On the origin of recrystallization textures in aluminium

Acta metalL mater. Vol. 39, No. 7, pp. 1377-1404, 1991 Printed in Great Britain. All rights reserved

0956-7151/91 $3.00 + 0.00 Copyright @ 1991 Pergamon Press plc

OVERVIEW NO. 93 ON THE ORIGIN OF RECRYSTALLIZATION TEXTURES IN A L U M I N I U M J. H J E L E N t, R. O R S U N D l and E. N E S 2 ~SINTEF, Division of Metallurgy, N-7034 Trondheim, Norway and 2The Norwegian Institute of Technology, Department of Metallurgy, N-7034 Trondheim, Norway (Received 19 July 1990) Absa-act--The electron back scattering pattern technique (EBSP) in SEM has been applied to follow in detail the orientation aspects of the nucleation and growth of recrystallization in cold rolled aluminium. The investigation includes both high purity metal and a commercial grade. The cube- and Goss-oriented recrystallized grains nucleate from transition bands. In the cube case these bands are found in the Cuand ND-rotated-Cu deformation texture components while the Goss bands have been identified in the brass texture component. Shear bands in the S-deformation component have been identified as nucleation sites for recrystallization. The orientation inside the bands is also S (complementary to the matrix), causing the S-orientation to reappear as a recrystallization texture component. These types of deformation heterogeneities and others are frequently associated with a 40 ° (111) orientation relationship to the surrounding matrix, making this specific orientation relationship primarily a property to be associated with the oriented nucleation concept. Isolated examples of a true 40 ° (l 11 ) growth selection due to the rapid growth of (111)-tilt-boundaries have been observed. A characteristic aspect in these cases is a growth selection from a planar transformation front. From a careful literature review as well as from experimental observations no direct evidence has been found in support of the hypothesis that 40 ° ( 111 ) grains, when uniformly distributed in space, have a transformation rate potential exceeding that of grains belonging to other texture components. No growth selection has been found in the special case of the growth of cube oriented grains. R6sum~4:)n applique la technique des 61ectrons r6trodiffus6s en MEB pour suivre en d&ail les probl6mes d'orientation de la germination et de la croissance dans la recristallisation d'un aluminium lamin6 zi froid. L'6tude comprend ~i la fois des m6taux de haute puret6 et de qualit6 commerciale. Les grains recristallis6s d'orientation cubique ou d'orientation de Goss germent fi partir des bandes de transition. Dans le cas cubique, ces bandes sont trouv6es dans les composantes de la texture de d6formation du cuivre tourn6es ou non autour de la normale fi la t61e tandis que les bandes de Goss sont identifi6es dans la composante de texture du laiton. On a montr6 que des bandes de cisaillement dans la composante de d6formation S sont des sites de germination pour la recristallisation. L'orientation ~i l'int6rieur des bandes est aussi S (compl6mentaire de la matrice), faisant r6apparaitre l'orientation S comme une composante de la texture de recristallisation. Ces types d'h&6rog6n6it6s de d6formation ainsi que d'autres sont fr6quemment associ6s fi la relation d'orientation 40°( 111 ) avec la matrice environnante, faisant essentiellement de cette relation d'orientation sp6cifique une propri&6 qui dolt &re associde au concept de germination orient6e. On a observ6 des exemples isol6s d'une vraie s61ection de croissance 40°(l 1l> par suite de la croissance rapide de joints de torsion (111). Une s61ection de croissance fi partir d'un front de transformation plan est un aspect caract6ristique de ces cas. Malgr6 une 6tude soigneuse de la litt6rature ainsi que des observations experimentales, on n'a aucune preuve pour confirmer l'hypoth~se que les grains 40°(l l 1), lorsqu'ils sont distribu6s uniform6ment dans l'espace, ont un potentiel de vitesse de transformation qui d6passe celui des grains appartenat aux autres composantes de la texture. Dans le cas sp6cial de la croissance de grains d'orientation cubique, on n'a trouv6 aucune s61ection de croissance. Zusammenfassung--Die Orientierungszusammenh/inge bei Keimbildung und Wachstum w/ihrend der Rekristallisation von gewalztem Aluminium wurden im Rasterelektronenmikroskop mittels der Elektronenrfickstreuung ausfiihrlich untersucht. Sowohl hochreines Aluminium wie auch solches kommerzieller Reinheit werden verwendet. Wiirfel- und Goss-orientierte K6rner bilden sich aus I~bergangsbfindern. Im Falle der Wiirfelorientierung werden diese B/inder in den Komponenten der Kupfer- und NDrotierten-Cu-Verformungstexturaufgefunden,bei der Goss-Orientierung dagegen in der Komponente der Messingtextur. Scherbfinder in der S-Verformungskomponente sinde die Keimorte der Rekristallisation. Die Orientierung innerhalb der B/inder ist ebenfalls S (komplement/ir zur Matrix), wodurch die S-Orientierung als Texturkomponente der Rekristallisation wieder auftritt. Diese Arten der Verformungsinhomogenit~it und andere sind h/iufig begleitet yon einen Orientierungszusammenhang 40°(111) mit der umgebenden Matrix. Dadurch wird dieser spezifische Orientierungszusammenhang vorwiegend eine Eigenschaft, die mit dem Konzept der orientierten Keimbildung zusammenh/ingt. Einzelne F/ille einer wahren 40°(111)-Wachstumsselektion durch rasches Wachstum von (111)-Kippkorngrenzen wurden beobachtet. Ein charakteristischer Aspekt dieser F/ille ist eine Wachstumsselektion von einer ebenen Umwandlungsfront. Eine sorgf//ltige Literaturdurchsicht und experimentelle Beobachtungen ergeben AM39/7--^

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HJELEN et al.: OVERVIEW NO. 93 keinen direkten Hinweis fiir die Stiitzung der Hypothese, dab 40°(lll)-K6rner bei gleichf6rmiger Verteilung im Raum ein Umwandlungspotential aufweisen, welches das yon K6rnern anderer Texturkomponenten iibersteigt. Keine Wachstumsselektion wurde beobachtet in dem speziellen Fall des Wachstums wfirfelorientierter K6rner.

1. INTRODUCTION The origin of recrystallization textures during annealing of a deformed metal has been a matter of controversy for decades, with the discussion focusing on the rival theories, oriented nucleation and oriented growth. At the 8th International Conference on Textures in Metals (ICOTOM 8) a special panel was assembled to assess the current status of this scientific controversy, The panel report [1] is sad reading in the sense that the current state of affairs appears to have changed only marginally compared to that of nearly 40 years ago [2, 3]. The reasons for this lack of progress are easily identified: (i) Growth selection experiments as well as controlled growth rate experiments using deformed single crystals have clearly demonstrated that the grain boundary mobility is sensitive to the orientation relationship between the new grain and the deformed matrix. However, to translate this observed mobility-orientation relationship into a theory capable of predicting the evolution of recrystallization textures in a deformed polycrystalline metal in an unambiguous way has turned out to be a difficult task, (ii) The idea of oriented nucleation was for many years a concept without well founded experimental and theoretical support. The true nature of the nucleation event was hidden in the dark and has remained so until more recently. Since the nature of the relevant nucleation and growth reactions which take place upon annealing of a deformed polycrystalline metal have successfully escaped direct experimental examination over all these years, the alternative left upon was to follow the evolution of recrystallization textures on a more global scale using x-ray pole figures and ODF's. This approach has supplied a massive amount of valuable empirical insight into the evolution of recrystaUization textures. Further, over the last 10-15 years a much clearer picture has also emerged regarding the microstructural evolution during deformation of metals. In a recrystallization context, the importance of deformation heterogeneities in supplying sites for nucleation of recrystallization is well documented. That the orientation characteristics of these sites play an important role in the evolution of recrystaUization textures is now generally accepted. But, since the fundamental question relating to the dominating

mechanism responsible for the origins of recrystallization textures still remains partially unresolved, this simply illustrates that an adequate experimental technique has been missing. Such a technique now exists. With the electron back scattering technique (EBSP) in SEM [4], the textural evolution can be followed on a local scale. This opens for direct examination of both the nucleation and growth aspects of recrystallization textures, as demonstrated in some recent short communications from the present authors [5-7]. In the present paper the evolution of recrystallization textures has been considered in a broader context, experimentally as well as theoretically. The new results found have been discussed in terms of both of the rival mechanisms. However, as a background for subsequent discussion some comments on the oriented growth theory will be presented on this introductory level. The basic concept of oriented growth is that grains with a special orientation relationship to the deformed matrix will have a growth advantage compared with grains of random orientation. In f.c.c. metals grains with a 40 ° (111) orientation relationship seem to be of this special character. But these 400 (111) grains have not only been observed to have a higher growth rate, another characteristic aspect is that the growth of these grains is associated with a strong growth rate anisotropy, i.e. the mobility of the 400 (111) twist boundary is much lower than that of the orthogonal tilt boundary. The result is that 40 ° (111) grains grow with a lenticular shape having aspect ratios frequently as high as 50:1 [8]. An important question now becomes: to what extent does this growth anisotropy influence the growth selection mechanism? Or in other words, in comparing the volume transformation rate of 40 ° (111) grains to that of randomly oriented grains, is the migration rate of the tilt boundary sufficiently high to compensate for the slower twist boundary? It is surprising that this vital question has been ignored for all these years. 1.1. Growth anisotropy and the oriented growth mechanism The objective is to analyze the transformation kinetics during recrystallization of a deformed metal where nucleation and growth result in two populations of grains: (i) grains with a random orientation relationship to the deformed matrix which grow with an isotropic growth rate GR and (ii) 400 (111) grains which grow with an anisotropic growth rate characterized by Gm (tilt) > G m (twist). For simplicity these growth rates are assumed to be constant during the

