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Transformations and structures in the uranium-rhenium system I. Metastable α′, α“ and γ0 structures
JOURNAL
OB NUCLEAR
21 (1967) 263-276.
MATERIALS
TRANSFORMATIONS
AND STRUCTURES I.
Metastable R.
for Atomic
Institute
Research
The compositional
J. JACKSON+
and W.
temperature
on producing
the sequence metsstable
the
crystal
effect
Finally, states
are
structures
of cooling
OI” and
presented
rate
were
tetwgonal
to
and
metastability
et
les phases stables
et metastables,
de refroidissement
entre
observes Es
and
innerhalb
wurden
mation
de structures
transformations phases
alternatives
de phase
de transition
les phases
ont
bildet
werden
zeugung
quadratique
y”. On presente
les arguments
d’expliquer
les structures
cristallines
1.
monokline
Work was performed Present
Address:
Dow
oti
ont
Bte
est discutee.
wurde
Der Einfluss der
die
Er-
st,udiert.
der Phasen-
Zustilnde
fiir die orthorhombische
wurden
Phase
a’, fur die
Phase LY”und fur die tetregonale
Phltse ~0.
Kristallstrukturen
und es wird weit’erhin
such den Einfluss der Abkiihlungtgeschwin-
digkeit und der Zusammensetzung zu erlllutern.
Abschliessend
tilt anderer
Systeme
auf diese-Strukturen
wird
von Uran-Rhenium-Legierungen
die
Metastabilittit
mit der Metastabili-
verglichen.
A literature survey revealed only three significant research studies dealing with U-Re alloys. Hatt 10) described the structure determination of the URes intermediate phase. Rough and Bauer 11) list results of tentative phase studies. Jackson et ~~1.12)presented a phase diagram and reported that /3 and y phases of uranium could be retained to room temperature by rhenium additions and rapid cooling. The occurrence of metastable transition states at certain cooling rates and compositions was also reported. The work reported here had the following objectives :
of the U.S. Atomic
Company,
auf
wurde
Metastabile
ge-
der Abkiihlungs-
die A.ufeinanderfolge
zu erkl&ren,
les
Phasen
Temperatur
Strukturen
versucht,
a” et
observees,
und
bestimmt,
und metastabile
und Gefiigebilder
permettant
Chemical
systemes
des
Btaient
in the Ames Laborat*ory
The
alliages
Les
Introduction
t
la
les
die beobachteten
The presence of retained high temperature phases and metastable transition phases in binary uranium systems in which the y miscibility exceeds a few atom percent has been well documented. Metastable transition phases of the martensitic 01 types have been reported in the U-MO 1, s), U-Nb 3, a), U-Ru 5), U-Ti 6) and U-Zr 7) systems. Metastable transition phases based on the cubic y structures have been reported in the U-MO s), U-Nb 314) and U-Ti 9) systems. Because of the similarity in the nature of solute atoms, similar metastable structures can be expected in the U-Re system. *
Enfln,
dsns
Es wird versucht,
sur la for-
et la succession
konnen.
bestimmt.
de la vitesse
observees
d’autres
de refroidisse-
celles-ci.
Konzentrationsgrenzen
bestimmter
beobachtet
~11’)monoclinique
sur
me&stables
stabile
geschwindigkeit
se forment
et,4 determinees.
m&astables
orthorhombiques
de la vitesse
mktastabilite et
die
derer
Ausserdem
l’effet
la
des et&s
the
in uranium-
et de la temperature
et l’effet
uranium-rhenium
thereon.
erure lesquelles
Ames, Iowa, USA
University,
la composition
The
is discussed.
de composition
de
relation
wechsel Los limites
State
and
alloys and ot,her systems in which metastable
are observed
Iowa
ys
explain
composition
SYSTEM
1966
ment
the
and microstructures
the relation between
rhenium
structures,
observed
(x’, monoclinic
Arguments
observed
states
CO., AMSTERDAM
LARSEN
microstructures
rate and
of phase changes were determined. transition
orthorhombic phases.
alternative
L.
25 July
which the stable and
the effect of cooling
PUBLISHING
IN THE URANIUM-RHENIUM
of Metallurgy,
and Department
limits within
phases form,
NORTH-HOLLAND
01’, OI” and yo structures *
Received
metasttlble
0
Rocky 263
Flats
Energy
Commission.
Division,
Golden,
Contribution Colorado.
no. 1921.
264
R. J. JACKSON A9n
1. Determination
of the compositional
limits
within which the stable and metastable phases are produced, the cooling rate being constant. 2. Investigation
of the effect of different cooling
rates in producing 3. Determination of changes.
al~rnative structures. the sequence of phase
2.
2.1.
under
noneq~ibrium
approximately
cooling
con-
Experimental ALLOY PREPARATION,HEAT TREATINGAND QUENCHINo
The source of the constituent metals and their chemical analyses were given in an earlier report 12). All the alloys were prepared by melting under a zirconium-gettered helium atmosphere in a non-consumable tungsten electrode arc furnace. A master alloy of URes was crushed and used as the source of rhenium. The use of a lower melting rhenium source of small particle size shortened the melting time considerably. Spectrographic analysis indicated no contamination from the copper hearth or the tungsten electrode. The weight changes on melting were such that in all cases the apparent uncertainty was less than 0.1 at o/0 Re. The intended composition was thus assumed to be correct. The as-cast alloys were homogenized under vacuum in a tantalum container for 4 days at 825” C and then cooled to room temperature at 25” C/h. Specimens approximately &” x &” x &” (1.6 x 4.8 x 2.4 mm) were machined from the homogenized buttons. The small size tended to reduce thermal gradients and property variations within samples. All the experiments which required heating were conducted at pressures less than 1 x 10-s Torr using a tantalum holder or shield where necessary to avoid contamination. Water quenching refers to a cooling rate of
1000” C/set.
The
procedure
is
described elsewhere Is), Vacuum quenching refers to removal of the furnace from around an evacuated specimen.
fused silica tube containing the Cooling rates of approximately
10” C/sea were achieved 2.2.
To achieve these objectives, a study was made of the time-temperature-composition dependency of crystal structure and microstruct~e ditions.
W. L. LARSEY
on
in this manner.
OPTICAL AND X-RAY
METALLOGRAPHY
Specimens were subjected to one minute each 320, 400, 500 and 600 grit SiC paper.
Polishing consisted of 5 min each on rotary laps of 600 grit Sic on 10 ounce canvas duck, oil lapped 6 pm diamond on nylon, and 1 pm diamond on a medium nap rayon. The worked surface was then removed by electropolishing at 40 V dc for 15 set in a chromic-acetic acid bath composed of 1 part CrOs, 1 part H20 and 4 parts CHaCOOH (glacial). The alloy was then examined under polarized light. To verify the structures revealed by polarized light, it was desirable to etch the polished surfaces. A sometimes satisfactory electrolytic etch was obtained by lowering the potential across the chromic-acetic bath to 25 V. A chemical etchant consisting of a I:1 mixture of nitric and acetic acid was found to work on all the alloys. ~tehing times for the HNOa :HAc mixture varied proportionally from 25 min for the 1 at y. Re alloy to 10 set for the 14 at y. Re alloy. Vickers hardness numbers were obtained using a 10 kg load and averaging at least four readings. The grain size of each specimen was estimated using the ASTM comparative method. In instances where the grain size was larger than the standards, it is reported as the average diameter in inches. The DI phase was often present as platelets and its grain size is reported as plate width in millimeters. The average grain size estimated by visual integration is reported for duplex structures. X-ray diffractometer patterns taken immediately after each heat treatment served as a quick means of phase identification and as a check for isothermal transformations at room temperature. After quenching and prior to
TRANSFORMATIONS
examination,
the
AND
sample
STRUCTURES
was
IN
mechanically
polished and electropolished. The latter step was necessary to avoid line broadening, especially for the sofikr alloys. Filtered CuKa radiation was used at a chart speed of 1” (Ze)~min. After quenching, specimens for Debye-Scherrer photographs were ground to a 0.5 mm dia. cylinder and then electrolytically thinned to a needle of 0.15 mm dia. Filtered CuKa radiation was used with a 114.6 mm dia. camera. Lattice parameters were determined by leastsquares extrapolation against the Nelson-Riley function using an IBM-7074 computer 1% 14). 3. 3.1.
THE
VACUUM
a
The core experiments were those in which series of alloys were water quenched, vacuum quenched and slow cooled from the y and @ phase fields. More specifically, they were: (1) a series of alloys water quenched from 975” and 850’ C after Q and 4 h at temperature, respectively ; (2) a series of alloys vacuum quenched from 975, 850, 700 and 662’ C after $, 4, 12 and 6 h at temperature, respectively; and (3) a series of alloys cooled at 25” C/h from 825” C after 4 days at temperature. These alloys were examined for crystal structure, microstructure, grain size and hardness. The crystal structure of the uranium phase present is shown s~~ematieally in fig. 1. The results of the grain size and hardness measurements are shown in figs. 2 and 3, respectively. The significance of
QUENCkED
SYSTEM,
975O
C
a*‘
2
4
VkUI)M
6
I
8
‘QUfNCiiEb
850dC
R
I
I
I IO
t
1 12
’
’
’
’
B 0
4
6
8
QUENCHED
IO
2
vAw&
I
YI ’
6bOg’C
12 I
’
a 0
265 I
I_ 14 ’
Yp
2 WATER
I
,
’
B-
0
Results and discussion CORE EXPERIMENTS
URANIUM-RHENIUM
4
6
, IO
a
‘QlJ$NCkEo’
700°1
‘
14 I
I
Y1 I 12
‘ 14
12
14
.
C
a” 0
L
I 2
1
I I 4 ATOM
I , 6 8 PERCENT
IO RHENIUM
Fig. 1. Effect of cooling rate and quenching temperature on the crystal struct#ureproduced in the uranium-rhenium system.
IO
x
WATER
+
VACUUM
l
WATER
OUENCHED
975
QUENCHED QUENCHED
975 850
‘C “C OC
0
VACUUM
QUENCHED
950
‘C
b
VACUUM
QUENCHED
662
‘C
v
COOLED
25
+C/tiR
FROM
625°C
IO
12
these data is discussed later along with results from supplementary experiments. 3.2.
METASTABLE
y AND yo STRUCTURES
Inspection of fig. 1 shows that 8.5 at y0 Re was sufficient to cause capture of the y phase on vacuum quenching from either 850 or 975” C. Similarly, about 9.5 at o/0 Re was sufficient to retain the y phase on water quenching from 850” C. X-ray examination of the retained y alloys showed the structure to be body-centered cubic. Lattice parameters of 3.452, 3.448 and 3.448 A were obtained for the 10, 12 and 14 at o/0 Re
0
2
4
6
ATOM
Fig.
2.
6
PERCENT
14
6
RHENIUM
Comparison of grain size meas~emen~,s after various
thermal
treatments.
alloys, respectively, which were water quenched from 850’ C. Klepfer and Chiotti 15) report a value of ao= 3.538 i$ for pure uranium at 800” C. Correcting this value to room temperature using the linear coefficient of thermal
266
R.
x
700
_I
WATER
PUENCHED
+
VACUUM
.
