (XRD)
methods. A supercooled fl phase was found to retain in the alloy as a result of rapid cooling. Subsequent
isothermal treatments resulted in a fl ---+~2 transition which took place through a complex sequence of
phase transitions. Initially it proceeded by ordering transitions fl (disordered b.c.c.)--, B2 (ordered
b.c.c.) ---}T (ordered tetragonal) and was then followed by T ~ O (ordered orthorhombic) ---}~2 (ordered
hexagonal) transitions. Plate-like O and a2 phases were found to form at the early stages of both the T ~ O
and O---~a 2 transitions. The T/O interface (habit) plane was determined to be {223}r, and the O/~ 2
interface (habit) plane was determined to be {lI0}o. Both the T----}O and O---} ~2 transitions can be
explained by a shape deformation mechanism involving a homogeneous lattice distortion and a rigid-body
rotation. In addition to the shape deformation, diffusion of Nb away from the O/~ 2 interface also is needed
for the O----,~2 transition to progress.
R~um~--On 6tudie les transformations de phase isothermes d'un alliage Ti-24Al-11Nb, rapidement
solidfi6 fi des temperatures infrrieures /t 900°C, par microscopie 61ectronique en transmission et en
balayage (MET et MEB) et par diffraction des rayons X. On trouve une phase fl surfondue qui reste dans
I'alliage par suite du refroidissement rapide. Des traitements isothermes ultrrieurs provoquent la transition
fl ~ ~2 qui se produit fi la suite d'une srquence complexe de transitions de phases: Au drbut, il se produit
des transformations d'ordre fl (c.c. drsordonnr)---}B2 (c.c. ordonnr)---}T (quadratique ordonnr), suivies
par les transitions T---~O (orthohombique ordonnr)---}~2 (hexagonal ordonnr). On trouve que les phases
O e t ~2, en forme de plaquettes, se forment aux premiers stades des transitions T---}O et O--}a 2. Le plan
d'interface (plan d'accolement) T/O est {223}r et le plan d'interface O/~2 est {1T0}o. Les deux transitions
T--*O et O---}~2 peuvent s'expliquer par un m~canisme de drformation de forme qui implique une
distorsion homogrne et une rotation rigide du rrseau. En plus de la drformation de forme, la diffusion
du Nb hors le l'interface O/~2 est aussi n~.essaire pour que la transformation O---~a2 s'effectue.
Z~-,ammenfassung--Die isothermen Phasenumwandlungen in der rasch erstarrten Legierung Ti-24AI11Nb werden bei Temperaturen unterhalb von 900°C elektronenmikroskopisch (Raster- und Durchstrahlungselektronenmikroskopie) und mit Rfntgenbeugungsmethoden untersucht. Als Ergebnis der
raschen Erstarrung bleibt eine unterkfihlte fl-Phase in der Legierung zuriick. Nachfolgende isotherme
W/irmebehandlung fiihrt zum ~bergang fl---}~2, der fiber eine komplexe Folge von Phaseniiberg/ingen
l~iuft: zu Anfang l/iuft sic fiber Ordnungsumwandlungen fl (entordnet k.r.z.)---}B2 (geordnet k.r.z.)--+T
(geordnet tetragonal), gefolgt von T---}O (geordnet orthorhombisch)---,a 2 (geordnet hexagonal). Plattenertige O- und ~2-Phasen bilden sich im frfihen Stadium der Umwandlungen T---}O und O---~a2. Die
T/O-Grenzflfichen-(Habit)-Ebene wird zu {223} bestimmt, die O/~2-Ebene zu {li0}o. Sowohl der
Obergang T--,O, wie auch O---*~2 krnnen mit einem Form-Deformationsmechanismus, der eine
homogene Gitterverzerrung und eine Festkrrperrotation einschlieBt, erkl~irt werden. Zus/itzlich zur
Formdeformation wird die Diffusion von Nb weg von der O/a2-Grenzfl/iche ffir den l~bergang O---~:c2
benrtigt.
I. INTRODUCTION
Ti3Al-base aluminide alloys are candidate materials
for high temperature structural applications because
of their low density, good elevated temperature
strength and creep resistance. Advanced intermetallicmatrix composites (IMCs) that use the Ti3AI
aluminide alloys as a matrix can offer even better
performance in stiffness and strength over their
monolithic counterparts [1]. Ti3A1 aluminide alloys,
however, are difficult to be fabricated into structural
components by conventional processing methods due
to limited ductility and low toughness, To overcome
this, researchers are exploring the use of inductivecoupled plasma-deposition (ICPD) [1, 2] to produce
composite unidirectional monotapes followed by hot
3035
3036
HSIUNG eta/.:
RAPIDLY SOLIDFIED Ti-24Al-11Nb
isostatic pressing (HIPing) or vacuum hot pressing
(VHPing) to make a multilayer net shape component.
