Phase transformations and equilibrium structures in uranium-rich niobium alloys

JOURNALOFNUCLEARMAT~RIA~l2,No.3(l964)

PHASE

291-364,NORTH-HOLLANDPUBLISHINGCO.,AMSTERDAM

TRANSFORMATIONS URALS-~CH C. D’AMATO,

SNAM,

Laboratori

The microstructures

obtained

their distribution, microscopy

in uranium 0.2-7.2

de sa persaturation.

and

meme region

1964

structures;

une

Un refroidissement

donne

des structures

lent ri partir de la

(a + jjr), les phases a tem-

perature ambiante quel que soit le mode de refroidis~ment

the lower the trans-

apres tran~ormation

the finer the structure.

(@ + rr) and (a + rr) structures

Italy

obtenues &ant les mdmes que celles qu’on obtient

of the y phase above the aCr, tempera-

temperature

form 28 February

Milano,

structure form&e de phase o! mais dans deux Btats different6

and X-ray diffractometry.

ture results in two-phased

S. Donato Milanese,

Une trempe a partir de la region (/3 + rr) produit

wt”/b

commencing

The phases obtained,

IN

and T. B. WILSON

1963 and in revised

were studied using optical and electron

~ansformation formation

Studi e Ricer&e,

15 October

Nb alloys by a variety of heat treatments from the y phase, are described.

F. S. SARACENO

Riuniti

Received

AND EQUILIBRIUM STRUCTURES NIOBIUM ALLOYS

de la phase y dans la region (a f

yr).

Les phases mktastables a,‘, aa“ et y residue1 apparaissent

are obtained

suivant la concentration

in their

en niobium

par trempe a l’eau It

respective fields, while (a + yr) structures are formed in the

partir de la phase y. Le vieillissement

(a + ~a) field. The jjl phase, which is metastable,

dans le domaine

(a + ye) est typique,

y2

phase d’equilibre

y2 &ant precedee par celle de la phase r,

in

m&a&able.

poses by rejection

decom-

of excess uranium to the equilibrium

phase on extended

annealing at elevated

~rn~ratu~s

the (a + y2) field. Quenching structures, different

from

both

the

phases

degrees.

being

Slow cooling

(a + yi) structures, ~rn~rat~e

(/? + yr)

depending

quenching

Die Mikrostrukturen,

duplex

supersaturated

to

from this field results in

of the method

at room

of cooling

after

upon

the niobium

content,

the formation

being preceded

by water

of the equilibrium

On decrit les microstructures contenant

de 0,2 a 7,2%

yz phase

wurden mit

und dem Elektronen-Mi~oskop

sowie mit

ergab

Strukt~en

suite de traitements

thermiques

obtenues

der y-Phase

oberhalb

Zweiphasen-Strukturen.

dans lea alliages

iiberschiissiges

a la

zerfallt

Durch

Abschrecken

par microscopic

verschieden

gebildet

de la phase y au-dessus de la tempera-

~ber~ttigt

beides

sind.

sen, die bei Raumtemperatur

auftreten,

der y-Phase in den (a + y,)-Bereich.

est, plus basse.

Die metastabilen

Les structures

(/l + rr) et (a + ~1) sont observkes dans

Ieur regions respectives

Abh~n~igkeit

alors que les structures (a + ;j,) se

dans la region (a + ya). La phase F1 metastable

decompose

se

en rejetant l’uranium en exoes et donne la phase

ya d’equilibre

a-Phasen,

Langsames

est d’autant

de transformation

par chauffage prolong6 B temperature

dans la region (a + yz).

vom

im

wobei ausgewerden

die

aber

Abkiihlen

aus

dieselben Phaunabhingig

von

nach der Umwandlungsmethode

aa’-, ab’*- und Rest-y-Phasen werden in Niob-Gehalt durch dbechrecken in

Wasser aus dem y-Bereich erhalten. Das Altern der a,‘- und as”-Phaaen

im (a + ya)-Bereich

Hartungserscheinungen,

&levee

ihren

wurden. Bei

aus dem (/l + y,)-Bereich

erreicht,

der Abktihlungsmethode

plus fine que la temperature

in

(a + y,)-

$,-Phase,

diesem Bereich ergibt (a + y,)-Strukturen,

La transformation

die

Temperat.uren

die metastabile

ture i!f, produit une structure a deux phases; la structure

forment

wurden w&rend

in der Gleichgewichtsphase

Uran

Duplex-Strukturen

des rayons X.

niedriger

desto feiner ist die Struktur.

schieden wird.

8, partir de la phase y. Les et par diffraction

der .&f,-Tem-

Je

Anlassen bei bestimmten

phases obtenues ont et& Btudiees air& que leur ~stribution optique et l’electronique

untersucht.

in dem (a + l+Bereich

ausgedehntem

en poids de niobium

werden beschrieben.

dem optischen

(a + y,)-Bereich U-Nb

werden,

(B + rr)- und (a + y&Strukturen entsprechenden Bereichen erhalten,

F, phase.

