OBSERVATIONS OF THE INTERACTION OF LATTICE DISLOCATIONS WITH HIGH ANGLE GRAIN BOUNDARIES

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OBSERVATIONS OF THE INTERACTION OF LATTICE DISLOCATIONS WITH HIGH ANGLE

GRAIN BOUNDARIES

P. Pumphrey

To cite this version:

P. Pumphrey. OBSERVATIONS OF THE INTERACTION OF LATTICE DISLOCATIONS WITH

HIGH ANGLE GRAIN BOUNDARIES. Journal de Physique Colloques, 1975, 36 (C4), pp.C4-23-C4-

33. �10.1051/jphyscol:1975403�. �jpa-00216308�

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JOURNAL DE PHYSIQUE ColloqueC4, supplkment au no 10, Tome 36, Octobre 1975, page C4-23

OBSERVATIONS OF THE INTERACTION OF LATTICE DISLOCATIONS WITH HIGH ANGLE GRAIN BOUNDARIES

P. H . PUMPHREY Universitat des Saarlandes,

Werkstoffphysik, Bau 2, 66 Saarbriicken,

W.-Germany.

ResumB. - Les observations par microscopic klectronique de l'interaction entre les dislocations du rCseau et les joints de grains a forte dksorientation donnent de nombreux renseignements concernant non seulen~ent les propriktks, mais aussi les structures des joints de grains.

Les rksultats expkrimentaux suggkrent qu'h basse tempkrature, B la limite klastique, les joints de grains agissent d'abord comme sources de dislocations plut6t que comme barrikres de glissement. A haute tempkrature, I'incorporation par le joint de dislocations des rCseaux, se traduisant par un changement de structure du joint, est importante pour le glissement du joint et pour la recristallisation.

A I'exception de quelques cas particuliers, le modkle simple qui consiste

a

considkrer la superposition d'un rkseau de dislocations et de la structure d'un joint de coi'ncidence semble insuffisant pour comprendre ces interactions.

Abstract. - It is shown that direct observations by electron microscopy of the interaction of lattice dislocations with high angle grain boundaries can give valuable information about not only properties but also interfacial structure. The experimental results suggest that during low temperature yielding boundaries act primarily as sources of dislocations rather than as barriers to slip. At high temperatures, the incorporation of lattice dislocations by the rearrangement of boundary structure is important in boundary sliding and recrystallization. Except in a relatively few special cases, the simple boundary model of a dislocation network embedded in a Coincident Site Lattice interface appears to be inadequate for the understanding of these interactions.

1. Introduction. A number of important proper- ties of polycrystals are very much affected by the ways in which lattice dislocations interact with high angle grain boundaries. For example, the nature of the onset of low temperature yielding will depend upon whether the boundaries act primarily as sources of dislocations o r a s barriers t o slip. At higher temperatures, lattice dislocations may b e of importance in the process of sliding parallel t o the boundary. Also in this temperature range, grain boundaries may migrate and interact with lattice dislocations. This is especially important during recrystallization when relatively dislocation free grains grow t o consume the high density of dislocations in the deformed matrix.

T h e object of the present paper is to show that electron microscope observations of the interaction of lattice dislocations, with high angle grain boundaries can give valuable information about both the above properties and boundary structure.

2 . H i g h a n g l e b o u n d a r y s t r u c t u r e .

-

2 . 1 MODELS.

-

The earliest suggestion a s t o the structure of high angle boundaries was that they consisted of amorphous films similar t o supercooled liquids [ I ] . However, it is now well established [2] that the properties of high angle boundaries vary and recent models [3-51 emphasize

their periodic nature. Boundaries with relatively low interfacial energies are especially important, as their structures may be maintained by networks of dislocations when the misorientation deviates from the exact relationship [6j. Geometrical matching models have been quite successful in identifying such boundaries but not in predicting their energies [7]. Of particular interest are boundaries with misorientations such that a relatively high fraction (112) of the lattice points in the two grains lie on a Coincident Site Lattice (CSL) [3, 8, 91.

Electron microscopy of boundaries deviating slightly from C S L misorientations has confirmed the existence of the expected dislocation networks [ I 0- 131.

