Journal of Nuclear Materials 59 (1976) 137-148
0 North-l-ioliand Publishing Company
THE PHYSICAL METALLURGY OF HIGH STRENGTH Zirconium
ALLOYS
L.H. KEYS, G. JOHANSON and A.S. MALIN
School of Metallurgy, University of N.S. W., Kensington, Australia
Received 22 July 1975
The requirements of materials for use in the core section of unenriched uranium type nuclear reactors are, to a large
extent, fulfilled by zirconium alloys. However, the operating temperature of such reactors is circumscribed by the inability
of the presently available alloys to maintain their properties at temperatures exceeding 350-4OO’C. This paper is concerned
with the selection of alloying elements for the improvement of the high temperature strength of zirconium alloys and reports the results obtained from a study of the phase transformations in dilute Zr-Cr alloys.
Les exigences des matiriaux utilisks dans le coeur des rdacteurs du type a uranium non enrichi sont en grande partie satisfaites par les alliages d’uranium. Cependant, la temp&ature de fonctionnement de tels rdacteurs est limitie par l’impossibilit6
pour les al&ages disponibles actuellement de maintenir leurs prop&t& i des tempiratures exc$dant 350-400°C. Ce m&moire
est relatif i la sdlection d’dldments d’addition en vue d’ambliorer la risistance i haute tempdrature des alliages de zirconium
et rend compte des rdsultats obtcnus au tours d’une &ude des transformations de phase des alliages diluds Zr-Cr.
Der Bedarf an Material fiir den Kern der mit nicht angereichertem Uran betriebenen Reaktoren wird zum grossen
Teil von Zirkon-Legierungen gedeckt. Die Betriebstemperatur derartiger Reaktoren ist jedoch dadurch eingeschrLnkt, dass
die gegenwgrtig verfiigbaren Legierungen ihre ~igenschaften bei Temperaturen oberhalb 350-4OO’C nicht beibehalten k6nnen. Diese Arbeit beschiftigt sich mit der Auswahl von Legierungselementen zur Verbesserung der Hochtemperaturfestigkeit
von Zirkon-Legierungen; iiber die Ergebnisse einer Untersuchung zur Phasenumwandlung verdiinnter Zr-Cr-Letierungen
wird berichtet.
1. Introduction
of zirconium alloys used in water-cooled reactors [2].
An increase in strength of these alloys would allow a
reduction in wall thickness of the components and a
minimisation of the neutrons absorbed by the material
in the core. Also the thermal efficiency of these reactors is improved significantly if the operating temperature is increased to 450-500°C.
The alloys in
present use, based on the Zr-Sn and Zr-Nb binary systems, have good corrosion resistance and low hydrogen
pick-up but relatively low strengths. At the present
operating temperatures of water-cooled reactors
(300-35O”C), the strengths attained in these alloys
by cold working can be retained. However, as recovery
and recrystallisation of these alloys occurs in the vicinity of 400°C [3], work-hardening is not a significant
source of high temperature strength and much more
stable microstructures are required for the operating
temperatures envisaged.
There are a number of restrictions on the selection
of alloying elements for use in nuclear applications.
Since the development of nuclear reactors there has
been a continual search for materials suitable for the
manufacture of various reactor components, particularly fuel element cans and pressure tubes. The principal
requirements of these items, especially in reactors using
natural uranium fuel [ 11, are low neutron capture
cross-section, corrosion and oxidation resistance at
elevated temperature, low hydrogen pick-up from
aqueous coolants, high mechanical strength and creep
resistance at elevated temperature and high ductility
for ease of fabrication. Of the four pure metals beryllium, magnesium, zirconium and aluminium which
have acceptable capture cross-sections for thermal
neutrons, zirconium is the most attractive because the
mechanical properties and corrosion resistance are
superior under the reactor conditions which presently
apply.
