STUDY OF DISLOCATIONS IN CYCLICALLY TRANSFORMED β - PHASE IN Cu-Zn-Al ALLOYS

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STUDY OF DISLOCATIONS IN CYCLICALLY TRANSFORMED β - PHASE IN Cu-Zn-Al ALLOYS

D. Rios Jara, M. Morin, C. Esnouf, G. Guénin

To cite this version:

D. Rios Jara, M. Morin, C. Esnouf, G. Guénin. STUDY OF DISLOCATIONS IN CYCLICALLY

TRANSFORMED β - PHASE IN Cu-Zn-Al ALLOYS. Journal de Physique Colloques, 1982, 43 (C4),

pp.C4-735-C4-740. �10.1051/jphyscol:19824120�. �jpa-00221971�

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JOURNAL DE PHYSIQUE

ColZoque C4, supple'men L au n o 12, Tome 4 3 , de'cembre 7982

STUDY O F D I S L O C A T I O N S I N C Y C L I C A L L Y TRANSFORMED R - P H A S E I N Cu-Zn-A1 A L L O Y S

D. Rios 3ara'

,

M. Morin

,

C. Esnouf and G. Gu6nin

Groupe d 'L'tudes de Wtrr ZZurgie Physique e t de Physique des ;bfaLe'riaux, L.A.

341, P d s t i t u t PIationaZ des Sciences kppliqv.des, Ubciment 5 0 2 , 6 9 6 2 1 7i Z Zeurhatz~ne Ccdex, Pyanee

(Revised text accepted 11 October 1982)

Abstract.- The main features of dislocations in Cu-Zn-A1 f, phase exhibiting thermoelastic martensitic transformation have been analysed by means of electron microscopy and computer electron micrograph techniques. Two types of samples were studied : samples which were thermally cycled in the bulk state around the transformation range and two way memory samples which were obtained by temperature cycling with an applied stress. Xn both cases, most of dislocation line directions were near or coincident with <Ill>directions of the f phase. In thermally cycled samples, <I002 type directions were also observed and in all cases dislocations has .:I00 >type Burgers vectors. In the thermally cycled under stress observed dislocations are in fact often dislocation dipoles in {110\ planes and they are close to <Ill> screw orientation.

Introduction.- The knowled~e ofmechanisms producing dislocations in the reverse martensite transformation of thermoelastic alloys has received an increasing interest from people in the last few years. This is mainly because the presence of these defects is probably associated with some characteristics of the shape memory effect and could be at the origin of the two way memory effect in those alloys.

Very recently, S. Kajiwara and T. Kik~ichi (I) studied in a very detailed way the b.c.c.t3R, 9R transformations in Cu-Zn alloys ; from the first cycle a high density of dislocations in the b.c.c.t,3R transformation was observed. The b.c.c.=9R transformation gives very few dislocations in the first cycle, but a higher density is obtained after several transformation cycles (9). In both transformations, parti- cular arrangements of dislocations with <Ill> type directions were reporter. A formation mechanism is also proposed'in their work ; but, because of the great anisotropy of the alloy, Burgers vector determinations were not performed to prove their model.

In the present work, a similar analysis of defects in f Z 9 R transformation of ternary Cu-Zn-ALalIoy have been performed, and Burgers vector determinations are in- cluded ; but, as results are preliminary (and not in agreement with expected ones), no formation model is proposed for the moment. In addition, dislocation characteristics in thermally cycled under stress f, phase are presented.

Experimental.- Two alloys : 1) Cu, 14.1 At. % Zn, 17.1 At. 4 A1 and 2) Cu, 15.4 At.%

Zn, 17.1 At. % Al, were preparated by melting high purity Cu, Z n and A1 in an induc- tion furnance under argon atmosphere at IIOO°C. The obtained ingots were encapsulated in an evacuated silica tube and single crystals were obtained by a modified Brigdman method. Encapsulated single crystals were homogeneized for 5 hours at 650°C and

quenched in water at room temperature. The Ms temperatures for alloys I and 2 were -10'~

and

-

1 0 7 " ~ respectively (dilatometric

-

measurements).

The single crystal I (Ms = -IO0c) was thermally cycled 16 times in bulk state between room temperature and liquid nitrogen temperature. It was then oriented by back Laue technique and cut by a slow diamond blade sax? to obtain I001

1

- type surfaces in order to avoid the surface martensite formation (2).

+on leave of absence from Instituto de Investigaci6n de Materiales, U.N.A.M., MQxico D.F. (WEXICO)

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:19824120

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c4-736 JOURNAL DE PHYSIQUE

The single crystal 2 (Ms =

-

107'C) was cut to obtain a tensile test sample (6x4~40 mm). This sample was thermally cycled 20 times between room temperature and liquid nitrogen temperature under constant tensile stress application. A martensite single crystal was observed at low temperature in the central polished region of the sample while the stress was applied ; when the stress was removed, only two sets of parallel martensite variants were observed at low temperature.