The Burgers vectors of the dislocations in these networks, in general, differ from those of iat- tice dislocations and vary from one CSL t o another [lo]. However, lattice Burgers vectors may be always decomposed into combinations of the special boundary Burgers vectors. Thus, if high angle boundaries consist of dislocation networks embedded in C S L interfaces, lattice dislocations should undergo reactions with them, similar t o those with low angle boundary networks of lattice dislocations. High angle boundaries, therefore, should act a s sources or sinks of lattice dislocations according t o the forces acting.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1975403

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C4-24 P. H. PUMPHREY

2 . 2 OBSERVATIONS OF INTERACTIONS. - Evi-

dence .for the type of dislocation interaction predicted above, has been obtained [14] by electron microscope observations of welded, oriented thin film bicrystals of gold. The boundary shown in figure 1 is of the [001] twist type oriented near the 36.9",

2

= 5 CSL misorientation. The square cross grating pattern has the orientation and spacing predicted for a network of screw dislocations with B u r g e r s v e c t o r s b6 = a110 [310] and

b7 = a110 [i30]. The additional wide lines such as E F are extraneous lattice dislocations. According to normal dislocation theory, the lattice dislocation E F with b = a12 [10i] should decompose as follows :

plane of the boundary. It has been shown [15] that such sessile extrinsic defects can climb in the plane of the boundary when an excess of interstitials, produced by ion bombardment, is present in the boundary region. Straining stage experiments [I51 have shown that extrinsic lattice dislocations may be pulled from the boundary but that they are rather tightly bound to it by their interactions with the network.

Care should be taken in generalizing the results of these observations as, usually, rather few boundaries contain misfit dislocation networks resolvable in the electron microscope. It is not known to what extent networks have physical meaning at smaller spacings. For oriented [001]

twist bicrystals in gold [lo], only about 13 % of the angular range was covered by high angle boundaries containing resolvable networks. In polycrystals, the proportion is usually even lower [16, 171. Thus, in

.

practice, the interaction of lattice dislocations with boundaries without resolvable networks should be more important.

An example of a boundary in a Ni 5 % A1 solid solution, deformed at roo,m temperature is shown in figure 2. Lines along theydip plane traces on the boundary can be seen. These images are different from those of dislocations in the lattice but this is probably a contrast effect, as they lie at the surface of a strongly diffracting wedge. This is confirmed by observations of dislocations lying partly in the (a) boundary and partly in the matrix [IS] and is consistent with computations of dislocation images 119, 201. There is a possibility that the lattice dislocations have reacted with an unresolved network but that the products are sessile and would have to climb apart before they can be resolved. To obtain more information about this, thin foils have been heated in an electron microscope heating holder [IS].

The procedure used was to heat the foils at increasing temperatures for 30 s. periods and to cool to room temperature for photography. In a FIG. 1.

-

Courtesy : Schober and Balluffi [14].

-

(a) Extra-

neous lattice dislocations (e. g. EF) in boundary in thin film [001] twist gold bicrystal, oriented near the 36.9O, P = 5 CSL.

The background structure is a network of grain boundary screw dislocations. (b) Schematic diagram of configuration at extraneous dislocation EF. The numbers refer to the subscripts of the

Burgers vectors given in the text.

The first product dislocation should be absorbed by the network and the second one should undergo a further reaction :

b8

+

^b7⁼b9 = a110 [1251.

The resultant configuration of the network should be as shown in (b) which is confirmed by inspection of the micrograph (a). The effective Burgers vector of the defect has a strong component out of the

FIG. 2.

-

Dark field micrograph of grain boundary in Ni 5 % Al, recrystallized (30 min at 600 OC) and then defcirmed

at room temperature.

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OBSERVATIONS O F T H E I NTERACTION O F LATTICE C4-25

FIG. 3. - Dark field micrographs of boundaries in nickel.

(a)-(c) show a general high angle boundary at room temperature and-after 30s in situ heating at 150 "C and 190 OC respectively.

given series of photographs, the diffracting conditions were kept constant.

The effect of heating boundaries in nickel specimens is shown in figure 3. F o r a general high angle boundary, the dislocation images become wider and fainter and finally disappear (I). N o evidence was obtained for splitting into discrete partials. T h e disappearance temperature in the case

( I ) The disappearance of such dislocations during in situ

heating was noted first by Ishida et al. [21].

