On the development of the goss texture in iron-3% silicon

~~rrr meroll. Vol. 32. No. 12 pp. 2l85-2201. 1984 PrInted in Great Britain. All rights reserved

Copyright 0

Oool-6160/84 S3.00 + 0.00 1984 Pergamon Pms Ltd

ON THE DEVELOPMENT OF THE GOSS TEXTURE IRON-3% SILICON S. MISHRAt, Institut fiir Allgemeine Metallkunde

C. DARMANN

und Melallphysik Hochschule Aachen. (Receiued

9 March

IN

and K. LUCKE of the Rheinisch-Westfiilische F.R.G.

Technische

1984)

Abstract-A simple laboratory technique was developed to produce Goss-oriented specimens starting with commercial hot band of F+3.3yOSi. The evolution of microstructure and texture including-its throughthickness variation was followed through the subsequent stages of processing until final secondary recrystallization; the underlying mechanisms are discussed. In addition to pole figure determination the more quantitative and sensitive Orientation Distribution Function (ODF) analysis was applied. The results indicate that the structure of the hot band itself, consisting of large recrystallized grains of Goss orientation in the surface layers, exerts a strong influence on the eventual development of a sharp {01 I}( 100) Goss texture. Shear bands were observed after the first stage of cold rolling, and after the subsequent intermediate anneal, second cold rolling and primary recrystallization, strong (111) 11ND fibres were observed. Through ODF analysis, a small number of (01 I}( 100) oriented grains were revealed in the texture after primary recrystallization, again predominantly near the surface so that by chemical etching many of them could be removed. During the subsequent secondary recrystallization, these “Goss-nuclei” grew at the expense of all other matrix grains. R&smn&Nous avons d&veloppC une technique de laboratoire simple pour fabriqu&r des Clchantillons ayant I’orientation de Goss en partant de bandes commerciales de Fe-3%Si. L’Cvolution de la microstructure et de la texture, y compris leur variation en fonction de I’Cpaisseur, a iti suivie au tours des diffircnts stades du traitement jusqu’d la recristallisation secondaire; nous discutons les mccanismes sous-jaccnts. En plus de la d&ermination de la figure de tiles, nous avons appliqub I’analyse plus quantitative et plus sensible de la fonction de r&partition de l’orientation (FRO). Nos titats montrent que la structure de la bande chaude elle-m2me, qui consiste en gros grains recristallii d’orientation de Goss dans les couches supe&ielles, exerce une influence importante sur le d6veloppement &ventuel d’une forte texture de Goss {Oll}( 100). Nous avons observ6 des bandes de cisaillement apr&s le premier stade de laminage (i froid et de fortes fibres (I 11) 11ND apr&s le recuit intermtiiaire ultieur, le second laminage a froid et la recristallisation prima&. L’analyse FRO a r&ilC un petit nombre de grains orient&s {Ol l}( 100) dans la texture apr&s recristallisation primaire, de nouveau essentiellernent pr&s de la surface si bien que I’on pouvait Climiner bon nombre d’entr’eux par une attaque chimique. Au tours de la recristallisation secondaire uldrieure, ces “germes de Goss” croissaient aux dtpens de tous les autres grains de la matrice. Zmmenfaasung-Ausgehend von kommerziellem Fe-3.3%Si-Warmband wurde eine einfache Labortechnik zur Herstellung Goss-orientierter Proben entwickelt. Die Gefige- und Texturausbildung einschlielllich der Texturvariation iiber die Probendicke wurde wiihrend verschiedena ProzeDstadien bis zur abschlieBenden SekundiimLristallisation verfolgt, und die zugrundeliegenden Mechanismen wurden diskutiert. N&en der Bestimmung von Polfiguren wurde such die quantitative= und empfindtichere Methode der Orientierungsverteihmgsfunktion (ODF) angewandt. Die Ergebnisse weiscn auf den starken EintluD des WarmbandgefIiges, das sich durch groDe nkintallisierte Goss-Khmer in den AuBens&ichten auszeichnet. auf die nachfolgende Entwicklung einer scharfen {01l}(100)-Gosstextur bin. Nach dem ersten Kaltwalzen wurden Scherbiinder und nach dem anschlieBenden Zwischengliihen, dem zweiten Kaltwalzen und der Prim&rekristallisation starke (I 11) II ND-Fasem beobachtet. Durch ODF-Analyse lieI sich au&&m tine kleine Zahl {Oll}(lOO)-orientierter Khmer in der Textur nach Primiirrekristallisation nacbweisen, wiederum vorwiegend nahe der ObertlPche, so daB durch chemisches Abiitzen viele von ihnm mtfemt we&n konnten. Wiihnnd der anschlieBenden Sekundiirrekristallisation wuchsen diese “Goss-Keime” auf Kosten aller anderen Matrixk8mer.

1. INTRODUCTION Textured silicon-steel sheets have been of considerable fundamental and technological interest ever

ton sabbatical leave from the Research and Development Centre for Iron and Steel, Steel Authority of India, Ranchi, as a Humboldt Visiting Scientist.

since a process was invented in 1935 by Goss [l] for the production of material with a sharp {Oll}
2186