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transformation. In order to calculate the transformation kinetics the nucleation behaviour needs to be defined and for obvious reasons site saturation (~r = ~ , t = 0 and ~r = 0 at t > 0, N is the nucleation rate) will give rise to the strongest growth selection effect. In the following we define the rate advantage of the 40° (111) grains by the ratio-parameter n = Gul (tilt)/GR, and the shape of the grains by the aspect ratio q = Gm(tilt)/Gm (twist). Further, we assume that the 40 ° (111) grains grow with an ellipsoidal shape [small half axis, Gm(twist)-t and long half axis, Gm(tilt).t)], and finally we assume that the nucleation sites for both types of grains are randomly and uniformly distributed in space. The transformation kinetics is now obtained by solving the following two differential equations dV = (1 - V) dV ext and

dVll I ~--(1 -- V) dV~]~

(1)

where V is the total volume fraction transformed at the time t and Vnl the volume fraction of the 40 ° (111) grains. The extended volume fractions can be written

NeXt= -4n ( Nlll n3 + N R ) G ~ t 3 3

q

4 n3 ext = ~ n N l l l -q- G 3 t 3 Vii1

(2)

where N m and N R are the density of grains with 40 ° (111) and random orientation relationships respectively. Solving the differential equations give

V1,l = I1 + NRq ]-1 V nNI- -u~ J

(3)

which shows that the Vm/V-ratio is time invariant during the transformation period. For a more extensive analysis of similar type transformations, see recent work by Doherty [9]. It follows from this equation that an oriented growth effect requires

n3/q > 1. This analysis does not apply if the nucleation sites are non-randomly distributed in space. Extreme cases in this context are those covered in numerous single crystal investigations which rely on artificially stimulated nucleation of recrystallization [8, 10-14], in practice obtained by mechanically rubbing one surface of a deformed single crystal. In such cases the transformation ratio Vm/Vwill be time (or position) dependent, and a strong growth selection effect may occur as long as n > 1 even if n 3/q << 1. Of primary interest to us is, of course, what is found in more realistic cases, i.e. during annealing of deformed polycrystalline metals. In heavily deformed metals (~ > 1) the assumption of site saturated nucleation kinetics appears to be a sound one both in high purity metals [15] and in commercial alloys [16].

OVERVIEW NO. 93

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Practical experience also show that the nucleation sites for recrystallization in general are reasonably uniformly distributed in space. Although deviations from randomness in the spatial distribution may occur, such deviations are not expected to have any significant effect on the basic requirement for an oriented growth effect, namely that the extended volume of the 40 ° (111) grains have to increase during transformation at a higher rate than that of the randomly oriented grains, i.e. n3/q > 1. Now, it follows that to apply the results from single crystal growth selection experiments as a basis for the 40 ° (111) oriented growth theory, a necessary requirement becomes that the migration characteristics of the grain boundaries involved convincingly can be demonstrated to be in accordance with n3/q > 1. Of all the single crystal investigations performed over the last 40 years only two are found to be of relevance in this context. That is the famous work by Liebmann et al. [10] and the more recent investigation by Gottstein et al. [8]. Liebmann et al. carefully investigated how the migration rate of 40 ° (111) type boundaries varied with deviation from the ideal orientation relationships. It is convenient (as recommended by Liebman et al.) to separate this angular deviation into two categories: (i) an angular deviation ~b from the ideal 40 ° around the common (111) axis (because of lack of symmetry, effects of + q~ and - ~ need to be considered); (ii) an angular deviation ff which defines the deviation from "axiality", i.e. the common (111)-axis in the two crystals do not exactly coincide but define an angle ~ with respect to each other. The diagrams in Fig. 1 summarize the findings of Liebman et al. [10] and Gottstein et al. [8] by plotting the parameters n, q and n3/q as functions of the angular deviation from 40 ° ( I 1 1 ) , and both ~b- and ~k-type angular deviations from ideality have been included in the diagrams. In order to derive the migration-rate-ratio, n = Gl,(tilt)/GR, the migration rate of a "randomly misoriented" boundary, GR, needs to be quantified. While Liebmann et al. [10] measured the migration rate of boundaries between grains which were oriented close to the ideal orientation very accurately (~b and ~b in the 0-10 ° range), the speed of more "randomly oriented" boundaries is according to their results less well defined. However, boundaries separating grains which were reported to deviate substantially from the ideal orientation were measured to migrate at rates which were approximately a factor 0.7 slower than the ideally oriented tilt boundaries, and n = 1.5 (for ~b = ~k = 0) has been selected in plotting the curves in Fig. 1. However, which value we select for GR has a strong effect on the n3/q-parameter as indicated by the dotted line in Fig. 2 which correspond to n = 2. It follows from the results given in Fig. 1 that the speed of the (111 )-tilt boundary is strongly sensitive to deviation from the ideal orientation relationships. That is especially so for the ~b-type angular deviation.

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HJELEN et al.: OVERVIEW NO. 93

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Fig. 1. Effect of angular deviation from the ideal 40° (111) relationship for the normalized growth rate: n = Gm(tilt)/GR, the aspect ratio: q = Gm(tilt)'t/Gm(twist).t and the transformation parameter: n3/q. Oriented growth only if n3/q > 1.

As can be seen, q~ _+ 10° results in about 50% reduction in the migration rate. Deviations in the VJ-angle is less critical, for ¢ > 10° the boundary approaches a mobility level similar to that of a randomly oriented boundary. On the effect of variations in ~b and @ on the aspect ratio of the growing 40 ° (111)-grains, Gottstein et al. [8] reported that the q-value is insensitive to variations in q~ except for the value of ¢b which corresponds to a twin orientation relationship, where the aspect ratio approaches 100. However, the q-value is strongly dependent on the -angular deviation. As can be seen from Fig. 1, q has dropped to around 25 at VJ = 5 ° and is less than 1.5 at V/= 10°. The variation in q with ~ is well documented in the work by Liebman et al. From their observations an increase in ~ from 0 to 10° causes GHl(tilt) to decrease from 4.6 to 3.9 mm/min while the corresponding change in the twist boundary migration rate [Gin (twist)] is from about 0.1-2.7 mm/min, or the q value decrease from about 50-1.4.

The geometric parameter n 3/q represents the volume-growth-rate-ratio between 40 ° (111 ) grains and randomly oriented grains. An oriented growth effect requires n3/q > 1, and it follows from the data plotted in Fig. 1 that no unambiguous conclusion in favour of oriented growth can be drawn from these results. In essence, what the single crystal experiments have solidly documented is the very special anisotropic growth behaviour of 40 ° (111 ) grains, but in comparing their transformation potential, it has not been firmly established that the growth rate of these grains are sufficiently high compared to that of randomly oriented grains. It may be that the growth rate of randomly growing grains have been overestimated in the plots presented in Fig. 1 (although according to the Liebman et al. results it has not). The Liebman et al. results pertain to recrystallization growing into a relatively lightly deformed single crystal (20% by wire drawing). However, similar data obtained from single crystals deformed to larger strains are not available, except for the q-data reported by Gottstein et al. [8] which refers to single crystals (S-orientation) deformed 70% by rolling. It is interesting to note that these workers report the same q-values as those which can be deduced from the Liebmann et al. investigation, see Fig. 1, The growth selection experiments performed by Kohara et al. [11] and Parthasarathi and Beck [12] on S-oriented single crystals deformed 80% by rolling may indicate, qualitatively, that the n-value is larger than that which can be deduced from Ref. [10]. Unfortunately, no quantitative data in terms of relative growth rates between 40 ° (111 ) grains and randomly oriented grains can be derived from these works [11, 12]. In summary, the objective of this simple transformation analysis is to focus the attention on the fact that based on the experimental data presently available from single crystal experiments on the growth characteristics of 40 ° (111 ) grains, the conclusion has to be drawn that no unambiguous evidence exists in support of any oriented growth effect. That the artificial nucleation experiments [10-14] results in such a growth effect follows from the special case of having a "planar-growth-selection-situation" from a two dimensional site distribution, after all this is a trivial

Fig. 2. Electron back scattering pattern (EBSP).

HJELEN et al.: OVERVIEW NO. 93 observation which follows as a necessity as long as n > It, irrespectively of the value of q. We will continue this oriented growth/oriented nucleation discussion in view of the present experimental results in the appropriate section below. As a final introductory remark it is pointed out that the present analysis is highly idealized in the sense that the grains are assumed to grow in a deformed matrix having a uniform orientation. A deformed polycrystal represents a more interesting situation, and in this case, depending on the precursor grain size and the rolling reduction, a growing grain may rapidly become exposed to two or more differently oriented matrix components. In the oriented growth theory the idea is that a growth selection still applies in situations where alternating layers all satisfy the 40 ~' (111> orientation relationship, only the axis of rotation changes from component to component. (An important example being the case of the cube oriented grains growing in a polycrystalline deformation matrix of predominantly the S-type.) Such a situation will complicate the above analysis, but not change the basic physics related to the intrinsic properties o f 40 ° (111) tilt and twist boundaries. Simple inspection shows that these complications will go in a direction disfavouring the growth of 40 ° < 111 > grains. So the conclusion becomes that although the analysis above is idealized, if this approach does not convincingly support the idea of a growth selection, it becomes premature to extend this oriented growth idea to more complicated situations. Exceptions being cases where rather extreme deviations from non-randomness in the spatial distribution of nucleation sites applies.