( t l ) - ( j ) show a twin boundary at room temperature and after 30s

in situ heating at 350°C and 750°C respectively.

of the boundary shown in figure 3a-c was 190 OC.

This varied from one boundary t o another and was a s high a s 500 O C in boundaries nearer special misorientations. However, f o r the coherent twin boundary, a s shown in figure 3d-f, no spreading occurred even at t h e highest test temperature, although some rearrangement and loss of dislocations t o the foil surfaces did occur.

As the diffracting conditions in a given series of micrographs were constant, the observed spreading must represent physical changes in the dislocations.

3

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C4-26 P. H . PUMPHREY

It is difficult to account for the lack of discrete partials in terms of the CSL model. Although many of the product dislocations may have Burgers vectors too s m a l l t o ~ - b e resolved in the electron microscope, some should be comparable in size with lattice Burgers vectors [22] and, therefore, should be resolvable. It has been suggested [23]

that the observed widening of a dislocation image results from displacements of the boundary dislocation network as the additional dislocation is incor- porated. If this is so, the width of the image should decrease a s the network spacing decreases because of the increase in the interaction forces between the members of the network. However, the experimental observations [23] show exactly the opposite effect of an increase in image width as the network spacing decreases.

T o argue in terms only of normal dislocation reactions, neglects the possibility of changes occurring in the core of the incoming dislocation. It is reasonable to suppose that as the density of material is lower in the boundary than in the lattice, the core energy will be lower in the boundary. It has been argued further [24, 251 that the relatively low density of material in the boundary should result in lower restoring forces there. If this is so, its shear strength should be reduced and the equilibrium core width of the dislocation should increase. This argument is based upon the not necessarily valid assumption that the potential wells for atoms in the boundary are flatter than those for atoms in the lattice. Even if this assumption is valid, the effect on the core width is likely to be much smaller than that due to the fact that the potential wells for atoms in the boundary are usually much shallower than those for atoms in the lattice. This means that the boundary structure can be changed with a rather small effect on the interfacial energy. For example, figure 4 [26]

M I s o r I e n t a t I o n , Bdeg

FIG. 4. - Courtesy : Hasson and Goux 1261. - Relative interfacial energies as a function of misorientation for [OOI]

symmetrical tilt boundaries in aluminium.

shows that high angle [001] tilt boundaries in aluminium have a rather small variation in energy.

Therefore, if the dislocation core is spread, as is shown schematically in figure 5, the boundary energy is hardly affected. That the energy barriers to rearrangement are small, is indicated by the relatively low disappearance temperatures of the dislocations. The driving force for the process is not simply the reduction in the core energy but the reduction in the total energy of the dislocation which includes that stored in the long range strain field. This is reduced to zero when the core is spread t o infinity.

FIG. 5 . - Courtesy : Pumphrey and Gleiter [18]. - Schematic diagram to illustrate the model of dislocation annealing in a grain boundary. A lattice dislocation with a localised core (circled) (a) approaches a grain boundary and introduces an extra half plane at the boundary (b) If the dislocation core remains localised (b), the strain energy is very high. The energy may be reduced by distributing in the boundary (c), the misfit associated with the extra half plane, if the boundary energy is relatively

insensitive to the number of terminating planes.

The extent to which individual dislocations will spread in boundaries with energy minima will depend upon the degree to which the decrease in dislocation strain energy on spreading the core can balance the local increase in the boundary energy.

T o the first approximation, the boundary energy will be raised locally to that of a general high angle boundary. In the case of boundaries such a s the coherent twin, with deep energy minima, no significant spreading should occur.

If the whole of an interface with an energy

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OBSERVATIONS OF THE INTERACTION OF LATTICE C4-27

minimum is to be converted by core spreading into a general high angle boundary, a sufficient density of dislocations run into the boundary during plastic deformation must be present to provide the energy. However, the dislocation spacing is limited to

2

100

A

by the repulsive forces between dislocations [27]. In nickel, for example, this cor- responds to a maximum available energy of

-

200 ergs cmm2 which is much less than the g e n e r a l h i g h a n g l e b o u n d a r y e n e r g y o f

-

1 000 ergs cm-' [28]. As the coherent twin boundary energy ?s only

-

50 ergs ~ m - ~ , [ 2 9 ] , there is insufficient energy available to convert it into a general high angle boundary. T o test this, a recrystallized Ni 5 % A1 alloy was deformed 15 % and annealed 30 min. at 600 O C . As figure 6 shows, the coherent twin boundaries contain dislocation networks, indicating that they have retained their special structure.