Before one can control and optimize the matrix
microstructures (and thus the IMC's mechanical
properties) by controlling the deposition and subsequent consolidation processes, a full understanding
of time-temperature-transformation (TTT) as well as
isothermal transformation behavior of the Ti aAl-base
aluminide alloys is needed. The TTT behavior
of different Ti3Al-base aluminide alloys have been
investigated by Weykamp et al. [3], Peters and Bassi
[4], Djanarthany et al. [5], and Hsiung et al. [6].
Nevertheless, inconsistency can still be found among
their results in terms of transformation sequence. In
addition, little has been mentioned about mechanisms,
and a systematic study of the TTT behavior of the
TiaAl-base aluminide system has yet to be reported.
A candidate Ti3Al-base aluminide system used as
the matrix material for fabricating IMCs is Ti3A1
with an addition of Nb. Nb is added to stabilize the
fl phase and improve the deformability of the Ti3A1base aluminide at lower temperatures [7]. Studies of
microstructure in quenched (or rapidly solidified) and
aged Ti 3A1 + Nb alloys have revealed the existence of
numerous phases [8-14]. These include fl (disordered
b.c.c.), B2 (ordered b.c.c.), to-like (ordered h.c.p.),
PREP
T (ordered tetragonal), O (ordered orthorhombic)
and a2 (ordered h.c.p.) phases. Three different modes
of isothermal phase transitions have been recognized
in the Ti3A1 + Nb system [4, 6]: (I) fl --, ~2 + fl,
(II) fl--, a2(+O) and (III) fl--* to. For the case of
Ti-24Al-llNb [6], mode (I) was found to occur
between 900 and 1150°C, mode (II) below 900°C, and
mode (III) was found to be a competition reaction to
the mode (II) reaction at temperatures below 650°C.
We report here on a detailed study of the mode (II)
reaction. The effort has focused upon elucidating the
mechanisms of the mode (II) reaction.
2. EXPERIMENTAL
A rapidly-solidified Ti3AI+ Nb aluminide alloy
with a nominal composition of Ti-24at.%Alll at.%Nb (Ti-14wt%A1-21 wt%Nb) was chosen
for this study. Two forms of this alloy were used:
powder and foil. The powder was produced from
ingot via the plasma rotary electrode process (PREP)
by Nuclear Metals Inc., Concord, Mass. The foil was
produced from powder of identical composition via
the inductively coupled plasma deposition (ICPD)
process at GE Aircraft Engines, Lynn, Mass. [2].
During the ICPD process, powder was melted by
Powder
110
(b)XRD
E
211
200
i
|
I
30
i
,
,
I
,40
I
i
~
I
50
I
I
!
I
|
60
I
!
i
I
I
70
2#
Fig. I. (a) SEM micrograph showing microstructure of the as-prepared PREP powder, (b) XRD pattern
(Cu K~) showing the existence of the B phase in the as-prepared powder.
I
a
HSIUNG et al.: RAPIDLY SOLIDFIED Ti-24AI-I 1Nb
passing it through a plasma. The molten droplets
were then immediately deposited onto a preheated
mandrel inside a vacuum chamber where they were
rapidly quenched to a solid state. This mandrel was
preheated to ~ 800°C and held at this temperature
during deposition. Thus, during the deposition
process the alloy first deposited was cooled to
~800°C, and was held at this temperature until
completion of the deposition, and was then cooled to
room temperature. The total deposition period lasted
approximately 1 h. Therefore the alloy first deposited
has already been preaged.
Both the powder and foil were examined in their asprepared and aged status using X-ray diffractometry
(XRD), electron microscopy (SEM and TEM), and
energy dispersive X-ray spectrometry (EDS). Prior to
aging, specimens were wrapped with tantalum foils
and encapsulated in cleaned and evacuated quartz
ampoules. Aging was performed for various times in
the temperature range 450 to 900°C. Following the
heat treatment, Vickers hardness of PREP powders
was measured using a MICROMET microhardness
B
i~°
XRD
ICPD Foil: As-sprayed
I
(a) SideA
B
o
~
o
B
o
20
O0 I ~0
=(
I
(b)Center
ool
~
0
;k
~12~
3037
indentor. TEM specimens were prepared from both
the as-prepared and aged ICPD foil. Microstructures
were examined using a Philips 400T transmission
electron microscope. Selected-area electron diffraction
(SAD), microdiffraction (MD) and convergent-beam
electron diffraction (CBED) methods were applied to
identify and distinguish the different phases.