Gew. %-Niob-

in der y-Phase beginnende

Phasen und deren Verteilung

Umwandlungstemperatur,

of the aa’ and ah”

by t,hat of the metastable

erhalten

de la

Die erhaltenen

peratur

phases in the (a + yz) field results in typical age-hardening phenomena,

Warmehandlungen

la formation

die in Uran 0.2-7.2

durch verschiedene

Die Umwandlung

aa” and retained y are ob-

from the y phase. Ageing

Legierungen

R~ntgenbeugungs-Methoden

of the y phase in the (a + yr) field.

The met&stable phases a,‘, tained

a, but

produces

the same phases occurring

i~de~ndent

transformation

field

des phases aa” et ab”

gewichtsphasen vorausgeht.

291

ergibt typische

wobei der Bildung

diejenige

der

Alterung-

der y,-Gleich-

metastabilen

y,-Phase

292

1.

C. D’AMATO,

F. S. SARACENO

Introduction

AND

T. B. WILSON

Since it is desirable to irradiate a structure which

In view of the present interest in uranium-rich

is known to be thermally

stable, most interest has

alloys for use as nuclear fuels, a joint research pro-

been centred on the heat treatments

gramme is being carried out by Agip Nucleare,

obtain

experimental

(the

the irradiation

work being done by the SNAM-Labo-

A number

ratori Riuniti Studi e Ricerche associated company), The Nuclear Power Group and the Comitato nale Energia

Nucleare,

uranium-rich

alloys

Nazio-

in which the behaviour

under

neutron

irradiation

As background

i.e. second phase from other pro-

element.

perties

for this work, detailed investiga-

of uranium-rich

either molybdenum1-3)

binary zirconium

present paper is introductory the investigation

alloys

pro-

containing

or niobium.

The

to a series dealing with

of uranium-rich

and is concerned structures

of the diagrams

niobium

alloys,

with the different types of micro-

obtainable

by various

heat treatments.

775oc U-Nb-

Oiayram to

Royrrs

et

will not be attempted,

but

affect the present work. Although

there

is general

these two diagrams, uranium-rich

notably

agreement

between

there are, particularly

end, a number of discrepancies,

are probably

a reflection

both

at the which

of the techniques

the transformations,

sluggish in uranium-niobium

these being alloys, and of

the different levels of impurity content in the alloys. There is some disagreement transformation

temperature

ture was 665 i

775

the /l + c(

therefore

as to

occurs by a peritectoid

reaction. For example Pfeil et al.*)

or by a eutectoid

found by bracketing

a17

regarding and

whether this transformation

780

according

dia-

will be made in so far as the differences

used to investigate

tions have been made into the metallurgical

equilibrium

recent ‘9 “) are shown in fig. 1. A critical apprecia-

are being irradiated to differentiate

perties of the alloying

range.

of uranium-niobium

tion

position and similar structures in samples of differ-

from those deriving

temperature

to

state in

is

Different structures in samples of the same com-

the effects due to microstructure,

necessary

the equilibrium

grams have been reportet14-*), and two of the more

comparison

ent composition

having

of

being studied.

distribution,

structures

techniques

that the tempera-

3 “C, but placed

it at 667 “C on

DC U-

,

Nb - Diagram

according

740

700

-

I

I I 62a

d+

6.2%

-56+04%

x2 J

0

2

Weight

4

Per

cent

6

8

62C

Wright

niobium

Fig. 1.

Uranium

- niobium equilibrium

I

I

2

diagrams.

4

PQP

cent

6 ntabium

a

PHASE TRANSFORMATIONS

AND EQUILIBRIUM

the basis of metallographic and Muellers), gave 668 f obtained

analysis results. Dwight

also using

bracketing

2 “C, while Rogers

rising temperature

STRUCTURES

techniques,

et aL7), employing

resistometry

and clilatometry,

The alloys fabricated

although

the transformation

to occur, in the range 662-670

appears

“C, it is important

dus temperature

sidered

1.3 “C! by Blumentha19)), ther the reaction

since this determines whe-

would

or eutectoidal.

A

lead to the formation

of

of y1 in a obtained

precipitate

distribution

from

/3, and

since

by slow

second

may affect the irradiation

this phenomenon

phase

behaviour,

in the amount

eliminate the formation

difference of niobium

between required

to

of the p phase, the extremes

et al. (4.6 wt y. Nb).

obtained

in both

made corrections

Similar values were

investigations,

but

et al.

Pfeil

to allow for the reaction

of nio-

This limit and the shape of the y (b $ JJJ boundary are important the

in that, for any particular alloy, they

quantities

of y1 and

B obtained

by

annealing in the (/l + rl) field. 2.