There are two main incentives to improve the strength
137
138
L.H. Keys et al. /Physical
metallurgy
First, it is essential that the alloying elements do not
nullify the intrinsic advantage of the low capture cross
section of zirconium. Thus dilute alloys only can be
used since most of the useful alloying elements have
high capture cross sections. Also at 500°C the increased rate of oxidation and corrosion requires a
more resistant alloy than those presently employed.
It is generally recognised that for zirconium alloys an
increase in strength leads to lower corrosion resistance.
However, techniques have been developed for cladding
high strength atloys with a corrosion resistant alloy [2]
and corrosion resistance is therefore considered a
secondary factor in the selection of alloying elements
for high strength alloys. The other main requirement
for high strength alloys is that they have sufficient
ductility for both successful fabrication and reliable
performance.
With the exception of the Zr-Nb alloy system [4,5],
very little work has been published on the effect of
individual alloying elements on the microstructure
and properties of zirconium alloys. In this paper, the
possible methods available for strengthening zirconium
alloys will be discussed and, in addition, the results obtained from a study of the phase transformations in
dilute Zr-Cr alloys will be discussed.
2. Group IVA metals
As very little work has been reported on the selection of alloying elements for high strength alloys of
zirconium some of the guides employed in the design
of other group IVA metal alloys, particularly those
of titanium which is on a well established basis, may
be used as a useful starting point.
The Group IVA metals titanium, zirconium and
hafnium undergo allotropic transformations;
the high
temperature body-centred cubic form (@) transforms
to the low temperature, close-packed hexagonal form
(01)at 882,862 and 1310°C respectively. The addition
of alloying elements may either raise or lower the
transformation temperatures of the metals. The alphastabilising elements result in a peritectoid reaction
whereas the beta stabilisers, in general, give rise to a
eutectoid reaction. A review of the physical metallurgy
of titanium 161, which would be expected to behave
similarly to zirconium, reveals that the alloy chemistry
differs and that many of the basic criteria used in the
of high strength zirconium
alloys
design of titanium alioys cannot be applied to the development of high-strength zirconium alloys. The solid
solubilities of metallic alloying elements in both alpha
and beta phases of zirconium are considerably less than
that of titanium because the ratio of atomic radii is less
favourable. In the case of titanium alloys with microstructures based on single phase alpha, alpha plus beta,
single phase beta as well as alpha plus an intermetallic
compound have been developed and are commercially
available. The alpha-beta alloys are by far the largest
group in commercial use. On the other hand, in all
binary alloys of zirconium, except those containing
titanium or hafnium which form continuous series of
solid solutions, or tantalum or niobium, which form
a monotectoid, the beta phase is unstable at ambient
temperatures and decomposes to form an intermetallic
compound. Therefore, the design criteria used for the
majority of commercial titanium alloys, which are
based on the formation of a structure containing
both the alpha and beta phases, give no useful guide
in the selection of alloying elements for high strength
zirconium alloys.
3. Strengthening
methods available in Zr alloys
Published data [3] indicate that zirconium alloys
may be strengthened by the decomposition of the
beta phase to form either martensite (c.p.h,), the
metastable omega phase (hex.), or alpha plus an intermetallic compound. Although the martensite and the
omega phase cause marked strengthening at room
temperature, both these phases are unstable at elevated
temperatures and decompose to form intermetallic
compounds. Therefore it would appear that the only
satisfactory method of improving the strength of
zirconium alloys for use at higher temperatures in
nuclear reactors is by the precipitation of an intermetallic compound 161. The properties of the precipitate considered of importance are (i) mechanical
strength, (ii) thermal stability (iii) morphology and
(iv) distribution. With respect to mechanical properties
of the precipitate, its ability to resist shear during
plastic deformation is the main criterion. Thermal
stability of the intermetallic depends on it having a
high soivus temperature and a low solubility in zirconium, The rno~holo~
and distribution of a precipitate is determined by the crystal structures of the
L.H. Keys et al. /Physical
metallurgy of high strength zirconium alloys
precipitate and of the matrix from which it is formed.