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The texture analysis was based on X-ray pole figures and on the SEM electron back scattering pattern (EBSP) technique developed by Venables and Harland [4]. In the latter technique a stationary electron beam hit the highly tilted specimen, and backscattered electrons become diffracted on their way out of the specimen. Diffracted backscattered electrons hit a transparent phosphor screen and through a leadglass in the specimen chamber wall, the pattern (EBSP) is viewed by a low light television camera and displayed on a monitor. The orientation of a specific grain is obtained by positioning a cursor on each of two different zone axis, usually the 112and the 114 zones or the 112- and the 111 zones. On the basis of these positions a microcomputer calculates the orientation, and the result is automatically plotted on the screen as shown in Fig. 2. To determine local textures the specimen has either to be translated mechanically from grain to grain or the electron beam has to be moved electronically. In the coarse grained high purity aluminium the specimen itself was moved, while in the commercial alloy, where the grain size is much smaller, the electron beam was manually scanned in a raster. To recognize the grains from which the different EBSP's originate, electron channeling contrast images in both the secondary and the back-scattered modes were recorded. In the secondary electron image, Fig. 3(a), the con-

2. EXPERIMENTAL The starting materials were: (i) a directionally solidified high purity (99.99) aluminium ingot with nearly parallel columnar grains. The ingot cross section was 4 × 4 cm and the diameter of the columnar grains were in the mm range: and (ii) a commercial purity aluminium alloy (0.3wt% Fe and 0.04 wt% Si) received as hot rolled slab 9 mm thick. Plates for cold rolling of the high purity ingot were cut normal and parallel to the solidification direction and the rolling reduction was 90% (e = 2.3). The commercial alloy was prior to cold rolling homogenized for 24 h at 580°C followed by a slow cooling in order to reduce the solid solution level of Fe as much as possible. The alloy was cold rolled 95% (E = 3). Annealing was carried out in a salt bath. tMore rigorously this condition should be written n cos (~ n/2) > I where ct is the angle between the common <111)-axis and the surface normal (from which the new grains are artificially nucleated). Since we have 4 <111> directions there is always one which satisfied 0.82 ~-direction parallel to the surface normal. -

Fig. 3. (a) Secondary electron channeling contrast image. Contamination spots indicate cube grains. (b) Back scattered channeling image.

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HJELEN et al.: OVERVIEW NO. 93

tamination from the electron beam shows the scanned raster. When the beam came into a cube grain, the beam was stopped for about 30 s leaving a contamination spot there. To increase the signal to noise ratio a micrograph of the same area was made by backscattered electrons, Fig. 3(b). The secondary electron image was used to recognize the positions of the cubes while the backscattered image was used for grain size distribution measurements. Recently a more sophisticated imaging technique has been added. A new detector has been developed in order to image the mirostructure when the specimen is in the tilted configuration. The digital image which is processed to increase the signal to noise ratio is stored on a monitor when the beam subsequently is switched to spot mode. The position of the beam is indicated by a cursor which is overlayed the store image. The stationary beam can now be moved (pixel by pixel) from position to position, and the EBSP's are displayed on a second monitor for orientation determination. In this way the area examined, i.e. the (sub-)grain size, shape, surroundings and orientation are all shown simultaneously.

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3. EXPERIMENTAL RESULTS 3.1. A general texture and structure characterization

Rolling and annealing textures are given in Figs 4--6. Cold rolling of the directionally solidified high purity aluminium to a reduction of 90% resulted in widely different textures in the two cases (i) the columnar grains normal to the rolling plane (in the following referred to as NRD) and (ii) the columnar grains parallel to the rolling direction (LRD). In the NRD-orientation cold rolling resulted in a strong {112}(11T) texture (the copper component), Fig. 4(a), no other components could be detected from the pole figures. Rolling parallel to the casting direction (LRD) resulted in a more complex texture as shown by the pole figure in Fig. 5(a), where the main components form a strong fibre from the vicinity of the S-position {123}(41~) towards the brass (Bs) orientation {011}(2T1). Annealing at 400°C resulted in a strong cube recrystallization texture in the N R D case as illustrated by the pole figures in Fig. 4(b) and (c). Two recrystallization pole figures have been included in Fig. 4, this is because the coarse grained starting material (columnar grain with diameters in the mmrange) result in rolling and recrystallization textures which may vary slightly from specimen to specimen. The cube texture is the dominating component but a second component (20°-rotated-cube corresponding to {001}(310)) is also frequently observed, Fig. 4(c). These two components are of special interest in the oriented growth/oriented nucleation context as while the cube component has no common (111)-axis with the Cu-component, the rotated-cube indeed is rotated 40 ° around such a common axis. In the LRD orientation, annealing at 400°C resulted in a strong Bs-

TD

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Fig. 4. (lll)-pole figures from high purity aluminium, NRD-case. (a) Cold rolled, (b) and (c) pole figures from different specimens both annealed 6 s at 450°C. Goss orientation bridge in addition to the cube and rotated Goss, {011}(0II>, components, Fig. 5(b). Metallographic examination of the specimen used to

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OVERVIEW NO. 93

RD

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(a)

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ing of the commercial grade aluminium resulted in rather typical textures as shown by the pole figures in Fig. 6(a) and (b). On a coarse scale the recrystallization reaction is rather inhomogeneous in both the high purity and the commercial purity metals, Figs 7 and 8. The two chanelling contrast micrographs in Fig. 7 are from the high purity metal and illustrates the fully recrystallized structure in the N R D orientation, Fig. 7(a), and a partially recrystallized structure in the LRD orientation, Fig. 7(b). The relatively coarse structure (sheet plane section) in Fig. 7(a) display large local variations in grain size, with this variation being on a scale typically reflecting the ingot grain structure. The transverse section micrograph in Fig. 7(b) show a lamella structure where the lamella thickness corre-

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Fig. 5. (200)-pole figures from high purity aluminium, LRD-case. (a) Cold rolled and (b) after 24 s at 400°C. obtain the pole figure in Fig. 5(b) revealed that this condition is only partially recrystallized. The strong Bs component in this pole figure reflects retained rolling texture as confirmed by EBSP analysis. The slight lack of symmetry and the absence of some of the "equivalent components" in the pole figures in Figs 4 and 5 are typical aspects when working with very coarse grained starting material. A more detailed texture analysis of rolled and annealed high purity aluminium in conditions similar to that used in the present investigation can be found in a recent paper by Hirsch e t al. [17]. Cold rolling (95%) and anneal-

t

D

(b) Fig. 6. (I 1l)-pole figures from the commercial purity alloy, (a) as-rolled and (b) after 6 s at 400°C.

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HJELEN et al.: OVERVIEW NO. 93

Fig. 7. Channeling contrast images showing (a) the recrystallized structure in high purity aluminium (NRD-case, sheet plane section) and (b) a partially recrystallized structure in the LRD-case, short transverse section. sponds to the initial grain size subjected to a 90% rolling reduction. This micrograph clearly illustrates the inhomogeneous nature of both the recovery and the recrystallization reactions. Some of the deformed grains are fully recrystallized, some are well recovered while others still have a substructure which on this scale of inspection is almost indistinguishable from that found in the as rolled condition. As opposed to the high purity conditions, recrystallization of the commercial grade aluminium resulted in a uniform and equiaxed grain structure, Fig. 8(c). In contrast to this final uniformity the recrystallization transformation in itself is very inhomogeneous. The micrographs in Fig. 8(a) and (b) are from the same specimen recrystallized to approximately 50% showing that some regions, Fig. 8(b), are almost fully transformed, while other areas, Fig. 8(a), show individual grains growing in untransformed matrix. 3.2. Nucleation o f recrystallization

Cold rolling of the directionally solidified high purity aluminium with the columnar grains oriented normal to the rolling plane (NRD) resulted in a sharp Cu-deformation texture as shown in Fig. 4(a) and subsequent annealing gave a strong cube and a weaker rotated-cube texture components, Fig. 4(b) and (c). In this case the columnar grains extend from

surface-to-surface of the rolled sheet, or, in other words, the old grain boundaries will remain more or less normal to the sheet surface. Therefore, by studying nucleation of recrystallization in the short transverse section, effects due to old grain boundaries can easily be avoided. EBSP-analyses of the deformed structure demonstrates that the rolling has caused the columnar grains to break up into a layer structure of alternating Cu-texture components, with the different layers separated by sharp transition bands. As shown by the channelling contrast micrograph in Fig. 9, upon annealing, cube (or slightly "off-cube") oriented grains nucleate preferentially from these transition bands. Figure 10 is from another NRD-specimen, which has been sectioned normal to the rolling direction. This channelling contrast micrograph shows a wide band of elongated grains all having the same 20°-ND-rotated cube orientation ({001}(310)). The deformation matrix has the same Cu-orientation above, below and in between these grains. Since these elongated grains are uniformly distributed within a band which is several hundred # m wide, they are obviously not formed by the same mechanism as that illustrated in Fig. 9. In addition, beside these two types of nucleation mechanisms which dominated in the NRD-case, a series of other different ones have also been identified in the various conditions and materials investigated. The various mechanisms will now be presented in separate sections below. 3.2.1. 20°-ND-rotated cubes and some associated components. The micrograph in Fig. ll(a) is taken

from a region close to that shown in Fig. 10 (i.e. the NRD-case sectioned normal to the RD-direction). The area covered in Fig. l l ( a ) is of more general interest as in this ease the elongated grains are not all growing with their long-axis parallel (as in Fig. 10), but are distinctively differently oriented. As outlined in Fig. 1 l(a), the grains can be separated into two main categories labelled A- and B-types being parallel to the (T01) and ( 0 I I ) directions respectively. A third group is also present, marked C and extended in the ( i l 0 ) direction. Only a few grains of this latter type can be identified in Fig. 1l(a). As these grains are all embedded in the same Cu-oriented deformation matrix, their directions of growth may simply outline the traces of the three different { 111 }-planes inclined to the (111) section defined by the micrograph. Or in other words, these grains may all be of the 40 ° (111) type growing with the typical ellipsoidal shape defined by the three different (111) rotation axes (see Section 1.1). The EBSP pole figures in Fig. 1 l(b) and (c) confirm this possibility for the A- and B-type grains, but not for the C-type. The A-grains are of the same type as those shown in Fig. 10, rotated approximately 40 ° around the [l]'l]-matrix-axis into the 20°-ND-rotated-cube orientation, {001}(310), (Fig. lib). The B-grains are also of the 40 ° (111) type rotated around [Ill]-matrix-axis. However, the rotation of the B-grains is of opposite sense compared to that of the A-grains, resulting in an orientation

HJELEN et al.:

OVERVIEW NO. 93

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20pro

Fig. 8. Channeling contrast micrographs from the commercial alloy. (a, b) 40% transformed and (c) fully transformed. close to {123}(0327. The few C-grains are of a different type, these are also, strangely enough, rotated 40 ° with respect to the matrix orientation, but the axis of rotation is in their case the TD-axis, and the C-grains are close to the {001}(110) orientation. The sharpness of the 40 ° (111) orientation relationship is illustrated in Fig. 12 where the orientation of a random selection of 110 grains (A, B types) have been classified in terms of their angle of rotation around either one of the axes [1T1] [Fig. ll(b)] or Jill] [Fig. ll(c)]. Except for the few Cgrains which deviate substantially from this orientation relationship, no grains in the micrographs shown in Fig. 11 were detected which deviated more than about 10°-14° from ideality. Note also that the shape of the A and B grains is consistent with the growth-anisotropy-characteristics of 40 c' (111) boundaries, i.e. the mobility of the (111) twist boundary is considerably less than that of the corresponding tilt boundary. In an effort to identify the nature of the nucleation sites responsible for the A-, B- and C-type grains, the untransformed matrix around and inbetween the grains in Fig. 11 has been subjected to a detailed

EBSP-orientation examination. By systematically stepscanning the beam (step length less than 0.5/tm), the local variation in orientation over relatively large distances can be mapped with a high degree of accuracy. A typical result from such a scan is presented in terms of a "continuous-orientation-changecurve" and pole figures in Fig. 13. This "zig-zag" curve represents a recording of the orientation changes from subgrain to subgrain, with reference to a mean orientation. As such a mean reference orientation we have selected a grain at the center of the orientation distribution of all subgrains (i.e. the center of the orientation spread of the texture component, in this case close to the Cu-component). Further, the angular deviation of a specific subgrain (which is recorded on the vertical axis in Fig. 13) represents the angle of rotation around a nonspecified axis which this subgrain in question will have to be rotated to coincide with the reference orientation. With respect to the reference orientation the curve in Fig. 13 simply represent a misorientation-plot. Regarding the orientation of the nonspecified axis of orientation, it follows from the EBSP examination that we have two main axes around

1386

HJELEN et al.: OVERVIEW NO. 93 DEF. SUBSTR

ND A

RECRYST. GRAINS RD

~

TD

lOOpm Fig. 9. Channeling contrast micrograph (short transverse section, NRD-case) showingcube grains growing out of transition bands. EBSP (200)-pole figures to the left and right give the orientation of the deformed structure find the recrystallized grains respectively(10 s at 400°C). which the subgrains are scattered in an alternating way. These axes are either one of the < 111 >-axis or the TD-axis. This rotation-pattern can easily be seen from the EBSP-pole figures in Fig. 13. It follows from these misorientation plots that the orientation spread is not statistical around a mean in a random way but is in the form of local accumulation of misorientation into sharp local orientation gradients which frequently amount to deviations from the mean of 30-35 ° or more. An important result which can be derived from line scan examinations such as those given in Fig. 13 is that the deformation components in general, on the scale of individually growing grains, cannot be precisely defined. An extensive examination of the orientation spread on a local scale in both the columnar grain material and the commercial grade gives a scattering standard deviation around a local mean typically in the range +5-10 °. 3.2.2. The transition band mechanism. The cube:

The channelling contrast micrograph in Fig. 9 is from the short transverse/longitudinal section of a specimen annealed for I1 s at 400°C. The EBSP-pole

figures show that the layers of ND-rotated cube grains are located either between Cu-components of complementary orientations (D1/D2 and D3/D4 in Fig. 9) or alternatively have approximately the same orientation on each side (D2/D3). Cold rolling with the columnar grains oriented parallel to the rolling direction (LRD), the deformation texture becomes more complex as shown in Fig. 5(a), and so does also the annealing texture, Fig. 5(b). It is interesting to note that the EBSP-analysis of partially recrystallized material revealed that cube oriented grains were only found within deformation texture components of the Cu- or ND-rotated Cu orientations. The micrographs and EBSP pole figures in Fig. 14 are from a specimen annealed for 7 s at 400°C. Figure 14(a) shows a sharp transition band where the texture components above and below the band are ND-rotated Cu-components which are mirror oriented with respect to each other. The band in Fig. 14(b) separate Cu-components which are only marginally different. While in the NRD-case the cube texture component is totally dominating this component rep-

HJELEN

et al.:

OVERVIEW NO. 93

DEF. SUBSTR

1387 RECRYST. GRAINS RD

TD

Fig. 10. 20°-ND-rotated-cube grains (EBSP-pole figure to the right) growing in a Cu-oriented matrix (pole figure to the left). High purity material, NRD orientation and shot transverse section normal to RD (10 s at 400°C).

resents only a minor volume fraction in the LRDcase. In both cases, however, the nucleation of cube (or close to cube) oriented grains are found exclusively in association with planar deformation heterogeneities (transition bands) preferentially located in the Cu- or ND-rotated Cu texture components. The same picture is found also in the commercial grade aluminium. Figure 15 shows cube oriented grains growing out of a transition band separating Cu-comportents of complementary orientations. A systematic EBSP analysis of cold rolled (95% reduction) and recovery annealed commercial purity aluminium revealed that such transition bands were a common feature within the copper and ND-rotated copper components. (The recovery annealing caused no sign of recrystallization but was performed to "clean" up the structure for EBSP-analysis.) A step scan over such a band is illustrated in Fig. 16. In traversing the band, the rotation is firstly around the rolling direction, then around the TD-direction through the cube orientation and finally around the rolling direction back into a new stable orientation. Note the sharpness of the transition as only one subgrain (no. 2 in Fig. 16) separates the Cu-component from the close to cube oriented subgrains (nos 3 and 4) at the core of the band. Further the transition from the cube oriented subgrain no. 4 to the slightly ND-oriented copper component (no. 5) below the band occurs over one boundary only. In general the about 50 ° rotation

in orientation from the stable matrix orientation to the core of the band occurred over a distance typically 1 - 2 # m long. In accordance with what was found in the high purity metal cases, cube type transition bands frequently separated Cu-components which were almost indistinguishable in orientation relationships. The cube oriented subgrains within transition bands in the cold rolled condition did not display any systematic rotation (NDrotation) as was some times found in the high purity metal, but were scattered around the ideal cube orientation in a more random way as illustrated by the EBSP pole figure in Fig. 17. As cube oriented transition band nuclei frequently appear as slightly ND-rotated or grow into NDrotated Cu-components, it follows that a (111 )-axisof-rotation orientation relationship between nuclei and matrix quite commonly occurs. However, this orientation relationship is not found in all cases, and if found, this relationship applies to only one of the deformation components involved in the transition, i.e. either the component above or the one below the transition band. Further, in the investigated cases the transition band nuclei are generally found to be rotated 8-10 ° away from the ideal 40 ° (111 )-relationship. This most probably explains why these transition band grains, as opposed to the 40 ° (111 )-grains described in Section 3.2.1, grow with an approximately equiaxed shape (see Fig. 9), even

1388

HJELEN et al.: OVERVIEW NO. 93

RD L

RD !

IX SUBaRAINS ~ +A GRAINS b

,'~" ~ + B GRAINS c

Fig. 11. (a) Channeling contrast micrograph normal to RD section in the NRD-case showing elongated grains of different orientation. Pole (b) refers to A-grains and pole (c) to B-grains. For more details see the text. within the deformation component where the < 111 >orientation relationship applies. The Goss-texture: in aluminium the Goss orientation {011}<100> appears as both deformation and recrystallization components. In the high purity case, LRD-orientations, the Goss component is easily identified in both pole figures in Fig. 5(a) and (b). In the channelling contrast micrograph in Fig. 18, taken from partially recrystallized LRD-material the EBSP pole figures show that both recrystallized grains as well as deformation substructure with the Goss orientation are present. Note that the recrystallized grains are preferentially located in the brass deformation component which in the micrograph in Fig. 18 is nearly fully recrystallized and located between layers of the Goss deformation texture. These Goss-oriented grains appear to be nucleated from the Bs-Goss transition regions and, as clearly seen; some of these grains are strongly elongated in the <110> direction. These grains, however, are not of the 40 ° < 111 > type and it follows that the elongated shape must be due to other mechanisms. In general, the Goss- and

Bs-deformation components are found to be formed in contact with each other. And further, the Goss oriented grains are most frequently found to nucleate and grow in the brass deformation component. The Goss grains nucleate either from the Goss-Brass transition region as shown in Fig. 18, or from true transition bands inside deformed regions having the brass orientation as shown in Fig. 19. In this case, the Bs component has broken up into a deck of alternating Bs-components, ND-rotated with respect to each other. The Goss orientation is at the center of these orientation transitions, and the new recrystallized grains in this region are also of this orientation, Fig. 19. The Goss-orientation is also frequently observed as an annealing texture component in commercial purity aluminium. In the step-scan-EBSP-analysis of such type material in the cold rolled (and partially annealed) condition transition bands similar to that observed in Fig. 19 were frequently observed within the Bs-components. Such a transition band is shown in Fig. 20. As was found for the "cube-transition-

HJELEN et al.: OVERVIEW NO. 93 110 G R A I N S 30

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• ALL GRAINS ® FREE nRAINS

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ROTATION A R O U N D < 111 > [ DEll ]

Fig. 12. Frequency- and size vs orientation distributions of a random selection of grains shown in the micrograph in Fig. 1I. bands" also the Goss-bands in brass either had orientation relationship between the layers on each side which was of a symmetry character or the orientations of these two layers were nearly identical as is the case shown in Fig. 20. The investigated Goss transition bands in the commercial purity aluminium (as cold rolled) seemed to be somewhat broader than the cube bands. Considering that the Bs-Goss transition requires a rotation of 35 ° around the sheet normal, it follows that the orientation gradient associated with the Goss band is much less steep than the case is for cube bands. This fact also explains that it was possible to scan along the rolling direction and remain inside the transition band in the Goss-case. A scan parallel to the rolling direction revealed very small misorientations between neighbouring subgrains inside the bands. 3.2.3. Nucleation o f recrystallization from shear bands. In aluminium, the S-orientation { 123} (41~) is one of the prominent stable rolling texture components. This component, however, is also frequently reported after recrystallization. The pole figure in Fig. 5(a) from LRD-oriented high purity aluminium shows that cold rolling in this case has produced a