Experimental

2.1. ALLOY

of induction-melted necessary,

induction

was sufficient

to eliminate

The level of the major metallic impurities

Homogenisation

of the alloys was carried out at by furnace

The heat treatments transformation structures

obtainable

1) Controlled

employed

behaviour

to investigate

of the

alloys

the

and

the

in them were as follows:

slow cooling

from the y phase field

to the (a + yz) phase field. After being held for were cooled at

4 “C/min to room temperature. 2) Isothermal

treatments

in the

,d range.

After

being held for 1 hour at 850 “C, specimens were to

a

prescribed

temperature

in

the

ing into a lead bath, or slowly by furnace cooling at a controlled

rate of 4 “C/min). The specimens at temperature cooled

powder,

either by water quenching

(see table 1) were pre-

the latter being cold compressed and minimise

melting. The major impurities

into pel-

loss during

of the raw materials

were as follows :

and

then

to room

for up to 50 temperature

or by furnace cooling.

N. B. The treatment involving

slow cooling both

from the y phase to the transformation

tempera-

ture and from the latter to room temperature

Fe 70 ppm, Al 15 ppm, Si 15 ppm, C 400 ppm

Niobium:

cooling.

2.2. HEAT TREATMENT

hours

Uranium:

in the

40 ppm Al.

pared by melting together uranium bar and niobium fusion

ma-

of the metals.

were maintained

lets to facilitate

of the

(p + rl) phase field, (either rapidly by quench-

Methods

alloys

action

con-

alloys was in no case greater than 90 ppm Fe and

cooled

PREPARATION

Uranium-niobium

Re-

was not

since the mixing

furnace

crosegregation

alloys

1 hour at 850 “C, the specimens

bium with the carbon present in their alloys.

define

and over-

the

given being those of Pfeil et al. (3.4 wt % Nb) and Rogers

of the 7 wt %

these were remelted

950 “C for one week followed

may be important.

There is also a significant diagrams

(given as 667.7 5

is peritectoiclal

reaction

transformation

because of the rapid rise in liqui-

turned six times to ensure macrohomogeneity.

transformation

a dispersed

up to 5 wt % Nb were

with niobium content, vacuum arc

Nb alloy buttons; melting

in pure uranium

293

ALLOYS

in the form of pins by vacuum induction

melting. However,

to know whether it occurs above or below the a//l

eutectoid

containing

NIOBIUM

melting was used for the preparation

664 “C.

Therefore

IN URANIUM-RICH

Ta 500 ppm, Zr 100 ppm. TABLE 1

Wt o/0 Niobium in the alloys investigated

has been termed step-annealing. The isothermal

treatment

temperatures

were

740”, 730”, 715”, 690”, 680” and 675 “C. treatments in the (a + yl) phase 3) Isothermal field. The treatments were carried out as in 2) above,

the temperatures

selected

being 660 “C

and 650 “C. ym 0.19 0.58 1.01 1.97 2.59 3.203.705.20 7.2

5

4) Isothermal treatments

treatments

in the (a + yz) field. The

were carried out as in 2) above,

temperatures

the

selected being 620 “C and 630 “C.

C. D’AMATO,

294 5) Quenching

F. S. SARACENO

from the y phase and ageing in the

AND

T. B. WILSON

particular

temperature

(ct + ya) phase field. Specimens were water quench-

degree of undercooling

ed after

holding

for

in the

(,!I’+ yl)

field the

of the y phase would dimi-

1 hour

at 850 “C, and

nish resulting in a decreasing rate of transformation

then aged for times increasing

from 30 seconds

with increasing

up to 500 hours at either 620”, 550” or 450 “C. 6) Thermal

stability

field. Long

tests in the

term annealing

(a + yz) phase

was carried

out at

niobium

Lateral growth of the /? phase during the transformation

caused

containing

1 wt o/0 or less Nb, to decrease progress-

620”, 550”, 450” or 250 “C for up to 1000 hours

ively in thickness,

on selected samples which had previously

of

been

heat treated as above.

rods

of different

samples

were

examined

the results of all the experimental

servations

are not reported,

most significant

only those considered

being given.

ever refers to the complete

ob-

The discussion

work.

clearer explanation

small

became

of the structure.

content,

to

(fig. 2 b) as the

the /3 phase

permit

of some of the structures

a ob-

from the y phase to room

temperature.