WATER
J.
975
OUENCHED
975
QUENCHED
650
JACKSON
OC OC
VACUUM
OUENCHED
850
OC
o
VACUUM
OUENCHED
662
OC
o
COOLED
25
FROM
W.
L.
LAESEN
His proposed structure is based on a block of 2x2 x 1 y cells in which the solute atoms arc restricted to the edge-center sites. The observed line positions and intensities for yb0 agree very
oc
o
OC/HR.
AND
825
well with the structure proposed by Tangri for a cell with ao= 6.92 A and CO= 3.38 A. Hence,
“C
600
z
it is proposed morphous
that
the yb0 structure
with the tetragonal
ordered
is isophase
observed by Tangri in the U-MO system. A comparison of the y and ybs microstructures is shown in fig. 4. In the 9 at o/o Re alloy vacuum quenched from 775” C coarse bands were observed in addition to the fine interwoven bands. The bands were similar to those observed next to hardness indentions in y alloys and are believed to be deformation bands. II
I
0
2
I
II 4 ATOM
Fig. 3.
I
I
6 PERCENT
I
I
8
I IO
I
I 12
I
IA I4
RHENIUM
Comparison of hardness values after various thermal
treatments.
expansion determined by Klepfer and Chiotti 15) of 22.5 x lo-6/“C yielded an CO at 23” C of 3.476 A. Hence, the addition of rhenium atoms causes a contraction of the bee y cell. This contraction would be expected on the basis of size considerations. Additional quenches were made in an effort to obtain a y-like structure showing definite tetragonality as reported by Williams 16). Compositions of interest were vacuum quenched from 700 and 775’ C after 12 and 8 h at temperature, respectively; and water quenched from 775” C after 8 h at temperature. Vacuum quenching the 9 at y0 Re alloy from 700 and 775” C, and the 10 at o/o Re alloy from 700” C did yield specimens whose X-ray diffraction patterns showed line splitting and weak additional reflections. This structure is hereafter referred to as yba where the subscript indicates the metallographically identified bands (crosshatch structure) and the superscript the corresponding alteration in crystal structure. Tangri s) determined the structure of a tetragonal ordered phase observed under somewhat similar conditions in the U-MO system.
Fig. 4. Photomicrographs contrasting the y and y C (a) The 10 at y. Re alloy water microstructures: quenched from 850” C showing retained y microstructure. Polarized illumination.
x
alloy
150. (b) The 9 at y0 Re
vacuum quenched from 700” C showing tho crosshatch y g microstructure. Polarized illumination. x 100.
TRANSFORMATIONS
3.3.
B~TASTABLE
AND
01' AND
STRUCTURES
n"
IN
THE
STRUCTURES
URANIUM-RHENIUM
MetalIographic
SYSTEM,
I
267
structures associated with the
Two metastable modifications of the normal orthorhombic a-uranium structure were ob-
a’ and OI’ phases are shown in fig. 5. The observed structures varied from granular to
served for rapidly cooled y alloys containing between 3.5 and 9 at yO Re. The a-phase
very fine crossed bands (crosshatch struct~e) Taomixtures of fine and coarse bands which often
modifications are designated (x’ and a”, the superscripts indicating a departure from the equilibrium structure as determined by X-ray diffraction techniques. The n’ modification is a result of a relative contraction of the b parameter with slight expansions along the a and c
appeared
axes of the normal a-uranium structure. The bn modi~cation occurs at higher rhenium content and is due to a relative contraction of the b parameter, together with a change from an orthorhombic a structure to a monoclinic structure.
Fig. 5.
Photomicrographs
showing the microstructural
acicular.
The type of microstructure
obtained appeared to be directly associated with cooling rate and composition and not with the crystallographic phase present, i.e., orthorhombic OI’ or monoclinic &“. Vacuum which at
quenching yielded microstructures low rhenium compositions were
granular and which with increasing rhenium content developed a crosshatch structure (eompare figs. 5a and 5d). Water quenching yielded much coarser structures which at all compositions were considered banded. At the higher
types associated wit,h the 0~’ and 01” phases:
(a) The
6 at e/0 Re alloy vacuum quenched from 85W C showing crosshatch and granular appearing a” microstructure. x 100. (b) The 6 at y/o Re alloy water quenched from 850” C showing n’ microstructure consisting of coarse and fine bands. x 100. (c) The 7 at o/0 Re alloy water quenched from 850’ C showing OLD microst,ructuro consisting of coarse and fine bands. x 150. (d) The 8 a.t o/0 Re alloy vacuum quenched from 975” C showing crosshatch a” microst,rueture. Note cruciforms around carbide particles. x 160. All polarized illuminr&ion.
rhenium
contents
determine termed
it
whether
banded
was
the
often
diffioult~ to
structure
or acicular
should
bc
(figs. ;ib and 5~).
Lattice parameters (10, 7,0, COand angle y of the 0~‘ and CX”structures were determined using specimens water quenched from S50° C. The results are represented by solid lines in fig. 6. The main features are: The values of cl0 and CO increase with increasing rhenium content.
slightly
The value of bo decreases markedly increasing rhenium content.
with
The angle y gradually increases from 90.0” at 6 at yi;, Re to 92.2’ at 9 at 7; Re. The effect of introducing rhenium atoms into the uranium lattice may be explained in the following manner. The unequal interatomic distances in a uranium give an ellipsoidal shape to each uranium atom rather than a spherical shape. Assuming that each atom behaves as a hard particle, the rhenium atoms may be represented as hard spheres of radius 1.380 !J when in a coordination of twelve, and uranium atoms may be represented as hard ellipsoids *
I
s
.
EXPERIMENTAL
o
THEORETICAL
I
,
,
I
t
1
h
z
4.98
.. _____._ I-w-m
2 4.94 2.90 z ;
--
.
.
* ___ ..---
2.86
-----
.
0
-
5
6
l
I-
0
I
2
3 ATOM
Fig.
4
PERCENT
7
8
9
IO
RHENIUM
6. La&ice parameters of the ,x’ and 1” structures developed on wlttcr quenching from 880” C.
having and
the dimensions 1.660
K
along
of
1.3S1, l.426,
principal
1.640
inter-atomic
directions. The substitution of rhenium spheres into t,he lattice of ellipsoids results in the caleulatPed changes shown as open circles in fig. 6. The agreement between experimental and calculated lattice parameters is very good considering that the electronic interact,ion n-as ignored for the purpose of this rough calculation. 3.4.
DIRECT
y +x"
(A') TRANSFORMATION
hwez 17)has shown that rates of cooling up to 10 000” C/see horn the y phase do not prevent the formatio~l, in order? of /j’ and oi in pure uranium. This sequence of phase changes observed by Duwez at ext,remely fast cooling rates is also observed by vacuum quenching U-Re alloys containing less than about 3.5 at ‘:/A Re. Some observations supporting this sequence of phase changes are:
1. For alloys containing less than 3.5 at y0 Re which were LY at room temperat,ure, two thermal arrests were detected on cooling. 9 -I Irregular cy grains, similar to those obtained on quenching pure uranium, were observed on vacuum quenching the 1 at “/o Re alloy. 3. Duplicate alloy specimens! with the same thermal and mechanical histories, cont,aining 1 at :I0 Re were vacuum quenched from within the B and y fields, respectively. The ~~licrostructures as well as the hardness were nearly identical in both the p and y quenched specimens. Thus, it would seem that under these conditions the formation of in~rmediate /3 is not suppressed. It is shown in II. The ,!JFE B transformation, t,hat as t’he rhenium content increases the b --f ti transformation temperature is lowered and the stability of the ,I3phase increases with increasing rhenium cont,ent. On either water quenching or vacuum quenching from the y region Dhe 3 at 76 Re alloy was ,8 while the 4 at o/0 Re alloy was z‘. If the p phase had formed in the 4 at y;b alloy, one would expect it to be present at room temperature. The absence of the 1 phase
TRANSFORMATIONS
AND
STRUCTURES
IN THE
suggests a direct y --f LX’transformation
at fast
cooling rates. This viewpoint
by the
following
is further
observations
supported
URANIUM-RHENIUM
SYSTEM,
I
269
:
A 3.8 at y0 Re alloy water quenched from the y region had an ix’ structure while one cooled at a slower rate (400’ C/h) had a f!l structure. The 4 at y0 Re alloy cooled at 25” C/h had retained fi present, while as mentioned above the rapidly converted
cooled
alloys
were completely
to a’.
The microstructure of the LX’or 0~”alloys all showed straight grain boundaries (fig. 7a) which were those of the original y grains. All the retained /l alloys showed very irregular grain boundaries (fig. 7b). Likewise, any o( formed from /3 had highly irregular grain boundaries (fig. 7~). Had the 01’ or OL’ formed via a double transformation, i.e., would be y + p + a’, the microstructure expected to be irregular. The as-cast 1, 2 and 3 at y0 Re alloys contained numerous small cracks and voids which formed during the y --f @ transformation. These imperfections were readily seen metallographically and led to low gravimetrically determined densities. For example, the 3 at o/o Re as-cast alloy was retained ,8 with a density of 18.0% g/cm3 compared with a theoretical room temperature @ density of 18.80 g/ems. For alloys containing more than 3.5 at o/o Re, the ascast buttons were free of cracks and had densities of approximately 19.0 g/cm3 which compares favorably with the expected room temperature density for the distorted 01 structures. Had any /3 formed during the rapid cooling of these compositions, cracks would be expected in the as-cast buttons. This is further evidence for the direct y + 01’ or y + an transformation. Thus, at cooling rates obtained by water or vacuum quenching, the postulation of a direct y --f cy’ transformation is necessary to explain the observed behavior.
Fig.
7.
direct
Some microstructural y -P (x’ kansformation:
evidence suggesting a (a) The
4 at
YJ, Re
alloy water quenched from X50 C showing the banded microstructure boundaries. quenched
of bhc LX’ st,ate amidst
prior y grain
x 100. (lo) Tho 4 at y0 Re alloy vacuum from
662” C showing
highly
irregular
/3
grains.
x 150. (c) The 1 at. O/o Re alloy vacuum quenched from 662” C showing highly irregular OL grains.
3.5,
x 150. All polarized illuminat,ion.
SEQUENCE
OF PHASE
CHANGES
In the preceding sections it has been shown that during the transformation of y to (Y, the metastable structures a’, 0~” or yb0 may occur.
270
E. J. JACKSON
The mode of formation
AND
of these structures, and
their relationship to cooling rates explained in the following manner.
may
be
It has been proposed by Tangri and Williams2) that increasing additions of molybdenum progressively stiffen the uranium lattice, thus making it more and more resistant to shear, and that the effect of increased cooling rate is to favor arises
the shear process.
from
increased
caused by faster cooling
This presumably
hydrostatic
pressures
rates and also from
increased transformation stresses due to faster rates of transformation, once shear has started. Tangri also proposed that an optimum value of the ratio of quenching stresses to lattice stiffening, which for convenience of discussion was called the “stress/stiffening ratio”, is required to generate any given martensitic phase. Hatt and Roberts 18) have shown from X-ray diffraction work on a Zr-50 at y0 U single crystal that the 01 structure can be generated from a bee y structure by a shear on (112} y planes in a
W.