1389

strong S-component. An example of recrystallization of S-oriented grains within an S-deformation matrix is shown in Fig. 21. The recrystallization texture [Fig. 21 (d)] is a complementary "mirror" component of the matrix texture [Fig. 21(b)] (the mirror plane parallel to the rolling plane). Notice a high density of faintly visible bands in the deformation substructure. These bands have a 35 ° angle of inclination to the RD, and they are most likely shear bands. A selected area of the partly recrystallized structure is further magnified in Fig. 21. Note that a shear band is localized between the recrystallized grains A and B. The subgrain texture in the shear band [Fig. 21(c)] is identical to the recrystallization texture, showing that the S-component is nucleated in shear bands. It is interesting to note that the orientation relationship between the subgrains inside the shear band and the surrounding S-matrix is of the 40°(111 ) type, as illustrated by the (111)-pole figure, Fig. 21(e). Nucleation of S-oriented grains due to this mechanism was regularly observed in the LRD-case. Also in the commercial purity aluminium did the EBSP-analysis show that S-oriented grains are nucleated within the S-deformation component, obeying the same orientation relationship. The micrograph in Fig. 22 shows a 1 mm thick layer of partially transformed material (LRD-case). This layer can be subdivided into alternating layers of brass- and Goss orientations. Within the thick Goss layer in the top half of the micrograph, recrystallization has resulted in scattered equiaxed grains with a TD-rotated Goss texture, as can be seen from the EBSP-pole figure in Fig. 22. These grains are most likely nucleated from shear bands in the Goss oriented deformation matrix. 3.3. Growth o f recrystallization 3.3.1. Growth o f 40 ° (111) grains. Whether oriented growth is a mechanism of relevance to the 40 ° (111 ) rotated grains described in Section 3.2.1 or not is an interesting question which will be further discussed in a subsequent section. However, in the cases shown in Figs 10 and 11, the only grains present are those (111 ) rotated grains, and it follows that their dominance cannot be the result of a growth competition selection in a classical sense. Nevertheless, a case like the one illustrated in Fig. 11 provides a unique opportunity to evaluate the orientation dependence of the growth rate of (111 ) rotated grains. A random selection of 110 such grains have been studied and their orientation distribution has already been presented in Fig. 12. The size of these grains has also been measured, and the variation in size, d, with orientation is also included in the diagram in Fig. 12. [The grain size is given as the geometric mean d = (d, d2) u3 where dl and d2 are the short and long diameters of the ellipsoidal grains.] In order to separate the effect due to early impingement, the grain sizes in Fig. 12 are given both as an average of all grains (the averaging referring to the average of grains inside

1390

HJELEN et al.: OVERVIEW NO. 93

30 r

0

10urn

J

2O

Z O

~- I0

STEP-SCAN DIRECTION,ND

O Q1 Ig

-10 -20

RD

RD

TD

TD

Fig. 13. Step scan illustrating the variation in orientation from subgrain to subgrain using the Cu-orientation as a reference, taken from untransformed regions in Fig. I 1. The rotations associated with some of the misorientation peaks are illustrated by the EBSP-pole figures. each individual orientation class in the frequency distribution) and as an average of free-to-growgrains. As free-to-grow-grains is understood grains which have more than 50% of their periphery in contact with the untransformed matrix. Although the total number of grains investigated is too low to give a reliable statistical representation, the results in Fig. 12 illustrate a clear tendency, namely that while the orientation selection in terms of frequency is strongly peaked around a 40° rotation, no similar variation in the size distribution of these (111)-rotated grains is evident from Fig. 12. Two cases of oriented growth selection in the true meaning have been observed in this investigation. One case is illustrated in Fig. 23 where the two rows of close to {011}(01T) oriented grains (rotated Goss) are growing from above and below into a Bs-oriented matrix. These grains have a common 40° (111) axis of rotation with the Bs-matrix as shown by the EBSP pole figure in Fig. 23. The growth selection is consistent with that the faster 40 ° (111) tilt boundary is oriented for migration in the vertical direction in the micrograph. Another case is found in the micrograph presented in Fig. 22. Below the wide Goss-oriented layer in this micrograph is a more heavily transformed brass layer, which in addition to the more commonly observed Goss oriented grains marked a in Fig. 22 (Section 3.2.2), also contain grains having the Sorientation. Some of these S-oriented b-grains are

elongated in the (110) direction of the brass oriented matrix, a growth pattern consistent with the 40 ° (111) orientation relationship between the S-oriented b-grains and the Bs-matrix. 3.3.2. Growth o f cube grains. One of the claimed successes of the oriented growth theory is that it explains the origin of the cube texture. However, experimentally there have been few attempts to directly measure and compare the growth rate of cube oriented grains to that of grains belonging to other texture components. The reason for this has been lack of an adequate technique as also mentioned in the introduction section above. Recently, however, attempts have been made in this regard by applying a STEM channeling pattern technique in order to follow the evolution of recrystallization texture. These investigations by Juul Jensen et al. [18] and Nes and Solberg [19] pertain to commercial l l00-series alloys and it follows from both of them that the average size of the cube grains increases during recrystallization at a higher rate than grains belonging to the other texture components. Nes and Solberg concluded that the more rapid increase in the volume fraction of the cube texture component was not due to a higher growth rate but to a non-random distribution of nucleation sites. The non-cube nuclei were clustered in colonies and so impinged on each other earlier than did the cube grains. In the present paper we have followed the evolution of texture in the same alloys as studied by Nes and Solberg, by applying the

HJELEN et al.: OVERVIEW NO. 93 new and more powerful electron back scattering pattern technique (EBSP). The possibility for a growth selection of the cube component during recrystallization of the high purity aluminium has also been investigated. Commercial alloy: almost fully recrystallized material (volume fraction recrystallized 86%) had an average grain size of the non-cube grains of 13/~m while the cube grains were 85% larger, or 24#m. The volume fraction of the cube texture component was about 12%. To find out whether this size difference is due to a faster growth rate of the cube grains or not, partially transformed material has been studied in great detail. As was reported in Ref. [19], recrystallization of the commercial material is very inhomogeneous. This is illustrated in Fig. 8(a) and (b) showing that after 4 s some areas are fully transformed [Fig. 8(b)], while other areas show individual grains growing in untransformed matrix [Fig. 8(a)]. It is evident that with such an inhomogeneous structure the average grain sizes of the different texture components can only be compared in kinetics terms as long as these different components represent equal fractions of grains belonging to transformed regions and more free to grow grains. This important aspect is born in mind in the quantitative analysis of partially transformed material. After two seconds annealing at 400°C about 11% was transformed. At this stage only freely growing grains were measured and such grains were found to be about 12/~m in diameter irrespectively of if they were cubes or noncubes. After four seconds annealing, the fraction recrystallized was 40%, of which 32% was in the form of island of completely transformed material while the remaining 8% was due to nearly freely growing grains. As also mentioned above a nearly free to grow grain has been defined as a grain with more than 50% of its periphery (as seen in a metallographic section) in contact with untransformed matrix. Counting all the recrystallized grains, the average size of the non-cubes was found to be 11 # m while the cube grains were 17 #m. Selecting only the nearly free to grow grains, the average sizes of the non-cube- and cube grains were found to be 15 and 18/~ m respectively. Table 1 lists all the important transformation data for both the partially- and the neary fully transformed conditions. Note that while 85% of the area fraction of cube grains in the 40% transformed material is due to free cubes, the corresponding

1391

fraction for the non-cube grains is only 15%. Or in other words, while the overwhelming majority of all non-cube grains were found in fully transformed regions, less than i of the cube grains were found in such places. Classification problems: in obtaining the distribution data given above one is faced with some classification problems in terms of deciding what is a transformed region or not. The channeling contrast micrographs given in Fig. 8 illustrates this point. Figures 8(a) and (b) are both from material which is about 40% recrystallized. The micrograph in Fig. 8(a) shows isolated grains and colonies of fully transformed materials embedded in untransformed matrix. In this case, untransformed material is relatively easy to identify due to the somewhat blurred microstructure from which sharp diffraction patterns are difficult to obtain. But the microstructure in Fig. 8(b) is less well defined, here local regions look completely transformed with an average grain size of about 4 # m . Whether this is the case or if such regions only are very well recovered subgrain structures is difficult to decide unless a careful investigation of boundary misorientations are undertaken. It follows that in calculating average grain sizes in partially transformed material, one has to apply a subjective judgement in defining what is recrystallized or not, making the data given in Table 1 somewhat uncertain. Super purity material: as seen from Fig. 4(b), annealing for 6 s at 450°C results in a strong cube texture, with pole figure maxima of about 13 times random. These cube grains grow in a Cu-type deformation matrix [Fig. 4(a)], accordingly they do not satisfy any 40 ° (111) orientation relationship. However, in view of the findings in the commercial alloy, it is of general interest to see if the cube grains also in the high purity metal case deviate in size compared to grains of other texture components. The micrograph in Fig. 7(a) shows a typical grain structure in the same specimen as used to determine the pole figure in Fig. 4(b). Note that the grain size varies periodically from area to area. Within relatively large areas the grain size is rather uniform, but with sharp changes from area to area. The size of the "uniform-grain-size-areas" seems to reflect the size of the as cast columnar grains. For a more detailed EBSP- and channeling contrast examinations, two areas were selected, one with a relatively small grain size and one coarse grained area. For each area the

Table 1. Area fractions recrystallized and grain sizes in the commercial alloy after 4 and 8 s at 400°C Area fraction (%) Annealing time

(s) 2 4 8 60

Recrystallization

Cube grains

11 40 86 100

. 2.4 12 11

G r a i n sizes (#m) Free grains

.

Free cubes .

