Plastic deformation of foil copper crystals—II Electron microscopical study

PLASTIC DEFORMATION OF FOIL COPPER CRYSTALS-II

ELECTRON MICROSCOPICAL STUDY *

K. SUMINOtf, Y. KAWASAKI:, M. YAMAMOTOt and M. P. SUMINO§

In order to study the relation between work-hardening characteristics of 10.0 and 50.8 p copper foil crystals and dislocation phenomena occurring in them, the distribution patterns of residual dislocations in the crystals have been observed as a function of strain, orientation, and specimen thickness.

Dislocation tangles were found at an early stage in the deformation of a specimen that showed the lowest value of the work-hardening rate. Tangles developing only along one kind of slip planes were observed in specimens showing low work-hardening rates. On the other hand, in specimens with high hardening rates, the development of cell structure is found without exception. The structures of the cell boundaries become more complex and, at the same time, the sizes of cells decrease as the strain increases. Configurations which are thought to show the formation of sessile dislocations, and pile-ups against them were also observed.

It is suggested that the cell structure is formed because of irregular and complex motions of dislocations determined by the complex internal stress fields originating from dislocations accumulated during deformation.

DEFORMATION PLASTIQUE DE CRISTAUX DE CUIVRE PELLICULAIRES-II

ETUDE EN MICROSCOPIE ELECTRONIQUE

Dans le but d’&udier les relations entre les caractkristiques de consolidation des cristaux pelliculaires de cuivre de 10.0 et 50.8 p et les dbplacements des dislocations, les auteurs ont observe la distribution des dislocations r&iduelles dans ces cristaux en fonction de la dbformation, de I’orientation et de 1’8paisseur de 1’6chantillon.

Des 6cheveaux de dislocations ont &B observes tout au debut de la deformation dans un Bchantillon montrant unet&s faible taux de consolidation. Des Qcheveaux se dbveloppant uniquement le long d’une seule catbgorie de plans de glissement ont 8tB observes dans des Qchantillons montrant des taux de consolidation peu 81evBs. D’autre part, dans des Bchantillons p&sentant des taux de consolidat,ion Blev&, on trouve sans exception une structure cellulaire. La structure des limites de cellules devient plus complexe quand la deformation augmente, et simultanbment. la dimension des cellules diminue. Les auteurs ont observe des configurations qui sont vraisemblablement libes B la formation de dislocations sessiles, ainsi qua des empilements contre ces configurations.

Les auteurs suggirrent que la structure cellulaire se forme B la suite de mouvements irrbguliers et complexes de dislocations, mouvements d&ermin& par les champs de contraintes internes complexes provoqu& par les accumulations de dislocations qui se produisent pendant la d&formation.

PLASTISCHE VERFORMUNG VON FOLIENFORMIGEN KUPFERKRISTALLEN-‘II

ELEKTRONENMIKROSKOPISCHE UNTERSUCHUNG

Urn die Beziehungen zwischen den Verfestigungseigenschaften van folienfiirmigen Kupferkristallen mit 10.0 fi und 50.8 ,u Dicke und den in ihnen ablaufenden Versetzungsvorgiingen zu studieren, wurcle die Verteilung der Rest,versetzungen in den Kristallen als Funktion von Dehnung, Orientierung und Probendicke beobachtet.

Versetzungskn&uel fanden sich im friihen Stadium der Verformung einer Probe. die den geringsten Wert des Verfestigungsanstieges zeigte. Kniiuel, die sich nur auf einer Schar van Gleitlinien entwickelten, wurden bei Proben mit niedrigem Verfestigungsanstieg beobachtet. Andererseits wird bei Proben mit hohem Verfestigungsanstieg ohne Ausnahme die Entwickhmg einer Zellstruktur gefunden. Mit zuneh- mender Dehnung wird die Struktur der Zellgrenzen komplexer, gleichzeitig nimmt die GrGsse der Zellen ab. Es w&den Konfigurationen beobachtet, die nach unserer Meinung die Bildung van nioht gleitfiihigen Versetzungen zeigen, und ebenso Aufstauungen gegen diese.

