Scripta METALLURGICA Vol. 27, pp. 1471-1475, 1992 Pergamon Press Ltd. et MATERIALIA Printed in the U.S.A. All rights reserved

CONFERENCE SET No. 1

NUCLEATION OF RECRYSTALLISATION

Bevis Hutchinson

Swedish Institute for Metals Research Drotming Kristinas v~ig 48, S-114 28 Stockholm

(Received September 16, 1992)

Introduction

It has long been reeognised that nucleation in the classical meaning of the term does not take place during recrystallisation, eg. (1). Unlike most other transformations, recrystallisation is an irreversible structural change with a driving force that is almost independent of temperature, and no change of phase is involved. However, most workers find it useful to talk in terms of nucleation and growth as separate stages in the process for several reasons. Firstly, an incubation time prior to growth has been identified at least in some cases such as those studied by Anderson and Mehl (2). Direct observations in the high voltage electron microscope such as those of Ray et al. (3) and Humphreys (4) show that certain crystallites* start to grow with a high rate at a stage which gives a practical definition of nucleation. Secondly there is a close association between the point of origination of recrystallised grains and recognisable features in the deformation substructure such as shear bands, prior grain boundaries, second phase particles etc. It is therefore convenient to talk about nucleation taking place within different environments, especially as there is evidence that various environments promote specific orientations of the new grains and may bias the resulting textures accordingly (5). The present paper alms to survey the main concepts which have become accepted in the literature on nucleation and point out some new possibilities of interpretation arising from recent experimental observations.

Criteria for nucleation

Doherty (6) has emphasised that the driving force for recrystaUisation is invariably very small with the result that the critical nucleus size is quite large (~ 1 pm), containing of the order of 1010 atoms or more. The probability that so many atoms could be thrown together by random thermal processes is far too low for classical nucleation processes to be valid in general (7). A possible exception is represented by the creation of twin oriented crystallites in low stacking fault energy metals as studied by Gottstein (8) and Bcrger et al. (9), which may be considered as heterogeneously nucleated against the matrix lattice and which presumably originate by statistical fluctuations (growth accidents).

In reality, however, the lack of thermally created nuclei is of no great concern since the deformed structure of the metal consists of a contiguous mass of crystallites providing an embarrassment of ricbes as regards potential nuclei. It does not really matter whether subgralns or cells are generated during deformation, as seems to be most commonly the case, or whether they form by polygonisation during the early stages of annealing (10). The result is a structure consisting of crystallites* separated by quite sharp boundaries at the stage where nucleation will take place.

* The word 'crystallite' is used here to mean a volume having a virtually perfect lattice, which may be separated from its neighbours by either low or high angle boundaries.

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1472 NUCLEATION Vol. 27, No. 11

The modern views that a process of growth within such an aggregate of crystallites leads to the creation of a recrystallised grain was first expressed clearly by Beck in a paper provocatively entitled 'Do metals recrystallise?' (11). The residual stored energy of deformation is localised within the crystallite boundaries and the process is "fundamentally akin to a type of grain growth which occurs in substantially strain-fr~ matrix with strong single orientation texture". In view of the plethora of potential nuclei available in such a structure, the central question becomes why so few of these normally develop into recrystallised grains. Two criteria must be fulfille, d concerning stability and mobility respectively.

The nucleus must be thermodynamically stable at all stages of its development, i.e. experiencing a positive driving force for growth. The Volmer stability condition is here a function of the crystallite's size and the energy of its boundary reladve to those in the surrounding matrix, but also of its geometrical form as shown by Dillamore et al. (12) and Ray et al. (3). Long lath-like crystallites or large plate-like ones are favoured for stability and growth even though their volume may not be great. A large crystallim may be developed during annealing by the process of rotation and coalescence of two or more subgrains as proposed by Hu (13). How frequently this process is involved in nucleation is still a matter of dispute although convincing evidence has been seen in some cases (eg. (14)). It is worth emphasising, however, that the coalescence process relies on there being a very small misorientation between the subgrains such that there is a very steep gradient in the energy-misorientation relationship. Since the final orientation must lie in between those of the two original subgrains, coalescence cannot create new orientations which were not originally present within the deformation texture, as has sometimes been claimed (15).

With or without the help of coalescence there can be expected to exist a very large number of crystallites which fulfill the condition for thermodynamic stability in a typical deformed metal. The second and presumably more restrictive criterion is that the growing crystallite has a high mobility boundary so that it can grow rapidly. Most of the crystallites have small misorientations with their neighbours and low angle boundaries are well known to be immobile. It is therefore essential for a successful nucleus to have or to rapidly achieve a misorientation above the critical value to confer high mobility.