3. RESULTS
3.1. As-processed microstructures
3.1.1. PREP powder. The microstructure of the asprepared powder was studied using XRD and SEM.
The results are shown in Fig. 1. In contrast to the
formation of martensitic microstructures reported for
Ti3AI (without a Nb addition) PREP powder [15],
a cellular-dendritic microstructure was observed in
this alloy [Fig. l(a)]. The XRD result shown in
Fig. l(b) indicates that the/7 phase was the only phase
to form in the as-prepared powder. Since the Ms
temperature of the b.c.c. ~ h.c.p, martensitic transformation in the Ti 3AI + Nb alloys containing more
than 5 at.% Nb is known to be well below the room
temperature [9], the retention of the /7 phase in the
PREP powder prepared by rapid cooling from the
stable fl phase regime is to be expected. Similar
results have been reported in Jominy end-tested
Ti-24AI-I 1Nb specimens [3]. There, the fl phase was
found close to the most rapidly-cooled end of the
sample.
3.1.2. ICPDfoil. While the fl phase was the only
phase found in the as-prepared PREP powder, an
initially confusing variety of different phases including
~/B2, T, O and ~t2, were found to exist in the asprepared ICPD foil. We have been able to understand
this by considering the thermal histories created
during the deposition. XRD results generated from
different foil depths are shown in Fig. 2. The r/B2
phase was found to exist mainly in a region close
to top foil side (side A hereafter) that was the last
0
042
0
431
20
(c) Side B
":'
a2
I"=
=e
o~' {°9.
a,
Ill
4o
a2
=o.=
so
2e
a~
==..o =o'=1
do
"Io
Fig. 2. XRD pattern (Cu K=) generated from the as-sprayed
ICPD foil, (a) side A (last deposited), (b) center portion and
(c) side B (early deposited).
Fig. 3. Dark-field TEM image showing the formation of a
network of thermal anti-phase boundaries within a B2 grain,
Z (zone axis) ~[100], g = 001.
3038
HSIUNG et al.:
RAPIDLY SOLIDFIED Ti-24A1-11Nb
HSIUNG et al.: RAPIDLY SOLIDFIED Ti-24AI-11Nb
deposited (least preaged), and had been almost
immediately cooled after deposition [Fig. 2(a)]. The
O phase existed mainly in the center of the foil
[Fig. 2(b)], whilst the ~t2 phase was found as one
approached the first deposited side (side B hereafter)
of the ICPD foil [Fig. 2(c)]. Note that the deposit on
side B was directly attached to the mandrel, and
received a substantial amount of preaging at ~ 800°C
during deposition. TEM studies of the microstructures
for different depths (and therefore different preaging
status), within the ICPD foil are shown in Figs 3-5.
Figure 3 is a dark-field image (obtained near to side A)
showing the formation of a network of thermal antiphase boundaries (APBs) formed within a B2 grain.
The existence of the ordered b.c.c. (B2) phase in this
region was confirmed using microdiffraction (MD)
and CBED methods• M D and CBED patterns of
the <001 >B2, < 111 >s2 and < 110>a 2 zones are shown in
Fig. 4. Superlattice reflections of the B2 phase can be
found in the M D patterns, and the cubic symmetry
of the B2 phase is displayed in the CBED patterns.
Thus, a fl --~ B2 ordering reaction clearly occurred on
side A during the cooling that followed deposition.
Detailed TEM studies of the T, O and ct2 phases
in isothermally aged foils will be demonstrated later.
First we consider the morphology of the microstructures containing the T, O and ct2 phases in the
3039
as-prepared foil. Figure 5(a) is a dark-field image
showing a typical microstructure observed near the foil
center. A plate-like O phase was found to form within
a matrix of the T phase. A typical microstructure observed in a region near to side B is shown in Fig. 5(b).
Plate-like ct2 phase was found to form within a matrix
of the O phase. From these distinct microstructures
observed in the different portions of the as-prepared
foil, it was realized that the isothermal/~ ---, ct2 transformation must take place through a complex reaction
sequence fl --. B2 ~ T---* O ~ ot2 during foil deposition at ~ 800°C. To confirm this reaction sequence,
a series of isothermal aging treatments and analytical
experiments using XRD, TEM and electron diffraction techniques was carried out, and are presented
below.
3.2. Isothermal phase evolution sequence
3.2.1. PREP powder. Powder samples were isothermally aged at 450, 650 and 850°C. X R D results
obtained from samples after aging for different times
at 450°C are shown in Fig. 6. The T phase was found
in a sample after aging for 30 min [Fig. 6(a)]. The O
phase was found in a sample after aging for 24 h at
450°C [Fig. 6(b)]. Since the reaction of the O--4 ct2
transition was sluggish at 450°C, the ct2 phase was
e
o •
PREP
Powder
•
T
XRD
r
• ~ _ ~ 0 "
,~"~
(a)450°C,
~e"
¢
30 min
- ..