Nach unserer Meinung bildet sich die Zellstruktur infolge unregelmiissiger und komplexer Versetzungs- bewegungen, die bestimmt sind durch die komplexen inneren Spannungsfelder, die von den wiihrend der Verformung angesammelten Versetzungen ausgehen.

1. INTRODUCTION

In a foil crystal, an active slip plane has only one

macroscopic dimension, so the motions of dislocations

* Received January 9, 1963; revised April 22, 1963.

IllLo: U.S.A. ow temporarily at,: University of Illinois, Urbana,

$ Tde Research Institute for Iron, Steel and Other Metals, Tohoku University, Sendai, Japan.

5 Department of Applied Science, Faculty of Engineering, Tohoku University, Sendai, Japan.

contributing to its deformation can be regarded as

nearly one-dimensional. This reduces the variety of

dislocations that move and also the kinds of dis-

location interactions. Hence, it becomes easier to

find relations between dislocation interactions and

macroscopic work-hardening characteristics in foils

than in bulk crystals. We previously investigated(l)

t’he plastic properties (especially the work-hardening

behavior) of recrystallized copper foil-crystals having

ACTA METALLURGICA, VOL. 11, NOVEMBER 1963 1235

1236 ACTA METALLURGICA, VOL. 11, 1963

various thicknesses and tension orientations, and ob-

tained fairly clear relations between theoretically ex-

pectable dislocation interactions and work-hardening

characteristics. However, no direct observations were

made of dislocation phenomena occurring in the speci-

mens. Hence, the interpretations given in the previous

paper remain speculative. In the present paper, trans-

mission electron microscopy has been used to observe

distributions of residual dislocations in foils after thin-

ning them by electrolytic polishiI~g. These data

together with those obtained in the previous work are

expected to be useful for revealing the mechanism of

work-hardening in f.c.c. crystals.

2. EXPERIMENTAL PROCEDURE

Tensile specimens were prepared as in I. Thin film

specimens for microscopic examination have been pre-

pared by the Bollman method’2) after various deforma-

tions in an Instron tension instrument. Observations

were made on 10.0 and 50.8 p specimens whose

orientation-dependent plastic properties have been

investigated in detail. Electrolytic polishing was

carried out in 33% nitric acid + 67 “/h methyl alcohol(a)

at 30~4O’C; the voltage and the current density

being 4 N 7 V and 1 - 3 A/cm2, respectively. Elec-

tron microscopy was done initially with a Japan

Electron Optics Laboratory JEM-5, and later with a

Hitachi HU-10 microscope.

It is most important in this work to observe patterns

of residual dislocations developing over large areas,

However, it is rather difficult to prepare suitable thin

film specimens from heavily deformed 50.8 p crystals.

Moreover, not all of the existing dislocations give con-

trast in the microscope. Also, the distribution of dis-

locations is not necessarily uniform throughout a

specimen. Therefore, one must be careful not to draw

an incorrect conclusion from observations of patterns

that exist only locally. By taking as many as 1500

electron mierographs with low magnification (5,000 x ) at an operating voltage of 100 kV, we have tried to

record the most general features of the patterns corre-

sponding to each stage of deformation.

Certain factors may prevent the observation of true

facts. For example, long-range stresses may be re-

lieved by the thinning process; and hence, dislocation

~nfigurat,ions dependent on them would disappear.

Thermal stresses induced by electron bombardment

might also affect the dislocation distribution. Finally,

new dislocations might be introduced duringhan~ing.

According to our observations, most of the observed

dislocations are in rather immobile forms consisting of

tangles or cell boundaries, and their characteristics are

shown to be closely reIated to the work-hardening

characteristics. In specimens that were not deformed

before thinning, the dislocations were found to be

quite few and to show no definite patterns; also, they

were moved easily by the action of the thermalstresses.