However the advantages of particular structures in
determining the morphology and distribution of
precipitates in zirconium alloys are, at the moment,
difficult to assess. Since strengthening by precipitation
of an intermetallic compound in zirconium alloys must
occur by the decomposition of either the martensite
or the omega phase the properties and microstructures
of these two phases must be considered.
The alloys which are based on beta-stabilising elements are considered to show the greatest development potential and the most important are those which
have the highest eutectoid temperatures and lowest
solubilities in the alpha phase. A similar degree of
strengthening may be possible in some zirconium alloys
containing alpha stabilisers but their lack of response
to heat treatment precludes their consideration in this
paper. There are insufficient data available to determine whether the characteristics of the precipitate or
the characteristics of the matrix are the most important
in producing the optimum strength. On the basis of the
above requirements of a high eutectoid temperature
and low solubility the details of the alloying element
systems which may produce intermetallic compounds
with desirable properties are shown in table 1. The
zirconium-niobium
system is included for comparison.
The zirconium-chromium
system was chosen for a
detailed study because chromium fulfils the requirements of a high eutectoid temperature and low solubility
in alpha phase. In addition it has been reported [7,8]
that the omega phase as well as martensite may be produced on quenching dilute zirconium-chromium
alloys.
Thus it is possible in this system to make a direct comparison of the structures produced by tempering the
Table 1
The zirconium-binary
stabilising elements.
Element
Ni
Cr
cu
Fe
MO
Nb
Eutectoid
temp. (“C)
845
835
822
800
780
610
alloys which contain strong beta
Zr-Nb is added for comparison.
Eutectoid
camp.
at. %
LYSolub. at
eutectoid
temp.
Inter
metallic
phase
2
1.7
2.2
4
7.2
17.2
very small
<0.2
-0.2
very small
<0.2
6.4
Zrz Ni
ZrCr2
Zrz Cu
ZrFez
ZrMoz
PNb
139
martensite and the omega phases. A study of the effect of heat-treatment on the hardness and structure
of a series of zirconium-chromium
alloys was carried
out.
4. The zirconium-chromium
system
The phase diagram for the zirconium-chromium
system as given by Domagala et al. [9] is shown in
fig. 1. The eutectoid occurs at a temperature of 835°C
and a composition of approximately one weight percent (1.7 at. %). The eutectoid reaction results in the
transformation
of the beta phase to OL+ ZrCr,. The
solid solubility of chromium in (Yzirconium is very
small and is less than 0.28% at the eutectoid temperature.
In the present investigation a series of alloys ranging
from 0.5 to 2.5% chromium was chosen for study; the
compositions and their respective solution treatment
temperatures in the beta range are shown in table 2.
The effect of cooling rate on microstructure and hardness was studied by furnace cooling, vacuum cooling,
helium quenching and water quenching from the beta
range. The structures produced were examined using
optical and electron microscopy.
5
E
p 900=+P
L
\ -_ A
835’C
800 .
700
Q-=028
LT+ZQ
0
I
1
I
2
WEIGHT
PERCENT
Fig. 1. The Zr-Cr phase diagram
I
3
I
4
5
CHROMIUM
(Domagala
et al. [9]).
Table 2
The beta solution treatment temperatures of the zirconiumchromium alloys.
Composition (wt%)
~_
__
--0.5
1.0
1.5
I.8
2.0
2.5
p Solution Temperature (c”)
1000
1000
1000
1100
1150
1250
4.1. Qmnched structures
The effect of cooling rate from the beta range on
hardness as a function of chromium content is shown
in fig. 2. Furnace cooling produced a coarse structure
of alpha plus ZrCr, intermetallic and these specimens
indicated the eutectoid composition of these alloys to
be approximately 1.5% chromium. Despite the fact
that the specimens were only held ten minutes at temperature the beta grain size for all the alloys was large
and ranged from 0.5 to 1.O mm. As can be seen from
fig. 2 there was little difference in hardness between
these alloys as the chromium content was increased,
However as the cooling rate was increased the eutectoid
transformation was suppressed and the structures became finer with an accompanying increase in hardness.
