full text

455

ISSN 0031-918X, The Physics of Metals and Metallography, 2009, Vol. 108, No. 5, pp. 455–465. © Pleiades Publishing, Ltd., 2009.Original Russian Text © I.V. Gervas’eva, V.A. Zimin, 2009, published in Fizika Metallov i Metallovedenie, 2009, Vol. 108, No. 5, pp. 482–493.

INTRODUCTION

The nonoriented electrical steel is a magneticallysoft material, which is widely used in electric genera-tors and motors. In the general world production oftransformer steels, the nonoriented steel occupies about70% [1]. Approximately half the electrical energy pro-duced in the world is used in motors; therefore, thenecessity of reducing the energy loss in these steels iseconomically substantiated and stimulates studies onthe improvement of the quality of materials [1–6]. Themagnetic properties of nonoriented electrical steel arecontrolled to the greatest degree by two characteristicsof structure, namely, by the final texture and by the finalaverage grain size. The dependence of steel propertieson the crystallographic texture is caused by the influ-ence of the energy of magnetocrystalline anisotropy onthe magnetic permeability and losses for the magneti-zation reversal. In single crystals of iron, the magneticpermeability is maximum and the losses are minimumwhen the external magnetic field is oriented along thedirections

⟨

100

⟩

. The nonoriented electrical steel is usu-ally employed when the magnetic field is applied in alldirections in the sheet plane. For this case, an ideal tex-ture would be an axial

{100}

⟨

0

v

w

⟩

texture, in whichtwo

⟨

100

⟩

directions in each crystallite are parallel tothe surface. At present, there is no economical methodof the production of such a texture. Therefore, in theentire world there is used an “isotropic” “nonoriented”steel with a multicomponent texture, in which the mag-netic properties are determined as being averagebetween the values observed in the direction of rollingand in the transverse direction. The difference in thesevalues must not be too great and the absolute values ofthe magnetic induction should be as high as possible. In

Russia, the value of the anisotropy of magnetic inductionis indicated in the State Standard (GOST 21427.2-83). Inforeign steel, this difference is specified in the limits of20%; as to the texture, of a common trend is obtaining aminimum fraction of the {111} component and a maxi-mum fraction with a Goss orientation {110} [2]. It isobvious that the cube component also must contributeto an improvement of magnetic properties. An increasein the grain size leads to a decrease in the hysteresislosses and an increase in the eddy-current energylosses. It is assumed that an optimum average size ofgrain is approximately 150

μ

m [2].In this work, we undertook an attempt to study the

influence of the initial chemical composition, structuralstate of the hot-rolled sheet, and method of cold rollingon the texture and the grain size of finished steel.

EXPERIMENTALAs an initial material, we used samples of hot-rolled

strips (HRS) of iron–silicon alloys of two different man-ufacturers (No. 1 and No. 2) 2.2 mm thick. The chemicalcomposition of the samples is given in Table 1.

The usual scheme of production of the isotropicsteel consists in a single cold rolling of a HRS to athickness of 0.5 or 0.35 mm and annealing in a contin-uous furnace. The latter includes the stage of a decar-burizing annealing at a temperature of approximately

860°ë

and the stage of normal grain-growth at a highertemperature of

950 °ë

.The experimental treatments consisted in a change

in the scheme of deformation. In the first variant, therewere compared samples obtained according to the sin-gle-pass and two-pass rolling; in the second variant,

Textural and Structural Transformations in Nonoriented Electrical Steel

I. V. Gervas’eva and V. A. Zimin

Institute of Metal Physics, Ural Division, Russian Academy of Sciences, ul. S. Kovalevskoi 18, Ekaterinburg, 620990 Russia

Received December 26, 2008; in final form, March 3, 2009

Abstract

—The structure and texture in sheets of nonoriented electrical steel of different chemical compositionhave been studied at different stages of technological processing. It has been shown that the use of prerecrys-tallization annealing between the cold-rolling passes leads to an increase in the average size of grain in the fin-ished steel.

Key words

: texture, recrystallization annealing, transformer steel

PACS numbers: 81.40.Ef, 81.40.–z

DOI:

10.1134/S0031918X09110052

STRUCTURE, PHASE TRANSFORMATIONS, AND DIFFUSION

456

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GERVAS’EVA, ZIMIN

samples after usual single-pass rolling and rolling to thesame degree but with an intermediate prerecrystalliza-tion annealing at

450°ë

for 15 min were used.

The samples of the first variant were rolled in theplant laboratory on a mill with a diameter of the work-ing rolls of 180 mm. The degree of deformation

ε

wasmeasured by the relative change in the sheet thickness;the deformation scheme is indicated in Table 2. In thecase of the two-pass rolling, the intermediate annealingwas performed in the same plant laboratory at a temper-ature of

800°ë

in a nitrogen–hydrogen mixture.

The samples of the second variant were rolled on aKvarto-150 laboratory mill with a diameter of theworking rolls of 45 mm at the Institute of Metal Phys-ics, Ural Division, Russian Academy of Sciences, using

single rolling with an intermediate prerecrystallizationannealing in a vacuum at a temperature of

450°ë

for15 min after reaching the degree of deformation of70%. For comparison, samples rolled and annealedwithout an intermediate prerecrystallization annealingwere used (control samples).