MISHRA er cd.: GOSS TEXTURE IN IRON-3% SILICON

discussed by Beck [4], the general pre-condition for secondary re~~stalli~tion is a small primary grain size arrived at through the inhibition of normal grain growth either by second phase particles (“impurity inhibition”) or through the presence of a sharp primary recrystallization texture (“texture inhibition”). Detert (51and Walter and Dunn [6] demonstrated a third mechanism for secondary re~~st~i~ation in high purity Fe-Si alloys which is characterized by the inhibition of normal grain growth because of the thickness limitation effect and by surface energy differences as the driving force [I. For commercial grain-oriented silicon steel May and Turnbull [8] showed that secondary rec~stalli~tion was due to irnpu~~ in~bition caused by a fine disper#on of MnS precipitates. A significant technological advancement was made by Taguchi and associates [9-l l] who developed a process which was based on the inhibiting effect of fine precipitates of Al N assisted by particles of MnS and which led to an extremely high degree of perfection of the (01 l)( 108) texture. Imanaka and Goto et 01.[12,13] developed a highly oriented material by taking advantage of the inhibiting effects of Mn, Se and small amounts of Sb. Following a suggestion by Grenoble [14] that solutes such as S and N are major sources of grain growth i~ibition in low-impu~ty FeSi, Fiedler [ls] used a combination of B, S and N for producing highly textured silicon steel. All these new products are characterized by an average deviation of only 3-4” from the (100) axis in the rolling direction, as compared to about 7’ for conventionally textured material using MnS as the inhibitor. In spite of the many advances in the technology of producing oriented silicon steels, the basic mechanism of the development of the remarkably sharp {011}(100) texture has so far not been unambiguously established, even though it is generally accepted that a minor {Oll>( 100) component must be present in the primary recrystallized matrix [161to serve as the nuclei for secondary recrystallization. Several explanations have been proposed. According to the oriented growth theory by Beck [4,17] and Ibe and Liicke [18], the boundaries of the {Oll)(lOO) oriented grains have on the average a higher mobility than the boundaries of other major components of the matrix structure. According to Dunn [19], oriented nucleation produces large primary grains of the (01 i}( 100) orientation which grow selectively at the expense of other grains (oriented nucleation-selective growth theory). According to Decker and Harker PO], (01 l)
smaller neighbours (geometric grain coalescence theory). A large number of fundamental studies of the origin of the Goss texture has been published in the past. The reason for undertaking a further investigation of this subject is three-fold: (i) Surprisingly, even though it is well established that metallurgic parameters in the early stages of the production process can strongly influence the final texture, the evolution of microstructure and texture through various stages of processing has not been reported adequately; most authors only started with an investigation of the texture of primary recrystallized specimens. Therefore the first purpose of the present work is to carry out a detailed examination of the microstructural and textural changes during all the successive stages of rolling and annealing starting with the hot band. (ii) Although there are strong indications that the texture in a specimen of iron or steel can vary through the thickness of the sheet (e.g. 122-251)and that particularly for the formation of the Goss texture the surface layers are important, in most of the earlier investigations texture was determined either only at the surface or at the half-thickness levei of the specimen: Therefore it is the second purpose to study texture fo~ation at all levels of thickness. (iii) The third purpose of the present study is to apply the rather new Orientation Distribution Function (ODF) analysis technique (e.g. Bunge [26]) to obtain more quantitative and precise information on the textural components present in a specimen. In the last few years this new technique has been extensively and successfully applied in f.c.c. metals and alloys (e.g. (271).Its application to b.c.c. systems and particularly to Fe-%, however, has been limited so far to a very small number of investigations [28-321. In the present paper, discussion of the voluminous complex information contained in the ODF’s is restricted to the more obvious features, particularly to those related to the formation of the Goss texture. Detailed and comprehensive analyses of the structure of selected ODF’s and their implications from the point of view of the general theory of texture development in cold-rolled and annealed steels will be presented in a subsequent paper [33]. The starting material for the present investigation was commerical silicon-steel hot band with MnS as the grain growth inhibitor. The result of a similar investigation for an AIN-inhibited material will be reported elsewhere ]341. 2. MATERIAL AND METHODS 2.I. Materid and meli~ilur~ic~I

freutment

Commercial silicon steel of the following nominal chemical composition (by wt%) was used: 3.32x%, 0.031 C, 0.060 Mn, 0.010 P, 0.023 S, 0.0034 N and

MISHRA et cd.: trace

GOSS TEXTURE

Al. Silicon steel slabs of this variety are nor-

mally preheated

to about 14Oo”C, then hot rolled and

at the end of rolling

operation

700°C

4min

in

vestigation

less than

quenched

[35].

The

to below

present

in-

started with such a 2 mm thick hot band

strip which

was further

processed on a laboratory

scale

in a manner similar to that described in [8] and

[l6].

This simulates

the processing of methods

nor-

in industrial production and was found to produce full secondary recrystallization. Here the hot band was first annealed for I h at 900°C in 10e4 torr vacuum, then cold rolled by 70% to 0.58 mm (roll diameter 25 cm), again annealed for 15 min at 850°C and cold rolled by 50% to a final thickness of 0.28 mm. Subsequently this strip was heated for 30min at 800°C to produce primary recrystallization and then annealed for 32 h at 1050°C in a dry hydrogen atmosphere for secondary recrystallization. The heating rate to the annealing temperature was %5O”C/h. At no stage was the strip chemically etched; neither were the edges sheared after rolling. The cold rolling operation was oriented parallel to the original hot rolling direction and was carried out in light passes using palm oil as the lubricant. Between the passes the strip was reversed in order to achieve symmetry with respect to the rolling and transverse directions. The cold rollin was always homogeneous, i.e. the f,/d ratio ( = P- rAd/d [36], IO= length of contact) was always greater than 1.2 (r is the radius of the roll, d the thickness of the strip and Ad the reduction per pass). Special metallurgical treatments carried out by us will be mentioned wherever relevant. mally applied

2.2. Metallography and X-ray measurements Specimens for metallographic examination were polished mechanically and etched either with nital (3% nitric,acid in methanol) or with a freshly prepared dilute solution of ferric chloride to give better contrast. Vickers microhardness (300 g load) measurements were Performed in special eases. X-ray texture measurements were carried out in the reflection mode (0-8S”) using the Lficke pole figure goniometer [37,38] with Co-K, and an Fe filter; specimens were in the form of rectangular coupons 25 x 15 mm in size. To determine the throughthickness variation of texture in a specimen, it was mechanically ground from one side and etched with nital. The thickness level at which a pole figure was determined was specified by the ratio s between the distance from the specimen centre to the level in question and the half thickness. Since pole figures measured on the actual surface are normally smeared out and not representative (e.g. [39]), in the following for s = 1 not the “absolute” surface but a layer approximately 0.01 mm below it was chosen. For all texture specimens { 1lO}-pole figures were obtained, and additionally (200) and {211} pole figures for the ODF analysis of selected specimens.

IN IRON-3’;/, SILICON

2181

The intensity data were normalized with respect to a random specimen which was prepared from high purity iron-carbonyl powder by pressing in a die (3r/cm’) and sintering in vacuum (lO-4 torr) for 60 min at 850°C. The randomness was checked by determining the pole figures of the specimen which showed constant intensity over the whole orientation

range. The orientations

of secondary

grains were determined

from

2.3.

X-ray

recrystalliqxi Laue

patterns.

Determination qf ODF’s

In the present work, the series expansion method by Roe [40] and Bunge [26,41,42] was used. Here the orientation g of a crystallite is specified with respect to the specimen coordinate system (i.e. rolling direction RD, transverse direction TD and normal direction ND) by the three Eulerean angles cp,, 4, and (p2, and the orientation distribution function f(g) is calculated from the measured pole density distribution by means of series expansion into spherical harmonics using the program system of Pospiech and Jura (43,441. Series truncation was effected at L = 22. Liicke et al. [45,41 recognized that a major drawback of the series expansion method is the appearance of “ghost” intensities. These were shown by Matthies [46] to be due to the omission of the odd terms of the series expansion of the ODF which occur whenever the ODF is reproduced from pole figures. A discussion of the appearance and form of ghosttype errors as well as a method of correcting them was recently given by Liicke et al. [471. Because of the complexity of the present textures, however, such quantitative corrections were carried out only for a few examples, and for the rest it was only qualitatively checked that the measured maxima are not due to ghost intensities. In any case, even the uncorrected information contained in the ODF’s is still much more extensive and more detailed than that which can be derived by the inspection of pole figures. [Also whereas the negative intensities to be seen in Fig. 10(a)-(c) below are due to ghost errors, they are sufficiently small not to affect the discussion.] Due to the cubic crystal and orthogonal sample symmetry, a genera1 orientation possesses 96 symmetrically equivalent positions in Euler angle space. The Euler angle range used here: 0 5 cpI , 4, (p2s 90”, still contains each orientation three times and thus can be divided into the three “basic ranges” shown in Fig. 8(c). The measured ODF f(s) will be presented in the form of contour lines in sections ‘p, = constant (at p, = O”, 5”. . . . 90”) through this Euler angle space, with intensity levels of 1.5,3,5,8, and 12 times random. For certain symmetric orientations 2 or 4 of the symmetrically equivalent positions fall together so that their multiplicity reduces to 48 or 24. For these cases also the observed intensity f’(g) must be reduced by a factor of 2 or 4, respectively, in order to be comparable to that of a general orientation (cf. Table 1).