Water quenching was necessary to suppress ZrCrZ precipitation particularly in the lower chromium contents.
Maximuin hardness values were obtained on the
water-quenched specimens and as can be seen from
fig. 2 the higher the chromium content the greater
the increase in hardness over that of the furnace cooled
specimens. This was particularly true of the higher
chromium alloys where there was a dramatic increase
in hardness at approximately 1.8% chromium. Optical
metallographic examination of the low chromium
water-quenched specimens revealed an acicular structure usually associated with the martensitic transformation in these alloys (fig. 3), however, in the alloys
containing greater than 1.8% chromium as the chromium content was increased the martensite was gradually replaced by what appeared to be retained beta and
the structure of the 2.5% alloy appeared to be large
600
500
1
or
8
2 LOO
E
$w 300
iIj
!!
200
100
WEIGHT
PERCENT
Fig. 2. The effect of composition
CHROMIUM
and quench rate on the hardness of dilute zirconium-chromium alloys.
Fig. 3. Microstructure of the water quenched O.S% chromium
alloy (X 500).
L.H. Keys et al. /Physical metallurgy of high strength zirconium alloys
Fig. 4. Microstructure of water quenched 2.0% chromium alloy
(X 500).
141
Fig. 5. Microstructure of water quenched 2.5% chromium z3lloy
(X 250).
(a)
Fig. 6. Thin foils of the martensitic structures produced in zirconium-chromium
lath martensite.’
l.O%Cr showing lath plus twinned martensite.
alloys (X 25 000). (a) O.S%Cr showing
142
L.H. Keys et al. /Physical metallurgy of high strength zirconium alloys
grains of single phase beta (figs. 4 and 5).
Examination of thin foils of these alloys in the electron microscope showed that in the composition range
OS 1.8%Cr fully martensitic structures were produced
on water quenching. The OS%Cr alloy consisted of
lath martensite containing a high dislocation density
(fig. 6a). On the other hand, the 1.8%Cr alloy contained well-defined plates of fully twinned martensite
with few dislocations. The martensitic structures in
alloys with intermediate compositions showed mixtures of the lath and twinned structures; the proportion
of twinned plates increased with increasing alloy content (fig. 6b).
The structure of the water-quenched 2%Cr ahoy
consisted of a number of large primary plates of
twinned martensite in a matrix of small equiaxed
grains (1500-4000
A diam.) of the omega phase. The
2.5% Cr alloy contained only the omega phase with a
grain size of IZOO-2000 A (fig. 7). Diffraction data
indicated that the individual grains of omega were
crystallographically
related 161. In addition, the 2 and
Fig. 7. Thin foil of the omega phase structure
chromium alloy (X 25 000).
in the 2.5%
2.5% Cr alloys contained a uniform distribution of
oriented, fine precipitates (50 A) of ZrCr, throughout
the omega matrix. No retained beta could be detected.
Hence the extremely rapid increase in hardness from
340 H.V. for the 1.8% ahoy to 550 H.V. for the 2.0%
alloy and 575 H.V. for the 2.5% alloy (fig. 2) can be
explained by the formation of the omega phase on
rapid quenching from the beta region.
4.2. Tempered structures
4.21. M~rtens~t~c
st~ctu~es
Tempering of the alloys at temperatures between
350 and 700°C showed that the effects produced were
strongly dependent on the prior structure. The alloys
which formed martensitic structures on water quenching
(0.5, 1.O and 1.5% chromium) in general, showed rapid
softening in the first 2-4 hours of tempering temperatures exceeding 500°C and then remained at the same
hardness value for up to 24 hours. There was some
evidence to suggest that the 1.O and 1 S%Cr alloys agehardenedat the lower temperatures (350-45O’C) but
the degree of hardening was only minimal. These features are shown in fig. 8.
The precipitation process in specimens aged for 2
and 24 hrs at 600°C was studied using thin foiI electron
microscopy techniques. In the O.S%Cralloy, after 2hrs
at 600°C, the dislocation structure of the martensite
laths had recovered and precipitation had commenced.