8 ---

. 2 ---

All non-cubes

All cubes

Free non-cubes

Free cubes

17 24 26

12 15 ---

12 18 ---

. 11 13 21

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HJELEN et al.:

OVERVIEW NO. 93 Table 2. Orientation and size ofgrainsin the high purity ma~rial Small grains Large grains Dev. ~om Size Size cube, 0 Number (#m) Number (#ml) 0-9° 26 54 50 125 10~19" t18 52 81 105 20-29° 87 60 27 106 30-39" 70 65 32 107 8 >40 ° 35 72 47 117

grains were separated into classes with respect to their angular deviation from the true cube orientation. The m e a n grain size for each such class was determined a n d the numerical results are listed in Table 2 as well as displayed graphically in Fig. 24. These results are conclusive in the sense that there is no correlation between grain size a n d orientation. The average size of the cube grains is equal to the average size o f any o t h e r texture c o m p o n e n t . It follows from the n u m b e r o f grains vs o r i e n t a t i o n curve in Fig. 24(a) t h a t the peak is n o t at the expected true cube o r i e n t a t i o n but 10 ° off. This slight asymmetry is

also reflected in the pole figure in Fig. 4(b), a n d such asymmetries are typical o f X-ray pole figures from very coarse grained materials.

RECRYST. GRAINS

DEF. SUBSTR

~

'TD

(a) RECRYST. GRAINS

DEF. SUBSTR

~ RD

I

ll}lj/jill RD (b)

Fig. 14. Channeling contrast micrographs showing cube oriented grains growing out of transition bands. LRD-case, short transverse orientation. Pole figures to the left gives the matrix orientation above and below the bands and the ones to the right gives the orientation of the recrystallized grains (7 s at 400°C).

'ID

HJELEN et al.:

OVERVIEW NO. 93

1393

DEF. SUBSTR

RECRYST. GRAINS

TD

Fig. 15. Cube grains growing out of a transition band in the copper deformation texture. Commercial aluminium (3 s at 400°C). 4. DISCUSSION A most conspicuous finding in the present investigation can be summarized as follows: Cold rolling of aluminium to strain in the range from 2 to 3 results in deformation heterogeneities of various kinds (transition bands, shear bands and others) which frequently have been found to have the following feature in common: The substructure located in the core of each of these heterogeneities have a 40 ° (111) orientation relationship with respect to the surrounding matrix. Referring to the oriented growth/oriented nucleation debate, an alternative view now becomes that a 40 ° (11 l ) orientation relationship rather confirms an

oriented nucleation mechanism, than that of oriented growth. It follows that the 60-year-old idea of Burgers and Louwerse [20] may still be of good value. However, what definitely can be concluded is that a < 111)-orientation relationship frequently pertains both to the nature of the nuclei in itself as well as to the subsequent growth behaviour. Most likely, it is this intertwining of the nucleation and growth reactions around a common orientation relationship which has fuelled the scientific debate for so long, and made an entangling of mechanisms so difficult. Just to realize this should bring this debate an important step forward. In the following, the implications of the present findings on our understanding of both the formation mechanism behind the various deformation heterogeneities as well as on subsequent growth of nuclei will be discussed in further detail.

RD BOUNDARY MISORIENTATION [DEC] 54

1 2

5 49

3

37

4 5

.4

T ND

TD

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A

{112~<11T>

Fig. 16. Lattice rotations associated with an EBSP-step scan crossing a cube-oriented transition band located in the Cu-texture component. Commercial purity aluminium, cold rolled and recovery annealed. AM39/7--B

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HJELEN et al.: OVERVIEW NO. 93

4.1. Nucleation o f recrystallization 4.1.1. Nucleation o f "20°-ND-rotated-cubes'" The micrographs presented in Figs 10 and 11 illustrate that within relatively large regions recrystallization can be totally dominated by grains obeying the 40 ° (111) orientation relationship. The question then becomes: Is this a result of the special mobility properties of the 40 ° (111) boundary, or does the sharp frequency distribution in Fig. 12 simply reflect a corresponding frequency vs nucleation-site-orientation distribution? Detailed examination of the orientation distribution in the untransformed matrix did show that local orientation gradients are produced during deformation. As a salient aspect these gradients revealed a common (111) axis of rotation between the heterogeneity and the surrounding matrix. What has not been established so far is the mechanism responsible for such local fluctuations in orientation. One might argue that these heterogeneities most likely are associated with a rather broad scattering range in terms of orientations around the common axis of rotation. The sharp frequency vs orientation distribution of the recrystallized grains (Fig. 12) may then simply reflect a corresponding boundary mobility effect. In this way, arguments supporting oriented DEF. SUBSTR

RD

TD

Fig. 17. (200)-pole figure from a selection of cube-oriented transition band in cold rolle, commercial purity aluminium. growth would be associated also with the nucleation stage, a view most recently proposed by Duggan et al. [21], who labelled this "micro-growth-selection". The problem with their interpretation in the present case RECRYST. GRAINS RD

TO

Fig. 18. Goss-oriented grains growing in a brass deformation component which is located inbetween layers of deformation structure having the Goss-orientation. High purity aluminium, LRD-case, short transverse section mierographs. Pole figures to the left and fight refer to deformation substructure and recrystallized grains respectively (12 s at 400°C).

HJELEN et aL: OVERVIEW NO. 93 is the lack of a corresponding migration rate effect in the grain size vs orientation distribution. In our opinion, the sharp distribution in Fig. 12 is a characteristic aspect associated with the mechanism responsible for the formation of the nucleation sites. Further arguments in support of such a view will be presented in Section 4.1.1 below, which deals more specifically with the possible connection between the 40 ° (111) orientation relationship and the microstructure of the nucleation sites. 4.1.2. Nucleation f r o m transition bands. A transition band separates different parts of an old grain which has split during deformation and where the individual new parts have rotated towards different, but stable, end orientations [22]. A characteristic feature of a transition band accordingly becomes a sharp orientation gradient bridging the two neighbouring texture components. Dillamore and Katoh [23] predicted that during plane strain deformation of f.c.c, metals (high stacking fault energy), transition bands might develop which contained the cube orientation {001}(100) at their center of orientation rotation. Working with cold rolled copper, Ridha and Hutchinson [24] demonstrated for the first time that cube oriented grains did originate from transition bands similar to those described by Dillamore and Katoh. Later, both direct and indirect observations [17, 25] confirmed that a similar type of mechanism was responsible for the nucleation of the cube orientation also in aluminium. The present results complement these earlier findings, making it now possible to draw a more comprehensive picture related to both the conditions under which transition bands are formed, of what type they are, and what they look like.

a

b

c

1395

As a general comment, a reasonable assumption is that transition bands which form in different rolling texture components will display characteristic differences in terms of frequency of occurrence and in the nature of the lattice rotations involved. So far only two types of "crystallographically" different transition bands have been identified, namely the cube type, which appears to form preferentially in the Cuand ND-rotated-Cu-texture components and the Goss type, which is associated with the brass deformation component. The cube transition band: in the present investigation this type of transition band has been found most frequently within the Cu-deformation component or variants of the copper components, most notably ND-rotated copper. Occasionally cube bands have been found also in components more of the Cu/S-type. This is in contrast to a very recent investigation on copper by Duggan et al. [21], where the cube type transition bands in copper were found in regions having the S-orientation. In the present investigation we have been studying conditions where the deformation texture have been of various types, i.e. in the high purity metal, NRD-orientation, the rolling texture was almost exclusively copper, in the LRDcase the S-component dominated while the commercial material had a normal fl-fibre type rolling texture. In all these cases cube bands were easily detected and preferentially located in the Cu-component. The reasons for this difference between copper and aluminium needs further investigation. The findings of Duggan et al. [21], however, is important in showing that transition bands with the cube as the core-orientation is possible also in the S-component.

d

e

RD

Fig. 19. The micrograph shows a partially transformed region where Goss oriented grains grow in a brass texture. Note that the brass component has broken up into a layer structure of alternating orientations, see individual pole figures over the micrograph. Pole figure to the right of the micrograph represent recrystallized grains. For details see the text (10 s at 400°C).

f

1396

HJELEN et al.: OVERVIEW NO. 93 RD BOUNDARY MISORIENTATION [DEG] 1

8

32 10 10 14

2 3 4

~ ND TD

5 6

19

17

7 E

RD

<1"T2> 0 |110t<001> Fig. 20. Lattice rotations associated with an EBSP-step-scan across a Goss-oriented transition band inside the brass deformation texture. Commercial alloy, cold rolled and recovery annealed.

Comparing the results of Ridha and Hutchinson [24] and the present ones, a characteristic difference between the transition bands in copper and aluminium is that while in copper the transition in orientation is spread out extensively over a "stack" of sub-boundaries, in aluminium this transition has collapsed into a few high angle boundaries (see Fig. 17). In a way the transition bands in aluminium recrystallize on a micron scale in the sense that regions of new orientations separated by high angle boundaries are formed during processing even at room temperature. This can only be understood as being a result of extensive dynamic recovery. The situation in copper becomes a little unsettled, however, as in contrast to the findings of Ridha and Hutchinson, Duggan et al. [21] reported that in their case, cube transition bands in copper were separated from the surrounding S-oriented matrix by sharp orientation gradients much like what we find in aluminium. The evolution of deformation texture is deterministic in the sense that given the initial orientation and the mode of deformation, then the subsequent path (or alternative paths) through orientation space is well defined. An important question then becomes: What are the possible precursor orientations which upon subsequent rolling will develop cube-oriented transition bands? Of special interest in this context are the results obtained with the directionally cast high purity aluminium in the NRD-case where the starting texture is a strong, well defined, (001) fibre parallel to the ND-direction. Two extreme orientations in this fibre can be represented by {001}(100) and {001}(110). It is interesting to note that controlled plane strain deformation of aluminium single crystals starting from these orientations [26, 27] cause the crystals to split into bands with the different bands rotating as follows: (i) for the {001 }(110) case

lattice rotations are around the TD-direction towards the two complementary (112) ( I I 1) and (II2) (111 ) [25, 26]. (ii) In the cube case the rotation is firstly around TD reaching {10~}(~01) at a strain of about one, followed primarily by a RD-rotation reaching the vicinity of the S-orientation at a strain of approximately 1.5 [27]. Such rotations towards Cu or S (S/Cu) we expect as a general pattern for the whole fibre of {001 }(hk0) starting orientations, and such a behaviour is clearly born out in the present NRDcase as shown by the pole figure in Fig. 4(a). From this pole figure, only one component is visible, namely the Cu-component. A more detailed ODF analysis of similarly cast and cold rolled (95%) high purity aluminium by Hirsch et al. [17] gave the following volume fractions for the main components: Cu (CuNo)= 50%, Cu/S = 20% and {10~}(~01)= 15%. The {001}(100) cube texture is an especially interesting starting orientation since in this case the cube may remain at the center of symmetry for transitions towards several possible sets of Cu components. Driver and Akef [28] have modelled the texture evolution during plane strain deformation for this special case of starting orientation. In accordance with the experimental observations [27] the modelling predicts that from the initial cube orientation, the lattice rotates firstly around the TD-direction towards {102}(~01) at intermediate strains and then proceeds by a rotation generally around RD, eventually ending up in the stable S or Cu-components. From this initial {001}(100) orientation, by selecting different sets of possible slip planes, the crystal lattice may rotate through orientation space along four different routes towards the four equivalent and complementary {112}(11I) orientations, two of which eventually overlap at large strains to produce two twin related Cu-components. These four different