MISHRA

2188

et ul.:

GOSS

TEXTURE

RESULTSANDDXSCUSSION OF MICROSTRUCIURAL AND POLE FIGURESTUDIES

3.

3.1. Hot band Figure l(a) shows the microstructure of the hot band. It is representative of silicon-steel hot bands in general, with either MI& or AlN grain-growth inhibition. One recognizes that the outer layers are recrystallized whereas the interior has largely retained a deformation structure. Since the average Vickers microhardness values measured near the surface and at the centre were similar (i.e. HV 220 and 246), one can conclude that in the interior considerable recovery has occurred. The Role figures (Fig. 2) indicate a pronounced through-thickness variation of the texture. At the surface (s = I) the large recrystallized

IN IRONJ%

SILICON

grains yield a sharp ~Oll}(l~> texture (with some rotation around RD and ND). Then, especially from s = 0.7 to s = 0.4, a rotation around TD occurs and at mid-thickn~s (s = 0) the texture is characterized by strong (112)( 110) and (OOl)( 110) components. {112}(110) and (OOl}(110) are the usual components for heavily rolled b.c.c. metals [39]. Dunn [48] has shown that a (112x1TOJ-orientedand Hu 1491 that a (OOl)[lIO]-oriented deformed FeSi single crystal recovers completely without resolution. The microstructure, hardness value and the texture observed for the s = 0 level of the hot band are therefore fully self-consistent. Textural rotation through the strip thickness has previously been reported for hot rolled low carbon steel [24] and iron [25]; however, the microstructure of hot rolled low carbon steel is in general much more homogeneous through the thickness, e.g. [50], than noted here. The Goss texture observed at the surface of hot roBed steels can be interpreted [51] 8s a shear texture. The f8Ct th8t the surface layers are fully recrystallixed would suggest that some process of simuhaneous shearing and dynamic in siru recryst8llization takes place as siticon steel is hot rolled.

3.2. First annealing andfitst stage cold rolling Upon annealing the hot band at 900% the recrystallized gr8ins in the outer layers grew to very large sizes, some of them to as much as 0.2 mm, and

up to a depth of s = 0.7-0.5 from the surface. The central portions were par&By recrystalliit there was some indication of many grains having nucleated at former grain boundaries. Analysis of the pole figure results was complicated because of the large grain sizes encountered. Nevertheless, they indicated that here the through-thickness variation from surface to centre is essentially that for the hot band in the as-received condition, i.e. the large grains in the surface exhibiting 8 sharp Goss orientation and ~~1~(110~~~12~(110) components in the centre part. This indicates again that this annealing treatment causes more recovery than genuine recrystallization. After the first stage of cold rolling by 700/,, the microstructure [Fig. l(b)] has an appearance characteristic of shear bands [52] with incIinatio~ of % 35” to the rolling plane. The occurrence of shear bands in deformed silicon steels has not been reported beforehand. In fact, shear banding in b.c.c. metals and alloys appears not to have been adequately documented in the literature. Mathur and Backofen did report 1531the presence of shear bands in a low carbon Al-killed steel at z 32” but claimed that they are not

Fig. I. (a) Microstructure of hot band, cross section, FcCI, etch, 75 x . (b) Microstructure after tirst stage cold reduction, longitudinal section. nital etch. 100 x .

observed until 90% deformation by rolling. Willis and Hatherly 1541, on the other hand, observed shear bands in a Nb-stabilized steel relied 72% at x35”. The deviation from the angle of 45” at which the maximum resolved shear stress lies is rather generally observed also for f.c.c. metals 1521. The reasons for

MISHRA e, ul.:

s= 0.5

GOSS TEXTURE IN IRON-3% SILICON

2189

s=Ok not Fig. 2. {1IO} pole figures

such deviations are not entirely clear, though Dillamore et al, 1521have proposed that the “geometric softening factor” may favour an angle of about 35”. The through-thickness variation of the texture also occurs in the 70% cold rolled sheet (Fig. 3). The centre ischara~te~~againby{~l~(l~O)and~llZ~(llO> ~om~nen~, with the spread and intensity being

for hot

band.

almost identical to that in the hot band and annealed hot band specimens. This can be described as some sort of “texture inheritance” which probably only means that hot band annealing essentially leads to recrystallization in situ SO that prior to the first cold rolling a “rolling texture” still exists in the centre. The texture of the outer layers is similar; it shows strong ~~l~(llO), but only weak {112~(110) and addi-

s=O,8

First

Cold

Roiling

Fig. 3. {I IO} pole figures after first stage of cold rolling.

MISHRA et al.: GOSS TEXTURE IN IRON-3% SILICON

2190

tionaily weak ( ill I(2 i i ). For the in-between layers the pattern is more complex. At s =0.7 a doublet f 1I 1)(21 I} was observed but at the s = 0.7-W levels only one of the two symmetrical {111j(2 Ii) components was detected. As is known from single crystal work (e.g. Dunn [19]) the (Oll)( 100) orientation is unstable during deformation and rotates into the (Ii 1)(211) orientations. Therefore, the large recrystallized grains of Goss orientation in the annealed hot band material might be the origin of the {111}(211) component in the outer and in-between layers of the cold roiled strip. The appearance of only one component of this orientation may be caused by the fact that the Goss grains are very large and that due to deviations from the ideal Goss position, one of the two equivalent {lli}(211) orientations may be preferred. Thus it appears that the inhomogeneity of texture in the rolled sheet is largely due to the inhomogeneity of the texture and the microstructure in tire annealed hot band strip. It is less iikeiy that it is caused by inhomogeneous deformation during cold rolling. This is supported by the work of DHrmann and Liicke [55] on (001)[110] oriented Fe-3.5%Si single crystals showing that deformation by roiling is inhomogeneous only for I,# < 0.2 whereas in the present case tbe ratio was always > 1.2. Moreover, shear deformation during inhomogeneous cold rolling should stabilize the Goss orientation. 3.3. Intermediate annealing and second stage cold rolling After intermediate annealing the microstructure indicated full recrystallization throughout the sheet

thickness, with a grain size of 15-i8pm as is normally observed in tire case of Mn~inhibit~ silicon steel [S]. It might be presumed tirat the shear bands in the cold roiled material aided in bringing about this complete and rather homogeneous recrystallization. The texture of tire specimen was, all the same, somewhat inhomogeneous through the thickness (Fig. 4) and also rather diffuse. The main components at the surface and the central layers were {111)(211) and (001)(110>. The most interesting feature, however, was the presence of the Goss component at s = 0.8. Along with our results for the s = 0.8 level of the cold rolled strip, this is consistent with the single crystal result f19] which showed that the Goss o~enta~on upon cold rolling and annealing recrystallizes back into the original orientation (oriented nucleation-selective growth). After the second stage of cold reduction by SO%, the microstructure, probably due to the smaller degree of roiling, did not show any evidence of shear band fo~ation; the grains were merely elongated along the rolling direction. There was also very littie through-thickness variation of tire texture. The pole figures at ail s values [Fig. 5(a)] were practically identical, exhibiting major (111}(211) and minor (OOl)( 1IO>components as often observed after light cold rolling. 3.4. Primary and seconahry recrystaihkation After primary recrystallization the microstructure was again homogeneous and consisted of grams with a diameter of 1518 pm. As seen in Fig. S(b) the pole figures at s = 1, 0.4 and 0 levels were quali~t~veiy similar; only the sharpness increased from the surface towards the interior. The major component was

s=O.8

s=O.6 Intermediate

Annealing

Fig. 4. (I IO} pole figures after intermediate

annealing.