•
i
j
. . . .
200
o
. . . .
.o" . . . . . .
so
~ ' ' 'o ~ '
2o
(b)450°C,
24 h
e
'I
o
Fig. 5. (a) Dark-field TEM image showing a morphology
of the plate-like O phase formed within a T matrix, Z
[I 11ITII[110]o. (b) Dark-field TEM image showing a morphology of the plate-like ct2 phase formed within an O matrix,
Z ~ [T14]o II [T216L2.
2#
Fig. 6. XRD pattern (CuK~) generated from a powder
sample after aging for various times at 450°C, (a) 30 min,
(b) 24h.
3040
HSIUNG et aL: RAPIDLY SOLIDFIED Ti-24Al-11Nb
PREP Powder
o
o~
.~
oo
~
(a) 650oc, lomin
V
T
=Tp'
0o
o
02
.~
(b) 650°C, 2 h
*g
.
~.
~0
Fig. 7. XRD pattern (CuK,) generated from a powder
sample after aging for various times at 650°C (a) 10 rain,
(b) 2h.
barely detected even after aging for a week. XRD
results obtained from samples after aging for various
times at 650°C are shown in Fig. 7. Both the T and O
phases were found after aging for 10 min [Fig. 7(a)].
However, the ~2 phase was found after aging for
2 h at this temperature [Fig. 7(b)]. The XRD result
generated from a sample after aging for 24 h at 850°C
is shown in Fig. 8. Only the ct2 phase was then
detected.
Microhardness measurements were made as
function of aging time from powder samples. Typical
results obtained from the samples after aging for various times at 650 and 450°C are shown in Fig. 9(a,b).
An age hardening effect was observed for both temperatures. A hardness increase occurred during the
fl-+ O transition stage which we attribute to the
formation of plate-like O phase within the parent
matrix [Fig. 5(a)]. The hardness reached a maximum
when samples contained only the O phase. While the
hardness decreased during the O--~ ~t: transition at
650°C and approached a low plateau when the powder
contained mainly the ~q phase [Fig. 9(a)], the hardness
drop was relatively insignificant at 450°C even though
the aging time was extended to a week [Fig. 9(b)].
This again suggests the sluggish nature of the O --+ ~t2
transition at 450°C. Note that similar hardness
distributions also has been measured from Jominy
end-tested Ti-24Al-llNb specimens [3], where the
hardness was measured as a function of distance from
the quench end. These results clearly point to the
existence a time-temperature-transformation behavior
in the Ti-24Al-llNb alloy.
3.2.2. ICPD foil. Foil samples were isothermally
aged at 650, 800 and 900°C. TEM specimens were
prepared from the portion near to the side A, where
the original microstructure contained only the B2
phase (Fig. 3). The T phase was found after aging for
l0 min at 650°C. Figure 10(a) is a dark-field image
showing the morphology of the T phase. Notice the
existence of antiphase boundaries (marked by arrows)
within the T phase. MD patterns generated from the
[10017, [111]x and [01 lit zones are shown in Fig. lO(b).
In order to compare these patterns with those of
o
850Oc,24 h
*
u2
o
o
o
o
0
AI
30
•
•
40
2O
Fig. 8. XRD pattern (Cu K~) generated from a powder sample after aging for 24 h at 850°C.
HSIUNG et al.:
700 ~
....... ,
. . . . . . . . , , ....... ,
........ ,
RAPIDLY SOLIDFIED Ti-24AI-11Nb
........ ,
.......
700 -~
la)
.......
,
........
,
........
3041
,
........
'
........
,
........
'
.......
,
........
~-
.....
(hi
600
>
"r
v
:~>' 500
500
(/)
(/)
40Q
t~
-r"
-r
300
300
Temperature= 6 5 0 " C
200
i?? . . . . . . . i
........
101
0
,
........
102
Temperature = 450°C
,
........
103
Time
,
........
10 4
,
.....
20o
10 5
10 s
x)~ ....... '
0
........
101
'
........
102
(sec)
103
Time
10 4
10 5
10 6
(sec)
Fig. 9. Vickers hardness (load: 100 g) plotted as function of aging time measured from powder samples
after aging at (a) 650°C and (b) 450°C.
the B2 phase shown in Fig. 4, the reflection spots in
the Fig. 10(b) were indexed on the basis of the B2
reflections. Notice that the forbidden reflections can
be found on various ½(i10)B2, ½ ( 1 1 2 ) B 2 ,
¼(i12),2
and ~(i12)B2 positions [14]. The tetragonality of the
T phase has been determined (from the MD patterns)
to be ~ 1.02 [14]. When the aging time was extended
to 40 min at 650°C, a plate-like O phase was found
to form within the T matrix [Fig. 1l(a)]. Two different
orientation variants of the O plate can be found in
the Fig. ll(a). The T/O interface (habit) plane was
measured to be ~ 4 4 ° away from (001)r, i.e. parallel
to (~23)x and (2~3)T [Fig. l l(b)]. Selected area
diffraction patterns generated from the (T + O) two
phase region is shown in Fig. 1l(c). The orientation
relationships between the T and O phases were
derived from the SAD patterns to be (i]-0)r II (001)o,
and [li0]T ^ [010]o = [001]T A [100]O ~ 4.5 °.