FIG. 6. - Dark field micrograph of coherent twin boundary in Ni 5 % Al, recrystallized, deformed 15 % in tension and

annealed 30 min. at 6 0 0 ° C .

3. Properties.

-

3.1 LOW TEMPERATURE YIELD- ING.

-

At low temperatures, it is well known that the yield stress u,, of a polycrystal is related to the grain size d by an equation of the form

where go, k and n are constants and n

-

^112.

As the yield stress increases with decrease in the grain size, it was first suggested [30, 311 that grain boundaries act as barriers to slip originating within the grains. u,, is the lattice friction stress and k a measure of the barrier strength. The observed yield stress dependence on grain size can be predicted from dislocation theory, if the stress to initiate slip in the second grain is applied at the boundary by a dislocation pile up. In simple terms, the stress at the head of a pile up is proportional to the number of dislocations n in the pile up. Now naL-"2 where

L is the length of the pile up. This is limited by the grain size. The exact expression for k depends upon the model assumed and the method of calculation.

These theoretical studies have been reviewed recently by Li and Chou [32]. A detailed treatment of these would be inappropriate here as there are a number of experimental observations on the yielding process .which cast doubt upon the validity of the pile up model.

(i) Dislocation pile ups on grain boundaries are sometimes seen in alloys [33] but not in pure metals. However, both alloys and pure metals show a similar yield stress dependence on grain .size.

(ii) The stress at the head of a pile up can be relieved at rather low temperatures by the spread- ing process described in section 2. Figure 7 shows how a grain boundary in Ni 5 % Al can consume the dislocations in a pile up a s long as the stress is high enough to push them into the boundary (2). If dislocation pile ups are important in yielding, there should be a discontinuity in the yield stress at the temperature at which pile up absorption becomes significant. In fcc metals and alloys, there is no such discontinuity.

(iii) Dislocations have been shown t o emerge from grain boundaries in the absence of stress magnifying pile ups by optical microscopy [34, 361 and electron microscopy 137, 391. As it may be difficult to distinguish between relaxed pile ups and sources, dynamic experiments in the electron microscope showing dislocations emerging from boundaries 137, 391 are especially convincing. Even if grain boundaries act a s sources, some pile ups resulting from dislocations gliding right across the grain should be visible.

An alternative approach is to assume that grain boundaries can emit a certain number of disloca- tions per unit area [40]. Simple geometry gives the dislocation density pad-'. When taken with the experimental relationship aap"', this gives the Hall-Petch relationship. The exact form of the expression for k depends upon the assumptions made in the calculation [32].

Much effort has been directed towards the identification of these sources. The first model, proposed independently by Mott [41] and Li 1401, is shown in figure 8. It was suggested that pre- existing boundary faults could be desorbed into the lattice, under the influence of an applied shear stress. Some electron microscope observations of linear boundary defects, similar to those in figure 2 and 3 a [2 1 , 27, 42-45] have been interpreted 1321 a s Li ledges. However, it appears that these are the result rather than the cause of dislocation genera-

(*) The disappearance temperature was unusually low in this case because the sample was rapidly quenched from just below the melting point, to obtain high densities of vacancies at the grain boundaries (Pumphrey and Gleiter, to be published).

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C4-28 P. H . PUMPHREY

____)

'r

-2

-

^r-

FIG. 8. - Courtesy : Li [40] Grain boundary ledge source of lattice dislocation.

FIG. 9. - Bright field micrograph of grain boundaries in recrystallized (30 min. at 600 "C) Ni 5 % Al.