3. WORK-HARDENING CHARACTERISTICS OF COPPER FOIL CRYSTALS

The work-hardening characteristics of the foil

specimenso) are summarized in this section prior to the

description of the microscopic observations. Many of

the deformation characteristics of single crystals are

maintained in pseudo-single crystals.

The stress-strain curves of specimens thicker than

50 ,u (with favorable orienta,tions) show three deforma-

tion stagosas in bulk crystals, and the work-hardening

rate in stage II decreases with decreasing specimen

thickness, finally becoming indistinguishable from

that in stage I when the thickness is about 10 p.

Specimens thinner than this do not show stage II for

any tensile direction. The magnitude of the work-

hardening rate for 10.0 ,u specimens increases as the

orientation changes from Loll] to [OOl] becoming

fairly large near the [OOl] orientation. The work-

hardening rates for 50.8 ,u specimens in both stages I

and II show orientation dependences similar to those

of the 10.0 p specimens. It is noteworthy that the work-hardening rate in the well-developed stage II

for orientations near [Oil] is much smaller than for

those near [OOl]. Slip lines formed in the latter

specimens had a clustered distribution showing a

strong localization of deformation, whereas those ob-

served in the former were uniformly distributed.

4. EXPERIMENTAL RESULTS

(A) Distribution of residual dislocations

in the 10.0 p spepeeimerbs

(A 1) Specimens elongated in the A direction. Figs.

1 and 2 show typical distribution patterns of

residual dislocations in specimens with the A orien-

tation of Fig. 4 in the previous paper, after 12 per cent

shear strain and 100 per cent shear strain (fracture),

respectively. Most of the residual dislocations take the

form of tangles developed along a (110) direction. The film surfaces are nearly (100) planes, so a (110) direc-

tion is an intersection of a slip plane with the surface.

Thus, the dislocation tangles lie mostly along a (111)

slip plane. A few tangles that are not along (110) are

also observed. The configuration within each tangle

becomes more complex as the strain increases. Also, in advanced stages of deformation, the tangles are some-

times composed of the sides of rectangles or parallelo-

grams elongated in (110) direction, their major sides

being approximately parallel to (110) in most eases.

SUMINO et&.: PLASTIC DEFORMATION OF FOIL COPPER CRYSTAL&-II 1237

FIG. 1. Typical pattern of residual dislocations in 10 y specimens elongated in the A direction (12% resolved

FIG. 3. Typical pattern of residual dislocations in 10 ,u

shear strain). x 8000. specimens elongated in the C direction (lOyO resolved

shear strain). x 8000.

These rectangles or parallelograms are thought’ to be a

kind of cell; their sizes are approximately 2 ,u x 5 ,u

or larger. Such elongated distribution patterns along

one direction are not influenced by the strain. Long-

range dislocation motions are often observed in

tangle-free regions within the electron microscope.

(A 2) Specimen~s elongated in the C direction. Figs. 3

and 4 show typical patterns of residual dislocations in

specimens having the C orientation of Fig. 4 in the

previous paper. The former was taken after 10 per

cent shear strain and the latter after 60 per cent

(fracture). In this case: dislocation tangles tend to

develop along (110) direction for small strains but this

FIG. 2. Typical pattern of residual dislocations in 10 p specimens elongated in the A direction (after fracture).

x 8000.

FIG. 4. Typical pattern of residual dislocations in 10 /J specimens elongat,ed in the C direction (after fracture).

x 8000.


tendency is not so marked as in (A 1). For larger deformations, cells of a square or rhombic form develop, one side of them tending to be composed of tangles developed along the (110) directions. The cells are about 1.5 ,u x I.5 p in size at smallest. Both the dislocation density, and the complexity of cell bonndaries increase with increasing strain.

(A 3) Specimens elongated in the E direction. Figures 5 and 6 show typical patterns of dislocations in speei- mens having the E orientation of Fig. 4 in the previous paper. The former was taken after 4.7 per cent shear strain and the latter after 25 per cent (fracture).