When the aging time was further extended, more
O plates were found to nucleate within the T matrix.
(b)
e
"
"
•
"
•
•
•
O
I
[loo]
0
I
I
0
•
O
tj
[_111]
B
e
w
"
•
•
Ot
tlt
It
[011]
Fig. 10. (a) Dark-field TEM image showing a morphology of the T phase observed in a foil sample after
aging for 10 min at 650°C, Z ~ [110]. g = lI10, (b) MD patterns generated from the [100IT, [111]T, and
[011]T zones.
3042
HSIUNG et al.:
RAPIDLY SOLIDFIED Ti-24Al-I 1Nb
//'/
li0
(b)
--
OOl
v~t
variant 2
~
52a ......,,.,
iio
1...-%
/
/
trace of T/O interface
(habit) plane
0
!
Fig. 11
HSIUNG et al.:
RAPIDLY SOLIDFIED Ti-24AI-I1Nb
3043
(a
e
•
4)
s
•
0
0
•
f
e
@
#
A
5S.3°
Fig. 12. MD and CBED patterns generated from the (a) [010] and (b) [001] zones of an O grain showing
a 2 mm symmetry.
Eventually, the T phase disappeared and the morphology of the O phase evolved from a plate to an equiaxed grain with a plate-like ct2 phase formed within
the O matrix. M D and CBED patterns generated
from the [010] and [001] zones of an O grain formed
in a sample after aging for 2 h at 650°C are shown in
Fig. 12. A 2 mm symmetry is displayed in both the
[010] and [001] patterns [10]. A typical observation of
the (O + ct2) microstructure made from a sample aged
for 4 h at 650°C is shown in Fig. 13(a). The O/~t2
interface (habit) plane was found to be parallel to
(ll0)o. Selected area diffraction (SAD) patterns
generated from the (O + ~2) two phase region is shown
in Fig. 13(b). M D patterns generated from the [001]o
and [0001]~2 zone are shown in Fig. 13(c). The orientation relationships between the O and ~t2 phases can
be derived from the SAD patterns: (001)o II (0001)~2,
[100]o A [11~0]~2= [1310]o A [I 100]~2~ 1.5 °. The ~2 plates
eventually grew and coalesced to an equiaxed ~t2
grain as the aging time was extended to 24 h at 650°C
[Fig, 14(a)]. Notice that the plate-like structure disappeared and only a small number of dislocations
were left within the ct2 grain. CBED patterns generated from the [0001] zone of a coarsened ~t2 phase are
shown in Fig. 14(b). A 6 mm symmetry is displayed
in the CBED patterns. Occasionally, a small amount
Fig. 1I. (Opposite) (a) Dark-field TEM image showing formation of the lath-like O phase within a T matrix
(after aging for 40 min at 650°C), (b) a [-ll0]r stereographic projection showing the orientation of the
T/O interface, (c) SAD patterns generated from the (T + O) two-phase region in (a), Z ~ [II0]T II[001]o.
3044
HSIUNG et al.:
RAPIDLY SOLIDFIED Ti-24AI-11Nb
e
Ii
l.s °
0
"
...............
G 2
1
,
lii
.
e ............. :~
•
1120ae
(c)
,,
S
•
•
gO
•
Q
"
•
"
•
•
•
•
•
•
•
•
•
•
•
•
t
•
c) 6)
•
t
t
•
•
6
•
•
,l~f
'
m
•
•
•
•
•
"
qD
•
•
•
•
•
•
•
qt,
o
t
•
•
•
('
•
•
•
•
•
0
o
.-\
/-.
60 °
Fig. 13. (a) Dark-field TEM image showing formation of the plate-like ct2 phase within an O matrix
(after aging 4 h at 650°C), Z ~ [001]o II[0001L2, (b) SAD patterns generated from the (O + ~t2) two-phase
region in (a), (c) MD generated from the [001]o and [0001]~2 zones.
o f retained O phase was observed within the ~t2 matrix
even after prolonged aging (Fig. 15).
The above results confirm that the isothermal
fl---, ct~ transition occurs through a complex phase
evolution sequence fl ~ B2 ~ T---* O ~ ~t2 at temperatures below 900°C. The phase evolution resulted
in a change of microstructure morphology as well as
hardness of the alloy.