Electron microscopy samples were prepared from both alloys by electro-polishing using a double jet method with 20 % HN03 - methyl alcohol solution. The transmission electron microscope was a Jeol 100 C used at 100 KV. Both alloys were confirmed to transform into 9R martensite by in situ low temperature observation in the T.E.M.

For contrast simulation,a C.1.1.IristiOcomputer was used. Computation program was a littlemodified version of the Iiead et al; program (3). In all cases, determinagions were made in the P-phasewhichis wnsideredto be a B2 order structure (a = 2.928 A

(4)) ; so all index in this work are in tegms of this b.6.c. B2 structuge. Calcula-

ted ~xtinction distances are :

51

10 = 324 A,

5200

= 466 A

, 5

= 589 A and

5130

=

842 A. Adopted anomalous absorption coefficients were : AC(l lbf2= 0.07, AC(200) = 0.08, AC(II2) = 0.09 and AC(130) = 0.10. Elastic constants values are (5) : CI 1 = 1185, C12 = 1074 and C44 = 832 x

lo9

dyne/cm2. In all this work : - B = beam direction, F = foil normal,

g

= diffraction vector and E = dislocation line direction.

Experimental results and discussion.-

Dislocations in untransformed @- phase.- As already reported (6), a very low dislocations density was present in the untransformed B - phase. Dislocations line directions of the < I l l > and

<loo>

types were observed and Burgers vectors of the <loo>

type were determined. A detailed study of the dislocations in the untransformed 6 phase will be published elsewhere.

Dislocations in thermally cycled

6-

phase.- As expected, no surface martensite was present in this samples (F = c101)).

Fig. I-a shows a typical zone on the sample. A high density of dislocations can be observed ; most of them having <ill> line directions (determined by stereogra- Sic analysis). Rows of dislocations in <110> directions were observed ; these rows could be related to the boundary between martensite plates at low temperature (as their geometry, separation and direction highly suggest).

As representative examples three different isolated dislocations were chosen for Burgers vector determinations by computer simulation and they are shown in Fig. I-b.

Fig. 2 shows a series of seven electron microscopy photographs together with their corresponding computer simulated contrast for dislocation 2 in Fig. I-b. The Burgers vector is concluded to be b = [OIO]. In fig. 2 : B =

[loll

for a ; 3 = [OOI] for b, c, d and e ; B' = [li3j for f ; 3 = [I3251 for g.

As can be seen, agreement of-contrast is good. For this dislocation, line direction was considered to be 5 =Clll],so we tested all the other possibilities of Burgers vectors of the <Ill>, <100>

,

_<110> and <112> types for each one or the three

( I 10

1

planes containing the [I I 1 ] direction. This means that we consider a

(

1101 plane as the slip plane. i4one of these Burgers vectors gives a good match between observed contrasts and computed ones.

We have also tested the

b

=

[I

¹⁰¹possibility even if this suppose a (i12) slip plane ; this was made, because after reference ( 7 ) , if we look for symetries of electron microscopy images with

z

// ⁱⁱ projection (fig. 2-c) and g lii - projection (fig. 2-b) under

B

// conditions, we would obtaine (related to the dislocation line) a left-right and a top-bottom symetries, respectively. This geometries (always following (7)) can be produced only when the extra half-plane of an edge

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Fig.l.(a) Low magnification photograph showing dislocations induced by 16 transformation cycles i n bulk 0 - phase. ii projections are shown.

(b) Enlargement of the framed zone In (a), showing the threedisioca.

tions for which Burgers vector deter~llinat~orls were perfornrcd : dislocation 1 has b:I1001. 2 has 'b I0101 and 3a b has 6 ^10011.

F1g.2. Seven electron microscope photographs and their corresponding computer simulated - -

image for dislocation 2 i n Fig.1-b. &r- gers vector is 6~10101. Fig.2-g was obtained by a . ' t w o - h a l f " simulation to increase details in the simulated image. Diffraction vectors are indicated.

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dislocation contains the incident beam, and because

h 1

8 , then % =

[I

101. The simulated images for % = [110] gave well lef t-right and top-bottom contrast symetries, but their shape was very different to the observed ones. In fact, very careful observation of electron microscopy photographs shows that they do not have the exact left- right and top-bottom symetries : see (again) for example fig. 2-c and note that the first black lobes arc not exactly in face to face position across the dislocation line (a situation which is well reproduced by the corresponding computer simulated image). However, this is a highly misleading situaiion and much care must be taken when using this method (7) of identification, because even a % = [010] Burger vector which is 45" from the ideal symetric situation, may give a contrast symetry which is very close to the ideal symetric one.

