On order twinning in AuCu

ON

ORDER

TWINNING

B. HANSSONt

and R.

IN

AuCu*

S. BARNES:

Electron microscopic examin8tion of thin films from ordered AuCu I hes revealed feulted regions which have been interpreted 8s twins in which only the super-lattice is rotated. The development of these “order” twins can be inferred, 8nd it is suggested thet they are generated by the movement of &(lOl) dislocations on successive {lOl} pl8nes. This twinning, 8s 8 mode of stress relief, is discussed. SUR LES MACLES

D’ORDRE

DANS

AU CU

L’examen en microscopic Qlectroniquede pellicules minces de Au Cu B 1’Qtatordonnb 8 mis en Qvidence l’existence de rbgions perturbbes qui ont Bti interpn%es comme &ant des ma&s dans lesquelles le super-r&eau seul 8 subi une rotation. On peut en dbduire le m&anisme de developpement de ces “macles d’ordre”: les auteurs estiment qu’elles sont engendrees par des mouvements de dislocation t(lO1) sur des plans {lOI} successifs. Cette formation de m8cle est discutbe en t8nt que mode de lib&&ion des contraintes. ZWILLINGSBILDUNG

DURCH

ORDUNG

IN AuCu

Elektronenmikroskopische Untersuchungen diinner Folien 8us geordnetem AuCu I zeigen fehlgeordnet,e Gebiete, welche als Zwillinge gedeutet werden, bei denen nur die Uberstruktur gedreht wird. Ferner kBnnen Riickschliisse auf die Entwicklung dieser “Ordnungszwillinge” gezogen werdcn. Es wird gefolgert, da0 die bei der Bewegung von +(lOl)-Versetzungen auf aufeinander folgenden {lOl}-Ebenen gebildet werden. Diese Zwillingsbildung als eine mBgliche Art der Entspannung wird diskutiert.

1. INTRODUCTION

Above

a critical

temperature

together

of about

410°C the

alloy of50 at.% copper and 50 at.% gold forms a f.c.c. structure

(a = 3.878)

with a random distribution

copper and gold atoms.

Below ~410°C

orders to form a super-lattice planes composed

specimen

of

enabled

(002)

thickness

either of all copper or all gold atoms

(see Figs. 1 and 2(a)).

vacua

was then

atmosphere

the structure

with alternating

in

in

an

alumina

heated

by cold-rolling

to about

mm thick, were obtained.

by

in

[OOl]

at 325°C for 72 hr in vacua.

contraction

[OlO] directions

an (the

so that the structure

a-

In

this

ordered

in

and

the

b-axis

becomes

ratio of 0.92 (a = 3.96& boundaries

the

expansion

tetragonal there

and

are

ordered by annealing the foils

The foils were then thinned

respectively),

c = 3.67&.(l)

structure

direction [loo]

This treatment 5076 its original

After a final quench from

750°C the structure’was

and

The

Quenching and rolling were repeated until foils, -0.03

Because of the different sizes of is accompanied

slight

form

with no sign of cracking.

ordering

a

to

copper.

to 750°C in a hydrogen

and quenched into water.

it to be reduced

the copper and gold atoms, (the c-axis)

boat

an alloy of 50 at. o/o gold and 50 at. “/

an electrolyte

consisting

and three parts acetic acid.

with a f a

then examined

electrolytically

of one part perchloric

using acid

Suitable thin films were

in a Siemens Elmiskop

I operating

at

100 kV.

anti-phase

at which the (002) layers are out of step,

i.e. a (002) layer of gold atoms adjoins a layer of copper atoms and vice versa. these anti-phase

terval of five lattice axis.(l)

Between

boundaries

-380°C

and -410°C

occur with a regular in-

parameters

along

the a- or b-

This structure, known as AuCu II, will not be

considered

in this paper.