L.
LARSEK
of any given rhenium given
content,
and thus of a
restrict the explain the amount of generation of the a” structure on water quenching the !I at o/o alloy from 850’ C which on vacuum quenching under identical conditions yielded
stiffness?
is to relatively shear. This would
a y structure.
It can be seen in fig. 1
that, in each case, the effect of rapid quenching is to produce a structure further removed from the parent bee phase than that obtained
with
slower cooling, indicating that severe quenching causes t’he transformation to go towards completion. This effect is well illustrated by the photomicrographs of a !I at o/o Re alloy vacuum quenched (fig. 8a) and water quenched (fig. Xb)
(111) y direction. Furthermore, they developed a crystallographic model which explains the generation of the 01’, OI”and yo structures in the uranium-molybdenum system by progressively restricting the amount of shear. On water quenching alloys containing 3.5 to 6.5 at o/o Re, the optimum stress/stiffening ratio is realized, thus producing the critical amount of shear required to generate the a’ structure. With increasing rhenium additions, the lattice
is steadily stiffened until with 7.0 at o/o Re the lattice is too stiff to permit the critical amount of shear. Only a limited shear is produced resulting in an angle y of 00.8’, thus yielding a monoclinic 0~”structure. Further additions of rhenium progressively restrict the amount of shear, thus causing angle y to steadily increase to a value of 92.2”, as observed in the 9 at $” alloy. In highly stiffened lattices, as for example in alloys containing 10, l:! and 14 at oh Re, no shear or only a very limited amount takes place resulting in the generation of the y or yo structures, respectively. The effect of slower cooling rate in a lattice
Fig. 8. Comparison of microstructures of the 9 at) % Re alloy water quenched and vacuum quenched from 975” C:
(a) Thr 9 at y0 Re alloy vacuum quenched
from 975” C showing clean equi-axed grains of y structure. x 150. (b) The 9 at o/0 Re alloy water quenched from 975’ C showing single-phase, marten&o-like N” struct,ure. x 100. Both polarized illumination.
TRANSFORMATIONS
975” C.
from
undergone
AND
The
STRUCTURES
rapidly
cooled
IN
alloy
a shear transformation,
has
while the
THE
URANIUM-RHENIUM
water quenched
SYSTEM,
microstructures
271
I
disproves
this
more slowly cooled alloy appears untransformed. Work by Hills et ~~1.19) on end-quenched
and indicates the softening is inherent in the transition structure. Another possibility is that as the rhenium content is increased, theitendency
U-MO rods showed similar structural variations.
for precipitation
The slowly cooled end of a 10 at y0 molybdenum
y + 01’, an reaction occurring at a lower temper-
rod so quenched showed y” whereas the structure
ature);
of the
stage
rapidly
quenched
pattern
showed
end was
LX”. An
7.5 at y0 molybdenum
01” structure X-ray
cooled
only
throughout
rod had an
its length
of the rapidly
but
quenched
a little line splitting,
endthe
region
indicating
that its monoclinic angle was very near 90”, i.e.? the structure was approaching 0~‘. 3.6.
METALLOGRAPHICOBSERVATIONS
Metallographic observations have been presented in earlier sections to substantiate certain arguments. In this section, metallographic observations not previously discussed are presented and their relationship ments is developed. 3.6.1.
to earlier argu-
Hardness
Fig. 3 contrasts the differences in hardness for the vacuum quenched, water quenched, and slowly cooled alloys. The hardness-composition curves for vacuum quenched and water quenched y alloys are quite similar. The general shape of the curve is explained by the following hypothesis. The initial OLhardening is a combined solid solution and precipitation hardening effect. The difference in hardness at 1 at o/o Re on water quenching and vacuum quenching is due to the higher inherent hardness of the /? solid solution which is retained on water quenching. The general increase in p hardness with composition is a combined solid solution (due to supersaturation) and precipitation hardening effect. The linear decrease in hardness of the 01’ and 01” structures with increasing rhenium content above 4 at o/o is puzzling. At first glance it might appear that the decrease is associated with the coarsening of the structure, i.e., the appearance of bands and large lenticular plates. However, comparison of vacuum quenched and
is lessened (possibly due to the
and at the lower compositions of precipitation
tectable
occurs
by normal X-ray
techniques.
which
an early is unde-
and metallographic
It is well known that precipitated
phases that are undetectable by normal X-ray and metallographic techniques can markedly affect hardness. In any event, the continuous and gradual hardness change in passing from y’, to a” to (Y’ lends credence to the argument that y“, a’ and a’ are intermediate transition states in the formation of LYfrom y. The hardness minimum of the y quenched alloys is always associated with the first appearance of the y structure. The increase in hardness at higher rhenium contents is due to the grain boundary network of URez phase. The higher hardness of the 4 and 5 at y. Re alloys vacuum quenched from the y region is due to the formation of a detectable second phase. Water quenching of these compositions from the y field resulted in apparent single phase structures. The difference in hardness at the minimum for y quenched alloys is due to the effect of cooling rate on structure produced, i.e., CL”or y. The similarity in hardness at higher rhenium contents of the /? quenched and slowly cooled y alloys is interesting. The uranium phase present was /l in one case and OLin the other, thus the hardness is probably governed by the UzRe phase distribution which was similar in both cases. 3.6.2.
Grain size
The results of the grain size measurements are shown graphically in fig. 2. For the 0~’ and DL”structures the grain size shown is that of the prior y grain boundaries. The dotted symbol for the series quenched from 662” C is also for the prior y grain size. It is seen that all heat treatments used resulted in similar y grain sizes ,that were quite large except at relatively high
272
R. J. JACKSON
AND
rhenium contents. The reason for the restricted grain growth at high rhenium contents is undoubtedly
the grain boundary
network
of
URes particles. An interesting alloys obtained
grain size trend for retained /3 on y quenching
is also readily
seen. The fi grain size of water quenched
alloys
W.
I,. LARSEN
with increasing
rhenium
content
developed
a
crosshatch structure. The reason for the difference
in appearance
of
and
LY’ and
pendency not
LY” microstructures on cooling
readily
explains
the
their
rate and composition
evident.
One
hypothesis
observed
microstructures
deis that
is as
decreased with increasing rhenium content whereas on vacuum quenching the /3 grain size
follows. The tetragonal ybo phase having a crosshatch
increased with increasing rhenium content, This
microstructure
effect is probably
atures low in the y region and this phase yields
associated
with the greater
degree of supercooling of the y + p reaction on water quenching and the resulting larger number
of nucleation
sites.
3.6.3. Hicrostructures
The y quenched
pure uranium
was charac-
forms on quenching from temper-
a granular structure after transforming
to 01on
aging. This behavior possibly explains the a’ and CY~microstructures seen in the vacuumquenched alloys. Presumably, on vacuum quenching the cooling rate is slow enough that the alloys order to form yb0 as an intermediate
terized by irregular grain boundaries and highly twinned grains. A similar microstructure was observed in the 1 at o/0 Re alloys vacuum quenched from 850 and 975" C. The 1 at y0 Re alloys water quenched from 850 or 975’ C were retained /3. The retained /3 phase at these compositions decomposed at room temperature by the isothermal growth of LY.The 2 and 3 at o/o Re alloys either vacuum quenched or water quenched from 850 or 975' C were
transition state. The characteristic crosshatch mi~~rostructure of the ybo phase is thus given to the resulting 0~‘ or 0~” phase. The 01’ or a” platelets would thus be too small to be resolved in the optical microscope due to the small size of the bands making up the yt,O crosshatch structure. The observed change from a crosshatch to a granular structure is due to the presence of a second phase which in the case of the 4 and 5 at 0/o alloys was detectable by
retained p. The retained ,8 at these compositions was quite stable and no decomposition at room temperature was detected. The grain boundaries in these alloys were often characterized by a
normal X-ray techniques. On water quenching from the y region, there is not sufficient time for ordering and the rr,s structures does not form. Also, the thermal gradient and quenching stresses developed on water quenching are more severe than during vacuum quenching. Hills et aZ.20) and Butcher and Hatt 21) have shown that retained y0 alloys in the U-MO system deform by twinning mechanisms which results in coarse-banded ye microstructures. Thus, on water quenching, coarse twin bands may form followed by transformations to a’ and ocw by the growth of lenticular plates. The size of the resulting a’ or ~l# platelets would be governed by the size of the twin bands, or if twinning was absent by the size of the parent y grain. Thus in fig. 5b, the large bands would be twin bands and the smaller bands lenticular plates of cx’. In fig. 5~3, twinning of the y phase was restricted resulting
sawtooth appearance. For water quenched
alloys containing
4 to 9
at o/0 Re, the Iw, temperature for the direct y + ac’ (ol”) transformation lies above that for the y -+ /3 transformation and only DC’or (x” is observed in the microstructure of these alloys. These alloys were distinguished by smooth y grain boundaries and coarse microstructures which at all compositions were banded and/or acicular in appearance. For vacuum quenched alloys containing 4 to 8 at o/o Re a similar direct y -+ 00 (a”) behavior was observed except for the 4 and 5 at o/0 alloys where a second phase was detectable. Vacuum quenching yielded microstructures which at low rhenium compositions were granular and which
TRANSFORMATIONS
in the formation followed
AND
STRUCTURES
of large lenticular plates of a”
at a lower temperature
by the neces-
sarily smaller lenticular plates of (Ye. This scheme with some modification interpretation microstructures.