2191

MISHRA er al.: GOSS TEXTURE IN IRON-37: SILICON

behaviour has been reported by Taguchi CI al. [ 1I] for silicon steel stabilized with AIN. and very recently by Pease ef a/. [59] inMnS-inhibited steel, All this evidence suggests that most of the Goss nuclei are present either in the surface layers or in their immediate vicinity. Such an idea was first tentatively put forward

by Philip and Lenhart [60]. Recently Inokuti

et al. [61], using trancmission Kossel technique, also concluded inantly

that the Goss nuclei are present predom-

close to the surface

stabilized

with Mn,

layers of silicon

Se, and Sb. Matsuo

reported similar conclusions

ef al.

steel

[32]

for an AIN-containing

silicon steel.

4. ODF RESULTS AND DISCUSSION

4.1. General

s=o Second Cold Rollinq

s=o Primary Recrystallization

Fig. 5. (a) {llO} pole figures after second stage of cold rolling; insert shows positions of some ideal orientations. (b) { 110) pole figures after primary recrystallization. {111}(211) and the minor components {111}(110) and {OOl}( 110). Interestingly, { 110) and also (200) pole figures failed to indicate the presence of the Goss orientation. Other investigators (Koh and Dunn [56] and Detert [57’j) have also noted the absence of a clearly defined Goss component in primary recrystallized silicon steel stabilized with MnS, whereas others did find its presence (e.g. May and Turnbull [8] and Fiedler [16]). On annealing at 1050°C full secondary recrystallization occurred as evidenced in Fig. 6(a). The orientations of the grains were verified by the Laue method to be very close to the Goss orientation, i.e. deviating up to about 10” from the ideal position as is commonly observed in industrially processed strip as well. This secondary grain growth was not due to surface energy differences, since our material contained sulfur and (001) oriented grains would be favoured to grow instead of (011) grains, if surface energy is the driving force (see Walter [SS]). On etching the surface of the primary recrystallized strip (0.03mm from each side) before applying the secondary recrystallization treatment, a most interesting result was obtained in that only partial secondary grain growth took place [Fig. 6(b)]. Similar

The ODF’s for various specimens are presented in Figs 8 and 9. They exhibit maxima which are listed in Table 1. Their positions in Euler angle space are summarized in Fig. 7 and the corresponding Euler angles are given in Table 2. While the exact orientations of the maxima are listed in the tables, in the text approximate orientations are used (indicated by x ) if the indices are clearly simpler. Additionally, in Table 1 the intensities of the maxima as observedf(g) and reduced with respect to their multipticityf,(g) (cf. Section 2.3) are given. One recognizes that some components (e.g. {001}(110)) which, due to their symmetry, appear to be the largest by visual inspection of the pole figures or ODF’s [and as indicated also by&I)] correspond in reality [as indicated byf,(g)] to rather small volume fractions.

(4

W Fig. 6. (a) Surface macrograph of secondary recrystallized strip, approx 3 x . (b) Surface macrograph of strip that was chcnkally etched before secondary recrystallization trcatment, approx 2 x

2192

MlSHRA

et cd.: GOSS TEXT ‘URE IN IRON-3% SILICON

10 .20 30 LO 50 60 70 80 90

0

10 @ 20 o30 -

4Ol21h001

All

40. ~+-lollMoOl 50 -'rzt 61+.-_(02i1hOO1

Fig. 7. Positions

of most

pqE;!$ HllOMOl mlNl~o1 -Ia) ,,/lrn)IfOl ~'-4112)11f01 )X1 ml)' / :c-_(lll)~liol

60

of the important orientations (b) cp, = 90”~section.

As will be shown, the ODF data on the whole are consistent with the pole figure results. What is more significant, they enable one to unambiguously resolve many textural components and derive quantitative information as to their relative intensities. This makes it possible, to a fairly large extent, to directly follow the sequence of texture development through successive metallurgical treatments. As discussed by Lilcke (271for f.c.c. metals, here two concepts for describing the ODF’s will be applied: (i) by peak-type components, i.e. more or less isotropic scatter around certain ideal orientations and (ii) by orientation tubes, i.e. by complete or limited fibres. 4.2. Peak-type description The main maxima one finds in the ODF’s (Table 1)are{OO1}(110)and ~{112}(110)whicharechar-

Fig. 8(a)

40 50 60 70 60 90

in Euler space; (a) 9, =O”-section,

acteristic of the cold-rolling texture in steels and result from plane strain compression (Section 3.1). Although {001}(110) exhibits the higher observed intensity, %{11Zj( 110) always has the larger volume fraction due to its higher multiplicity. It is to be noted that this orientation may be shifted from the exact position {112}(110) by up to 4” to (337) (110). It appears doubtful, however, that a discussion of this rather small effect is meaningful. After the first cold rolling and also on subsequent treatments the orientations z{ lllf(211) and {11 l)( 110) appear as strong maxima for both surface and interior. The components %{Oll}(lOO) and (4 4 11) (11118) on the surface of the hot band also exhibit strong maxima (Table 1). The (Oll}(lOO) orientation has been found by many investigators (e.g. [22,25,62]) in the surface texture of hot rolled iron

Fig. 8(b)

MISHRA er al.:

GOSS TEXTURE IN IRON-3% SILICON

2193

Fig. 8(c)

Fig. 8(d)

Fig. 8(e)

Fig. 8(f)

Fig. 8. (a) ODF for s = 1 level of hot band; s = 1 denotes a texture at s z 0.97 (the corresponding pole figure is similar to that of s = 0.9). (b) ODF for s = 0 level of hot band. (c) The three basic ranges of Euler space for ‘p, const. (d) ODF for s = 0 level of annealed hot band. (e) ODF for s = 1 level after first stage of cold reduction. (f) ODF for s = 0 level after first stage of cold reduction.

and steels, but here, interestingly, the maximum intensity for this component is identified as occurring at {011}(911) which is about 9” away from the ideal position, rotated around ND. In contrast to {011}(100) the presence of (4411}(11118) component in the surface texture has never been reported before, as far as we know. Quite obviously the ODF technique utilized in the present investigation, being the first such analysis for hot rolled steels, has aided in resolving the two components.