Perhaps the most fundamental question relating to nucleation concerns the critical misorientation for high mobility of a grain boundary. This is a difficult matter to investigate experimentally and only a few systematic attempts have been reported (eg. 16 - 18), giving a fairly wide range of values (10 ° - 20 °) for the critical misorientation. It seems probable that the axis of misorientation and the boundary plane will influence the critical value but little is known about these. Neither is it certain whether the transition in mobility is sharp or spread out over a range of misorientations. The mechanisms discussed by Doherty (6) could be interpreted as giving a fairly sharp transition. It remains to be seen whether theoretical approaches will provide reliable answers to these questions before convincing experimental data become available.

As well as the lxansition in mobility between low and high angle boundaries there is also a transition between general and special high angle boundaries. Special boundaries such as the 40 ° <111> type in fee metals have unusually high mobility (19) and may in favourable cases dominate the later stages of growth and, accordingly, the resulting texture. Bunge (20) has used the expression 'texture relevant nuclei' to distinguish such cases from other grains which have nucleated but failed to achieve a significant volume by growth.

Develonment of hizh anzle boundaries on annealinz

It has generally been considered that the microstructure existing after plastic deformation comprises a mixture of the prior high angle grain boundaries and low angle subboundaries formed from accumulated dislocations. This type of mixed substructure provides several different opportunities for nucleation:

(i) Nucleation at prior high angle grain boundaries That a pre-existing grain boundary can bulge out creating a strain-free volume which can continue

to grow as a recrystallised grain was demonstrated by the elegant metallography of Sperry and Beck (21). Growth occurs into the invaded grain but not in the reverse direction because the misorientation is almost nil in that case. Belier and Doherty (22) showed, however, that bulges could form in both senses along a single grain boundary, reproducing the original grain orientations on the opposite sides of the

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prior boundary. Other types of nucleation also seem possible in association with prior grain boundaries. Sperry and Beck also provided evidence for new recrystallised grains which did not share either of the parent grains' orientations. Although these observations could have been an artifact of sectioning, the same phenomenon has been shown unambiguously in bicrystals of iron by Hutchinson (23), and is believed to play a significant role in the development of recrystallisation textures in steels.

Leaving aside the effects arising from pre-existing high angle boundaries, it is well known that recrystallised nuclei can develop within the deformation substructure, for example in single crystal specimens. Although individual cell or subgrain boundaries arc generally of low misorientation, a sequence of these may constitute a long range curvature in the crystal lattice. In any such case, growth of a crystallite must inevitably lead to its acquiring an increasing misorientation and eventually a high mobility. Nucleation of recrystallisation will therefore be especially favoured in microstructural hetcrogcneities having sharp lattice curvatures.

(ii) Transition bands As shown by Hu (13) and subsequently analysed by Dillamore et al (12), the interface between two

deformation bands having a large mutual misorientation is usually composed of rows of low angle boundaries providing a cumulative misorientation. A crystallite which grows in such an environment can maintain its thermodynamic stability since its size and boundary energy increase simultaneously until a sufficient misorientation is achieved to confer high mobility. The orientations of nuclei which form in such transition bands may be well defined and so contribute systematically to the final recrystallisation texture. Examples of important textures which are believed to nucleate preferentially within transition bands are the Goss texture in iron alloys (12) and the cube texture in copper (24).

(iii) Deformation zones at particles Humphreys (25) described how dense substructures or deformation zones develop around second

phase particles to accommodate the misfit resulting from flow in the softer surrounding matrix. There is a progressive lattice rotation from the remote matrix to the particle surface which increases in magnitude with particle size and with the applied strain. Such a structure can be thought of as a transition band wrapped around the particle as pointed out by Nes (26) and is well known to provide recrystallisation nuclei on annealing. Although the crystallites which become nuclei can probably exist anywhere within the deformation zone (4, 27), the ultimately successful ones will normally lie close to the particle surface since these have the high misorientation necessary for Continued growth into the remote matrix. Detailed measurements by Humphreys and Kalu (28) showed that recrystallised grains nucleated at particles in fact have preferential orientations with respect to the deformation axes and matrix orientation. Statistically, however, the result of such particle stimulated nucleation is to produce an almost random distribution of new grain orientations and accordingly a weak final texture (eg. 29).