FIG. 7 . - Dark field micrographs of grain boundary in Ni 5 % AI. Part of the dislocation pile up in ( a ) is consumed by the boundary ( b ) following 30 s in situ heating at 270 O C .

tion, a s boundaries in carefully prepared thin foils, such as those in Ni 5 % Al shown in figure 9 show no such line defects. It has been shown also [46]

that boundary segregation affects uo and k in the Hall-Petch equation more than does the density of linear defects.

A major difficulty with the original Li ledge mechanism is that it allows only one dislocation to be generated from a given source. Experimental observations [35, 391 indicate that sources have multiple strengths. In general, this requires shear of the boundary. However, this is not necessary if both halves of a generated loop can be emitted on different slip planes. This process has been

observed by electron microscopy [37]. A number of mechanisms for the generation of multiple loops on a single slip plane have been proposed. For example, many screw dislocations can be emitted as the result of the repeated shear of a boundary ledge lying parallel to a lattice slip direction 1471.

The boundary shear is achieved by the gliding of the second halves of the screw loops in the plane of the boundary (Fig. 10). In situ electron microscope observations [37] have shown that boundaries can generate multiple screw dislocation loops (Fig. 1 I ) but have not demonstrated convincingly the nature of the boundary shear. Another possible mechanism for the generation of lattice dislocations is illustt'ated in figure 12. It was suggested [37] that special boundary dislocations with Burgers vectors bl in the plane of the boundary could pile up against a kink P (or triple point) in the plane, where each could split into a lattice dislocation with Burgers

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OBSERVATIONS O F T H E INTERACTION O F LATTICE

dislocations d ~ o c o t i w FIG. 10. - Courtesy : Price and Hirth [47]. The operation of a

grain boundary ledge source of lattice screw dislocations.

FIG. 11.

showing

- Courtesy : Baro et al. [37]. Bright field micrographs the generation of lattice screw dislocations from grain

boundary, during in situ straining.

vector bo and a new boundary dislocation with Burgers vector b2 (bl ^-+bo

+

b*) which could glide on the new boundary plane PB. The degree to which this mechanism is limited t o certain combinations of planes A P and PB will depend upon the restrictions on the dislocation Burgers vectors allowed. Although it is known that lattice dislocations can be generated a t kinks and triple points

FIG. 12. - Courtesy : Baro et al. [ 3 7 ] . Model for the generation of lattice dislocations at grain boundary kink P.

(e. g. [38]), there is less evidence for the mobile boundary defects required [37, 481.

The exception to this is the coherent twin boundary, in which glissile dislocations are well known [49-521. Of special importance is the emission of these dislocations on the lattice slip plane continuation of the coherent twin boundary at steps in the twin plane. The stacking fault energy has been shown t o be lower than in the matrix for dislocations on the coherent twin plane and nearby slip planes [52]. Unequal forces on the pairs of partial dislocations can lead to the emission of large non equilibrium stacking faults [5 I].

3 . 2 GRAIN BOUNDARY SLIDING.

-

At elevated temperatures, it is well established that sliding parallel t o the grain boundary can take place. The role of lattice dislocations in this process is not clear. Except perhaps at the very early stages, interactions with lattice dislocations will certainly occur. Some of those involved in the, deformation of the grains will impinge upon the boundaries.

Lattice dislocations may be generated from the boundaries at obstacles such as ledges, kinks and triple points as a relaxation process.

By analogy with the lattice, it may be argued that as a grain boundary is a periodic structure, sliding should occur by a dislocation mechanism rather than by viscous shear. Proposed mechanisms suggest that sliding can occur by the motion of networks of grain boundary misfit dislocations [53, 541 or by the motion of extrinsic boundary dislocations originating from within the boundary [37] or from the lattice [43, 551. It is difficult to obtain experimental evidence for such mechanisms as the images of most boundaries following sliding are free of the dislocation images seen following low temperature deformation [25]. Only in the case of some boundaries near special misorientations, which usually contain misfit dislocations, are defor-

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C4-30 P. H. PUMPHREY

mation induced dislocations visible. An example of a slid boundary in a 33.4°/[110] tilt bicrystal in aluminium is shown in figure 13 [57]. However, slid [001] tilt bicrystals show only arrays of misfit defects [56]. The general lack of dislocations following sliding is hardly surprising in the light of the observations of dislocation spreading described earlier. Even if localised defects were present under stress a t high temperature, these would delocalize very rapidly when the stress was removed. There has been, however, a sugges- tion [58] that if sliding is carried out at a suffi- ciently low temperature, the dislocations may be retained.