IN THE
of all the observed For
example,
SYSTEM,
I
273
In the 4 and 5 at y. Re alloys the UsRe phase
was present
in the
form
of
platelets
(often parallel) which were nearly continuous. allows
01’ and 0~” the
URANIUM-RHENIUM
micro-
The 4 and 5 at y. alloys also showed a continuous
jl network
boundaries.
outlining
the
prior
In the 4 at y. alloy,
y grain
the matrix
structure of the 7 at y0 Re alloy water quenched
consisted
from 850” C shown in fig. 5c could be interpreted
Presumably
‘in the following MS temperature
due to the non-selectivity of the etchant. On cooling the hypoeutectoid alloys, uranium-rich ,!l nucleates along y grain boundaries. On further
manner. On cooling below the the large lenticular a” plates
(mainly at 45”) form a network in the supercooled y matrix. The resulting transformation and quenching stresses cause the “boxed-off” y to deform by twinning giving rise to the vertical bands (some horizontal). The vertical ,bands then transform to an and the resulting ,transformation stresses cause motion along the twinned surfaces. This results in deformation of the LX”plates (white) that were not part of L the original 0~” network. As the rhenium content increases, the M, temperature falls below room temperature. Thus, for vacuum-quenched alloys containing 9 at y0 Re or more and for water-quenched alloys containing 10 at o/o Re or more, the y phase is stabilized at room temperature. Fig. 4a shows the characteristic equi-axed microstructure. A word is in order concerning the micro‘structure of the series of alloys cooled at 25” C/h ‘after 4 days at 825’ C. The 1 at oh Re alloy was characterized by twinned and very irregular LXgrains which would be expected from the double transformation (y + ,9 + CL).The 2 at o/o Re alloy consisted of a matrix of ,!3and oc uranium and a discontinuous precipitate of UaRe. The micrograph of this alloy at low magnification showed what appeared to be small equi-axed grains of a in the coarse /l matrix. This a structure is similar in appearance to a “pseudo (x” structure reported by Hills et aZ.22) for U-MO alloys. They report that this phase forms in the high temperature o( region and presumably consists of a duplex lamellar structure of (x+ UaMo within the visible “pseudo 0~” grains. Both the 2 and 3 at y. Re alloys showed a discontinuous UaRe precipitate.
mainly
of p with
some
01 present.
the a-j3 boundaries were not visible
cooling, coring of the ,!l phase results giving rise to concentration gradients within p. Precipitation of UaRe will occur at the rhenium-rich /l interfaces. This retards lateral growth and gives rise to transverse growth of the p phase. The regular decrease in the amount of retained fi with increasing rhenium content is interesting. The 2 at o/o Re alloy was heavily ,3 while only a small amount was present in the 5 at oh Re alloy. Possibly, the increasing amounts of UaRe promote formation of 01 from /l by acting as nucleation sites. Another possibility is that two reactions are competing for y, i.e., the hypoeutectoid y -+ ,!? and the direct y + LY reactions. The latter possibility seems the more plausible. The 6 at y. Re alloy was free of retained /? and its microstructure was more acicular in appearance than those of lesser rhenium contents. As rhenium content increased further the acicular appearance lessened and was nearly absent in the 9 at o/o Re alloy. The acicular appearance was replaced by optically active irregularly shaped apparent grains. Electron micrographs of replicas from these irregular shaped apparent grains showed them to consist of a very fine duplex lamellar structure of 01i- UaRe. A complete series of photomicrographs from which the preceding arguments were developed is shown in a thesis by Jackson 23). 3.7.
RELATION
TO OTHER
SYSTEMS
Metastable transition phases are found in a number of alloy systems in which the bee structure is stable only at high temperature.
274
R.
The
tendency
phase
towards
formation
the
interest
a-uranium
JACKSON
metastable
is most
W.
like
LARSEN
Burgers’ shear mechanism 24) for the bee + hcp The
This group is of
is
L.
transformation.
during
in this investigation
structure
AND
transition
pronounced
the bee --f hcp transformation. particular
J.
since
hexagonal
nature
of
phases
observed
mation
is shown
the
metastable
transition
for the bee + hcp in table
1. At
transfor-
low
solute
metals except that successive layers of atoms in the basal plane are skewed back and forth
concentration the transformation occurs /3 (bee) --f 01’ (hcp) by Burgers’ shear mechanism. The
in the (010)
effect of increasing solute content is to stiffen the lattice making it more resistant to shear.
uranium
direction.
By comparing
the a-
cell with the orthohexagonal
cell for
a hcp structure, one obtains b/u= 2.06 and c/a = 1.73 for 01 uranium compared with b/a = 1.73 and c/a= 1.63 for a hcp structure.
This
The
solute
TABLE transition
Metastable
comwosition Very
concentrations
transition
transformation phase?
Met&stable phase* bee + ol uranium
bee --f hcp
low solute
a? (orthorhombic) (forms from /3) /I
(*’ (hcp)
concentration
(supersaturated
(tetragonal)
(w) aa’ (orthorhombic)
Low solute concentration
o(s’ (orthorhombic)
Low solute concentration Medium
solute
w (C32 hexagonal)
0zb”(monoclinic)
concentration High
solute
concentration
diffuse o (faulted p
(tetragonal
w)+
c/n >
yso (ordered tetragonal)
1)
or /?’ (solute rich) High solute
y
All found in Zr-Nb alloys 25). LY’, W, and diffuse o in Zr-U alloys Is, 26). LY’, COfound in Ti-V, Ti-Cr, Ti-Mn, Ti-Fe, Ti-Mo, systems. 0~’ found in Ti-Cu plus many
*
other systems.
All present in U-MO system 1.2, s). All present in U-Re All except
(bee or tetragonal c/u w 1)
concentration t
/I in U-Nb
system
phase
a metastable
1
concentration Low solute
a
phases observed during the bee + hcp and
bee -+ 01 uranium Relative
in
from the hcp structure.
further
At medium transition
phase forms via a mechanism that requires shorter atom movements. At higher solute concentrations the product may be diffuse w (the X-ray reflections are diffuse ; hence, the name diffuse W) accompanied by a supersaturated ,6’ or bet /l” phase. The diffuse CJ
addition of rhenium atoms to the B-uranium lattice has little effect on the c/a ratio while the b/a ratio decreases ; thus, making a-uranium more “hexagonal like”. It also should be pointed out that the shear mechanism proposed earlier for the U-Re system is merely an extension of
Metastable
results
removed
(this investigation).
system 3.4).
a’, aa’, a’s, yo in U-Ti system 6, 9). a’, OCR’and possibly CQ,” in U-Zr system 7).
Ti-Co
and Zr-Mo
TRANSFORMATIONS
structure
AND
presumably
STRUCTURES
arises from
IN
incomplete
TEE
URANIUM-RHENIUM
bears
a similar
relation
SYSTEM,
for
the
I
U-MO
275 and
solute
U-Nb systems, i.e., at about b= 5.74 8. The a’-&” and LY”-y boundaries in the U-Re system
concentrations the p phase is stabilized and aging is required for the appearance of the
occur at about 5.80 and 5.75 A, respectively, on water quenching from 850” C, thus sub-
metastable transition phases. As mentioned earlier, metastable
stantiating the above hypothesis. These size correlations are based on a limited
glide
due to lattice
heavily
phases
faulted
have
uranium formation transition
structure.
been
alloy
stiffening At
observed
systems.
A
resulting higher
in a
transition
in a number summary
of
of the
conditions for these metastable phases in several uranium-systems
is presented
in table
1.
Comparison of columns 2 and 3 of table 1 shows that the transition phases occurring during the bee -+ hcp and bee -+ a~ uranium transformation are quite similar. Inspection of column 3 of table 1 shows that the behavior of the U-Re system is not unlike those of other highly y-miscible ~anium systems. The greatest similarity lies between U-Re and U-MO alloys ; the difference being the absence of the acicular a’ microstructure at low rhenium contents. The sequence of phase changes for the two systems is in general the same. The U-Nb system is also quite similar except the @ phase is not encountered. Lattice parameters are available for the distorted a structures in the U-MO 122), U-Nb 39 4) and U-Ru 5) systems. These lattice parameters along with w’-CQ,” and ~‘-y composition boundaries on water quenching have been discussed by Anagnostidis et aZ.3) and Tangri and Chaudhuri 4). They point out that if one compares the lattice parame~rs for the three systems, a correlation with size of the solute atom becomes obvious. As the solute size increased the change in 6 parameter decreased as would be expected on the basis of substituting hard spheres for ellipsoidal uranium atom. The ab’-~” composition boundary occurred at about b= 5.80 ,k for the two systems in which a” formed. This finding indicates the monoclinic (XII phase corresponds to a given lattice distortion and that elements having a small atomic diameter favor the formation of the Olbnphase. Likewise, the &b”-y composition boundary
number of observations, of the known
but do correlate most
data. As others 3*4) point
out,
this could permit one to predict the extent of metastable phase regions on the basis of atomic size ; the elements of small size decreasing extent 4.
the
of the domains.
Summary
A study was made of the time-temperaturecomposition dependency of crystal structure and microstructure under non-equilibrium cooling conditions for uranium-rich alloys in the uranium-rhenium system. The y phase under certain conditions was found to transform to ar without passing through the intermediate ,3 structure. Accompanying this transformation was a number of metastable transition states. Two transition states were modifications of the orthorhombio ~-uranium structure. They were: 1. a supersaturated a structure showing a notable oontraction of the b parameter with slight expansions along the a and c axes, and 2. a supersaturated a structure that had distorted to a monoclinic cell in addition to the parameter changes noted above. At high solute contents
the y phase can be
rendered metastable to room temperature. By cooling certain compositions from low temperatures in the y region a metastable transition state (~0) based on the bee structure occurs. This was a tetragonal ordered phase based on a block of 2 x 2 x 1 y cells having a c/a< 0.5. The compositional limits within which the stable and metastable phases form, and the effect of cooling rate and temperature on producing alternative structures were determined. From this was deduced the sequence of phase changes. A crystallographic and microstructural model
276
R.
was developed
J.
JACKSON
to explain the observed
structures
and microstructures
of cooling
rate and composition
AND
crystal
and the effect thereon.
Finally, the relation between metastability uranium-rhenium alloys and other systems which
metastable
states
are
observed
in in was
W.
Common
13) M.
14) D.Hansen,
argonno
(USA)
(1961)
St,ructure of YNi, Ames,
Williams,
Iowa,
226
M.S.
USA,
Mat.
4,
K.
5,
(1965) 278 K. Tangri,
Colombie
5 (1962)
and
H.
109
29)
Tangri and D. Chaudhuri,
J. Nucl.
Mat.
Thesis (Iowa
1964)
15 (1965)
and C. Rao,
288
22)
Tangri,
Mem.
A. Harding,
AERE
Acta
Cryst.
Met.
58 (1961)
and C. Knight.
M/R
2673A
14 (1961)
‘9 11) F. Rough and A. Bauer, Constitutional diagrams of uranium and thorium alloys (Addison-Wesley,
Reading,
Mass.
1958) p. 60
2558/g
24 (1953)
152
Harwell (UK)
B.
Butcher and J. Heywood,
R. Hills,
B. Butcher
11 (1964)
Report
(1960)
Metals
3 (1961)
J. Less-
155
and B. Howlett,
J. Nucl.
149
B. Butcher and B. Hatt, R. Hills,
B. Howlett
Common
Metals
Jackson,
system, Ames,
Harwell
119
Phys.
J. Nucl. Mat. 11 (1964)
and B. Butcher,
5 (1963)
J. Less-
369
Structures
and
transformation
kinetics of phases in the uranium-rhenium
469
(1958)
J. Appl.
R. Hills,
R.
23)
305
Sci. Rev.
M. Waldron
Report
B. Hatt,
50 (1958)
Fulmer
163
J. Nucl.
9 D. Douglass, Trans. Am. Sot. Metals 53 (1961) 307 7) D. Douglass, L. Marsh and G. Manning, Trans. Sot. Metals
1
communication, (1963)
Common Mat.
15 21
D. Chaudhuri
)
Monti,
(UK)
and J. Roberts,
AERE/X/PR 19
M.
B. Hatt
9
private
Inst.
P. Duwez,
J. Nucl.
(UK)
443
Heaton.
I51 H. Klopfer and P. Chiott,i, Ames (USA) Report G.
(1961)
9
L.
6156
State Univ.,
Research
3) M. Anagnostidis,
9
and
ANL
and W. Larsen, J. Less-
5 (lY63)
IHC-X93 (1957)
‘1 J. Lehmann, J. Nucl. Mat. 4 (1961) 218 2, K. Tangri and G. Williams, J. Nucl. Mat. 4
K.
Metals
Mueller
Report
References
Sm.
LARSEN
R. Jackson, D. Williams
9
discussed.