The precipitates were approximately 100-300 A diameter and were uniformly distributed throughout the
martensite with no preferred nucleation at plate interfaces. The precipitates showed no appreciable change
in size after ageing for 24 hrs at 6OO’C(fig. 9a). In the
I .O and 1S%Cr alloys nucleation tended to be heterogeneous; nucleation was favoured at plate interfaces
and, in the case of twinned martensite, was very prevalent at twin interfaces (fig. 8b). After ageing at 24 hrs
the precipitates had assumed a rod-like morphology
(fig. SC) and in the 1.5%Cr alloys areas of the plate
structure had begun to recrystallise (fig. 9d). The distribution density of precipitates was very similar in
the 0.5, 1, 1.5% Cr alloys but the size of the precipitates increased with chromium content. In contrast
to the above behaviour, in the higher chromium alloys
nucleation was homogeneous at low temperatures
froming a fine, uniform distribution of precipitates
(50 A diameter), (fig. 10).
143
L.H. Keys et al. / Physicd metallurgy of high strength zirconium alloys
I
I
0.5 Percent
I
Chr omi u m
(4
1 *O Pctceht
Chromium
I
(b)
1.5 Percent
Chromium
(cf
8
12
16
20
24
AGflNG TIME (Hours)
Fig. 8. The effect of ageing on the hardness of the water quenched alloys. (a) 0.5%Cr. (b) l.O%Cr. (c) 1.5%Cr.
4.2.2. 77~ o~egast~c~re
The ageing of the omega phase was studied in the
2.5% Cr alloy so as to avoid the complication of the
martensite present in the 2%Cr alloy. The ageing pro-
CUSS was followed by tempering at varying times and
temperatures in the ranges, I-48 hrs, 350-7OO’C.
Significant changes in hardness occurred on tempering
the 2.5%Cr alloy. The effects produced are shown in
LX
Keys
et al. / physical
metallurgy
o~high strength z~~~o~~~~rn
allqys
(a)
(4
Fig. 9. The effects of ageing at 6OO’C cm the precipitation of the ZrCrz intermetallic from the martensitic structures (X 25 000).
(a) O.S%Cr aged for 24 hours. fb) I.O%Cr 2 hrs. 600°C showing nucleation at twin interfaces. (c) 1.5%Cr 24 hrs. 600°C showing
rod-like morphology of ZrCra. (d) l.S%Cr 24 hrs. 600°C showing areas of recrystallisation.
145
L.H. Keys et al. /Physical metallurgy of high strength zirconium alloys
fig. 11. At the lowest temperature used (350°C) an
initial slight decrease (575 + 550 HV) in hardness occurred in the first hour of heating; the hardness then
remained unaltered for up to 48 hours exposure. At
450°C the initial rapid decrease in hardness to 500 HV
was followed by a period of approximately one hour
where the hardness remained constant. A second decrease to 325 HV occurred over a period of 8 hours
and the hardness then remained steady for up to
48 hrs. At higher temperatures the as-quenched hardness of 575 HV rapidly decreased in the first hour of
tempering and then decreased more slowly until a steady
state was reached and no further significant softening
occurred after prolonged heating. As stated previously,
after quenching from the beta field, the 2.5%Cr alloy
consisted of a matrix of omega phase with a fine grain
structure and a uniform distribution of fine particles
of the ZrCr, intermetallic. Ageing of the alloy resulted
in the progressive transformation of the omega phase
to alpha. At 350°C no apparent change in structure
occurred after 48 hr., whereas complete transformation
Fig. 10. Thin foil of 1.5%Cr alloy aged 8 hours 350°C, showir ‘g
uniform distribution of precipitates (X 25 000).