In alloys of

since the ,B/yr ratio was

the

joining

up

y +

(/3 +

obtained

at temperature

yr) transformation

was

EXISTENCE

OF THE

FIELD

4.6 wt. %Nb,

,!?PHASE

alloys

eliminated

in alloys

less than

of ,!? phase being entof higher niobium

con-

tent.

a) The temperature the

higher

isothermal

treatment

carried

(/I + yr) field, both by quenching from the y phase, transformation

showed

occurs

boundaries then

into

in the

that the y + (j3 + yI) and growth.

at the original y grain

and at the non-metallic

extended

out

and slow cooling

by nucleation

The ,Q phase was nucleated

the y grams

lamellae having a Widmanstatten

structure

temperature

and the

smaller

the the

coarser

content

c) The time of holding

of the alloy ; the matrix

of the non-matrix

inclusions

and

in the form

of

type of distribu-

temperature differences different

the coarser the structure.

produced cooling

the structures

Al-

this may have been due to an associated

was more equilibrium

likely

a result

transformation

with increasing

niobium

in the y phase, it

of the decrease temperature content

in the

y/(/I + yl)

(fig. 1). At any

by

treat

were, however, very slight.

of cooling

to room temperature

of

formed in the (p + yl) field had on

the contrary a significant effect upon the structures observed at room temperature. Cooling under equilibrium

(a + yr) temperature

content.

The

in the final structure

rates to the isothermal

ment temperature The method

treatment

in the (B + yJ field; the slower the

according

niobium

in the (,!? + yr) field; the

phase.

decreased

decrease in the rate of diffusion

of I wt y0

d) The rate of cooling to the isothermal

ture in the (/I + yr) field the speed of transformation though

of /j’

longer the time the greater the spheroidisation

duce an increase

increasing

the

quantity

formed.

tion (fig. 2 a). At any particular treatment temperawith

de-

of holding in the (p + yl) field;

the

rate of cooling

The

aft,er the

contents.

containing

the formation

B

and less, and the y1 phase for higher niobium

OF

y + (p + yr) is permissible only

in uranium-niobium irely

TO THE

to the diagram of Rogers et al. (fig. 1)

According

the transformation

the

complete

being the ,B phase for concentrations RELATING

of

lamellae, the y1 phase remained the matrix (fig. 2~).

b) The niobium

3.1. STRUCTURES

p

the

pended upon :

first, as this facilitates

by slow cooling

globules and

matrix

niobium

in the alloys

and finally fragment into strings

and

The structure

how-

Structures relating to the field of existence of the @ phase are considered

too

the y1 lamellae,

coalesced

continuous higher

Since some hundreds

tained

small

lamellae

Results

3.

content.

conditions

in the quantity

to the phase diagram is reached,

should pro-

of the fl phase

until the (,!!I+ yI) at which tempe-

rature the p phase transforms to y phase. According to Rogers eutectoid

et al. the transformation

occurs

by a

y1 phase being precipitated

in

the a phase as a result of the lower solubility

of

niobium

reaction,

in a than in ,9 at this temperature,

while

PHASE TRANSFORMATIONS AND EQUILIBRIUM STRUCTURES IN URANIUM-RICH NIOBIUM ALLOYS

a) U-lwt%Nb

690 “C - 15 min

0) U-2wtyoNb

Fig. 2.

675 “C -

675 “C -

3 hours

3 hours

Uranium - niobium alloys cooled rapidly from the y phase (850 “C) to the /I + y1 field, held for a predetermined time and then water quenched. (500 x )

according to Pfeil et al., the reaction is not eutectoidic but peritectoidic and should result in a in the quantity of y1 in the structure after the reaction is completed. Further cooling should result in rejection of a from y1 until the latter arrives at the monotectoid composition. At this temperature the monotectoid reaction should take place with decomposition of the y1 phase into a and yz phases. Although both the y1 a,nd ya phases have bee structures they differ in that at the monotectoid temperature the former has a lattice parameter of 3.48 d and contains approximately 6 wt o/o Nb while the latter has a decrease

b) U - 1 wt% Nb

2%

lattice

parameter

of 3.34 A and contains approxi-

mately 50 wt oh Nb. The most recent phase diagrams’? 8, indicate that the quantity of bee phase present as yz in hypomonotectoid alloys, after the reactions is complete, should be only one-eighth by weight of that existing as y1 before reaction. Assuming that the densities of uranium-niobium alloys in equilibrium obey a simple mixtures law, it can be calculated that the above ratio is one sixth by volume. A considerable decrease in volume of bee phase should therefore be expected as a result of the monotectoid reaction.

C. D’AMATO,

296

F. S. SARACENO

Furnace cooling at 4 “C/min of the structures ob-

AND

T. B. WILSOK

diagram, it was deduced that the monotectoid

reac-

t,ained in the (p + yl) field was too rapid to allow

tion was very sluggish and had not gone to comple-

all of t’he above

t,ion during cooling from t’he (,8 + yl) field to room

fact

increase

changes to go to completion.

in the

amount

of fi phase

cooling to the p + a transformation to the extent

expected.

did not occur

Precipitation

noted in the form of very fine particles dia.