HSIUNG et al.:
RAPIDLY SOLIDFIED Ti-24Al-11Nb
3045
Fig. 14. (a) Dark-field TEM image showing a coalesced ct2 grain observed in a foil sample after aging for
24 h at 650°C, (b) CBED patterns generated from the [0001] zones of an ~: grain showing a 6 mm
symmetry.
4. DISCUSSION
We have demonstrated that the isothermal fl--, ~2
phase transition in a rapidly solidified Ti-24Al-I 1Nb
occurs through a complex reaction sequence at
temperatures below 900°C. Initially this proceeds
by ordering transitions (fl--~ B2--, T) followed by
T--* O--* ct2 transitions. We next propose mechanisms for these isothermal phase transformations.
4.1. The fl --~ B 2 --~ T ordering transitions
The fl ~ B2 ~ T ordering transitions are illustrated
in Fig. 16 based on the stoichiometric composition of
TisA12Nb (Ti-25 at.%Al-12.5 at.%A1), i.e. 12.5 at.%
of Ti in the TiaA1 compound was substituted by Nb.
Note that the alloy composition of Ti-24 at.%Al-11
at.%Nb can still be maintained with existence of
vacancies in the Ti 5A12Nb compound. In the fl lattice,
the Ti, A1 and Nb atoms were uniformly and randomly distributed at both (0, 0, 0) and t!
'~2' !2~ !~
2 J lattice
sites. The site occupancy is 62.5% for Ti, 25% for A1
and 12.5% for Nb. The fl ~ B2 ordering leaves the
(0, 0, 0) site of the B2 lattice Al-rich and the (1, ~,i)1
1
site Ti-rich. The probability of finding each atom type
at the (0, 0, 0) site is 43.75% for Ti, 6.25% for Nb
and 50% for Al; at the ~-2'
t! !2 , !~
2,' site is 87.5% for Ti
3046
HSIUNG et al.:
RAPIDLY SOLIDFIED Ti-24AI-11Nb
Fig. 15. Dark-field TEM image showing an observation of the retained O phase within an ~t2 matrix
(after aging for 72 h at 800°C), Z ~ [I2T6]=2.
and 12.5% for Nb. The B2---~ T ordering leaves the
1_ _1 1"I
4,4,4.s',
1 3 3, I
[ 3 ! 13
~.4,4,4/,
3 3 3
i-I 3 1~t
'..4,4,4/,
l
(_3 3 1"~ (1 1 33
'.4,4,4",
k4,4,4/,
( 1 _1 1~
[ 3 1 3"I
".],4,.4J,
4, 4, 4,, (~, i, ~), (0, 0, 5) and ~2,2, 2, sites of the T
lattice occupied by Ti; the (½, 0, 0) and (0, ½, 0) sites
occupied by Nb; and the (0, 0, 0), (~,
1 ~,
1 0), (½,0,91
and (0,½,½) sites occupied by Al. Notice that the
dimensions of the T lattice are doubled by the
chemical ordering. In addition, the atomic arrangement in the T lattice lead to a tetragonal distortion
'
[~
it ,,-,,~J
"1i 67.5%
Nb 12.5%
AI 25.0%
>
( + 2 % ) along the [001]r direction as a consequence of
the atomic size difference effect (the atomic radii (nm)
ofTi, AI and Nb are 0.147, 0.143 and 0.143 respectively [16], i.e. the atomic radius of Ti is 2% greater
than that of AI or Nb).
4.2. The T ~
0 transition
The lattice correspondence between the T and
O phases is shown in Fig. 17(a): [001]X---~[100]O,
T
B2
[*" a'-* I
@ Xi 87.5%
Nb 12.5%
•
"11 43.75%
Nb 6.25%
AI 50.00%
l-
a
LI
@ Ti
•
AI
•
Nb
Fig. 16. Schematic illustration of the 1 / ~ r2---* T ordering transitions.
c/a = 1.02
HSIUNG et al.:
RAPIDLY SOLIDFIED Ti-24AI-I 1Nb
[ilOl,r
[ololo
°
1110]T
(b)
3047
[o1o1o
L~loO{]op10IT
(a)
@o
[oo.q.r
[100lo
(001) T I (100) 0
CT
(c)
/
tlOOl T
IllO-lz
(OOi) o I (110) x
['i'iol.r
1oo11o
1
1
0~
[00110
---
c°
a'r I1
[1lilt
O = 63.4,°
B
[]{qo~
Fig. 17. (a) Schematic illustration of (a) lattice correspondence between the T and O phases, (b) lattice
deformation and (c) atomic shuffling during the T ---*O transition. Only one eighth of the T and O lattice
was drawn in (c).