{Oll}(lOO) and {4411}(11118) are in fact the ones predicted theoretically by Dillamore and Katoh [Sl] for b.c.c. metals deformed under the action of an imposed shear. In this context it is also important to note that Backofen and Hundy [63] and Williams [64] had observed (Oll}(lOO) and {112}(111) (close to {4411}(11118)) orientations in Armco iron upon deformation by torsion [63] and shear [64], res$ctively. Recently, &terle and Wever [65] have reported a strong {Oll}(lOO) component in the

2194

MISHRA et al.: GOSS TEXTURE IN IRON-37; SILICON

near f44 lIf(ll Its> move to (001)(110). As first observed by Inagaki and Suda [67] on ODF’s, these two orientations both rotate to some extent towards {11I}( 1IO) leading to broad scatter along these lines for medium degrees of deformation.

surface texture of Armco iron rolled with high deformation per pass, high total reduction in thickness and elevated rolling temperatures. The actual hot-rolling conditions in the mill are far more severe. Considering all the available information together it would be justjfiable to conclude that the {Sl I]( 100) Goss orientation and the {4 4 1I)( 11 I 18) orientation observed in the surface layers of silicon-steel hot band are shear texture components. The development of the surface texture during the subsequent first cold rolling can again be explained by single crystal ex~~rnen~ [66]: the oblations near (01 I)( 100) rotate around TD towards {111](211) (cf. Section 3.2), whereas orientations

As can be seen in the ODF’s (Figs 8 and 9), the maxima described above are not isolated peaks but more or less parts of extended orientation tubes. The positions of these tubes and the intensity ~st~butions along them will be thoroughly analyzed (also with respect to the theoretical aspects) in a subsequent

Fig. 9(a)

Fig. 9(b)

Fig. U(c)

Fig. 9(d)

4.3. Fibre-type description

MISHRA

er cd.: GOSS TEXTURE

Fig. 9(e)

IN IRON-3U,/, SILICON

2195

Fig. 9(f)

Fig. 9. (a) ODF for s = 1 level after intermediate annealing. (b) ODF for s = 0 level after intermediate annealing. (c) ODF for s = 1 level after second stage of cold reduction. (d) ODF for s = 0 level after second stage of cold reduction. (e) ODF for s = I level after primary recrystallization. (f) ODF for s = 0 level after primary recrystallization.

(i) The first tube corresponds to (110) 11 RD and runs along the line cp, = 0”. (p2= 45” from {OOl}( 1IO) over {112}(110) to about {lll}(l IO) (denoted as the “a-tube” in [33]). It contains the highest intensities of the ODF’s. As shown in Fig. 10(a) for &values ‘260” the orientation density decreases significantly, so that one could speak here of an incomplete fibre component. (The maxima near {441}( 110) are largely due to the ghost phenomenon resulting from the strong {OOl}(llO) and { 112}(110) components). On the other hand, in some cases the above maxima are rather pronounced so that their description as peak-type components is also justified. (ii) The second orientation tube corresponds to ( I I I ) IIND and runs along the line t$ = 55”. cpr= 45 from {lll}(llO) to {111)(211) (“y-tube” [33]). As seen in Fig. IO(b) this fibre contains medium intensities, but, in most cases, at a rather constant level. Furthermore, in a large number of specimens this line practically coincides with the skeleton line, i.e. with the line containing the orientations with maximum intensities in the different sections cp, = constant. Figure 1l(a) and 1l(b), respectively, give some examples. These facts mean that in all these cases one has nearly an ideal (111) I(ND fibre. As in the special case of the hot band surface texture [Fig. 8(a)], for the central layers of the as received hot band, the annealed hot band and after the first stage of cold rolling [Fig. 8(b), (d), (f)] the position of the skeleton

A.M. 32/12-E

line does not agree with that of the fibre. This is due to the fact that in the hot band and in the annealed hot band, very strong {OOl}( 110) and x{ 112}( 110) components were already present which are stable end orientations in silicon steels [48,49] and therefore this texture was essentially “inherited” even after 70% deformation by roiling. An important feature found in Fig. 10(b) is that the intensities of the {II l}( 110) and {lllj(211) orientations always increase upon cold rolling. However, whereas the maximum intensity value is reached for { 1 11}( 110) after the first stage of cold rolling, in the case of {Ill}{21 1) it is reached after the second stage. It is interesting to note that the strength of the (111) IIND fibre increases upon cold rolling whereas it decreases upon annealing [Fig. 10(b)]. This relationship will be discussed in Section 5. (iii) Further fibres can be found in the ODF’s as will also be shown in [33]. Here it shall only be mentioned that the {01 l)( 100) and (44 1l}( 11118) components are not isolated peaks but part of an orientation tube (%-tube” [33]). This can recognized, e.g. in Fig. 8(a) where the tube runs from (01 l}( 100) at {cp, ,& (p2}= {0°, 45”, 90”) through the orientation space with increasing q, and decreasing 4 and q2 to {4 4 1l}( 11 118) at {90”, 27”, 45”) including also the 9” shifted position (011}(911) at {9”, 45”, 90”}. Concerning the literature, the occurrence of a strong (110) 11 RD fibre has often been reported especially for the rolling textures of b.c.c. metals. A fibre of the type (111) IIND has been observed as a major component in various recrystallized steels by many investigators, e.g. [68]. There appears, however,

2196

MISHRA er al.:

GOSS TEXTURE IN IRON-3% SILICON

considerable dispute as to its presence or absence in cold-rolled low-carbon steels (cf. (331). The orientation tube (01 I}( lOO)-(4 4 1l}( 11118) with the intensity maximum at {011)(911) has not been described in the literature before. It would be worthwhile to compare the present ODF results with those reported in earlier ODF research [28-321 on Fe-3% Si. Unfortunately all this earlier ODF research was limited in one sense or another. Thus Morris and Heckler [28,29] and Pospiech et al. [31] used composite specimens and their

AnnealedHot Band

SecondColdRollin

AnnealedHot Band -6’

SetondColdRol\inQ

10

20

30

CO

50

60

70

69

1 J

0

99

Fig. IO(c)

Fig. 10. (a) Variation of orientation density along the line specified by ‘p, = O”,(p2= 45”. (b) Variation of orientation density along the line specified by I#I= SS”, (p2= 45”. (c) Variation of orientation density along the line specified by

Fig. lo(a) HotBond

f(g) 10 8 6

L

lo

20

30

40

54

60

70

60

90 'pt

results are therefore necessarily an average over the whole specimen thickness. The metallurgical treatments used by Pospiech ef al. [31] were, besides, rather different. Shinozaki er al. [30] determined the ODF only at the mid-thickness level of the primary recrystallized strip, Matsuo et al. [32] at a layer l/5 to l/4 of the thickness below the specimen surface. In any case, most. of these investigators observed {111}(211) and {111}(110) as important components in the pfimary recrystallized texture. Remarkably, however, nearly all the earlier investigators failed to document the presence of a (! 11) IIND complete fibre axis, especially for the cold rolled conditions, as very clearly established in the present study. 5. THE MECHANISM OF FORMATION OF THE GOSS TEXTURE

10

20

30

40

SO

Fig. IO(b)

60

70

&I9

1

90

When the microstructural, normal pole figure and ODF analysis results are considered as a whole, a fairly consistent picture emerges regarding the development of the sharp Goss texture in Fe-3’;/,Si. This is summarized in Section 6. One should realize (Table I) that in all cases, except for the hot band and annealed hot band surface, the Goss orientation is a

MISHRA er al.:

GOSS TEXTURE IN IRON-3% SILICON

minor component. In the primary recrystallization specimen its presence could not even be ascertained from the pole figures but only from the present ODF’s [Fig. 9(e) and (f)]. Besides, it is also seen that the intensity of the Goss orientation is always higher at s = I than at s = 0. No Goss orientation was detected after the cold rolling treatments. An ultimate sharp Goss texture obviously requires the preexistence of large recrystallized grains of Goss orientation in the surface layer of the as-received hot band which form there as the result of shearing of the surface during hot rolling and simultaneous dynamic recrystallization. Thus the fact that the shear texture

l

a-

O= 55*,C92=45~i Ill11 I ND FIBRE Ideal Positiirts Positions of Intensity Maxima

64

+

hsiths

ol Intensity Maxima @I

Fig. 11. (a) Skeleton line (i.e. line connecting the orientations with maximum intensities in the different sections q, = const.) for s = 0 level tier second stage of cold rolling. (b) Skeleton line for s = 0 level after primary recrystallization.