0.1 p.) in the

a phase,

/J + (a + yl) transformation

In

during

Nb alloys step-annealed

at 675 “C for 50 hours was

3.450 a

to

that

the

(fig. 9a).

urt,her particles

react’ion

was not,

during the cooling to room temperature

from the (p + yJ field. Monotectoid

took place by rejection

of the y1 phase, which

of a and enrichment

of the

remaining y1 phase in niobium, increased the degree

Kg. 3.

of

a niobium

content’

of

This

non-equilibrium

phase

having

a niobium

cont!ent much lower t)han t,hat of the equilibrium yz phase has been designated transforms

1/1 and as shown later.

slowly to yz on extended

ageing in t’he

region of the (a -I- yz) field.

It must be pointed out, that the a paramet,er given by X-ray

diffractometry

refers to that of the 7,

phase present as a result of the monotectoid

decom-

Uranium - niobium alloys cooled rapidly from the “Jphase (850 “C) to the/l + ^J~ field (675 “C), held for 3 hours and then cooled slowly to room temperature. (500 s )

fragmentation

and

spheroidisation

lamellae in the alloys containing (figs. 3a and 9a), and produced

of

the

y1

1 wt y0 or less Nb, a lamellar structure

of a and a bee phase around the islands of a (ex ,B) in the

corresponding

wt, y0

12 wt %.

high temperature

decomposition

which showed that

(maximum

indicating

that, hhe precipitation

completed

diffractometry,

by the

the a parameter of the bee phase in the 2.0-3.7

Annealing of t,his structure in bhe high temperature showing

results of X-ray

was confirmed

of y1 was

is eutectoidic

region of t’he (a + yz) field produced

t,emperature. This deduction

alloys

containing

2 wt

o/0 or more

Nb,

occupied

by

the bee phase in these alloys at room temperature was much greater than that predicted

duced

by the eutectoid

insufficient

reaction

to permit examination

by the phase

of bee phase pro/? + (a f

rl)

is

by conventional

techniques. For convenience,

however, the precipitate

ing from the latter reaction

(figs. 3b and 9b). Since it was found that the volume

of yl, since the quantity

position

result-

has also been referred

to as rl. The phase changes involved during the isothermal transformation

of the y phase of uranium-niobium

PHASE TRANSFORMATIONS

AND EQUILIBRIUM

STRUCTURES

IN URANIUM-RICH

NIOBIUM

ALLOYS

297

alloys in the (p + ri) field and subsequent slow

be suppressed in the alloys containing 1 wt yO or

cooling to room temperature can be s~ma~ed

more Nb.

schematically as follows :

The transformat,ion when the @ phase was suppressed appeared to be of the nucleation and growth type, eharacterised by very high rates of both

Isothermal

/

transformation

processes. Consequently the structural characteristics were different from those when ,8 phase formation was involved (see sect’ion 3.1 above) in that

MonotecOoid reaction

Eutectoid reaction

a + y1 (precipitate)

1 & + a (lamellaef

Water quenching after isothermal transformation in the (p + ri) field suppressed the eutectoid and monotectoid reactions described above, and resulted in the transformation of the /l and y1 phases into supersaturated a phases (figs. 2 b and Zc), the degree of supersaturation depending upon the solubility of niobium in the p and y1 phases at the temperature from which the alloy was quenched. These supersaturat,ed phases have been termed u’. The scheme of phase changes occurring during this heat treatment can be indicated as follows:

Isothermal

1 transformation

Fig. 4.

Urssnium - 2

x-t% niobium alloy cooled rapidly

from the y phase (850 “C) to the a + y1 field (655 “C) held for 15 minutes and then

water quenched.

(500 x )

the struchures were very fine (fig. 4) with a tendency for the lamellae to become increasingly acicular as the niobium content increased. An important feature of these structures was that the continuous phase was always a independent of the niobium content. It is still uncertain whether this transformation occurs “directly” by the rejection of a phase from t,he y phase with consequent enrichment of the

~a~ensiti~ transformations I

1 I

a’

(slightly supersaturated)

(highly zupcrsaturated)

3.2. I~~THERAxAL TRANSFORMATION

OF y PHASE IN

THE (a -+ yl) FIELD

It was found t.hat, by cooling rapidly from the y phase field to an isothermal treatment temperature in the (M.+ yl) field, the formation of B phase could

latter in niobium to form yl, or whether it takes place “indirectly” analogous to the bainite reaction in steels and involves the rapid formation and decomposition of a supersaturated a phase. The alternative mechanisms of transformation are shown below. “Direct” transformation y + (a + yl) “Indirect” transformation y + supersaturated a + (a + rr). After the transformation was complete and the samples cooled to room temperature the phases

C. D’AMATO, F. S. SARACENO AND T. B. WILSON

298 present

were (a + yl) irrespective

of the cooling

denum

alloys,

the

different

types

depending upon the alloy content”).

rate. 3.3. ISOTHERMAL TRANSFORMATIONOF y PHASE IN THE (a + yl) FIELD When the transformation isothermally

in the

temperature, essentially thermal

the the

(a + yz) field above

structures

same

of

the

as those

transformation

3.2 above).