Burgers-vector of dislocations named 1 and 3 in Fig. I-b were also determined ; they are : b =

[loo]

and h = [OOI] respectively. L)islacation I has Ti =

[

^{1 1}. I] Dis- location 3 is composed by two segments : segment A has Ti =

[TI I]

and segment B has

-

u =

[loo].

This later appears very short (even if it is almost parallel to the sample surface), because it is very close the sample surface ; thence, a little devia- tion of its line direction from the exact [I001 direction makes it to intersect the surface. With this geometry, almost all of the contrast from segment A is obtained : a situation which favoured its identification.

As we can also see in Fig. I, some dislocations are associated by pairs. As can be deduced from their contrast, botlh dislocations in the pair have the same Burgers vector ; this one is very probably of the <100>type, because their contrast in different diffracting conditions has the same shape as that for the analyzed dislocations. The plane containing the pair is a ( I 10

1

type plane.

Dislocations in thermally cycled

e-

phase under stress.- As the foil normal was

--.-

F = [ 2 2 9 ] (17,4' from LOOI] ) , surface martensite was present in this sample (alloy

2).

Fig. 3-a shows an electron microscopy photograph of a typical zone. A high density of dislocations can be observed. Dislocations line directions were identified to be of the <lll>type ; and, as in the thermally cycled sample, dislocations arrangements in <IlO>directions were obtained. But, in contrast to that sample, most of dislocations appears as dislocations dipoles or more complex groups of dislocations.

We can see at the top of fig. 3-a, a ribbon of dark contrast, which in fact corres- ponds to a very high density of dislocations, with all the different line directions of the

<

I 1 ]>type. In si t11 transformation in the electron microscope, martensite have been observed to grow in this zone from the low-right corner to the top-left one, so it runned approximately parallel to the [011] dislocations rows in Fig. 3-a.

Martensite was observed to cross the dark ribbon of dislocations without any parti- cular effect, so we can conclude that all of these dislocations are perfectly stable in the martensite phase. As it was impossible to find an isolated dislocation for the Burgers vector analysis we have chosen the dislocation dipole named A in Fig.3-b.

This is because of two reasons: I ) It has a line direction close to ii=[l l llwhich is one of the majoritary types in this zone (the other one is [Till; both appear as parallel for this orientation of the sample because

E=[ooI])

and 2) it is isolated enought so that its contrast is not perturbated by other dislocations configurations.

Fig. 4 shows five electron microscopy photographs and their corresponding computer simulated images for dislocation 1 of dipole A in fig. 3-b. Dipole type contrast can be well identified in this series, as in all photographs, contrast of dislocations I and 2 are related hy a 180° rotation. The plane of dipole was identified as (1 10). Burgers vector is b = [ l l I], so dislocations 1 and 2 are closc to screw orientation.*

h he

existence of these dipoles is somewhat surprising, indeed they have ii =

5

=[I

II]

and a plane (110) which is a common slip plane, therefore such a dipole must be unstable. However, there is some inaccuracy in Lhe stereographic line direction determination and the dislocations are probably not pure screw dislocations. The same type of dipoles was found in deformed Xi-A1 alloys ( 8 ) , which is also a B2 ordered structure.

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Fig.3.ca) Low maglrlfication photograph Fig.4. Five electron microscope photographs and their corras..

showing dislocations in thermally cycled ponding computer simulated image f o r dislocation 1 of under stress 13 phase. (b)- Enlargement dlpole & i n Fig.3-b. Burgers vector 15 b : I l l l I . Diffrac ..

of the framed zone in ( a ) , dipole A w a s - tion vectors are indicated. Contrast of dislocations 1 and 2 used for Burgers vector determination. are related by a 180' rotation ; both of them are screw u-proiections are i n d ~ c a t e d . dislocations.

As before, all the possibilities of Burgers vector within the three (110} planes containing the [I 1 I] direction were tested. Unfortucately, in Fig. 4, the presence of surface martensite as well as the modification of contrast at the interse_ction region of dislocations I and 2, makes the correlation not very good. Anyway b =[Ill]

was the better possibility to explain the observed contrast. It is important to note that Burgers vector is a whole [Ill] vector and not a 112 [Ill] one. This latter possibility was also tested, but its corresponding contrast is too much faint to be compared with.

Coqclusions. Dislocations characterization in the 6 phase (transforming to 9R martensite at low temperatures) of Cu-Zn-A1 alloys has been made in this work. Results can be summarized as follows :

I) A large number of dislocations was obtained in bulk ,3 - phase after 16 transformation cycles. Arrangements o f dislocations in <110> type directions were observed.

Line directions of most dislocations are of the <Ill' type and all of them have b = <loo> type Burgers vector.

2) High density of dislocations is obtained in thermally cycled E phase under stress.

Arrangements of dislocations in <110> type directions were also observed. Line directions are of the <Ill> type. Most dislocations are associated as dipoles or morc complex groups. One dipole was completly identified as been very close to the <I112

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30URNAL DE PHYSIQUE

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