2197

of b.c.c. metals is given mainly by {01 I )( 100) is the ultimate reason for the final Goss texture. This mechanism also predicts that during the final secondary recrystallization the formation of the Goss texture should originate at the surface and proceed to the centre. This agrees with the present experimental observations that during all processing stages the Goss component occurs with higher intensity in the surface layers than in the inner layers and that chemical etching of the surface layers results in only

limited secondary recrystallization. It is most interesting to note that at the surface after intermediate annealing and primary recrystallization the Goss orientation has its maximum intensity about 9” away from its ideal position at the hot band surface. The centre layers of the hot band and the annealed hot band which show very sharp textures also exhibit some small intensity near the Goss orientation but with a skeleton line somewhat beneath the Goss orientation tube (i.e. 4 =45”, (p2= O/90”), whereas in the other specimens having Goss intensity the skeleton line always coincides with the Goss tube. In the centre layers of the intermediate annealed and primary recrystallized specimens the Goss maximum occurs at its exact position. The reason for the fact that in the surface layer of the hot band the Goss component is found about 9” from the ideal position is not entirely clear. Possibly the deformation mode is not pure shear but might be more complex. What is significant all the same is that in the surface the go-offset position is retained through successive annealing operations, thus emphasizing again the important bearing the initial texture has on texture development in later stages of processing. This fact is also of great importance for the final deviation between the rolling direction and (001) of the fully processed commercial product. It means that the sharpness of the final Goss texture iS already determined by this initial deviation in the hot band. This represents a sort of “texture memory” through a succession of processing operations. Whereas annealing of the hot band prior to the first stage of cold rolling has been known [16] to help produce a sharper Goss texture during secondary recrystallization than when this step is omitted, the reasons for such an effect have never before been clearly understood. For the first time, results from the present work offer a viable explanation. Annealing of the hot band causes considerable growth of the Goss-oriented grains already present in the surface layers. The effect of the subsequent cold rolling and annealing treatments on the texture of the surface layers then parallels, even though in a qualitative and limited sense, the already mentioned situation in single crystals [19] where the Goss orientation is lost on cold rolling, but reappears on annealing. The occurrence of such a sequence in the present experiments is also clearly seen in Fig. 10(c) which shows the intensity distribution along the line ‘p, = 90”, (p2= 45"(i.e. along (1 IO) 11TD). In particular rather

MISHRA

2198

et al.:

GOSS TEXTURE

IN IRON-3%

SILKON

Table I. Microstructure and texture components {hkl)(uuw) observed in the different layers of the specimen during the various processing stages. h gives the observed intensity of the peak g, and f, that obtained after reduction with rcspeet to its multiplicity (Table 2) Pioceasing stage Hot band (2 mm)

Microstructure

lrkf

fully recrystalliaed, large grains (zO.04 mm)

011

stl:

Surface (s = 1)

Cenlre (s = 0) &!A

hki

uvw

All,

911

15.0/7.5

001 331 III 149 011 015 Of1 011

110

110 001 rotated o-150 around TD

110 001 PI1 1033

24.316.1 I6.2/8. I 6.813.4 4.914.9 4.211.O 3.911.9 2.5/l .2 2.2/l.f

001 112 I49 013 011 Oil 011

110 II0 995 109 Oil 922 IO0

15.213.8 14.617.3 4.914.9 3.y1.6 3.010.7 3.Ofi.5 2.2/0.5

Texture variation qualitatively as above

(equal to 01 I 100 rotated 9” ND) 44 II II II 8

8.5j4.2

.r so.5 partly reerystallizcd r 4 0: deformed structure Annealed hot band (P@“C)

s zz I-0.6:

Mainly GOSS-

extremely large grains (0.2 mm) s 0 0: partly recrystallmed

In-between

ww

orient. (Wk

110 110 995

layUS

First cold rolhng (TO%* 0.6 mm)

Shear bands inclined a35” to rolling plane

III 001 5511 111 144 015

110 II0 II0 I12 011 IO0

%9/3.9 ?.?/I .8 7.?/3.8 7.0/3.5 2.811.4 l.PjO.9

5511 001 144 If1 227 I49 015

110 110 011 112 774 995 I00

15.H7.5 13.313.3 5.212.6 4.912.4 3.6fl.8 3.513.5 2.4j1.2

Primarily one component of the Ill 211 doublet

Intermediate anneal (SSO°C)

Fully

Ill 111 Ill 013 1 I10 I13 011 011

358 110 II2 t 110 911 1033

4.114.1 4.612.3 4.612.3 4.112.0 3.7il.S 3.4&l 3.2il.6 2.5/1.2

001 337 Ill Ill 111 017 011 011

110 II0 358 I12 110 100 922 100

7.2ll.8 4.912.4 4.914.9 4.612.3 4312.1 4.2j2.1 3.4jl.l 2.310.6

DitTuse texture with 011 100 and 013 100

structure,

grain size 15-18 pm

E

Second cold rolling (50% 0.3 mm)

Grains elongated along rolling dire&on

111 112 001 017

112 110 740 100

9.814.9 5.6/2.8 3.9fl.P 2.611.3

III 111 001 II2 Ill 017

358 112 110 110 110 100

8.4/8.4 7.913.9 7.lf1.8 6.713.3 6.7j3.3 2.411.2

very little texture variation

Primary recrystallized (~Oc)

F~IIY recrystallized structure, grain size IS-RI pm

Ill 111 013 5511 Oil 001 001

110 112 100 110 911

7.013.5 5.512.1 4.0/2.0 3.911.9 3.8/l .9 3.5/0.9 3.311.6

III III 337 1110 013 011

112 110 II0 110 100 100

5.512.7 5.2/2.6 4.412.2 3.3/1.6 3.t/t.s 2.llO.5

Texture: sharpness increases towards eultn layer

;:

markedly for the surface layers, on cold rolling the intensity of the (111}(211) component increases while that of the Goss component decreases and vice-versa on annealing. The mechanism for this behaviour can be understood in an approximate way. As in single crystals [ 19,661, during cold rolling Goss crystals rotate into two symmetrically equivalent 1111I(21 1) components. One can then assume that, in a subsequent recrystallization, the remaining parts of the original orientation can act as nuclei for the Goss grains and that these grains are distinguished by preferred growth. The latter is qualitatively consistent with the findings of Ibe and Lticke [IS, 691, again based on single crystal investigations, that in Fe-3%Si a preferred 27” rotational relationship around a (I IO) pole exists for growth selection. in the present case the {111}(211) orientation is related to {011~(100)

by a 35” rotation around the (110) transverse axis (and to (01 I](91 1) by 36” (16 10)). As known from f.c.c. metals, small deviations from the ideal rapid growth relationship (there ~~‘{lll~)should~of minor importance so that also in the present case a high growth rate can be expected. For the component {lll)(llO), in contrast, one finds the rather different relationship 46” (I 312) and 41” (23 12>, respectively. It should be noted that due to the alternation of {111)(211) and {Ollj(lOO) on cold rolling and annealing a growth of the volume fractions of these components can occur: cold rolling increases {II 1}(211) because this is an end orientation for the rolling of the Goss component which, in turn, arose from the prior anneal by growing at the expense of { 111)(21 I). But it must be noted that the degree of the second cold rolling must not be too high. Other-