of the y phase occurred the MS

alloys

obtained

were

by

iso-

in the (a + yl) field (see

The bee phase obtained

at temperature

after transformation

in this case was yl, provided

that the isothermal

holding

The

characteristics

were found in the 0.6-3.7

Acicular struccontent

wt o/o alloys, while the

4.0 and 5.2 wt y. alloys exhibited

banded

struc-

tures. The 7.2 wt o/o Nb alloy gave no response to polarised light. X-ray

diffractometry

showed that, with the ex-

ception of the 7.2 wt o/o alloy, the structures present, were all basically nium, modified

to yz.

the orthorhombic

by distortions

lattice of ura-

induced

by the pre-

in a supersaturated

condition

in

A0

of transformation

of the y

phase in the (a + yz) field were virtually with those of the transformation

structure

in fineness with niobium

sence of niobium

time was insufficient

to allow the 1/1 to transform

tures increasing

of

identical

in the (a + yJ field.

3.4. CONTINUOUSCOOLINQFROM THE y PHASE TO ROOM TEMPERATURE The

cooling

rate of 4 “C/min

phase formation

was such that

was suppressed

p’

in the alloys con-

taining 2.0 wt yO or more Nb, while in those alloys of less niobium decreased

with decreasing

The structures

observed

alloy content. in the alloys containing

2.0 wt yO or more Nb were of the same type as those

observed

0.2

0.6

1 I

content the degree of /? suppression

after

transformation

2 I

2.5

3.7

The lattice

contraction

of the b parameter,

1.0 wt o/o

lamellar

type

increasing

2.0-3.7

of the y1 phase in the

wt y. alloys, was 3.464 & there being no

corresponds

with composit’ion.

This value

to a niobium content of approximately

of

alloys

having

banded

the orthorhombic

(fig. 5).

ment

of the ab angle to values greater than 90”

in the step-annealed

lattice

alloys, reflecting

a closer ap-

by step-annealing.

AGEINC IN THE (CZ+ yz) FIELD All of the alloys in the range 0.6-7.2

wt y. Nb

with the result, that the

of these alloys was monoclinic.

These structures, which can be considered as modifications

3.5. WATER QUENCHINGFROM THE y PHASE AND

in polarised

Since the angles of

cell are all right angles, enlarge-

causes loss of symmetry,

obtained

by

contraction

of the b parameter

structures

light (4.0 and 5.2 wt y. Nb).

9 wt %, which is slightly lower than that obtained proach to equilibrium

content

to contraction

was caused the

enlargement of the ab angle, was found in the lattice

figs. 3a and 9a).

variation

distortion

with niobium

In addition

found after step annealing in the (,!l + rl) field (see

systematic

%N”

of y in the the lattice.

The average a parameter

Nb

Fig. 5. Variation of the b parameter of the orthorhombic lattice of G(uranium with niobium content in uranium niobium alloys water quenched from the y phase (850 “C).

above), while those in the alloys containing

section 3.1 above,

“!

J

10

;

(a + yl) or (a + yz) fields (see sections 3.2 and 3.3 or less Nb were of the fragmented

4

of the normal

a uranium

structure,

are formed

formation

of the y phase.

The lattice

orthorhombic

by a martensitic

type trans-

of the 7.2 wt y. alloy was perfectly

presented single phase structures under bright field

cubic showing that the y phase had been retained

illumination.

to room temperature

The

polarised

those obtained

light

structures

in water quenched

were similar uranium-molyb-

to

behaviour

and explaining

the isotropic

when examined metallographically

polarised light.

under

PHASE TRANSFORMATIONS AND EQUILIBRIUM STRUCTURES IN URANIUM-RICH NIOBIUM ALLOYS

299

Evidently the MS temperature of uranium-rich niobium alloys, which decreases with increasing niobium content, must intersect room temper&t~e between the limits 5.2 and 7.2 wt y0 Nb, confirming previous work12* 13). In agreement, with the nomenclature proposed by Lehmenn and Hillsll), the structures obtained have been termed a,‘, q,” and y, the composition limits of which are given in table 2. TABLE 2 by X-ray diffraction and polarized light??

0.6-3.7 4.0t-5.2 --

I

a’& U”b

I

t extra alloys were fabricated especially to determine the composition limit dividing the a’, phase from the U”b phase. tt the classific&tiongiven is &ccording to that used for uranium-molybdenum alloys by Lehmann and Hillslo), in which a’ refers to an orthorhombic structure with a eontraction of the b parameter, a” indicates 8 monoclinic structure and y a metastable bee phase. The subscripts “a” and “b” indicate the type of structure observed under polarised light, “a” referring to acicular and “b” to banded.