100’
0
I
I
I
4
8
12
of the omega to alpha phase was evident after 4 hr.
exposure at 450°C. The transformation of the omega
occurred by nucleation of aipha platelets at the omega
’
16
1
fsdc o_
t
I
550-C
6OO’C iz
t
7oo.c @_
‘36
48
AGEING TIME f Hoers)
Fig. 11. The effect of ageing on the hardness of the water-quenched
2.5% Cr alloy.
146
L.H. Keys et al. /Physical metallurgy of high strength zirconium alloys
(4
grain boundaries and subsequent growth of the alpha.
The initial structure after complete transformation
from omega to alpha consisted of a fine array of alpha
grains containing ZrCr, intermetallic which had not
changed in size or distribution (fig. 12a). Prolonged
ageing (48 hr) at 450°C resulted in the growth of the
alpha grains to form plates and at the same time the
ZrCr, particles grew from 50 a to a final size of 100 8.
The structure of specimens aged for 2 and 24 hr at
600°C were also examined. After 2 hr. at this temperature the structure consisted of an array of very
fine plates of alpha, containing precipitates. Continued
ageing caused the alpha plates and the precipitates to
grow (fig. 12b).
5. Discussion
(b)
Fig. 12. Thin foils of the structures resulting from the ageing
of the omega phase in the 2S%Cr alloy (X 25 000). (a) 4 hours
at 450°C. (b) 24 hours at 600°C.
The two factors which have the greatest effect on
the structure and hence the strength of zirconiumchromium alloys were chromium content and the rate
of cooling from the beta solution temperature, the
strength of these alloys was increased by increasing
the chromium content and also by increasing the
quench rate (fig. 2). The higher chromium contents,
however, were only an advantage if relatively high
quench rates were used, that is when either martensite
or omega phase was formed. By far the greatest
strengths were associated with omega phase formation.
The observation that the structure of the martensite
produced changed from lath to the twinned type with
increasing solute concentration agrees with that of
Williams and Gilbert [S] on the zirconium-niobium
system. The hardness of martensite in zirconiumchromium alloys was greater than that of the presently
used zirconium-niobium
alloys, for example, the hardness of the martensite produced in a 2.5% niobium
alloy was 264 H.V., whereas the same hardness was
measured on the 0.5% chromium alloy.
Thin foil electron microscopy of the Zr-Cr alloys
has confirmed the presence of omega phase in the 2.0
and 2.5% water-quenched specimens and that a rapid
quench rate is necessary for its formation. However
the beta to omega transformation in these alloys appears to be unique. In all other titanium and zirconium
systems studied to date the omega phase has been
present as a precipitate in a matrix of retained beta
phase [lo]. In the Zr-Cr system no retained beta could
be detected.
L.H. Keys et al. /Physical
147
metallurgy of high strength zirconium alloys
For high temperature applications the as-quenched
structures are of little value as they are unstable, hence
the structures of most interest are those produced by
tempering and in the particular case of nuclear applications where the envisaged temperature range is 450500°C the structures produced by tempering at 600°C
are of most interest.
The hardness variations observed on tempering can
be satisfactorily related to the structural changes, particularly in the higher chromium (2.0 and 2.5%) alloys.
In these alloys, where the as-quenched product is principally omega phase, the softening on tempering is
attributed to the transformation
of omega to alpha.
However, the relationship between hardness and structure in the martensitic alloys is less obvious. In the
martensitic alloys the microstructural changes which
may occur on ageing are (i) a reduction in dislocation
density (ii) precipitation of the intermetallic and
(iii) recrystallisation of the plate structure. The results
indicate that the dominating strengthening influence
is the martensitic structure. In the 0.5% Cr alloy where
no age-hardening was observed and the softening effects were small at higher temperatures, the effects of the
reduction in supersaturation and the decrease in dislocation density of the lath martensite outweighs any possible increase in hardness due to precipitation. On the
other hand, it was only in the 1 .O and 1.5% Cr alloys,
where the martensite is twinned and a change in the dislocation density has little effect, that age-hardening was
observed at low ageing temperatures. This was caused by
the formation of a uniform distribution of very fine
precipitates. During the initial stages at higher temperatures nucleation was heterogeneous and no agehardening effects were observed and the subsequent
marked softening of the martensitic alloys can be
directly attributed to the recrystallisation of plate
structure. Because of the unique nature of the omega
phase formation in the Zr-Cr alloys the structures
produced by tempering the omega phase also showed
a number of unique features. The drop in hardness on
tempering at temperatures greater than 400°C can be
attributed to the decomposition of omega to form
alpha. At 500°C the transformation to alpha was
extremely rapid. Continued ageing resulted in the
growth of the alpha into plates and at the same time
the ZrCr, particles grew in size. The structures produced by tempering the omega phase however had a
very much finer alpha grain size and contained a much
Table 3
The hardness of the water-quenched
Alloys tempered 24 hr. at 600°C.