Mat.
L.
Ph.D. Iowa,
24
.J. Bowles,
25
J.
)
US-A,
Trans.
Stiegler,
J. Nucl. 26
Dissertation
Mat.
J.
(Iowa
State
alloy Univ.,
1964) AIME
Houston 11 (1964)
B. Hatt, and J. Roberts,
191 (1951) and
M.
44 Picklesimer,
32 Acta Met. 8 (1960) 575
OB NUCLEAR
21 (1967) 263-276.
MATERIALS
TRANSFORMATIONS
AND STRUCTURES I.
Metastable R.
for Atomic
Institute
Research
The compositional
J. JACKSON+
and W.
temperature
on producing
the sequence metsstable
the
crystal
effect
Finally, states
are
structures
of cooling
OI” and
presented
rate
were
tetwgonal
to
and
metastability
et
les phases stables
et metastables,
de refroidissement
entre
observes Es
and
innerhalb
wurden
mation
de structures
transformations phases
alternatives
de phase
de transition
les phases
ont
bildet
werden
zeugung
quadratique
y”. On presente
les arguments
d’expliquer
les structures
cristallines
1.
monokline
Work was performed Present
Address:
Dow
oti
ont
Bte
est discutee.
wurde
Der Einfluss der
die
Er-
st,udiert.
der Phasen-
Zustilnde
fiir die orthorhombische
wurden
Phase
a’, fur die
Phase LY”und fur die tetregonale
Phltse ~0.
Kristallstrukturen
und es wird weit’erhin
such den Einfluss der Abkiihlungtgeschwin-
digkeit und der Zusammensetzung zu erlllutern.
Abschliessend
tilt anderer
Systeme
auf diese-Strukturen
wird
von Uran-Rhenium-Legierungen
die
Metastabilittit
mit der Metastabili-
verglichen.
A literature survey revealed only three significant research studies dealing with U-Re alloys. Hatt 10) described the structure determination of the URes intermediate phase. Rough and Bauer 11) list results of tentative phase studies. Jackson et ~~1.12)presented a phase diagram and reported that /3 and y phases of uranium could be retained to room temperature by rhenium additions and rapid cooling. The occurrence of metastable transition states at certain cooling rates and compositions was also reported. The work reported here had the following objectives :
of the U.S. Atomic
Company,
auf
wurde
Metastabile
ge-
der Abkiihlungs-
die A.ufeinanderfolge
zu erkl&ren,
les
Phasen
Temperatur
Strukturen
versucht,
a” et
observees,
und
bestimmt,
und metastabile
und Gefiigebilder
permettant
Chemical
systemes
des
Btaient
in the Ames Laborat*ory
The
alliages
Les
Introduction
t
la
les
die beobachteten
The presence of retained high temperature phases and metastable transition phases in binary uranium systems in which the y miscibility exceeds a few atom percent has been well documented. Metastable transition phases of the martensitic 01 types have been reported in the U-MO 1, s), U-Nb 3, a), U-Ru 5), U-Ti 6) and U-Zr 7) systems. Metastable transition phases based on the cubic y structures have been reported in the U-MO s), U-Nb 314) and U-Ti 9) systems. Because of the similarity in the nature of solute atoms, similar metastable structures can be expected in the U-Re system. *
Enfln,
dsns
Es wird versucht,
sur la for-
et la succession
konnen.
bestimmt.
de la vitesse
observees
d’autres
de refroidisse-
celles-ci.
Konzentrationsgrenzen
bestimmter
beobachtet
~11’)monoclinique
sur
me&stables
stabile
geschwindigkeit
se forment
et,4 determinees.
m&astables
orthorhombiques
de la vitesse
mktastabilite et
die
derer
Ausserdem
l’effet
la
des et&s
the
in uranium-
et de la temperature
et l’effet
uranium-rhenium
thereon.
erure lesquelles
Ames, Iowa, USA
University,
la composition
The
is discussed.
de composition
de
relation
wechsel Los limites
State
and
alloys and ot,her systems in which metastable
are observed
Iowa
ys
explain
composition
SYSTEM
1966
ment
the
and microstructures
the relation between
rhenium
structures,
observed
(x’, monoclinic
Arguments
observed
states
CO., AMSTERDAM
LARSEN
microstructures
rate and
of phase changes were determined. transition
orthorhombic phases.
alternative
L.
25 July
which the stable and
the effect of cooling
PUBLISHING
IN THE URANIUM-RHENIUM
of Metallurgy,
and Department
limits within
phases form,
NORTH-HOLLAND
01’, OI” and yo structures *
Received
metasttlble
0
Rocky 263
Flats
Energy
Commission.
Division,
Golden,
Contribution Colorado.
no. 1921.
264
R. J. JACKSON A9n
1. Determination
of the compositional
limits
within which the stable and metastable phases are produced, the cooling rate being constant. 2. Investigation
of the effect of different cooling
rates in producing 3. Determination of changes.
al~rnative structures. the sequence of phase
2.
2.1.
under
noneq~ibrium
approximately
cooling
con-
Experimental ALLOY PREPARATION,HEAT TREATINGAND QUENCHINo
The source of the constituent metals and their chemical analyses were given in an earlier report 12). All the alloys were prepared by melting under a zirconium-gettered helium atmosphere in a non-consumable tungsten electrode arc furnace. A master alloy of URes was crushed and used as the source of rhenium. The use of a lower melting rhenium source of small particle size shortened the melting time considerably. Spectrographic analysis indicated no contamination from the copper hearth or the tungsten electrode. The weight changes on melting were such that in all cases the apparent uncertainty was less than 0.1 at o/0 Re. The intended composition was thus assumed to be correct. The as-cast alloys were homogenized under vacuum in a tantalum container for 4 days at 825” C and then cooled to room temperature at 25” C/h. Specimens approximately &” x &” x &” (1.6 x 4.8 x 2.4 mm) were machined from the homogenized buttons. The small size tended to reduce thermal gradients and property variations within samples. All the experiments which required heating were conducted at pressures less than 1 x 10-s Torr using a tantalum holder or shield where necessary to avoid contamination. Water quenching refers to a cooling rate of
1000” C/set.
The
procedure
is
described elsewhere Is), Vacuum quenching refers to removal of the furnace from around an evacuated specimen.
fused silica tube containing the Cooling rates of approximately
10” C/sea were achieved 2.2.
To achieve these objectives, a study was made of the time-temperature-composition dependency of crystal structure and microstruct~e ditions.
W. L. LARSEY
on
in this manner.
OPTICAL AND X-RAY
METALLOGRAPHY
Specimens were subjected to one minute each 320, 400, 500 and 600 grit SiC paper.
Polishing consisted of 5 min each on rotary laps of 600 grit Sic on 10 ounce canvas duck, oil lapped 6 pm diamond on nylon, and 1 pm diamond on a medium nap rayon. The worked surface was then removed by electropolishing at 40 V dc for 15 set in a chromic-acetic acid bath composed of 1 part CrOs, 1 part H20 and 4 parts CHaCOOH (glacial). The alloy was then examined under polarized light. To verify the structures revealed by polarized light, it was desirable to etch the polished surfaces. A sometimes satisfactory electrolytic etch was obtained by lowering the potential across the chromic-acetic bath to 25 V. A chemical etchant consisting of a I:1 mixture of nitric and acetic acid was found to work on all the alloys. ~tehing times for the HNOa :HAc mixture varied proportionally from 25 min for the 1 at y. Re alloy to 10 set for the 14 at y. Re alloy. Vickers hardness numbers were obtained using a 10 kg load and averaging at least four readings. The grain size of each specimen was estimated using the ASTM comparative method. In instances where the grain size was larger than the standards, it is reported as the average diameter in inches. The DI phase was often present as platelets and its grain size is reported as plate width in millimeters. The average grain size estimated by visual integration is reported for duplex structures. X-ray diffractometer patterns taken immediately after each heat treatment served as a quick means of phase identification and as a check for isothermal transformations at room temperature. After quenching and prior to
TRANSFORMATIONS
examination,
the
AND
sample
STRUCTURES
was
IN
mechanically
polished and electropolished. The latter step was necessary to avoid line broadening, especially for the sofikr alloys. Filtered CuKa radiation was used at a chart speed of 1” (Ze)~min. After quenching, specimens for Debye-Scherrer photographs were ground to a 0.5 mm dia. cylinder and then electrolytically thinned to a needle of 0.15 mm dia. Filtered CuKa radiation was used with a 114.6 mm dia. camera. Lattice parameters were determined by leastsquares extrapolation against the Nelson-Riley function using an IBM-7074 computer 1% 14). 3. 3.1.
THE
VACUUM
a
The core experiments were those in which series of alloys were water quenched, vacuum quenched and slow cooled from the y and @ phase fields. More specifically, they were: (1) a series of alloys water quenched from 975” and 850’ C after Q and 4 h at temperature, respectively ; (2) a series of alloys vacuum quenched from 975, 850, 700 and 662’ C after $, 4, 12 and 6 h at temperature, respectively; and (3) a series of alloys cooled at 25” C/h from 825” C after 4 days at temperature. These alloys were examined for crystal structure, microstructure, grain size and hardness. The crystal structure of the uranium phase present is shown s~~ematieally in fig. 1. The results of the grain size and hardness measurements are shown in figs. 2 and 3, respectively. The significance of
QUENCkED
SYSTEM,
975O
C
a*‘
2
4
VkUI)M
6
I
8
‘QUfNCiiEb
850dC
R
I
I
I IO
t
1 12
’
’
’
’
B 0
4
6
8
QUENCHED
IO
2
vAw&
I
YI ’
6bOg’C
12 I
’
a 0
265 I
I_ 14 ’
Yp
2 WATER
I
,
’
B-
0
Results and discussion CORE EXPERIMENTS
URANIUM-RHENIUM
4
6
, IO
a
‘QlJ$NCkEo’
700°1
‘
14 I
I
Y1 I 12
‘ 14
12
14
.
C
a” 0
L
I 2
1
I I 4 ATOM
I , 6 8 PERCENT
IO RHENIUM
Fig. 1. Effect of cooling rate and quenching temperature on the crystal struct#ureproduced in the uranium-rhenium system.
IO
x
WATER
+
VACUUM
l
WATER
OUENCHED
975
QUENCHED QUENCHED
975 850
‘C “C OC
0
VACUUM
QUENCHED
950
‘C
b
VACUUM
QUENCHED
662
‘C
v
COOLED
25
+C/tiR
FROM
625°C
IO
12
these data is discussed later along with results from supplementary experiments. 3.2.
METASTABLE
y AND yo STRUCTURES
Inspection of fig. 1 shows that 8.5 at y0 Re was sufficient to cause capture of the y phase on vacuum quenching from either 850 or 975” C. Similarly, about 9.5 at o/0 Re was sufficient to retain the y phase on water quenching from 850” C. X-ray examination of the retained y alloys showed the structure to be body-centered cubic. Lattice parameters of 3.452, 3.448 and 3.448 A were obtained for the 10, 12 and 14 at o/0 Re
0
2
4
6
ATOM
Fig.
2.