Below

-380°C

the anti-

phase boundaries occur irregularly (AuCu I structure). This paper reports some observations AuCu I structure. 2. EXPERIMENTAL

Spectroscopically troscopically

of faults in the

TECHNIQUE

pure gold (75.62 wt.%)

pure copper

and spec-

(24.38 wt. %) were melted

* Received August 14, 1963. t On attachment from Fysiska Institutionen, Ups818 Universitet, Uppsala, Sweden. $ Metallurgy Division, Atomic Energy Research Esteblishment, Harwell, Berks. ACTA METALLURGICA, 4

VOL. 12, MARCH 1964

Fm. 1. The unit cell of AuCu I in the orientation of the structure in Fig. 2. 315

ACTA

316

METALLURGICA,

VOL.

12,

1964

The orientations

of the films were determined

from

.electron diffraction patterns of the selected areas. The diffraction distorted distorted

patterns

of Figs.

3 and 4 were slightly

112 patterns and that of Fig. 5 was a slightly 101 pattern,

indicating

that the planes of

those foils were not, quite parallel to the (1122 and the (101) planes respectively.

(The directions

marked on

Figs. 3, 4 and 5 are therefore,

strictly

projections

on the planes of the

foils.)

of those directions

speaking,

the

In some cases the superlatt,ice spots could ,be

used to distinguish

between possible directions.

The planes of the faults were determined graphic projection of 800-1500

methods.

by stereo-

Assuming a foil thickness

A it was found that the (101) plane was

the only low-index

plane compatible

with the orienta-

tion and the width of the faults. FIG. 2(a). A model of the AuCu I structure. 3. EXPERIMENTAL

RESULTS

Figures 3, 4, 5 and 6 are typical photographs of different

regions of a thin film extracted

ordered foil.

taken

from an

In some areas structural faults reminis-

cent of elemental twins were seen. Figure 3 shows such faults which are only about each surface of the film.

3 ,u long but extend

This is presumably

stage of the fault which has developed

to

an early

further in the

region shown in Fig. 4, where the faults in some instances completely in it.

cross the area, and in others end

In Fig. 5 many of these faults end within the

area giving the impression

that the faults are pene-

trating in from the left.

In many areas the structure is

heavily

6) in the

twinned

(Fig.

manner

observed

previously.(24) FIG. 2(c). The three top (111) layers of the model shown in Fig. 2(a) have been sheared by a *[lOi] dislocation slipping on each of them. 4. DISCUSSION

It, has been shown(3*4) already that the { lOl} planes bound twins in AuCu. structural

It is therefore thought that the

faults shown in Figs. 3, 4 and 5 represent

early stages in the formation of these twins. In an ordered structure it is possible to have a type of twin in which the super-lattice which we have called an “order”

is reoriented

and

t’win to distinguish

it from the “normal” twin. A normal twin can be produced, for example, by a partial dislocation with l/6

FIG. 2(b). The AuCu I structure after a *[lOi] dislocation has produced shear on each successive (101) plane above the one indicated. In order to show the change of the (001) face the projection has been cut off and replaced on the other side (the (111) face).

[112] Burgers vectors slipping on successive (111) planes. The passage of this imperfect dislocation places atoms in new positions each of which is a reflection in the the original crystal. only

dislocations

twin boundary of one in In an ordered structure

with Burgers

vectors

representing

HANSSON

ANI) BARNES:

ORDER

TWINNING

IX

AuCu

317

FIG.. 3. An early stage in the formation of order twins in AuCu I. The plane of the foil is close to (121). The faults are on (011).

FIG. 5. Advancing twins. The foil is nearly parallel to (101). The twins are on (011).

FIG. 4. A later stage in the formation of order twins in AuCu I. The plane of the foil is close to (112). The faults are on (011).

FIG. 6. The twinned

structure

of AuCu I.