MISHRA et al.:

GOSS TEXTURE

the first formed {I I I}(21 I) component would rotate in the direction of {I I I)( I IO) [67] so that no preferred growth of the Goss grains would be possible during subsequent annealing and no size advantage of the Goss grains for the secondary recrystallization would be obtained. These relationships seem to be quite similar in the case of the final secondary recrystallization. From the present results it appears that during secondary recrystallization the grains with the (01 I}( 100) Goss orientation grow primarily at the expense of a (lll)IIND fibre of which {ll1}(llO) and {I I I }(211) are the strongest components. The results of Morris and Heckler [28,29] also indicated that the selective growth of the Goss grains takes place at the expense of the {ll1}(2ll) and { 11I}( I IO) matrix orientations. Not quite clear is the role of the ~111}(110) grains. In contrast to {11 I}<21 l), they have no obvious preferred growth relationship with respect to Goss grains; but the consumption particularly of this orientation by Goss grains was observed by Matsumuto et al. [70] and Littmann and Dahlstrom [71] during secondary recrystallization in Fe+3%Si single crystals (containing AIN as the grain-growth inhibitor along with MnS). It appears probable that first the {lll}(211) grains will be consumed because of the favourable orientation relationship and then the { 11l}( 110) grains because of the size advantage the Goss grains have assumed in the meantime. It seems that Nielsen’s theory of geometrical coalescence 1211plays very little role in the formation of a sharp Goss texture as the Goss component is a rather weak component in the primary recrystallized matrix. For different reasons, Pease et al. [S9] have wise

IN IRON-3’:<, SILICON

2199

also ruled out geometrical coalescence as an important mechanism for secondary recrystallization in Fe-3%Si. In order to explain the formation of the remarkably sharp Goss texture, one may then suppose that the Goss-oriented grains which form during primary recrystallization, particularly in the surface layers, possibly have a size factor advantage over grains in other orientations. Houze (as reported by Datta [72]) has concluded that the Goss-oriented grains are in fact somewhat larger on the average than other grains. Recently Inokuti el al. [6l] came to a similar conclusion. The size advantage which could provide the necessary driving force for the subsequent preferential growth of the Goss-oriented grains could most likely be due to the orientation dependent selective growth [4, 17, 18,691 of the Goss component at the expense of the { 1I I}(21 I) component, particularly in the surface layers, during primary recrystallization itself. Indeed our ODF analysis results for the surface layers before and after primary recrystallization tend to support such a view. If, as proposed by Decker and Harker [20], the (01 l}( 100) orientation is the first to recrystallize during primary recrystallization, then the size advantage would even be enhanced.

6.

SUMMARY AND CONCLUSIONS

(a) Hot band as-receiued. In its centre layer the usual rolling texture {112}( 1lo)-{OOl}( 1IO) is observed, but at the surface an orientatiop tube with high intensity near the Goss orientation (rotated by 9” around ND) is found. This Goss orientation is formed during hot rolling at the surface by the action

Table 2. Euler anglea cp,, t$,cp2 and multiplicities m for components listed in Table I. While the Miller indices only give the approximate orientations. the Euler angles give exactly the observed positions

48 48 48 48 48 48 48 24

5511 II2 III 144 011 Ill III 4411 227 149 001

110 110 110 011 011 112 358 III18 774 995 740

z: 48 48 96 48

017

IO0

48

015

100

48

013

100

48

011 011 011 011

100 911 922 1033

24 48 48

48

0 845 0 25 45 0 31 45 0 33 45 0 35 45 0 55 45 0 80 45 0 90 45 30 55 45 22 55 45 42 71 20 43 75 16 41 04 24 CPI+ cpz= 30 0 cp,+cp,= 120 0 0 8 0 0 82 90 011 0 0 79 90 018 0 0 72 90 045 0 945 0 1745 0 23 45 0

45 84 a4 48 72 72 49 68 67 50 68 66 51 66 63 60 55 45 80 46 14 90 45 0 90 55 45 38 55 45 90 27 45 90 22 45 46 66 6 rp,+(Pz=6Q 0 cp, + cpl= 150 0 0 890 9090 8 0 II 90 90 90 II 0 18 90 90 90 18 0 45 90 9 45 90 17 45 90 23 45 90

90 45 90 a2 55 45

61 25 76

082 90 90 079 90 90 072 90 90 90 90 81 90 73 90 67 90

0 82 0 79 0 72 45 45 45 45

2200

MlSHRA et 01.: GOSS TEXTURE IN IRON-3% SILICON

of shear which is due to the high rolling reduction at elevated temperatures. Once the Goss orientation is formed, it recrystallizes by dynamic in situ recrystallization and subsequentiy grows towards the centre layer of the strip thus resulting in large {Oii}( 100) oriented grains. These large Goss grams in the surface layers strongly influence the eventual development of a sharp Goss texture on secondary recrystallization. (b) Annealedhot band. The main effect of annealing prior to cold roiling lies in the further growth of the Goss-oriented grams in the surface layers to very large sizes. In the centre layer recovery and recrystallization take place with only minor changes of the texture. (c) First coid roliing. Here near the surface a rolling with the fibre (iil)IIND (i.e. texture {ii I)(21 l)-{ Ii i}( 110)) is formed. Qualitatively speaking,. the large Goss grams behave like single crystals rotating.dming rolling into (11 I}(21 1) (and from there into ( 111I( 110)). Ad~tionaiiy, clear evidence of shear band formation was obtained. (d) Intermediate anneal. Here several weak orientation tubes were formed apparently due to nucleation in shear bands which leads to a broad range of orientations. Additionally, however, the 9” rotated Goss o~en~tion reappears in the surface layers. Similar to what is known to happen in single crystals, remnants of the Goss grains in the rolled matrix seem to recrystallize back to the original orientation mainly at the expense of {111}(211). (e) Second cold rolling. Here again the above roiling texture occurs but with an increased (II1)(211) component. This is assumed to be formed mainly from the Goss component in the way already described for the first cold roiling. cf) Primary recrystallization. As in the intermediate anneal a broad spectrum of orientations is introdu~d. This includes-as could only be shown by ODF’s-a small volume fraction of 9” rotated Goss-oriented grams which reappear predominantly in the surface layers growing again mainly at the expense of {111}(211). fg) Secondary recrystallization. Grams in the Goss orientation already present in the primary matrix and already having a size-factor advantage grow seiectiveiy at the expense of grains primarily in the {111}(211) orientation before consuming the other grains in the matrix as well. This leads to very large grains with a Goss orientation rotated a few degress around ND. ~ummari~ng it can be said that the ultimate reason for the final Goss orientation is that during hot rolling a shear deformation takes place at the surface of the strip and that the shear texture for b.c.c. metals is (01 I]( 100). During cold rolling the Goss grains rotate into {I 1I f(21 I> and during subsequent annealings reappear growing at the expense of { 111)(21 I). This teads to a size advantage for the subsequent secondary recrystallization.