6
PERCENT
14
6
RHENIUM
Comparison of grain size meas~emen~,s after various
thermal
treatments.
alloys, respectively, which were water quenched from 850’ C. Klepfer and Chiotti 15) report a value of ao= 3.538 i$ for pure uranium at 800” C. Correcting this value to room temperature using the linear coefficient of thermal
266
R.
x
700
_I
WATER
PUENCHED
+
VACUUM
.
WATER
J.
975
OUENCHED
975
QUENCHED
650
JACKSON
OC OC
VACUUM
OUENCHED
850
OC
o
VACUUM
OUENCHED
662
OC
o
COOLED
25
FROM
W.
L.
LAESEN
His proposed structure is based on a block of 2x2 x 1 y cells in which the solute atoms arc restricted to the edge-center sites. The observed line positions and intensities for yb0 agree very
oc
o
OC/HR.
AND
825
well with the structure proposed by Tangri for a cell with ao= 6.92 A and CO= 3.38 A. Hence,
“C
600
z
it is proposed morphous
that
the yb0 structure
with the tetragonal
ordered
is isophase
observed by Tangri in the U-MO system. A comparison of the y and ybs microstructures is shown in fig. 4. In the 9 at o/o Re alloy vacuum quenched from 775” C coarse bands were observed in addition to the fine interwoven bands. The bands were similar to those observed next to hardness indentions in y alloys and are believed to be deformation bands. II
I
0
2
I
II 4 ATOM
Fig. 3.
I
I
6 PERCENT
I
I
8
I IO
I
I 12
I
IA I4
RHENIUM
Comparison of hardness values after various thermal
treatments.
expansion determined by Klepfer and Chiotti 15) of 22.5 x lo-6/“C yielded an CO at 23” C of 3.476 A. Hence, the addition of rhenium atoms causes a contraction of the bee y cell. This contraction would be expected on the basis of size considerations. Additional quenches were made in an effort to obtain a y-like structure showing definite tetragonality as reported by Williams 16). Compositions of interest were vacuum quenched from 700 and 775’ C after 12 and 8 h at temperature, respectively; and water quenched from 775” C after 8 h at temperature. Vacuum quenching the 9 at y0 Re alloy from 700 and 775” C, and the 10 at o/o Re alloy from 700” C did yield specimens whose X-ray diffraction patterns showed line splitting and weak additional reflections. This structure is hereafter referred to as yba where the subscript indicates the metallographically identified bands (crosshatch structure) and the superscript the corresponding alteration in crystal structure. Tangri s) determined the structure of a tetragonal ordered phase observed under somewhat similar conditions in the U-MO system.
Fig. 4. Photomicrographs contrasting the y and y C (a) The 10 at y. Re alloy water microstructures: quenched from 850” C showing retained y microstructure. Polarized illumination.
x
alloy
150. (b) The 9 at y0 Re
vacuum quenched from 700” C showing tho crosshatch y g microstructure. Polarized illumination. x 100.
TRANSFORMATIONS
3.3.
B~TASTABLE
AND
01' AND
STRUCTURES
n"
IN
THE
STRUCTURES
URANIUM-RHENIUM
MetalIographic
SYSTEM,
I
267
structures associated with the
Two metastable modifications of the normal orthorhombic a-uranium structure were ob-
a’ and OI’ phases are shown in fig. 5. The observed structures varied from granular to
served for rapidly cooled y alloys containing between 3.5 and 9 at yO Re. The a-phase
very fine crossed bands (crosshatch struct~e) Taomixtures of fine and coarse bands which often
modifications are designated (x’ and a”, the superscripts indicating a departure from the equilibrium structure as determined by X-ray diffraction techniques. The n’ modification is a result of a relative contraction of the b parameter with slight expansions along the a and c
appeared
axes of the normal a-uranium structure. The bn modi~cation occurs at higher rhenium content and is due to a relative contraction of the b parameter, together with a change from an orthorhombic a structure to a monoclinic structure.
Fig. 5.
Photomicrographs
showing the microstructural
acicular.
The type of microstructure
obtained appeared to be directly associated with cooling rate and composition and not with the crystallographic phase present, i.e., orthorhombic OI’ or monoclinic &“. Vacuum which at
quenching yielded microstructures low rhenium compositions were
granular and which with increasing rhenium content developed a crosshatch structure (eompare figs. 5a and 5d). Water quenching yielded much coarser structures which at all compositions were considered banded. At the higher
types associated wit,h the 0~’ and 01” phases:
(a) The
6 at e/0 Re alloy vacuum quenched from 85W C showing crosshatch and granular appearing a” microstructure. x 100. (b) The 6 at y/o Re alloy water quenched from 850” C showing n’ microstructure consisting of coarse and fine bands. x 100. (c) The 7 at o/0 Re alloy water quenched from 850’ C showing OLD microst,ructuro consisting of coarse and fine bands. x 150. (d) The 8 a.t o/0 Re alloy vacuum quenched from 975” C showing crosshatch a” microst,rueture. Note cruciforms around carbide particles. x 160. All polarized illuminr&ion.
rhenium
contents
determine termed
it
whether
banded
was
the
often
diffioult~ to
structure
or acicular
should
bc
(figs. ;ib and 5~).
Lattice parameters (10, 7,0, COand angle y of the 0~‘ and CX”structures were determined using specimens water quenched from S50° C. The results are represented by solid lines in fig. 6. The main features are: The values of cl0 and CO increase with increasing rhenium content.
slightly
The value of bo decreases markedly increasing rhenium content.
with
The angle y gradually increases from 90.0” at 6 at yi;, Re to 92.2’ at 9 at 7; Re. The effect of introducing rhenium atoms into the uranium lattice may be explained in the following manner. The unequal interatomic distances in a uranium give an ellipsoidal shape to each uranium atom rather than a spherical shape. Assuming that each atom behaves as a hard particle, the rhenium atoms may be represented as hard spheres of radius 1.380 !J when in a coordination of twelve, and uranium atoms may be represented as hard ellipsoids *
I
s
.
EXPERIMENTAL
o
THEORETICAL
I
,
,
I
t
1
h
z
4.98
.. _____._ I-w-m
2 4.94 2.90 z ;
--
.
.
* ___ ..---
2.86
-----
.
0
-
5
6
l
I-
0
I
2
3 ATOM
Fig.
4
PERCENT
7
8
9
IO
RHENIUM
6. La&ice parameters of the ,x’ and 1” structures developed on wlttcr quenching from 880” C.
having and
the dimensions 1.660
K
along
of
1.3S1, l.426,
principal
1.640
inter-atomic
directions. The substitution of rhenium spheres into t,he lattice of ellipsoids results in the caleulatPed changes shown as open circles in fig. 6. The agreement between experimental and calculated lattice parameters is very good considering that the electronic interact,ion n-as ignored for the purpose of this rough calculation. 3.4.
DIRECT
y +x"
(A') TRANSFORMATION
hwez 17)has shown that rates of cooling up to 10 000” C/see horn the y phase do not prevent the formatio~l, in order? of /j’ and oi in pure uranium. This sequence of phase changes observed by Duwez at ext,remely fast cooling rates is also observed by vacuum quenching U-Re alloys containing less than about 3.5 at ‘:/A Re. Some observations supporting this sequence of phase changes are:
1. For alloys containing less than 3.5 at y0 Re which were LY at room temperat,ure, two thermal arrests were detected on cooling. 9 -I Irregular cy grains, similar to those obtained on quenching pure uranium, were observed on vacuum quenching the 1 at “/o Re alloy. 3. Duplicate alloy specimens! with the same thermal and mechanical histories, cont,aining 1 at :I0 Re were vacuum quenched from within the B and y fields, respectively. The ~~licrostructures as well as the hardness were nearly identical in both the p and y quenched specimens. Thus, it would seem that under these conditions the formation of in~rmediate /3 is not suppressed. It is shown in II. The ,!JFE B transformation, t,hat as t’he rhenium content increases the b --f ti transformation temperature is lowered and the stability of the ,I3phase increases with increasing rhenium cont,ent. On either water quenching or vacuum quenching from the y region Dhe 3 at 76 Re alloy was ,8 while the 4 at o/0 Re alloy was z‘. If the p phase had formed in the 4 at y;b alloy, one would expect it to be present at room temperature. The absence of the 1 phase
TRANSFORMATIONS
AND
STRUCTURES
IN THE
suggests a direct y --f LX’transformation
at fast
cooling rates. This viewpoint
by the
following
is further
observations
supported
URANIUM-RHENIUM
SYSTEM,
I
269
:
A 3.8 at y0 Re alloy water quenched from the y region had an ix’ structure while one cooled at a slower rate (400’ C/h) had a f!l structure. The 4 at y0 Re alloy cooled at 25” C/h had retained fi present, while as mentioned above the rapidly converted
cooled
alloys
were completely
to a’.
The microstructure of the LX’or 0~”alloys all showed straight grain boundaries (fig. 7a) which were those of the original y grains. All the retained /l alloys showed very irregular grain boundaries (fig. 7b). Likewise, any o( formed from /3 had highly irregular grain boundaries (fig. 7~). Had the 01’ or OL’ formed via a double transformation, i.e., would be y + p + a’, the microstructure expected to be irregular. The as-cast 1, 2 and 3 at y0 Re alloys contained numerous small cracks and voids which formed during the y --f @ transformation. These imperfections were readily seen metallographically and led to low gravimetrically determined densities. For example, the 3 at o/o Re as-cast alloy was retained ,8 with a density of 18.0% g/cm3 compared with a theoretical room temperature @ density of 18.80 g/ems. For alloys containing more than 3.5 at o/o Re, the ascast buttons were free of cracks and had densities of approximately 19.0 g/cm3 which compares favorably with the expected room temperature density for the distorted 01 structures. Had any /3 formed during the rapid cooling of these compositions, cracks would be expected in the as-cast buttons. This is further evidence for the direct y + 01’ or y + an transformation. Thus, at cooling rates obtained by water or vacuum quenching, the postulation of a direct y --f cy’ transformation is necessary to explain the observed behavior.
Fig.
7.
direct
Some microstructural y -P (x’ kansformation:
evidence suggesting a (a) The
4 at
YJ, Re
alloy water quenched from X50 C showing the banded microstructure boundaries. quenched
of bhc LX’ st,ate amidst
prior y grain
x 100. (lo) Tho 4 at y0 Re alloy vacuum from
662” C showing
highly
irregular
/3
grains.
x 150. (c) The 1 at. O/o Re alloy vacuum quenched from 662” C showing highly irregular OL grains.
3.5,
x 150. All polarized illuminat,ion.
SEQUENCE
OF PHASE
CHANGES
In the preceding sections it has been shown that during the transformation of y to (Y, the metastable structures a’, 0~” or yb0 may occur.