FIG. 6. Typical pattern of residual dislocations in 10 p specimens elongated in the E direction (after fracture).

x 8000.

The patterns described above are representative of those observed most frequently. However, there exist other kinds of patterns which are found less frequently. An example is the long straight lines lying along (110) as shown in Figs. 7 and 8. They are often

Pm. 8. Typical pattern of residual dislocations in 10 y specimens elongated in the E direction (4.70/o resolved

shear strain). X 8000.

Here, directed dislocation tangles are not observed even for small defo~ations. Instead, cells in the form of polygons or circles (polyhe~al or spherical in three dimensions) are found. The sizes of the cells decrease with increasing strain, finally attaining values smaller than 1 p; and the structure becomes more complex as the strain increases.

(A 4) Some rema&. The density of the residual dislocations is closely related not to the strain, but to the stress as long as the deformation is not too large. For identical Aow stress values, the dislo~tion density is found to be the same independent of the orientation. This is not the oase, however, for specimens after fracture. For them, the density of dislocations increases in the order A, G, and E in spite of the lack of corre- lation of the fracture stress with orientation.

FIG. 7. Long straight dislocation lines lying along (110) observed in a 10 ,u specimen elongated in the A direction

(early stage of d0fo~ation). x 8000.

SUMINO et al.: PLASTIC DEFORMATION OF FOIL COPPER CRYSTALS-II 1239

observed clearly in specimens with the A and C orien-

tations at any stage of deformation* and always lie

along (110); i.e., on a slip plane. These long straight

dislocations appear to be immobile. They remain in

specimens even after thinning; also, they do not move

under the thermalstressesproducedby themicroscope.

The immobile dislocations may be (1) sessile dis-

locations(4-6) formed by the reaction of AB(d) dis-

locations with BD(a) or CD(a) (the meanings of these

notations are given in the previous paper), (2) disloca-

tion dipoles”) made by the motion of AB(d) disloca-

tions, or (3) screw parts of CB(d) dislocations which

have been immobilized by a high density of jogs formed

on them. It is not expected, however, that the CB(d)

dislocations would be activated at small strains in spe-

cimens with the A or C orientation, and probability (3)

is improbable. Those straight dislocations observed in

early stage of deformation (Fig. 7) may be dislocation

debris. Threefold nodes are found quite often on the

straight dislocations in Fig. 8 and it is thought that

sessile dislocations form by dislocationreactionsduring advanced stagesof deformation. Thereaction between

AB(d) and BD(a) leads to the formation of Lomer-

Cottrell sessiles, while that between AB(d) and CD(a)

leads to sessiles proposed by Hirthc6) BD(u) dislo-

cations become less active upon the orientation change

[Oll] -+ [OOl]. There observed, however, no evidence

* They are also found in specimens with the E orientation, but here they are not so clear because of masking by the cell boundaries.

FIG. 8. Long straight dislocation lines lying along (110) observed in a 10 ,G fractured specimen elongated in the C

direction. x 8000.

FIQ. 9. Dislocation configuration showing, possibly, a pile-up in a 10 ,u specimen elongated in the E direction

(2.2% resolved shear strain). x 32,000.

that the long straight dislocations decrease upon

this orientation change. Furthermore, if the formation

of the original Lomer-Cottrell sessiles were most dom-

inant in the work-hardening, the orientation depend-

ence of the work-hardening rate in stage I should be

less remarkable in the thicker specimens compared to

the case of the thinner specimens. This contradicts to

the observed fact. Thus, it is thought that a consider-

able part of the observed sessiles formed by the reac-

tion between AB(d) and CD(u) as proposed by Hirth.