FIG. 13. - Courtesy : Kegg et al. [ 5 7 ] . Bright field micrograph of 33.4 "/[I 101 symmetrical tilt boundary in slid Al bicrystal.

The experimental evidence suggests that at least in the case of some boundaries near special misorientations, sliding occurs by the motion of extrinsic boundary dislocations. It has been suggested [59] that these originate from the decom- position of incoming lattice dislocations into special boundary dislocations which may then multiply in the interface. In view of the high temperature instability of localised dislocations in general high angle boundaries, their mechanism of sliding is less well understood. It is not clear, for example, to what extent the motion of spreading incoming lattice dislocations could contribute to sliding. The process may be more realistically described in terms of the propagation of local rearrangements of boundary structure. Such an idea is suggested by a computer experiment [60] in which it was shown that sliding' occurred by the slipping of the plane ends in one grain past those.in the other but not with equal ease for every plane. The boundary defects are similar to the dislocations that Baro et af. 1371 proposed could be mobile at low temperatures.

Irrespective of the details of the mechanism of

sliding, the generation of lattice dislocations from high angle boundaries at high temperatures should be rather easy. This is because, unlike the case at low temperatures, the accompanying boundary shear presents no problems.

3 . 3 RECRYSTALLIZATION. - During this high temperature process, relatively dislocation free grains grow to consume a deformed matrix. The driving force is the reduction in the strain energy of the dislocations. Typically, the density is reduced from

-

1012 to -lo6 lines cm". Although much is known about the nucleation of new grains and rate of boundary migration [61], rather little attention has been paid t o the mechanism of consumption of the lattice dislocations. The dislocation spreading process described earlier is clearly a possible mechanism. However, it must be shown that, at the recrystallization temperature, this process can consume sufficiently fast the high density of dislocations encountered in cold worked material.

On the basis that the activation energy for the process is that for boundary diffusion, it has been shown [I81 from the observed dislocation disappea- rance time and temperature, that individual dislocations can be consumed in

-

^lo4 s at the recrystallization temperature. Using the further assumptions that dislocations 1 000 apart may be absorbed simultaneously and that once a boundary has consumed a dislocation it can accept another, it was estimated [I81 that most grain boundaries can consume a dislocation density of a t least 1012 lines ~ m - ~ . This implies that this process is sufficiently fast not to be rate controlling in recrystallization.

As the original observations of dislocation spreading were made on boundaries in fully recrystallized material, it is of some importance to check that the process occurs also in the interfaces between growing grains and the deformed matrix.

Therefore, the annealing behaviour of dislocations in such interfaces have been studied in material, partially recrystallized and then slightly deformed at room temperature. Actual recrystallization is rather difficult t o study in 'thin foils because of boundary pinning effects [62-651 and local varia- tions in driving force [66].

The annealing behaviour of dislocations in a new grainldeformed matrix interface in Ni 5 % Al is shown in figure 14. This demonstrates that the spreading process does occur also in these interfaces. However, the temperature a t which the dislocations disappear in 30 s is 50-100 "C lower than the 350 "C observed for boundaries previously studied in the fully recrystallized alloy. The same low temperature of disappearance occurs for some of the new grainlnew grain interfaces near the recrystallization front.

In the case of the new grainldeformed matrix interface, it could be argued that the process of

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OBSERVATIONS OF THE INTERACTION OF LATTICE C4-3 1

FIG. 14.

-

Micrographs of new grainldeformed matrix inter- field micrographs taken at room temperature and after 30s face in partially recrystallized (30 min at 550 "C) Ni 5 % Al. in situ heating at 230 "C and 270 "C respectively.