270
E. J. JACKSON
The mode of formation
AND
of these structures, and
their relationship to cooling rates explained in the following manner.
may
be
It has been proposed by Tangri and Williams2) that increasing additions of molybdenum progressively stiffen the uranium lattice, thus making it more and more resistant to shear, and that the effect of increased cooling rate is to favor arises
the shear process.
from
increased
caused by faster cooling
This presumably
hydrostatic
pressures
rates and also from
increased transformation stresses due to faster rates of transformation, once shear has started. Tangri also proposed that an optimum value of the ratio of quenching stresses to lattice stiffening, which for convenience of discussion was called the “stress/stiffening ratio”, is required to generate any given martensitic phase. Hatt and Roberts 18) have shown from X-ray diffraction work on a Zr-50 at y0 U single crystal that the 01 structure can be generated from a bee y structure by a shear on (112} y planes in a
W.
L.
LARSEK
of any given rhenium given
content,
and thus of a
restrict the explain the amount of generation of the a” structure on water quenching the !I at o/o alloy from 850’ C which on vacuum quenching under identical conditions yielded
stiffness?
is to relatively shear. This would
a y structure.
It can be seen in fig. 1
that, in each case, the effect of rapid quenching is to produce a structure further removed from the parent bee phase than that obtained
with
slower cooling, indicating that severe quenching causes t’he transformation to go towards completion. This effect is well illustrated by the photomicrographs of a !I at o/o Re alloy vacuum quenched (fig. 8a) and water quenched (fig. Xb)
(111) y direction. Furthermore, they developed a crystallographic model which explains the generation of the 01’, OI”and yo structures in the uranium-molybdenum system by progressively restricting the amount of shear. On water quenching alloys containing 3.5 to 6.5 at o/o Re, the optimum stress/stiffening ratio is realized, thus producing the critical amount of shear required to generate the a’ structure. With increasing rhenium additions, the lattice
is steadily stiffened until with 7.0 at o/o Re the lattice is too stiff to permit the critical amount of shear. Only a limited shear is produced resulting in an angle y of 00.8’, thus yielding a monoclinic 0~”structure. Further additions of rhenium progressively restrict the amount of shear, thus causing angle y to steadily increase to a value of 92.2”, as observed in the 9 at $” alloy. In highly stiffened lattices, as for example in alloys containing 10, l:! and 14 at oh Re, no shear or only a very limited amount takes place resulting in the generation of the y or yo structures, respectively. The effect of slower cooling rate in a lattice
Fig. 8. Comparison of microstructures of the 9 at) % Re alloy water quenched and vacuum quenched from 975” C:
(a) Thr 9 at y0 Re alloy vacuum quenched
from 975” C showing clean equi-axed grains of y structure. x 150. (b) The 9 at o/0 Re alloy water quenched from 975’ C showing single-phase, marten&o-like N” struct,ure. x 100. Both polarized illumination.
TRANSFORMATIONS
975” C.
from
undergone
AND
The
STRUCTURES
rapidly
cooled
IN
alloy
a shear transformation,
has
while the
THE
URANIUM-RHENIUM
water quenched
SYSTEM,
microstructures
271
I
disproves
this
more slowly cooled alloy appears untransformed. Work by Hills et ~~1.19) on end-quenched
and indicates the softening is inherent in the transition structure. Another possibility is that as the rhenium content is increased, theitendency
U-MO rods showed similar structural variations.
for precipitation
The slowly cooled end of a 10 at y0 molybdenum
y + 01’, an reaction occurring at a lower temper-
rod so quenched showed y” whereas the structure
ature);
of the
stage
rapidly
quenched
pattern
showed
end was
LX”. An
7.5 at y0 molybdenum
01” structure X-ray
cooled
only
throughout
rod had an
its length
of the rapidly
but
quenched
a little line splitting,
endthe
region
indicating
that its monoclinic angle was very near 90”, i.e.? the structure was approaching 0~‘. 3.6.
METALLOGRAPHICOBSERVATIONS
Metallographic observations have been presented in earlier sections to substantiate certain arguments. In this section, metallographic observations not previously discussed are presented and their relationship ments is developed. 3.6.1.
to earlier argu-
Hardness
Fig. 3 contrasts the differences in hardness for the vacuum quenched, water quenched, and slowly cooled alloys. The hardness-composition curves for vacuum quenched and water quenched y alloys are quite similar. The general shape of the curve is explained by the following hypothesis. The initial OLhardening is a combined solid solution and precipitation hardening effect. The difference in hardness at 1 at o/o Re on water quenching and vacuum quenching is due to the higher inherent hardness of the /? solid solution which is retained on water quenching. The general increase in p hardness with composition is a combined solid solution (due to supersaturation) and precipitation hardening effect. The linear decrease in hardness of the 01’ and 01” structures with increasing rhenium content above 4 at o/o is puzzling. At first glance it might appear that the decrease is associated with the coarsening of the structure, i.e., the appearance of bands and large lenticular plates. However, comparison of vacuum quenched and
is lessened (possibly due to the
and at the lower compositions of precipitation
tectable
occurs
by normal X-ray
techniques.
which
an early is unde-
and metallographic
It is well known that precipitated
phases that are undetectable by normal X-ray and metallographic techniques can markedly affect hardness. In any event, the continuous and gradual hardness change in passing from y’, to a” to (Y’ lends credence to the argument that y“, a’ and a’ are intermediate transition states in the formation of LYfrom y. The hardness minimum of the y quenched alloys is always associated with the first appearance of the y structure. The increase in hardness at higher rhenium contents is due to the grain boundary network of URez phase. The higher hardness of the 4 and 5 at y. Re alloys vacuum quenched from the y region is due to the formation of a detectable second phase. Water quenching of these compositions from the y field resulted in apparent single phase structures. The difference in hardness at the minimum for y quenched alloys is due to the effect of cooling rate on structure produced, i.e., CL”or y. The similarity in hardness at higher rhenium contents of the /? quenched and slowly cooled y alloys is interesting. The uranium phase present was /l in one case and OLin the other, thus the hardness is probably governed by the UzRe phase distribution which was similar in both cases. 3.6.2.
Grain size
The results of the grain size measurements are shown graphically in fig. 2. For the 0~’ and DL”structures the grain size shown is that of the prior y grain boundaries. The dotted symbol for the series quenched from 662” C is also for the prior y grain size. It is seen that all heat treatments used resulted in similar y grain sizes ,that were quite large except at relatively high
272
R. J. JACKSON
AND
rhenium contents. The reason for the restricted grain growth at high rhenium contents is undoubtedly
the grain boundary
network
of
URes particles. An interesting alloys obtained
grain size trend for retained /3 on y quenching
is also readily
seen. The fi grain size of water quenched
alloys
W.
I,. LARSEN
with increasing
rhenium
content
developed
a
crosshatch structure. The reason for the difference
in appearance
of
and
LY’ and
pendency not
LY” microstructures on cooling
readily
explains
the
their
rate and composition
evident.
One
hypothesis
observed
microstructures
deis that
is as
decreased with increasing rhenium content whereas on vacuum quenching the /3 grain size
follows. The tetragonal ybo phase having a crosshatch
increased with increasing rhenium content, This
microstructure
effect is probably
atures low in the y region and this phase yields
associated
with the greater
degree of supercooling of the y + p reaction on water quenching and the resulting larger number
of nucleation
sites.
3.6.3. Hicrostructures
The y quenched
pure uranium
was charac-
forms on quenching from temper-
a granular structure after transforming
to 01on
aging. This behavior possibly explains the a’ and CY~microstructures seen in the vacuumquenched alloys. Presumably, on vacuum quenching the cooling rate is slow enough that the alloys order to form yb0 as an intermediate
terized by irregular grain boundaries and highly twinned grains. A similar microstructure was observed in the 1 at o/0 Re alloys vacuum quenched from 850 and 975" C. The 1 at y0 Re alloys water quenched from 850 or 975’ C were retained /3. The retained /3 phase at these compositions decomposed at room temperature by the isothermal growth of LY.The 2 and 3 at o/o Re alloys either vacuum quenched or water quenched from 850 or 975' C were
transition state. The characteristic crosshatch mi~~rostructure of the ybo phase is thus given to the resulting 0~‘ or 0~” phase. The 01’ or a” platelets would thus be too small to be resolved in the optical microscope due to the small size of the bands making up the yt,O crosshatch structure. The observed change from a crosshatch to a granular structure is due to the presence of a second phase which in the case of the 4 and 5 at 0/o alloys was detectable by
retained p. The retained ,8 at these compositions was quite stable and no decomposition at room temperature was detected. The grain boundaries in these alloys were often characterized by a
normal X-ray techniques. On water quenching from the y region, there is not sufficient time for ordering and the rr,s structures does not form. Also, the thermal gradient and quenching stresses developed on water quenching are more severe than during vacuum quenching. Hills et aZ.20) and Butcher and Hatt 21) have shown that retained y0 alloys in the U-MO system deform by twinning mechanisms which results in coarse-banded ye microstructures. Thus, on water quenching, coarse twin bands may form followed by transformations to a’ and ocw by the growth of lenticular plates. The size of the resulting a’ or ~l# platelets would be governed by the size of the twin bands, or if twinning was absent by the size of the parent y grain. Thus in fig. 5b, the large bands would be twin bands and the smaller bands lenticular plates of cx’. In fig. 5~3, twinning of the y phase was restricted resulting
sawtooth appearance. For water quenched
alloys containing
4 to 9
at o/0 Re, the Iw, temperature for the direct y + ac’ (ol”) transformation lies above that for the y -+ /3 transformation and only DC’or (x” is observed in the microstructure of these alloys. These alloys were distinguished by smooth y grain boundaries and coarse microstructures which at all compositions were banded and/or acicular in appearance. For vacuum quenched alloys containing 4 to 8 at o/o Re a similar direct y -+ 00 (a”) behavior was observed except for the 4 and 5 at o/0 alloys where a second phase was detectable. Vacuum quenching yielded microstructures which at low rhenium compositions were granular and which
TRANSFORMATIONS
in the formation followed
AND
STRUCTURES
of large lenticular plates of a”
at a lower temperature
by the neces-
sarily smaller lenticular plates of (Ye. This scheme with some modification interpretation microstructures.