Chromium Content (wt%)
-~
0.5
1.0
1.5
2.5
and tempered alloys.
______
Hardness (HV 200g)
_.
215
230
180
215
distribution
of finer ZrCr, precipitates
than those produced by tempering the martensite
(compare figs. 8c and 1 lb). This is reflected in the
hardness values for the different alloys tempered
24 hours at 600°C and shown in table 3. The minima
at 1.5% chromium is caused by partial recrystallisation of the matrix plate structure in this alloy. (fig. 8d).
As pointed out previously precipitation hardening
was thought to be the most likely method of producing
high strength at elevated temperatures and that the
desired distribution could be obtained by precipitation
from either the martensite or beta plus omega structures. The results of the zirconium-chromium
work
suggest that for this system, the structures produced
by the martensite and the omega transformations are
controlling and that the effect of precipitation is onl:
secondary. For example, the plate structure, which is
the major source of strengthening in martensite, was
stable at 500°C and was still the major influence on
the room temperature strength of the tempered alloys.
It is considered that the relationship between the influence of the plate structure and the precipitate structure on properties will still hold at higher temperatures.
more uniform
6. Conclusions
Zirconium alloys are not effectively strengthened
at room temperature by the formation of the martensite and omega phases. The phases can only be formed
by rapid quenching from the beta phase region. In the
Zr-Cr alloys it has been demonstrated that the beta to
omega transformation can be achieved without the
retention of beta. In these alloys the omega phase had
extremely high strengths but low ductility and is only
stable at temperatures below 400°C. The martensitic
structures have lower strengths and higher ductilities
148
L. H. Keys et al. /Physical metallurgy of high strength zirconium alloys
but they too are unstable at temperatures above 400°C.
Both the martensite and omega phases decompose on
ageing to form the equilibrium products alpha and
ZrCr, ~ntermetalli~. However, the structures produced
by tempering the omega phase are considered to be
superior for use at elevated temperature-to
those produced from the martensite. It is considered that these
tempered omega structures show greater potential
for high temperature applications than the commercial
alloys at present in use.
References
[l]
[2]
[3]
[4]
[5]
(61
[7]
[8]
[9]
[lo]
H.H. Klepfer, tl.S.A.E.C. Report G.E.A.P. 3979.
CD. Williams, Reactor Technofogy 13 (1970) (2) 147.
D.L. Douglass, Atomic Energy Review I (I 963) (4) 71.
J. Winton and R.A. Murgatroyd, Electrochem. Tech.
4 (1966) 358.
CD. Williams and R.W. Gilbert, Trans. J.I.M. 9 (1968)
(Suppl.) 625.
G. Johanson, Ph.D. Thesis, U.N.S.W. (1972).
H. Demars and J.P. Givord, Zirconium 68, Conf. On
the use of zirconium alloys in nuclear reactors, Marianske,
Lazne, Czechoslovakia, Oct. (1968) p. 255.
W.M. Rumba11 and F.G. Elder, J. Less-Common Metals
19 (1969) 345.
R.F. Domagala, D.J. McPherson and M. Hansen, J. of
Metals 5 (1953) 279.
B.S. Hickman, J. Mat. Sci. 4, (1969) 554.