R

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TE;

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ox,s-y

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rains +matrix e Fig. 21. S-oriented recrystallized grains nucleated from shear bands located in a deformation structure also having the S-orientation. High purity material, LRD-case. For details see the text (9 s at 400°C). 1397

HJELEN et al.: OVERVIEW NO. 93

1398 DEF. SUBSTR RD

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NO

TO

]

200pm

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0

Fig. 22. RecrystaUized grains of various orientations growing in the Goss- and brass-texture components. The grains marked b are due to oriented growth selection, for details see the text. routes are outlined in the stereographic projections in Fig. 25(a) and (b). Figure 25(a) corresponds to a transition band between two complementary oriented Cu-components, examples of such bands are given in Figs 9, 14-17. Alternatively, different parts of a deforming grain may select rotations as shown in Fig. 25(b), in which case the pair of components on each side of the cube band will be crystallographically indistinguishable. Examples of such cases can also be seen in Figs 9 and 16. It is interesting to note that the path of rotations described by Driver and Akef [28] are in accordance with the rotation pattern shown in Fig. 17, i.e. by traversing a band starting in one of the stable Cu-components, the rotation is firstly around the RD-direction, then aroundTD through the cube orientation and finally around RD towards the component on the other side which may be either one of the two possibilities described in Fig. 25(a) and (b). Now, this picture applies to cube grains which split upon deformation. One might argue that when starting from a polycrystalline material of random orientation, the number of such cubes is far from sufficient to account for the density of cube bands found in the as rolled condition. In response to this, it follows

from the Dillamore and Katoh [23] theory that any grain which prior to deformation belongs to the (001)-normal direction fibre, upon plane strain deformation, will split and rotate around the normal and rolling directions in such a way that the cube becomes stabilized at the center of transition. Such a splitting and rotation pattern has recently been reported by Norestad et al. [29] in an examination of copper single crystals cold rolled from a ND-rotated "cube" starting orientation. The present result also provides strong support for such a rotation pattern, as in the NRD-case, the observations of transition bands are far too numerous to be associated only with the relatively small number of grains in the (001) ND-fibre which initially were cube oriented. Goss oriented transition bands: the Goss recrystallization texture has been observed to be nucleated from two types of transition regions. One type is the transition between the Goss and the brass deformation components. In cold rolled aluminium these two components are frequently found to accompany each other (Fig. 22). In such cases new grains with the Goss orientation easily grow into the brass (Fig. 18), the opposite has never been observed. True Goss-

HJELEN et al.: OVERVIEW NO. 93

1399 RECRYST. GRAINS

DEF. SUBSTR

RD

TD

YD

c o m m o n x,:,~: axis -'~ ~

I

,/

- grains *matrix

111- pole tig.

Fig. 23. Oriented growth selection of slightly TD-rotated {ll0}(IT0) grains growing into brass deformation texture. The 40° (111 ) orientation relationship is illustrated by the (111)-pole figure below the micrograph (12 s at 400°C).

oriented transition bands are regularly observed inside the brass component, Fig. 19. These bands are most likely formed in grains which during deformation on their way to the brass orientation pass through the Goss-orientation which again subsequently splits into alternating layers of brass, NDrotated with respect to each other and with the Goss orientation at the center of symmetry between the layers as shown in Fig. 19. 4.1.3. Nucleation from shear bands. Shear bands, being a result of a macroscopic instability during rolling, are of an entirely different nature than transition bands. This is a trivial observation in mechanistic terms, of course, but their differences are of interest in evaluating (comparing) their capacities as nucleation rites for recrystallization. While transition bands are truly intrinsic in character reflecting the crystallographic nature of slip, the shear band activity during deformation is strongly dependent on a range of extrinsic parameters. During rolling, the density of transition bands will be uniquely defined by the polycrystalline nature of the initial state, here interpreted narrowly as the initial grain size and texture. In principle, the formation of transition bands will be independent of the temperature of deformation and metallurgical parameters such as solid solution con-

tent, particle structure etc. Shear banding, on the other hand, is strongly sensitive to state variables as well as a whole range of microstructural parameters. The picture is not clear as to which extent nucleation of recrystallization from shear bands in aluminium will contribute to specific preferred orientations. The present observations (Fig. 21) identifies for the first time in aluminium a specific orientation relationship between the orientation within the shear band and the surrounding matrix. Recrystallization confirms that the grains growing out of such shear bands indeed have this orientation relationship to the surrounding deformation structure, Fig. 21. It is interesting that this shear band orientation is of the same type, but a complementary component of the matrix which in this case is the S-orientation. It is further interesting to note that this orientation relationship between the two S-components is that of a 35 ° rotation (approximately) around the TD-direction. Pictorially, the subgrains within the shear band has rolled over, around the TD-direction as compared to their neighbouring matrix subgrain, a relationship which appears as intuitively plausible. It is interesting to note that such a "rolling-over-mechanism" is the basic idea in the original oriented nucleation mechanism according to

1400

HJELEN et al.: OVERVIEW NO. 93

Burgers and Louwerse [20]. Of more profound interest, however, is the observation that the substructure inside the shear bands display a 40 ° (111 ) orientation relationship to the matrix. Some further comments on this aspect is given in 4.1.4 below. This type of shear-band nucleation mechanism was frequently observed within the S-deformation texture component both in high purity and commercial grade material. One might expect that shear bands located within other deformation texture components would exhibit a similar orientation relationship, but the only evidence for such a mechanism outside the S-components so far is the TD-rotated Goss grains scattered around in a Goss matrix as shown in Fig. 22. 4.1.4. The nucleation sites and the 40 ° (111) orientation relationship. A remarkable result of the present

EBSP-analysis is the frequent observations of a 40 ° (111) orientation relationship between various types of nucleation sites and the matrix. Sites as different as shear bands and transition bands both may result in this orientation relationship. In line with these observations in aluminium, Duggan et al. [21] have also found that in copper, cube oriented transition bands may form within the S-texture components, and it is well established that the cube-S orientation relationship is of this 40 ° (111) type. It follows that for a selection of important recrystallization texture components in aluminium (and copper), the celebrated 40 ° (111) orientation relationship primarily becomes a property intrinsic to the nature of the nucleation site. To which extent the special growth characteristics of grains with this orientation relationship makes an additional contribution to the texture evolution is an aspect which will be discussed in Section 4.2.2. In a way, these observations close the circle in the meaning bringing us back to the old mechanism of oriented nucleation originally proposed by Burgers and Louwerse [20]. The present observations give, for

I

I

I

the first time, solid experimental support for Burgers idea (later refined by Burgers and Tiedema [30]). For an extensive discussion of the oriented nucleation mechanism, see the excellent review by Beck and Hu [31]. In this review, Beck and Hu discuss several aspects which in their opinion contradict Burgers' theory, the most salient objection being that 40 ° (111) recrystallized grains, corresponding to lowactivity slip planes during the preceding deformation, are frequently observed. With the new insight (due to the EBSP technique) into the turbulent orientation nature of individual texture components (Fig. 13, see also Refs [32] and [33]), large deviations in local orientation from that predicted from the macroscopic shape charge can be expected. In this context it is interesting to note that the common (111)-axis of rotation between the shear-band-site and the surrounding matrix [Fig. 21(e)] corresponds to a lowactivity slip plane. 4.1.5. Other mechanisms. In a recent paper by Berger et al. [34] the possibility of an entirely different nucleation mechanism based on twinning has been raised also for aluminium. This very impressive work clearly demonstrates that annealing twins may form at early stages of recrystallization in aluminium. The Grttingen investigation pertains to aluminium (and copper) single crystals deformed to a strain of 75% in tension. This is a large tensile strain which corresponds to only a plane strain of 0.6. In our context, however, this is a small strain at which level the different rolling components have still not reached their stable orientations, and more important, the deformation heterogeneities which have been identified as nucleation sites for recrystallization are far from developed. Cube oriented transition bands, for instance, requires a strain of at last 1.5 to reach a fully developed stage. As opposed to the tensile case studied by Berger et al., the present investigation shows that in aluminium, rolled to strains in the 120

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Fig. 24. (a) Grain size as a function of misorientation with respect to the cube orientation. (b) Misorientation distribution. High purity aluminium, NRD-case.

60

1401

HJELEN et al.: OVERVIEW NO. 93

[

TT2,I_,?I •-~,.-- 11121 [ 111]

--=,,,--(1 12)[11 1]

Q

Fig. 25. Stereograph projections illustrating the rotations of the ND- and the RD-directions on each side of a cube oriented transition band in a Cu texture component. The (001)(001 ) cube is used as a reference orientation. (a) A case of complementary orientations below and above the band, and (b) the matrix orientations above and below the band are crystallographically indistinguishable. range from 2 to 3, deformation heterogeneities with the appropriate orientation relationships are present in the as deformed state, and further, direct observations confirm that recrystallization indeed grow out from these heterogeneities. Accordingly, we feel that to interpret the evolution of annealing textures in aluminium and aluminium alloys which have been cold rolled to strains of 1.5 and more in terms of a mechanism based on twinning must be considered at the present level of modelling to be both speculative and unnecessary.