IN THE
of all the observed For
example,
SYSTEM,
I
273
In the 4 and 5 at y. Re alloys the UsRe phase
was present
in the
form
of
platelets
(often parallel) which were nearly continuous. allows
01’ and 0~” the
URANIUM-RHENIUM
micro-
The 4 and 5 at y. alloys also showed a continuous
jl network
boundaries.
outlining
the
prior
In the 4 at y. alloy,
y grain
the matrix
structure of the 7 at y0 Re alloy water quenched
consisted
from 850” C shown in fig. 5c could be interpreted
Presumably
‘in the following MS temperature
due to the non-selectivity of the etchant. On cooling the hypoeutectoid alloys, uranium-rich ,!l nucleates along y grain boundaries. On further
manner. On cooling below the the large lenticular a” plates
(mainly at 45”) form a network in the supercooled y matrix. The resulting transformation and quenching stresses cause the “boxed-off” y to deform by twinning giving rise to the vertical bands (some horizontal). The vertical ,bands then transform to an and the resulting ,transformation stresses cause motion along the twinned surfaces. This results in deformation of the LX”plates (white) that were not part of L the original 0~” network. As the rhenium content increases, the M, temperature falls below room temperature. Thus, for vacuum-quenched alloys containing 9 at y0 Re or more and for water-quenched alloys containing 10 at o/o Re or more, the y phase is stabilized at room temperature. Fig. 4a shows the characteristic equi-axed microstructure. A word is in order concerning the micro‘structure of the series of alloys cooled at 25” C/h ‘after 4 days at 825’ C. The 1 at oh Re alloy was characterized by twinned and very irregular LXgrains which would be expected from the double transformation (y + ,9 + CL).The 2 at o/o Re alloy consisted of a matrix of ,!3and oc uranium and a discontinuous precipitate of UaRe. The micrograph of this alloy at low magnification showed what appeared to be small equi-axed grains of a in the coarse /l matrix. This a structure is similar in appearance to a “pseudo (x” structure reported by Hills et aZ.22) for U-MO alloys. They report that this phase forms in the high temperature o( region and presumably consists of a duplex lamellar structure of (x+ UaMo within the visible “pseudo 0~” grains. Both the 2 and 3 at y. Re alloys showed a discontinuous UaRe precipitate.
mainly
of p with
some
01 present.
the a-j3 boundaries were not visible
cooling, coring of the ,!l phase results giving rise to concentration gradients within p. Precipitation of UaRe will occur at the rhenium-rich /l interfaces. This retards lateral growth and gives rise to transverse growth of the p phase. The regular decrease in the amount of retained fi with increasing rhenium content is interesting. The 2 at o/o Re alloy was heavily ,3 while only a small amount was present in the 5 at oh Re alloy. Possibly, the increasing amounts of UaRe promote formation of 01 from /l by acting as nucleation sites. Another possibility is that two reactions are competing for y, i.e., the hypoeutectoid y -+ ,!? and the direct y + LY reactions. The latter possibility seems the more plausible. The 6 at y. Re alloy was free of retained /? and its microstructure was more acicular in appearance than those of lesser rhenium contents. As rhenium content increased further the acicular appearance lessened and was nearly absent in the 9 at o/o Re alloy. The acicular appearance was replaced by optically active irregularly shaped apparent grains. Electron micrographs of replicas from these irregular shaped apparent grains showed them to consist of a very fine duplex lamellar structure of 01i- UaRe. A complete series of photomicrographs from which the preceding arguments were developed is shown in a thesis by Jackson 23). 3.7.
RELATION
TO OTHER
SYSTEMS
Metastable transition phases are found in a number of alloy systems in which the bee structure is stable only at high temperature.
274
R.
The
tendency
phase
towards
formation
the
interest
a-uranium
JACKSON
metastable
is most
W.
like
LARSEN
Burgers’ shear mechanism 24) for the bee + hcp The
This group is of
is
L.
transformation.
during
in this investigation
structure
AND
transition
pronounced
the bee --f hcp transformation. particular
J.
since
hexagonal
nature
of
phases
observed
mation
is shown
the
metastable
transition
for the bee + hcp in table
1. At
transfor-
low
solute
metals except that successive layers of atoms in the basal plane are skewed back and forth
concentration the transformation occurs /3 (bee) --f 01’ (hcp) by Burgers’ shear mechanism. The
in the (010)
effect of increasing solute content is to stiffen the lattice making it more resistant to shear.
uranium
direction.
By comparing
the a-
cell with the orthohexagonal
cell for
a hcp structure, one obtains b/u= 2.06 and c/a = 1.73 for 01 uranium compared with b/a = 1.73 and c/a= 1.63 for a hcp structure.
This
The
solute
TABLE transition
Metastable
comwosition Very
concentrations
transition
transformation phase?
Met&stable phase* bee + ol uranium
bee --f hcp
low solute
a? (orthorhombic) (forms from /3) /I
(*’ (hcp)
concentration
(supersaturated
(tetragonal)
(w) aa’ (orthorhombic)
Low solute concentration
o(s’ (orthorhombic)
Low solute concentration Medium
solute
w (C32 hexagonal)
0zb”(monoclinic)
concentration High
solute
concentration
diffuse o (faulted p
(tetragonal
w)+
c/n >
yso (ordered tetragonal)
1)
or /?’ (solute rich) High solute
y
All found in Zr-Nb alloys 25). LY’, W, and diffuse o in Zr-U alloys Is, 26). LY’, COfound in Ti-V, Ti-Cr, Ti-Mn, Ti-Fe, Ti-Mo, systems. 0~’ found in Ti-Cu plus many
*
other systems.
All present in U-MO system 1.2, s). All present in U-Re All except
(bee or tetragonal c/u w 1)
concentration t
/I in U-Nb
system
phase
a metastable
1
concentration Low solute
a
phases observed during the bee + hcp and
bee -+ 01 uranium Relative
in
from the hcp structure.
further
At medium transition
phase forms via a mechanism that requires shorter atom movements. At higher solute concentrations the product may be diffuse w (the X-ray reflections are diffuse ; hence, the name diffuse W) accompanied by a supersaturated ,6’ or bet /l” phase. The diffuse CJ
addition of rhenium atoms to the B-uranium lattice has little effect on the c/a ratio while the b/a ratio decreases ; thus, making a-uranium more “hexagonal like”. It also should be pointed out that the shear mechanism proposed earlier for the U-Re system is merely an extension of
Metastable
results
removed
(this investigation).
system 3.4).
a’, aa’, a’s, yo in U-Ti system 6, 9). a’, OCR’and possibly CQ,” in U-Zr system 7).
Ti-Co
and Zr-Mo
TRANSFORMATIONS
structure
AND
presumably
STRUCTURES
arises from
IN
incomplete
TEE
URANIUM-RHENIUM
bears
a similar
relation
SYSTEM,
for
the
I
U-MO
275 and
solute
U-Nb systems, i.e., at about b= 5.74 8. The a’-&” and LY”-y boundaries in the U-Re system
concentrations the p phase is stabilized and aging is required for the appearance of the
occur at about 5.80 and 5.75 A, respectively, on water quenching from 850” C, thus sub-
metastable transition phases. As mentioned earlier, metastable
stantiating the above hypothesis. These size correlations are based on a limited
glide
due to lattice
heavily
phases
faulted
have
uranium formation transition
structure.
been
alloy
stiffening At
observed
systems.
A
resulting higher
in a
transition
in a number summary
of
of the
conditions for these metastable phases in several uranium-systems
is presented
in table
1.
Comparison of columns 2 and 3 of table 1 shows that the transition phases occurring during the bee -+ hcp and bee -+ a~ uranium transformation are quite similar. Inspection of column 3 of table 1 shows that the behavior of the U-Re system is not unlike those of other highly y-miscible ~anium systems. The greatest similarity lies between U-Re and U-MO alloys ; the difference being the absence of the acicular a’ microstructure at low rhenium contents. The sequence of phase changes for the two systems is in general the same. The U-Nb system is also quite similar except the @ phase is not encountered. Lattice parameters are available for the distorted a structures in the U-MO 122), U-Nb 39 4) and U-Ru 5) systems. These lattice parameters along with w’-CQ,” and ~‘-y composition boundaries on water quenching have been discussed by Anagnostidis et aZ.3) and Tangri and Chaudhuri 4). They point out that if one compares the lattice parame~rs for the three systems, a correlation with size of the solute atom becomes obvious. As the solute size increased the change in 6 parameter decreased as would be expected on the basis of substituting hard spheres for ellipsoidal uranium atom. The ab’-~” composition boundary occurred at about b= 5.80 ,k for the two systems in which a” formed. This finding indicates the monoclinic (XII phase corresponds to a given lattice distortion and that elements having a small atomic diameter favor the formation of the Olbnphase. Likewise, the &b”-y composition boundary
number of observations, of the known
but do correlate most
data. As others 3*4) point
out,
this could permit one to predict the extent of metastable phase regions on the basis of atomic size ; the elements of small size decreasing extent 4.
the
of the domains.
Summary
A study was made of the time-temperaturecomposition dependency of crystal structure and microstructure under non-equilibrium cooling conditions for uranium-rich alloys in the uranium-rhenium system. The y phase under certain conditions was found to transform to ar without passing through the intermediate ,3 structure. Accompanying this transformation was a number of metastable transition states. Two transition states were modifications of the orthorhombio ~-uranium structure. They were: 1. a supersaturated a structure showing a notable oontraction of the b parameter with slight expansions along the a and c axes, and 2. a supersaturated a structure that had distorted to a monoclinic cell in addition to the parameter changes noted above. At high solute contents
the y phase can be
rendered metastable to room temperature. By cooling certain compositions from low temperatures in the y region a metastable transition state (~0) based on the bee structure occurs. This was a tetragonal ordered phase based on a block of 2 x 2 x 1 y cells having a c/a< 0.5. The compositional limits within which the stable and metastable phases form, and the effect of cooling rate and temperature on producing alternative structures were determined. From this was deduced the sequence of phase changes. A crystallographic and microstructural model
276
R.
was developed
J.
JACKSON
to explain the observed
structures
and microstructures
of cooling
rate and composition
AND
crystal
and the effect thereon.
Finally, the relation between metastability uranium-rhenium alloys and other systems which
metastable
states
are
observed
in in was
W.
Common
13) M.
14) D.Hansen,
argonno
(USA)
(1961)
St,ructure of YNi, Ames,
Williams,
Iowa,
226
M.S.
USA,
Mat.
4,
K.
5,
(1965) 278 K. Tangri,
Colombie
5 (1962)
and
H.
109
29)
Tangri and D. Chaudhuri,
J. Nucl.
Mat.
Thesis (Iowa
1964)
15 (1965)
and C. Rao,
288
22)
Tangri,
Mem.
A. Harding,
AERE
Acta
Cryst.
Met.
58 (1961)
and C. Knight.
M/R
2673A
14 (1961)
‘9 11) F. Rough and A. Bauer, Constitutional diagrams of uranium and thorium alloys (Addison-Wesley,
Reading,
Mass.
1958) p. 60
2558/g
24 (1953)
152
Harwell (UK)
B.
Butcher and J. Heywood,
R. Hills,
B. Butcher
11 (1964)
Report
(1960)
Metals
3 (1961)
J. Less-
155
and B. Howlett,
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149
B. Butcher and B. Hatt, R. Hills,
B. Howlett
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Metals
Jackson,
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Harwell
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kinetics of phases in the uranium-rhenium
469
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Sci. Rev.
M. Waldron
Report
B. Hatt,
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163
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9 D. Douglass, Trans. Am. Sot. Metals 53 (1961) 307 7) D. Douglass, L. Marsh and G. Manning, Trans. Sot. Metals
1
communication, (1963)
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15 21
D. Chaudhuri
)
Monti,
(UK)
and J. Roberts,
AERE/X/PR 19
M.
B. Hatt
9
private
Inst.
P. Duwez,
J. Nucl.
(UK)
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Heaton.
I51 H. Klopfer and P. Chiott,i, Ames (USA) Report G.
(1961)
9
L.
6156
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Research
3) M. Anagnostidis,
9
and
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and W. Larsen, J. Less-
5 (lY63)
IHC-X93 (1957)
‘1 J. Lehmann, J. Nucl. Mat. 4 (1961) 218 2, K. Tangri and G. Williams, J. Nucl. Mat. 4
K.
Metals
Mueller
Report
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LARSEN
R. Jackson, D. Williams
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