Acknowle&emenrs-The authors thank DipI.-lng. Martin Hiilscher for help in preparing the manuscript. They acknowledge the financia1 support of the Deutschc For~hun~g~~n~a~ and of the Bund~ministe~um fib Forschung und Technologie, Bonn, and they are indebted to the Computer Centre of the Rheinisch-Westf%lischeTechnische Hochschule Aachen for making available its facilities and for personal assistance. One of the authors (S.M.) expresses his gratitude to the Alexander von Humboldt Stiftung, Bonn, for the award of a Research Fellowship which made this work possible. BKPKRENCES 1. N. P. Goss, Trmrp. Am. Sot. Metals 23, 51 I (1935). 2. R. M. Bozorth, Trans. Am, SW. Mrtals 23, I 107 (1935). 3. C. G. Dunn, in CaId Working of Meta&, p. 113.Am. Sot. Metals, Cleveland,OH (1949). 4. P. A. Beck, A&. Phys. 3, 245 (1954). 5. K. Detert. Aczu metalt. 7, 589 (1959). 6. J. L. Walter and C. G. Dunn, J. MetaLr10,573 (1958). 7. P. A. Beck, J. C. Kremer, J. L. Demcr and M. L. Hofzworth, Trans. Am. Inrt: Min. Engrs 175, 372 (1948). 8. J. E. May and D. Turnbull, Tram=Am. Inst. Min. Ertgrs

212, 769 (1958). 9. S. Taguchi and A. Sakakura. U.S. Patent 3,159,511 (1964). 10. S. Taguchi, A. Sakakura and H. Takashima, U.S. Patent 3,287,183 (1966). 11. S. Taguchi, A. Sakakura, F. Matsumoto, K. Takashima and K. Kuroki, J. &fugn. Muter. 2, 121 (1976). 12. T. Imanaka, T. Kan, Y. Obata and T. Sato, W. German Patent OLS 2i351.141 (1974). 13. I. Goto, 1. Matoba, T. Imanaka, T. Gotoh and T. Kan, in Proc. Conf. Soft hlugn. Mater. (SMM-2, Cardi@, Vol. 2, p. 262 (1975). 14. H. E. Grenoble. I.E.E.E. Trans. Magtt. MAG 13, 142 (1977).

IS. H. C. Fiedier, Meta@. Tranr. A SA, 1307 (1977). 16. H. C. Fiedler, J. appi. Pkys. 29, 361 (1958). 17. P. A. &k and H. Hu, in Recrystailisation.Grain Growthand Textures (edited by H. Margolin), p. 393. Am. Sot. Metals, Metals Park, OH (1966). 18. G. Ibe and K. Liicke, Archs Eiscnhlirt.39,693 (1968). 19. C. G. Dunn, Acta metall. I, 163 (1953). 20. B. F. Decker and D. Harker, J. uppt, Pkys. 22, 900

(1951).

21. J. P. Nielsen, Trans. Am. Inst. Min. Engrs 200, 1084 (1954). 22. N. P. Goss, Trans. Am. Sot. Metals 45, 333 (1953). 23. J. F. Held, %uns. Am. Inst. Min. Engrs 239,573 (1967). 24. H. Takechi. H. Kato and S.- Naaashima. Trans. Am. Inst. Min. &grs 242, 56 (1968). ’ 25. D. S. H~dinott and G. J. Davies, J. Iron &eel Inst. 624 (1972). 26. H.-J. Bunge, Z. Metailk. 68, 571 (1977). 27. K. Liicke. in Textures of Materials (edited bv S.

Nagashima), Vol. 1. p. 14. iron Steel lnstjapan, Tokyo (1981). 28. P. R. Morris and A. J. Heckler, in A&ances in X-ray Analysis, Vol. 11, p. 454. Plenum Press, New York (1968). 29. 3. W. Flowers and A. J. Heckler, I.E.E.E. Trans. Mugn.

MAG 12, 846 (1976). 30. M. Shinozaki. 1. Matoba, T. Kan and T. Gotoh, Truns. Japan Inst. Metals 19, 8.5 (1978). 31. J. Pospiech. D. Schhifer and M. Betel, Akud. Wiss. DDR

16, 39 (1978).

32. M. Matsuo, T. Sakai, M. Tanino, T. Shindo and S. Havami. ihi&127).Vol. 2 D. 918. 33. C. D&m&n, S.‘M&hra and-K. Liicke. To be published.

MISHRA (‘I
2201

(edited by G. Gottstein

55. 56.

57. 58. 59. 60. 61.

62. 63. 64.

and K. Liicke), Vol. 1, p. 465. Springer, Berlin (1978). C. Darmann and K. Liicke, To be published. P. K. Koh and C. G. Dunn, Trans. Am. Inst. Min. Engrs. 218, 65 (1960). K. Detert. Metallurgica 12. 817 (1958). J. L. Walter, J. appl. Phys. 36, 1213 (1965). N. C. Pease, D. W. Jones, M. H. L. Wise and W. B. Hutchinson, Mefals Sci. IS, 203 (1981). T. V. Philip and R. E. Lenhart. Trans. Am. Inst. Min. Engrs 221, 439 (1961). Y. Inokuti, Y. Shimizu. C. Maeda and H. Shimanaka, in Recrysfallization and Grain Growrh of Multi-Phase and Parr Containing Materials (edited by N. Hansen, A. R. Jones and T. Leffers), p. 71. Denmark (1980). R. Jones, J. Inst. Mefals 93, 486 (1964). W. A. Backofen and B. B. Hundy, Trans. Am. Inst. Min. Engrs 197, 61 (1953). R. 0. Williams, Trans. Am. fnsr. Min. Engrs 224, 129

(1962). 65. W. &terle

and H. Wever, Z. Metal/k. 72, 230 (1981). 66. T. Taoka, E. Furubayashi and S. Takeuchi, Trans. Iron Steel Inst. Japan 6, 290 (1966). 67. H. Inagaki and T. Suda, Texrure 1, 129 (1972). 68. S. Mishra, C. Dirmann and K. Liicke, Metall. Trans A 14A, II (1983). 69. G. Ibe and K. Liicke, ibid. [II, p. 434. 70. F. Matsumoto, K. Kuroki and A. Sakakura, A.f.P Con/. Proc. 24, 716 (1974). 71. M. F. Littmann and N. A. Dahlstrom, J. appl. Phys.

49, 2034 (1978). 72. G. L. Houze, as reported by A. Datta, I.E.E.E. Trans. Magn. MAC 12, 867 (1976).