Ageing the alloys containing 0.6-5.2

wt y0 Nb

at elevated temperatures in the (a + yz) field resulted in the decomposition of the metastable phases, giving rise to the typical precipit&tion effects of age-hardening and over-ageing (fig. 6). X-ray diffraction results indicated that the formation of the equilibrium (u + ra) structure was always preceded by that of (a + y#*). 3.6. THERMAL STABILITY TESTS Although morphologically different in some cases, the structures obtained at room temperature, both by slow continuous cooling directly from the y phase and by sop-annea~g in the (p -+ yl) field, were always composed of the (u + j;,) phases. These structnres could not be considered to be in equili-

Fig. 6. Uranium - 2 wt% niobium alloy water quenched from the y phase (850 “C) and aged at 620 “C (500 x ) a) 4 min; b) 500 hours brium despite the very high stability exhibited by the r1 phase at room temperature.

The 7, phase had an a parameter in the range 3.450-3.475 A depending on the initial heat treatment ; these values correspond to a uranium content in the range 87-92 wt %. The equilibrium phase yz has an a parameter of approximately 3.350 is, and contains approximately 50 wt o/oU. Clearly, decomposition of r1 into yz would occur by rejection of excess uranium as u phase, accompanied by a decrease in the a parameter of the bee phase, To study this decomposition, samples which had been previously heat treated as stated above were subjected to extended annealing (up to 1000 hours) at 250”, 400” and 550 “C and subsequently examined

C. D’AMATO,

300 by X-ray

diffractometry,

optical

F. S. SARACENO

microscopy

and

electron microscopy. The X-ray

results from the 2.0, 3.2

figs. 7 and 8. As expected

1

in a diffusion

10 2

10 a)

C

-

2.0

lo3 wt

9;)

500h1000h 81

10"