[I 10]T "--*[010]o and [I i0]T---' [001]o. The structural
relationships between the T and O phases can be derived by assuming a shuffling of atoms on alternating
(I I0)T planes along [I 10]T accompanied by homogeneous lattice deformations along the [ 0 0 l I T , [I10]x and
[I I0]T directions. These are illustrated in Fig. 17(b,c).
The dimensions of the O lattice can be obtained by
contracting the T lattice by 8.8% along [001]T to create
[I00]o, expanding [I10] by 6.6% to create [010]o, and
expanding tli0IT by 2% to create [001]o. The atomic
shuffling makes the stacking sequence of the O phase
of 2H-type (ABAB stacking sequence).
The structural relationships between the T and O
phases including the T/O interface (habit) plane and
the orientation relationships between the T and O
phases can be successfully predicted on the basis of
a shape deformation mechanism [17]. An important
aspect of the T - - , O transition is the shape deformation (P) required to convert a T to an O lattice.
P is composed of a homogeneous lattice distortion B,
a lattice-invariant inhomogeneous shear S (corresponding to slip and twinning), and a rigid-body
rotation Q, i.e. P = iaSB [18]. The resultant shape
deformation must be consistent with the experimental
observation that the habit (invariant strain) plane is
essentially undistorted and unrotated. By measuring
the invariant plane strain (SB), the structural relationships between the T and O phases can be predicted.
From [17], the predicted T/O interface (habit) plane is
near to the {~23}T plane, and the predicted orientation
relationships are (I I0)T II(001)O and [II0]T ^ [010]o =
[001IT ^ [100]O = 4.3 °. These are in agreement with the
experimental results shown earlier.
4.3. The 0 --~ :t: transition
The O phase may be considered as a pseudohexagonal phase with a lattice distortion resulting
from a supersaturation of the Nb in the ~t2 lattice.
Note that the Nb saturation concentration in ct2 below
900°C is 9-10 at.% according to [7, 12], i.e. ~ 1.5 N b
atoms per ct2 unit cell (16 atoms/unit cell). To accommodate the excess Nb content, the supersaturated ~t2
lattice is no longer able to maintain its hexagonal
symmetry and becomes an O lattice. We illustrate this
in Fig. 18(a). Notice that the Nb atoms occupy the
(0,11
31~) sites of the O lattice. The dimen~, ~) and (0, ~,
sions of the O lattice are a ~ 0.605 nm, b ~ 0.98 nm
and c ~ 0.47 nm. The ~2 lattice can be restored by
diffusion of the excess Nb away from the O lattice
accompanied by a homogeneous lattice distortion
and an atomic shuffling on every third (100)o. Since
long range diffusion of the excess Nb away from the
O/0~2 interface is required for the O ~ a2 transition to
progress, this results in a sluggish O---, a 2 reaction at
lower temperatures (for instance at 450°C or below).
In a similar way to the T --~ O transition, the shape
deformation mechanism can also be applied to predict the structural relationships between the O and =2
phases. The lattice correspondence between the O and
a 2 phases is shown in Fig. 18(b): [l-f0]o--* [10T0]=2,
[110]o---* [01-1-0]=:and [001]o--~ [0001]=:. The a 2 lattice
was drawn here by assuming that Nb does not occupy
3048
HSIUNG ef al.: RAPIDLY SOLIDFIED Ti-24Al-11Nb
Wlo
lccc11,2
0
Ti
l00011 a
4
I
l
4
I
A’
0
Ti (Nb)
.
A’
*
Fig. 18. Schematic illustration of (a) lattice correspondence between the 0 and a2 phase, (b) unit cell of
the t12phase.
a specific lattice site. By redefinition of the tlz lattice
on the basis of an orthorhombic cell [shaded area in
Fig. 18(b)], the dimensions of the t~zlattice becomes
a ~0.58 nm, b z 1.Onm and c x0.465 nm. Thus, the
tiz lattice can be obtained by contradicting the 0
lattice by 4.1% along [1OO]oto create [loo],, , expanding [OlO], by 2% to create [OlO],,, and contradicting
[OOl], by 1.1% to create [OOl],,. Referring to the
directions [lOO],, [OlO], and [OO1]oas X-,y-, z-axes,
these can be expressed as follows
0
0.020
0
0
0
0.011
The principal strains, denoted by vii, of the homogeneous distortion B for the O+ tlz transformation
are given by
[
qz2 O=
o
0
0
0
0
133
0
=I,59
0.956
0
0
=
0
1.017 0 .
(3)
01
[ 0
I
where v is Poisson ratio (N 0.3). A theoretical prediction of the habit plane and orientation relationships
between the 0 and CI~phases can be made in the
manner shown in Fig. 19. The effect of the invariant
plain strain B’ on a spherical 0 crystal, viewed along
the z-axis, is illustrated. The spherical crystal was
deformed into an ellipsoid due to the strain B’.