An interesting configuration that is observed rather

exceptionally is shown in Fig. 9. This was formed in a

specimen elongated 2.2 per cent shear strain in the E

direction. It is thought to be a pile-up of dislocations

against a sessile dislocation. Ingeneral, even if pile-ups

form during deformation and are maintained during

unloading, they may rarely remain in thin film speci-

mens since they cross-slip rather easily as shown in the

previous paper, and because the stress distribution in

thin films is different from that in tensile specimens.‘*)

The pile-up in Fig. 9 is thought to have been stabilized

accidentally by the crossing of dislocations on another

slip system so that it remained throughout the

thinning procedure. In any case, it is believed that the

configurations in Figs. 7, 8 and 9 show that sessile

dislocations do form.

(B) Distribution of residual dislocations in 50.8 ,u

specimen8

(B 1) Specimens elongated in the A direction. Figs.

10, 11 and 12 show typical dislocation patterns


FIG. 10. Typical pattern of resi 50.8 ,u specimen elongated in the

x 32,000.

dual dislocat ions in a A direction (stage I).

FIG. 11. Typical p&kern of residual dislocations in a 50.8 ,u specimen elongated in the A dire&ion (st,age II).

x 32,000.

FIG. 12. Typical pattern of residual dislocations 50.8 ,u specimen elongated in the A direction (stage

x 32,000.

in a III).

SUMINO et cd.: PLASTIC DEFORMATION OF FOIL COPPER CRYSTALS--FI I241

Fra. 13. Typical pattern of residual dislocations in a SO,% p specimen elongated in the C direction (stage I).

s x000.

in the specimens with the A orientation of Fig. 8 in the previous paper at deformation stages I, II and III, respectively. In stage I, tangled dislocations developing along 8 certain direction (similar to those in the 10 y specimens) are observed. In stage II, the distributions and tangbng modes are more complex, and cells begin to be formed. A fairly distinct cell structure is sometimes found in specimens during stage III. But the cells are not so distinct its those found in the specimens of (B 2) and (B 3). In some cases, no distinct cell structure is found and the residual dislocations are distributed rather uniformly at random. However, this might not mean the absence of a cell structure, but simply the existence of large

Fra. 15. Typical pattern of residual dislooat~o~s in a 80.8 +u, specimen elongated in the C direction (stage III).

x SOOO.

cells. It is dif&ult to prepare thin pieces from tensile specimens of this thickness which allow electron trans- mission over sufficiently large areas, especially after large deformations~ and detailed observations over large regions were not made. The indistinct nature of these cells may be related to the low work-hardening rates in stage II for these specimens.

(B 2) ~~~~~~~~ ~~~~~~d in the C ~~r~~~o~. Figs. 13, 14, and 15 show typical d&&cation patterns in specimens with the C orientation of Fig. 8 in the previous paper, at stages I, 11 and III, respectively. The changes in the patterns with increasing strain are like those in the 10 p specimens of the same orientation. Long straight dislocations lying along (110)

Fm. 14. Typical pattern of residual dislocations in 8 50.8 ,u specimen elongated in the C direction (stage II).

x 8000.

WC. 16. Typical pattern of residual dislocations in a 50.8 ,u specimen elongated in the E direction (early stage

of deformation). x 8000.


FIG. 17. Typical pattern of residual disloelutions in & 50.8 ,a specimen elongated in the E direction (fract,ured).

x 8000.

may be seen in Fig. 13. A distinct eel1 structure is always present during stage II. As the strain increases, the cells decrease in size, and the cell structure becomes more distinct.

(B 3) ~~~c~~~~ e~~ga~~~ in the E d~r~e~~o~. Figs. 16 and 17 show typical dislocation patterns in specimens with the E orientation of Fig. 8 in the previous paper. The former was taken at an early stage of deformation, and the latter after the fracture. A clear cell structure developed early in the deformation. The changes of cell size and complexities of the cell boundaries with increasing strain are like those in 10 j,4 specimens.