ACTA

318

METALLURGICA,

displacements between the same kind of atom produce no change in the order when they move. Two thirds of the possible ) (101) Burgers vectors represent displacements between different kinds of atom (see Fig. 1) A dislocation with such a Burgers vector will therefore on moving create an anti-phase boundary in its wake. If a similar dislocation then followed the first, so that a super-lattice dislocation had passed, this anti-phase boundary would be eliminated. If such a partial super-lattice dislocation slips on successive (IOl)planes an order twin can be formed in which only the order is changed, i.e. gold atoms occur in positions which are reflections, in the twin boundary, of the positions of gold atoms in the original crystal. This is ~ustra~d in the model of AuGu shown in Fig. 2(b) where the rearrangement of the original structure (Fig. 2(a)) is equivalent to that produced by the slip of a dislocation with ) [ lOl%jBurgers vector on each of three successive (101) layers. The order, with alternating (002) planes, is preserved but the orien~tion of the super-lattice in space has been rotated 90” so that order twinning has occurred with the (101) plane as the twin plane. (In an actual crystal the rotation of the super-lattice will not be exactly 90” because the unit cell is tetragonal, but (101) will still be a mirrror plane). Equivalent structural changes take place after ~rres~nd~g slip on other (101) planes exept (X10) and (110) which contain only one kind of atom. Slip on these latter planes can therefore not alter the order, and (110) and (li0) cannot be order twin planes. The situation where several dislocations have slipped on successive (101) planes is illustrated in Eig. 7 where the dislocations of opposite sign are shown at the extremities of the twinned region. The rotation of the planes containing only one type of atom by 90” is illustrated by the broken line. Therre will be an anti-phase bounda~ between the outermost dislocation and its neighbour and again between the third dislocation and the fourth, and so on. The boundaries of the faults shown in Figs. 3 and 4 do not all he exactly on (101) planes because the faults are either wedge shaped or lentieular. Therefore they

/

/

/

FIG. 7. A schematic diagram of an arrangement of dislocations on successive (101) planes. The broken lines indicate the orientation of the alternating layera of Au and Cu atoms. These layers are perpendicular to the plane of the paper (010) (cf. Fig. 2(b)).

VOL.

12,

1964

are probably made up of arrangements of the type shown in Fig. 7. To resolve these individual dislocations one would probably need a spacing of at least 200 A between them. Since the distance between successive (101) planes is ~1.3 A this means that for resolution of the dislocations the wedge angle must be less than tan-l l/MO. Interference fringes appear where the twin boundaries overlap on the micrographs which makes it possible, in a few cases, to get a rough estimation of the wedge angles. Geometrical considerations show that several of the faults in Fig. 4 have a wedge angle greater than tan-l l/150. This may explain why we have not been able to resolve unambiguously the ~lo~tions concerned. As mentioned in the introduction AuCu changes its structure from cubic to tetragonal on ordering. The stresses (compression along the a- and b-axes and tension along the c-axis) caused by this change are thought to be the reason for twinning since it has been shownf2) that at low ~mpe~tures the twins appear a&r the ordering has taken place. Since twinning is equivalent to a systematic shear on each successive atomic plane it can relieve some of the stresses if the shear (i.e. the twinning) occurs on a plane not containing the direction of the stresses. Thus twinning on (100) planes would not relieve the stresses but twirming on (101) and (111) planes would. In addition to the stress relief afforded by the shear on a (101) plane occurring during twinning, the rotation of the c-axis by 90” resulting from twinning will also reduce the stress, since the stress in the twinned region will no longer augment that due to the matrix; their c-axes being perpendicular. Although the electron microscope observations suggest that twinning occurs by a shear in the $(lOl) direction on successive (101) planes, it is worth considering what would happen if these ~locations were to slip on successive (111) planes. The atomic arrangement after the passage of 4 [lOI] dislocations on successive (111) planes is shown in Fig. 2(c). The structure is altered so that it becomes ordered with alternating (lli)planes instead of alternating (002) planes. The behaviour is the same with one of the other + (101) dislocations slipping on the same plane, but of course that which represents the displacement between atoms of the same kind does not effect the order. This is true for all the (“II> planes. Thus twinning of this sort is probably prevented by the forces w?lich determine the ordering of the atoms. This behaviour may restrict normal slip on (111) planes to those dislocations which do not produce an anti-phase boundary. Barker@) has suggested as a possible means of stress

HANSSON

relief the formation evenly distributed

BARNES:

AND

ORDER

of lamelae where the c-axes are in the three principal directions

in

TWINNING

Dissociation (112)

partial

original cube axes in the sequence XYZXYZ

The occurrence

plane.