However, the planes OQ’ and OP’ are not distorted
by the strain, but were rotated from their initial
position OQ and OP. TO produce a unrotated as well
as a undistorted habit plane, a rotation about the
z-axis has to be added to B’ to return one of these
planes to the initial position (OQ’ to OQ for
instance). Let Q be the point (x, y). The coordinates
of Q’(x’, y’) can be determined by
’
(1)
1
0
VII
00 1+t,i
I[
-0.041
B=o
ing the lattice distortion B with a uniaxial strain
(tension 6) along the z-axis. This gives an invariant
plain strain B
1.020
0
0
1+ $2
0
0
1+ 633
0
0 ] .
1.011
(2)
The condition for there to be an invariant strain
(undistorted) plane is that one of the principal
strains be unity, and the other two be greater or
less than unity. Since qJ3= 1.011 is close to unity,
the amount of slip or twinning (corresponding to
shear deformation S) is quite small. Alternatively,
the invariant plain strain can be obtained by combin-
[;r]=[“%”
1.;17
;I[!
p
;I.
(4)
That is, x’ = 0.956x,y’ = 1.017~. Since OQ = OQ’,
x2 +y2 = (0.956~)’ + (1.017)‘. Thus, LQOY = tan-’
HSIUNG et al.: RAPIDLY SOLIDFIED Ti-24Al-11Nb
[01010
Y
[mO]o
Fig. 19. An ellipsoid developed from a sphere of the O
crystal by the invariant plane strain B'. The z-axis is parallel
to [001]o.
3049
~t2 phases were determined to be (IT0)T II (001)o, and
[II0]T A [010]o=[001]r A [100]o =4.3°; (001)o(0001), 2,
[100]o A [11~0]=2 = [010]o ^ [II00]= 2 = 1.6 °.
3. Both the T--~ O and O ~ ~t2 transitions can be
explained by a shape deformation mechanism involving a homogeneous lattice distortion and rigid-body
rotation. F o r the O ~ ct2 transition, besides the shape
deformation, diffusion of the excess N b away from
the O/ct2 interface also is needed. This causes a
sluggish O ~ ct2 reaction at lower temperatures.
Acknowledgements--This research was co-sponsored by
the Defense Advanced Research Projects Agency and the
National Aeronautics and Space Administration through
Contract Number NAGW-1692, and GE Aircraft Engines
through Contract Number MS-GE-5222-92. The authors
wish to thank E. S. Russell and D. Backman of GE Aircraft
Engines, Lynn, Mass. for providing the material used for
helpful discussions during the course of tile investigation.
REFERENCES
( x / y ) = 32.2 °. The Ot~t 2 interface (habit) plane therefore makes an angle 32.2 ° with the (100)o plane, this
is only 0.5 ° away from the ( l l 0 ) o plane ((110)o ^
(100)o = 31.7°). / Q ' O Y = t a n - l ( x ' / y ') = 30.6 °, thus
the rotation angle 0 = 1.6 °. This rotation makes
(110)o N(10-i'0)~:. The predicted orientation relationships between the O and ~t2 phases are (001)o II (0001)=2.
and [100]o ^ [11~0]=2 = [010]o A []'100]=2 = 1.6 °. These
are in agreement with the experimental results shown
earlier.
Finally, we make a comment here on the retained
O phase found after prolonged aging. Since diffusion
of the excess N b content away from the O/ct2 interface
is needed for the O ~ ct2 transition, the N b content
in the retained O phase should be greater than that
in the primary O phase. Note that the saturation N b
content of the O phase has been reported to be
~ 2 5 at.%, i.e. Ti2A1Nb [10]. That is, a decomposition reaction O---~ ct2 + retained O(Ti2AINb) might
have occurred during a final stage of the fl ~ ~2
isothermal transformations.
5. CONCLUSIONS
The isothermal phase transformations of a
T i - 2 4 A l - l l N b alloy at temperatures below 900°C
have been studied using X-ray diffraction, electron
microscopy (SEM and T E M ) and electron diffraction
techniques. We find that:
1. The isothermal fl ~ ~t2 transition took place
through a complex phase evolution sequence, first by
ordering transitions fl ( b . c . c . ) ~ B2 (ordered b.c.c.)
--~ T (ordered tetragonal), and then by T ~ O (ordered
orthorhombic)---, ct2 polymorphic transitions.
2. Plate-like O and ~t2 phases were found at early
stages of both the T ~ O and O ~ g2 transitions. The
T / O interface (habit) plane was determined to
be {~23}T, and the O/~t2 interface (habit) plane was
determined to be {110}o. Orientation relationships
between the T and O phases, and between the O and
AMM 40/I I - - P
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