4.2. Growth of recrystallization The only indisputable observations of oriented growth found in the present investigation are the cases shown in Figs 22 and 23. The rotated-Gosscomponent visible in the pole figure in Fig. 5(b) is most likely due to an oriented growth selection of {011}(01i) grains growing into the Bs component. It follows from simple inspection that the growth selection mechanism in the cases illustrated in Figs 22 and 23 have a close resemblance to that accomplished in the single crystal experiments [10-12], i.e. selection from a planar recrystallization front. This may be an important mechanism for the evolution of recrystallization textures, especially in metals and alloys which have a coarse grain structure prior to rolling. In such cases, examination of partially transformed conditions frequently reveal that some texture components recrystallize completely, while at the same time no sign of recrystallization can be found in neighbouring texture components, see Fig. 7(b). A planar growth selection of 40 ° (111) grains then becomes possible due to the high mobility of the ( 111 )-tilt-boundary, a case which is nicely illustrated in Fig. 23. 4.2.1. The growth of cube grains. The salient experimental observations are: (i) in the commercial alloy the average size of the cube grains is much larger than

that of the non-cube grains; (ii) in the commercial alloy the sites for the nucleation of recrystallization are non-randomly distributed. In 40% transformed material about 70% of all cubes were classified as free to grow grains while most of the non-cube grains were found in fully transformed regions; (iii) in the 11% transformed material the free to grow cube and non-cube grains were of equal size while in the 40% transformed material the free to grow cube grains are 20% larger than the free non-cube grains which again are about 35% larger than the average size of all grains; (iv) in the high purity material the average cube grain size is about equal to the average size of all grains. The points (i) and (iv) are contradictory in the sense that if (i) is due to an enhanced growth rate of cube grains, no such effect is apparent in the high purity material. On the other hand, by comparing only free to grow grains in partially transformed material the cube grains are only marginally larger than the other grains, however, both categories of these grains are considerably larger than the average size of all grains in the partially transformed state. In conclusion, the larger mean grain size of the cube grains in the commercial alloy may, to the first order, be interpreted as being due to an inhomogeneous distribution of nucleation sites. However, from the small size difference between the free cube and non-cube grains, a slightly more rapid cube growth rate cannot be excluded. The high purity aluminium data on the other hand provides no support for such an oriented growth mechanism. However, the true important result revealed in the high purity metal case is that a sharp cube recrystallization texture will nucleate and grow in a deformation texture having the Cu-orientation. This reaction excludes any involvement of a 40 ° (111) mechanism and can only be understood in terms of oriented nucleation. Although it may be argued that the cube texture may be due to various mechanisms

1402

HJELEN et al.: OVERVIEW NO. 93

in various cases [21, 35], there seems to be no reason to include other mechanisms than nucleations from transition bands in the cases of pure aluminium and aluminium alloys. 4.2.2. On the transformations p o t e n t i a l o f 40 ° ( 1 1 1 ) grains. Although no systematic efforts have been

made in the present investigation to compare the growth rate of 40 ° (111) grains to that of grains belonging to other texture components, the frequent observations of such 40 ° (111) type grains provide an opportunity for some comments on their growth potential, at least in a qualitative sense. Especially the "20°-ND-rotated-cubes '' are of interest in this context. The orientation distribution of these grains have been carefully examined and shown to be strongly peaked around 40 ° (111), Fig. 12, but no similar variation in their size could be detected. It is interesting to compare these results to those of Liebman et al. [10] which were covered in detal in Section 1.1. However, before elaborating on this comparison, a comment on the orientation-relationship-issue in itself needs to be made. It follows from the detailed EBSP-examination of the deformed matrix, Fig. 13, that the orientation of a specific texture component cannot be sharply defined in terms of a precisely indexed orientation. The typical picture is that even on a micron level the subgrain orientations are scattered around a mean orientation with a standard scattering deviation in the range ___5-10°. Similar results have been found by Hjelen et al. [32] in channel die compressed aluminium single crystals and by Orsund et al. [33] in cold rolled commercial aluminium. It follows that in aluminium deformed to a strain of E = 1 or more, no texture component can be identified to a sharpness better than about ___7 °. In terms of the Liebman et al. investigation (see Section 1.1), this corresponds to cases where the 40 ° (111) grains would have angular deviations from ideality of ~k (or ~b) ~ + 7 °, which cause substantial changes in mobility compared to the ideal case. Liebman et al. demonstrated that by increasing the ff value to about 7, the aspect ratio of a growing 40 ° (111) grain (the q-parameter in Fig. 1) will drop from about 50 (at ~b = 0) to less than 10. This is in accordance with the shape of the "20°-ND rotated-cubes" in Figs 10 and 11 where no grains were found to have aspect ratios larger than 10. It is further interesting to note that it also follows from the Liebman et al. investigation that for 40 ° (111) grains with ~k values > 10° the transformation rate become comparable to that of more "randomlyoriented" grains. This result is also in accordance with the present findings as the transformation rate of the 40°(111) grains analysed in Fig. 12 over a scattering range of about 20 ° around the ideal orientation relationship did not display any measurable size variation. Some recent observations by N~ess [36], working with a commercial AIMnMg alloy are of special interest in this context. The alloy investigated by

N~ess was (following hot rolling) cold rolled to a strain E = 2.2, which resulted in a typical deformation texture as illustrated by the pole figure in Fig. 26(a). Subsequent annealing resulted in a recrystallization texture which in addition to a rotated cube component, {001}(310), also contained a rotated Goss variant, {011}(111), Fig. 26(c). In accordance with the present observations the rotated cube has a 40 ° (111) orientation relationship with reference to the Cu-deformation component, and so has also the rotated Goss component as shown in Fig. 26(d). As the Cu-component is one of the roiling texture components [Fig. 26(a)], this orientation relationship might be interpreted as supporting an oriented growth mechanism. However, N~ess has convincingly demonstrated that an oriented growth mechanism is highly unlikely in the present case. Inspection of the hot band revealed that the investigated alloy never recrystallized during hot rolling, resulting in a total accumulated strain after cold rolling of about 6.5. It follows that in this case the ingot grains (200-300/~m in diameter) after hot and cold rolling were reduced to a pan-cake-shape, approximately 0.5 tim thick, or each texture component would be, on average, less than 1 # m thick. This picture was confirmed by EBSP-analysis [37], Fig. 26(b). This EBSP-pole-figure shows that the whole fl-fiber is present within a sheet of cold deformed metal which is less than 5/~m thick. It follows that during subsequent annealing both the "rotated-cube" and the "rotated-Goss" grains, almost from the nucleation stage on, were exposed to a deformation matrix which in orientation terms covered the total B-fiber. In such a situation the basic idea of oriented growth has no meaning. It is also interesting to note that in the case of the "rotated-Goss" grains, the common (111)-axis of rotation is nearly parallel to the RD-direction, Fig. 26(d), i.e. a case which according to Beck and Hu [31] contradicts Burgers' theory (see Section 4.1.4). However, the observations by N~ess and co-workers [36, 37] invalidates these objections: since in their investigation (Fig. 26) an oriented growth selection does not apply, it follows that the origin of the rotated cube- and Goss components can only be given a rational interpretation in terms of oriented nucleation. The micrograph in Fig. 22 represents a case where the size of 40 ° (111) grains directly can be compared to other freely growing grains. It is true that the 40 ° (111) grains marked b are much larger than the other grains located in the row of generally Goss oriented grains from which these b-grains have grown out. But they are not larger than the strongly elongated Gossgrains (marked a) which also grow into the same Bs-compound as the S-oriented b-grains. Also the elongated Goss-grains which are growing in transition region between the Bs and the Goss deformation components seem to be able to grow at a rate comparable to the 40 ° (111) b-grains. Further, some

HJELEN et al.:

OVERVIEW NO. 93

of the equiaxed rotated-Goss-grains which are located in the Goss deformation texture component in the upper half of the micrograph are apparently the largest ones present. In summary, it was stated in Section 1.1 above that based on the results published so far, no unambiguous evidence exists that the 40 ° (111) grains (when uniformly distributed in space) are capable of transforming the deformation matrix at a higher rate than grains belonging to other texture components. This conclusion has not been altered by the present observations. On the other hand, as the 40 ° (111) grains of various kinds are abundantly present, accounting for a range of different recrystallization components, continued effort should be directed towards clarifying: (i) the nature of the deformation micromechanisms responsible for the creation of beterogeneities with this characteristic orientation relationship and (ii) whether or not the special boundary mobility aspects associated with these grains do facilitate either the nucleation and/or the transformation stage.

1403 5. CONCLUSIONS

1. The cube oriented grains in aluminium are nucleated from transition bands preferentially located in the Cu- or ND-rotated Cu-texture components. The 40 ° (11 l ) relationship has been found to have no effect on the nucleation and growth of cubes in aluminium. 2. Recrystallized grains having the Goss orientation are nucleated from transition bands in the Bs-component or from the transition region between Bs and Goss deformation components. 3. S-oriented grains nucleate from shear bands in the S-deformation component. 4. Several types of deformation heterogeneities capable of acting as nucleation sites for recrystallization, such as transition bands, shear bands and others are found to be associated with a 40 ° (111) orientation relationship to the surronding matrix. It follows that in such cases this orientation relationship becomes a property to be associated primarily with the concept of oriented nucleation.

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1404

HJELEN et al.: OVERVIEW NO. 93

5. F r o m a careful review of the literature and from the present experimental observations no direct evidence has been found supporting the idea that the 40 ° (111 ) grains (when uniformly distributed in space) have a transformation potential exceeding that of grains of other orientation relationships. 6. Isolated examples of a true 40 ° ( 1 1 1 ) g r o w t h selection due to the r a p i d growth of (I1 l)-tiltboundaries have been observed. A characteristic aspect in these cases is a growth selection from a planar transformation front. Acknowledgements--The authors thank The Royal Norwegian Council for Scientific and Industrial Research (NTNF) for financial support. Thanks are also due to Drs W. B. Hutchinson and S. E. N~ess, Professors J. Driver, H. Hu and F. J. Humphreys for stimulating discussions and useful comments. A very special thanks to Mrs I. G. Page for excellent secretarial work.

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