5. DISCUSSIONS

Dislocation tangles were found at an early sta*ge of deformation even in the specimen showing the lowest work-hardening rate (Fig. 1). This suggests that tangling is not directly related to the cause of high work- hardening rates. It is not clear from the present work whether or not these tangles form as a result of the meeting and the subsequent interaction of primary dislocations with those belonging to other slip systems. At any rate, it is concluded that the macroscopic work- hardening rate is low as long as the tangles develop along only one direction (or along one slip plane). In this case, most of the glide dislocaGons ean move rather freely in the tangle-free regions,and move out of a specimen without being trapped by tangles. A few dislocations remaining in a specimen because of tirapping by tangles, or because of the self-tangling

process(g) would be responsible for the so-called “basic-hardening”.(lo)

Two dimensional configurations of tangles are observed during stage II hardening in 50.8 ,H. specimens. The transition of work-hardening from stage I to stage II begins when these two-dimensional configurations appear. The absence of stage II hardening in 10 p specimens suggests that the formation of the two- dimensional configurations is markedly affected by specimen thickness. A possible interpretation of this was given in the previous paper.

A high work-hardening rate accompanies well- developed cell structures without exception. Since the formation of the cell structure is closely related to dislocation activity on non-primary slip systems, it is natural to think that cells form by cooperation between dislocations belonging to several slip systems. The boundary structures of cells are very complex at the advanced stages of defo~ation in specimens with high hardening ra-tes, so their formation can hardly be described in terms of simple gliding motions of dislocations on definite slip planes. Possibly, cells form as a result of irregular and complex dislocation motions induced by complex distribution of internal stress which originates from the dislocations accumulated during deformation. If the interpretation of the formation of sessile dislo~tions mentioned in the preceding section is correct, it seems likely that internal stresses at pile-ups will activate dislocations in other pile-ups or dislocation sources on other slip systems, and subsequent reactions between the dislocations will occur creating cell boundaries. Once simple cells had formed, they would trap subsequent glide dislocations. As the strain increases, the boundary structure becomes more and more complex, and the cell size becomes smaller and smaller, thus increasing the work-hardening rate.

The phenomenon that causes the t,ransition from stage II to stage III hardening was not detected in the present observation. Perhaps the trapping efficiency of glide dislocations changes at a certain state of the cell structure.

A distinct clustered distribution of slip lines is observed on the surfaces of specimens having well- developed cell structures in them, but the relation between these two phenomena is not yet clear. Since the individual slip lines in a cluster are longer than the cell diameters, the glide dislocations must break through the cell boundaries in order to form the surface slip lines; otherwise, one must think that cells form after deformation.

Work is now in progress to correlate the formation of clustered slip lines with the cell structure.

SUMINO et al.: PLASTIC DEFORMATION OF FOIL COPPER CRYSTALS-II 1243

ACKNOWLEDGMENTS 3. H. M. TOMLINSON, Phil. Mag. 3, 867 (1958).

The authors are indebted to Mr. K. Takahashi and 4. W. M. LOMER, Phil. Mag. 42, 1327 (1951). 5. A. H. COTTRELL, Phil. Mag. 43, 645 (1952).

Mr. H. Ota for their assistance in operating the elec- 6. J. P. HIRTH, J. Appl. Phys. 32, 700 (1961).

tron microscope. The authors are also grateful to Prof. 7. W. G. JOHNSTON~~~J.J.GILMAN, J. Appl. Phya.31,632

(1960). J. J. Gilman and Prof. S. Ogawa for reading the manu- 8. J. D. ESHELBY snd A. N. STROR, Phil. Mug. 42, 1401

script and for their kind advices and suggestions. (1951).

9. H. G. F. WILSDORF and J. SGHMITZ, J. AppZ. Phys. 33, 1750 (1962).

REFERENCES 10. A. SEECER, 2. Naturf. 9a, 758 (1954).

1. K. SUMINO and M.Y~~~~0~0,Act~~et.thisissuep.1223. 2. P. M. KELLY and J. NUTTING, J. Inst. Met. 87, 385

(1958-59).

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Plastic deformation of foil copper crystals—II Electron microscopical study