(iv) Shear bands When deformation becomes localised into a shear band, lattice rotations occur such that the

resulting structure can be considered formally as two transition bands back-to-back. In fact, since no orientation is truly stable in simple shear the possible range of spin angles of the crystal axes can be enormous. Such a structure should be ideal for nucleating recrystallisation which is in agreement with numerous observations since the first report by Adcock (30). Although the shear band represents a region of local high stored energy as compared to its surroundings, the crystallites which form nuclei within it must fulfill the condition of thermodynamic stability and so must still represent 'low energy blocks' with respect to their immediate environment. Having first consumed the dense shear band structure, such grains may become large enough for continued growth into the less highly strained matrix. As concerns the orientations of nuclei arising in shear bands, there seem to be different types of behaviour. In steel, Goss oriented grains have been identified by Ushioda et al. (31, 32) and the S-orientation was found in aluminium (33) whereas almost random textures were found in the case of copper (24) and in brass by Duggan et al. (34).

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1474 NUCLEATION V o l . 27, No. 11

Is a sub~ain arowth model necessary for nucleation,' ?

The concept of nucleus development by growth in size and misorientation of a subgrain has become deeply rooted as a result of the influential work of Beck (11), Cahn (10) and Cottrell (35). However, there is reason to reconsider the question now in the light of various experimental observations.

It is becoming increasingly clear that high angle boundaries as well as low angle boundaries can be created directly by moderate to high levels of plastic deformation. Crystallites having high misorientations acquired as a result of straining sbould be able to grow rapidiy and almost immediately on annealing, providing that they fulfill the condition for thermodynamic stability (suitable size and/or shape advantage).

The first work to demonstrate dearly that high angle boundaries could be produced by deformation, in single crystals of stable orientations, appears to be that of Schnell and Grewe (36). These workers made a very detailed study of the types of misorientations that existed and also showed a close correlation between orientations of nucleated grains at the early stages of recrystallisation and the misoriented crystaUites in the as- deformed state. More recent studies by Nes and coworkers (33) have shown that isolated crystaUites having a large misorientation to the matrix are not uncommon in 90 % cold rolled aluminium. Such erratic crystallites are frequently found to be rotated some 40 ° around a <111> pole and can accordingly help to explain why this relationship is so commonly observed in the formation of recrystailisation textures.

The stable orientation crystals studied by Schnell and Grewe required cold rolling reductions in excess of 95 % before high angle boundaries were detected in the deformation microstructure. In other orientations and in conditions where heterogeneities form more readily, the substructures can develop with high misorientations after smaller deformations. For example, the cube oriented transition bands observed by Ridha and Hutchinson (24) in 90 % cold rolled copper were sometimes comprised of an array of low angle boundaries but frequently consisted of a single high angle boundary to the matrix on one side. These may have been formed by some process of collapse from an earlier more gradual transition. A transition band structure composed of only two high angle boundaries was subsequently shown by Nes et ai. (37) to be typical of cube transition bands in cold rolled aluminium.

The availability of ready formed high angle grain boundaries in some transition bands would be expected to confer on these a great advantage as viable nucleation sites. In a recent study of cube transition bands in copper, Duggan and Liicke (38) observed that only some of these developed fully as recrystaLlisation nuclei whilst others grew only into the matrix on one side of the band or did not grow at aU. A possible explanation of such behaviour is that certain orientation topographies in the transition band are more prone to collapse into high angle boundaries than others and those that do so produce the successful nuclei on annealing. In the light of the findings of Nes et al (33) it might be expected that transition bands with a <111> rotation axis would most readily form high angle boundaries and therefore act as the most viable nuclei, as was in fact reported by Duggan and Liicke (38).

The deformation zones around second phase particles are also found in many cases to consist of highly misoriented crystaUites rather than just subgrains. Herbst and Huber (29) reported that some deformation zones in a 90 % cold deformed A1-Mg2Si alloy had almost a random polycrystalline structure. In a similar vein, the fine structure of shear bands in brass contains numerous high angle boundaries between adjacent crystallites even though these are only ~ 100 nm in diameter (39).

The above considerations suggest that nuclei do not always evolve out of a subgrain structure but, rather, that nucleation makes use where possible of mobile high angle boundaries which are created by the deformation process. The number of such boundaries can reasonably be expected to increase with progressive deformation which may be one reason why the ratio of nucleation rate to growth rate increases, leading to the weU known reduction in as-recrystaUised grain size. If this type of development continues up to very high strains it is not impossible that the vast majority of crystallite boundaries are either prior high angle boundaries or new ones created by deformation. Such a structure has recently been proposed by Oscarsson et ai (40) as the reason why heavily rolled strip-cast aluminium recrystallises 'continuously' on annealing. So many high mobility boundaries are present that virtuaily all crystailites are in a process of shrinking or growing and so the conventional process of recrystallisation by nucleation and growth is replaced instead by a continuous coarsening process akin to normal grain growth.

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