of this type,

advanced

1000 hours annealing were in the vicinity

1 1

I 10

~~~

3.37

1

I

I

I

102

lo3

104

b)

A0

I’

-

3.2

lh 2h

wt

O;,

1

2oh

Variation

IO5

SOOhlOOOh

102

U - 3.7 wtq,

1 103

1

10

I

I

103

lo*

I min

lo5

wt yc Nb 20h

50hlOOh

/

A 102

500h

lOOOh

“I

1'

1

3.35 10'

lh 2h

A0

min

105

_I

A

9-

/ 103

104

2oh

5$lWh

11

'I'

1

1

Nb

of the a parameter

I

3.47

I

c)

Fig. 7.

min

-.. _

min

lo5

b) U - 3.2 wt yo Pib

5oh100h

I

1 10

--*

Nb

3.47

3 35

-~

soohloooh

..__

102

3.35

I 10

20~ 5o"loo"

lh 2h I1 -1

A'

SOOh lOOOh

after

of that given for ya, while after

! I

3.35

at 400 “C;

at 550 “C the a parameters

a) U - 2.0 50hlOOh

The process was always

lh 2h

3.47

3 351

the rate of transformation

at 550 “C than

1

105

min

more

Nb 20h

lh 2h

A'

in

controlled

20h 5oh100h II 18,

,h 2h I'

I

process

T. B. WILSON

increased with temperature.

diffraction

and 3.7 wt o/0 Nb alloys are shown graphically

A’

AND

10

of the y phase with

annealing time at 250”, 400” and 500 “C, in uranium - niobium alloys previously cooled at 4 ‘C/min from the 7 phase (850 “C) to room temperature.

Fig. 8.

U

-

t--

=I

L

102 c)

500h1000h

3.7 wt”/

103

10'

1 min

lo5

Nb

Variation of the a parameter of the ^/ phase with annealing time at 250 ‘, 400” and 550 “C, in uranium niobium alloys previously step-annealed in the (9 + 13, field (675 “C) for 50 hours.

PHASE TRANSFORMATIONS AND EQUILIBRIUM STRUCTURES IN URANIUM-RICH NIOBIUM ALLOYS 1000 hours at 400 “C they had decreased 3.40 A, demonstrating not yet complete

to only

that the transformation

was

at the lower temperature.

The 1/r phase exhibited

high stability

only very slight parametric

variations

examination

at 250 “C, being found

revealed

transformation

of y1 into yz produced

in the volume

of bee phase present,

that

from

consideration

of

the

This may indicate niobium

system

the existence in the uranium-

of an hitherto

undetected

accompanied process.

Prior to annealing,

the j1 in each sample existed

in one of the three following tions depending

a) U - 1 wt y0 Nb step-annealed

b) U - 2 wt ‘$&Nb step-annealed

in the /J’ + y1 field (675 “C)

for 50 hours.

morphological

upon the niobium

the type of heat treatment:

cooled at 4 “C/min to room

temperature.

6

content

condiand on

in the /I + y1 field (675 “C)

for 50 hours.

c) U - 2 wt ‘$&Nb continuously

Uranium

phase

system.

It was noted however that the decrease in volume

Big. 9.

equilibrium

similar to the yll phase in the uranium-molybdenum the

a decrease

by a more or less marked spheroidisation

expected

less than that

diagram.

in the samples annealed at this temperature. Metallographic

of the bee phase was considerably

301

- niobium alloys heat treated from the y phase (850 “C). (7000 x )

PHASE TRANSFORMATIONS

generation

AND EQUILIBRIUM

STRUCTURES

of y1 phase, caused it to become discon-

tinuous where previously Since no significant observed

after

the

was complete,

as indicated

duced by extended equilibrium

in the (/3 + yl) field (i.e. fragmented The

changes were

by X-ray

diffraction

that the structures

pro-

annealing above 400 “C were in

and were morphologically

a much

in the structures

higher

uranium

present

(approximately

Conclusions

content

alloys

7.2 wt y0 Nb have

been

containing

subjected

to

of different types of heat treatment termine

the

structure

morphological

obtainable

A uniform in 0.2-l

(approximately

variations

in

microlamellae

treatment

“island”

structures

in

wt y. alloys, the u (ex ,!?) islands containing

very fine precipitate being occupied an

of 1/1 and the interisland

by very fine lamellae

interlamellar

spacing

of

space

of rl and a, the

order

of

0.15 t*. The presence

of very fine precipitate,

presuma-

bly pl in the a (ex /3) regions of the structures, leads to the conclusion

that either the p + GC transforma-

tion is of the eutectoid phase diagram significant uranium

type in agreement with the

in solubility

with decreasing

of niobium

temperature.

Water

of bee

quenching

after isothermal

treatment

the p + y1 field produces structures containing

in two

a phase, the degree of super-

saturation depending upon the solubility

of niobium

in the ,5 and y1 phase existing at the treatment

tem-

perature. The formation

of the metastable

phases in uranium-rich quenching

niobium

cc’ , CL” and y alloys

by water

from the y phase has been confirmed,

the results indicating

that the composition

for the 5c,’ phase are 0.6-3.7

limits

wt y. Nb, and for the

wt o/o Nb, while the Ms tem-

ab” phase are 4.0-5.2 perature

of these alloys which

decreases

creasing

alloy

room

content

reaches

with in-

temperature

between 5.2 and 7.2 wt y. Nb. Ageing of the uB’ and a,,” phases leads to typical and overageing

in u

precipit,ation

hardening

effects.

transformation

alloys containing

of the y phase in the

more than 1 wt y. Nb, at tempe-

ratures below that of the p + u transformation above the Ms temperature,

produced

but

fine lamellar

structures of a and Fl, the interlamellar

spacing de-

creasing from 1 p to 0.1 p with increasing alloy content. Similar structures are obtained in 2.0-5.2wt Nb alloys by continuous

cooling

y.

at 4 “C/min from

the y phase. Since this cooling rate is insufficient

to suppress

of p phase in alloys containing

or less Nb, the structures

obtained

1 wt y.

in these alloys,

by furnace cooling to room temperature tially the same as those obtained

Acknowledgements

This point

will be clarified in future work.

the formation

generally by spheroi-

of Rogers et al., or that there is a

decrease

Isothermal

of r1 to yZ. The

phase present.

in the (,!l + yl) field. This

produces

annealing

in the (a + yz) field re-

is accompanied

types of supersaturated

of fragmented

bee yz phase

disation and a slight decrease in the volume

in order to de-

.O wt o/o alloys when heat treated from the

2.0-3.7

with

a number

phase f1 in an u matrix is obtained

y phase by step-annealing heat

to

in these alloys.

distribution

of the metastable

up

above has

50 wt %), and extended

temperatures

decomposition

Uranium-niobium

at room

described

sults in the gradual decomposition 4.

lamellae of rl

90 wt %) than that of the equilibrium at elevated

stable.

bee TX phase

metastable

temperature

y1 + (a + yz) transformation

results, it was concluded

303

NIOBIUM ALLOYS

in an alpha matrix).

it had been continuous.

morphological

IN URANIUM-RICH

are essen-

by step annealing

The authors wish to thank the boards of directors of SNAM-LRSR, support to

CNEN,

publish

it. They

Dr. H. M. Finniston gestions.

TNPG

and IRD

for the

given to this research and for permission are particularly

grateful

They would like also to acknowledge

cooperation

to

for his helpful advice and sug-

in X-ray and electron microscopy

the work

given by Dr. M. Cesari and his staff.

References ‘)

R. L. Craik, C. Fizzotti and F. Saraceno, Nuclear Research Centre (Parsons, UK) Report NRC 60-68 2, R. L. Craik, D. Birch, C. Fizzotti and F. Saraceno, J. Nucl. Mat. 6 (1962) 13 3,

G. H. May, J. Nucl. Mat. 7 (1962) 72

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C. D’AMATO,

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H. A. Saller and F. A. Rough, Battelle

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