( ( I ) Bright field micrograph at room temperature. (b)-(& Dark

incorporating the high density of dislocations in the cold worked material gives rise t o a higher energy, more open structure. This could not, however, apply to the new. grainlnew grain interfaces. An important feature of these is that they .have rather small radii of curvature

-

¹^p.Therefore, a ratio- nalization that could account for the low disappearance temperatures 'of dislocations in both types of interface is that both suffer a high driving force for

migration

-

5 x lo7 dyne ~ m - ~ . In the case of the new grainldeformed matrix interface, this comes from the elimination of the stored energy of cold work

-

0.2 callg [67]. For the new grainlnew grain interface, it results from the small radius of curvature r

-

¹^p. and is given by 2 y l r where y i s t h e g r a i n b o u n d a r y s u r f a c e e n e r g y

-

¹000 ergs ~ m - ~ [28]. It would appear, therefore, that migrating grain boundaries have a more open structure than stationary ones, independent of the source of the driving force. A similar suggestion has been made [68] about the structure of interpha- se boundaries moving during massive transforma- tions. Under driving forces a n order of magnitude or more higher than those experience during boundary migration, it was concluded that the interface had a diffuse structure 50 b; or more wide.

In the case of grain boundaries, the present observations indicate that they can consume dislo- cations even faster than estimated previously [18].

A further point is that care should. be taken in applying equilibrium statical grain boundary models to such interfaces. They are not, however, disor-

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C4-32 P. H. PUMPHREY

dered like supercooled liquid films, as there are still interaction of lattice dislocations with high angle variations in the dislocation disappearance tempera- grain boundaries in specific properties :

ture and, as shown in figure 15, occasionally they (i) In low temperature yielding, grain boundaries contain arrays, of periodic defects. act primarily as sources of dislocations, rather than

as barriers to slip.

(ii) Grain boundary sliding is unlikely to be by the simple glidelclimb of lattice dislocations. The decomposition of incoming lattice dislocations may provide extrinsic special boundary dislocations for the sliding of special boundaries. Lattice dislocations may be generated from a sliding boundary as an accommodation process.

(iii) The high density of dislocations i n cold worked material may be consumed in grain boundaries during recrystallization by the diffusion con- trolled rearrangement o f . boundary structure.

Except in some special cases, it appears that the simple Coincident Site Lattice (CSL) plus dislocation network model is inadequate for understanding boundary interactions with lattice dislocations. The observations suggest that the small variation in FIG. IS. ^-Bright field micrograph showing periodic defect energy from one boundary structure to another is

structure in new grainldeformed matrix interface in partially

recrvstallized Ni 5 ,% AI. important and that migrating boundaries have more open structures than static ones.

Acknowledgments. - The author is grateful to 4. Conclusions. - From the electron microscope Prof. H. Gleiter for many helpful discussions and observations described in this paper, the following to the Deutsche Forschungsgemeinschaft for finan- conclusions may be drawn about the role of the cial support.

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DISCUSSION

D. WARRINGTON : Have you observed a broadening effect where dislocations have not intersected the foil surface ? Is it possible that either a ) diffu- sion of impurities from the foil surf ace or b) facetting of the grain boundaries may be involved in this broadening process ?

P. H . PUMPHREY : The answer to your front question is yes. Diffusion of impurities from the foil surface may occur but the effect is probably not very significant, as changes in dislocations similar t o those observed on heating thin foils were seen in specimens heat-treated in the bulk and then thinned. Facetting of the grain boundary is rather unlikely as the interfacial energy is usually rather insensitive to the boundary plane.

R. BALLUFFI : IS the dissociation of the lattice dislocations in the grain boundaries consistent in any way with formal CSL theory ? For example, in many cases the DSC-lattice could consist of extremely small primitive vector in all 3 dimen-

sions. Then the lattice dislocation might dissociate into a very large number of invisible GBD's possessing these vectors. However, in other cases the DSC-lattice could be highly elongated, and part of the lattice dislocation could then dissociate into a very large number of invisible GBD's and the remaining part could survive as a visible dislocation possessing the largest primitive vector of the DSC-lattice. Additional reaction with intrinsic network GBD's might also occur.

P. H . PUMPHREY : In some cases the answer to your question is yes. However, a s you point out, most DSC-lattices contain some rather large Bur- gers vectors. One would expect t o resolve these in the electron microscope in at least some cases.

However, no evidence was obtained for the forma- tion of discrete partials during the spreading process. A further point is that it is not physically meaningful to distinguish between splitting into a large number of dislocations with small Burgers vectors and a spreading of the dislocation core.