This

were parallel to a { 1111

arrangement

has not

been

found

319

of a 4 (101) dislocation

above-mentioned

where the lamelae interfaces

AuCu

dislocations

the crystal. He suggested an arrangement where the lamelae had mutually perpendicular c-axes along the

. . . and

IN

unfavourable

must be associated

into two l/S

ma’y be restricted

by the

change in order which

with at least one of the part,ials.

of slip on 1111) planes could thus be

reduced.

in

While concluding

that order twinning on the {lOI>

practice and the reason seems to be that it cannot be

plane appears feasible and can explain the observations

produced

we may ask why twinning

planes.

by the movement Instead

of dislocations

one would

on these

get the above

radical

change in order. In a given { 11 l} plane the partial dislocation

concerned vectors.

can have three possible

l/6 (112) Burgers

One of these Burgers vectors is perpendicular

to the rows of copper

(or gold)

plane and when such a partial successive

is favoured

atoms

in the (lllj

dislocation

but one

is more effective

by changing the c-axis direction, alter the c-axis direction. a less systematic

while slip does not

Also slip, by being generally

process,

is lrss likely to relieve the

stresses on a fine scale. To summarize

it is suggest,ed t,hat order twins in

AuCu are generated

by dislocations

having a 4 (101)

both a normal twin and an order twin.

The move-

Burgers vector moving on successive (101) panes.

ment of either of the other two possible

partial dis-

evidence

locations

will, however, produce an ordered structure

with alternating (002) planes. likely,

(11 l} planes instead

of alternating

Thus the former type of twin is more

and in addition

the direction

for Harker’s

this is because

are probably because

Faults on

of less importance

Their generation

high energy involves

is both a stacking-fault This boundary

creating

with them.

a boundary

will have considerably

is no dislocation

No

complete that

mechanism

can occur.

One of us (B. H.) is indebted Research laboratory tillst&ndets

higher energy

twin atoms

a,re

from their original lattice sites, while in its

order twin only the type of atom is changed.

to the Atomic Energy

Establishment, Harwell , for facilities and to Dolepationrn fysik,

Stockholm.

Rwedrn,

providing fi5r fasta

for fina;:c ial

support

which

as it is probable that the

since in a normal

gives

ACKNOWLEDGiMENTS

REFERENCES

boundary

energy is greater than that of the anti-

phase boundary displaced

and rarely formed

associated

and an anti-phase

than an order twin boundary stacking-fault

been observed

the three cutting across (Fig. 5) but they

of the

there

whereby the arrangement

to the rows

of copper (or gold) atoms are likely to exist. (for example

which

of the stress

perpendicular

(11 l> planes have indeed occasionally

model

stress relief, has been found and it is suggested

would favour it. On these grounds only those normal twins with shear directions

than

slip in that it reduces the stresses both by shear and

slips on

(111) planes it will generate a twin which is

t,o slip as a

This is not understood,

reason may be that twinning

However, normal twinning could also produce stress relief.

means of stress relief.

1.

C. H. JOHANSSON and

J. 0.

l~sne.

.-lnv.

I’hysiE

25, 1

(1936). 2. G. C. KUCZYNSKI, K. F. HOCH~IISY md 11. DOYAMA, J. Appl. Phys. 26, 871 (1955). 3. D. W. PASHLICY and A. E. 3. F’n~s~~rn, I’roc. IS’ur. Reg. Conf. cm Elemon Microscop/, Delft. 1, 4”9 (1960). 4. M. HIRABAYASHI and S. ~Vm~xxar.-:~. Actn Hct. 10, 2.5 (1963). 5. D. HARKER, Trrrns. Amer. Sot. Vntrrlr 32, 210 (1944).