The final annealings for conducting the primaryrecrystallization and grain growth in the samples aftertreatment according to the first and second variantswere carried out in a laboratory furnace in a vacuum of

10

–3

mmHg at temperatures of 800, 850, and

900°C

for15 min and at a temperature of

950°ë

using the schemeclose to the plant conditions of annealing (placing intoa furnace heated to

850°ë

, holding for 10 min, andheating to

950°ë

in 50 min).

Table 1.

Chemical composition of the samples, wt %

Variant of steel treatment Si Al C S N Mn P Cr Ni Cu Ti

1 1.54 0.543 0.025 0.006 0.006 0.26 0.062 0.04 0.08 0.08 0.009

2 3.03 0.390 0.005 0.005 0.005 0.19 0.010 0.02 0.01 0.04 0.004

Table 2.

Relative integrated intensity of texture components after annealing in steel No. 1 and No. 2 (type of treatment: I, asingle-pass rolling; II, two-pass rolling)

Sample Temperature,

°

C{110} {200} {222}

I II I II I II I II

No. 1,

0.48 mm,

800 0.34 2.12 3.38 0.73

80%

950 0.08 2.40 5.19 0.48

No. 2

0.48 mm, 800 0.22 1.83 3.12 0.65

80%

No. 2,

0.35 mm,

800 0.10 0.08 1.44 1.01 5.45 6.31 0.28 0.17

I – 84%, 900 0.13 0.14 1.33 1.48 5.47 7.34 0.27 0.22

II – 70 and 46%

950 0.09 0.08 1.25 1.20 7.43 7.03 0.18 0.18

110{ } 200{ }+222{ }

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TEXTURAL AND STRUCTURAL TRANSFORMATIONS 457

The microstructure was revealed using etching ofpreliminarily polished samples in a 4% solution ofnitric acid in alcohol. The preparation of the polishedsections before etching consisted in mechanical grind-ing and chemical polishing in a solution of hydrogenperoxide and orthophosphoric acid. The determinationof the average size of grain and of grain-size distribu-tion was carried out by the intercept method. The struc-ture of the samples was studied by optical metallogra-phy using a Neophot microscope.

The characterization of multicomponent texturesafter cold rolling and recrystallization was performedusing orientation-distribution functions (ODFs)according to the Bunge method [7]. The pole figures ofthe recrystallized samples with small grain were mea-sured in Co

ä

α

radiation on a DRON-2.0 X-ray diffrac-tometer with a GUR-5 goniometer and an automatedtextural attachment. The sample was discretelyrevolved in its plane and was inclined about the hori-zontal axis of the goniometer in steps of

5°

. The incom-plete pole figures were measured in the geometry “forreflection,” in the angular range of up to the inclinationangle of

65°

. The estimation of the quantity of texturalcomponents in the recrystallized samples with a coarsergrain was carried out by the method of determining therelative integrated intensity of radiation from the planes{110}, {100} and {111}. The measurements were per-formed on a DRON-3M X-ray diffractometer in themolybdenum radiation by scanning

θ

–2

θ

intervals ofangles in the vicinity of X-ray reflections {110}, {200}and {222}. The areas under the intensity curves (minusbackground) were normalized to the appropriate areasfor a textureless powder sample (standard).

The simultaneous analysis of structure and orienta-tions in the recrystallized and deformed samples wascarried out by the method of the analysis of Kikuchipatterns with the aid of a special attachment (EDAX) ina FEI Quanta 200 scanning electron microscope. Thedeformed samples rolled to the higher degree of defor-mation were preliminarily annealed at temperaturesfrom 300 to

500°ë

to partially eliminate the results ofwork hardening.

RESULTS AND DISCUSSION

Determination of Structure and Texture of the Hot-Rolled Strip

The study of the microstructure of the lateral sectionof the hot-rolled strips of steels produced at the plants 1and 2 showed their essential difference. In the near-sur-face layers of steel 1, the structure of equiaxed grains isobserved; in the central layers of the strip, a typicaldeformation polygonized structure is seen. In steel 2, adeformation polygonized structure is observedthroughout the entire section of the sheet, and theregion of equiaxed grains is absent. In the scientific lit-erature concerning the anisotropic steel, the large roleof equiaxed cube-on-edge grains in the subsequent tex-

tural transformations is emphasized. It is believed thatit is precisely the presence of a layer of cube-on-edgegrains in the initial hot-rolled strip that subsequentlyensures the presence of cube-on-edge nuclei requiredfor the development of a single-component cube-on-edge texture during secondary recrystallization [8, 9].Since for the isotropic steel it is also important toincrease the content of cube-on-edge grains after pri-mary recrystallization, it can be supposed that such astructure can be most favorable for this type of steels aswell. However, the study of texture in the near-surfacelayers of hot-rolled strips showed that in both steels thislayer has a scattered Goss texture (Figs. 1a, 1b); in thetexture, orientations are mainly observed that are closeto

{110}

⟨

112–113

⟩

, whose axes

⟨

001

⟩

are deviated fromthe rolling direction by

20°–30°

.

Texture and Structure after Single and Double Roll-ing.

In our previous work [10], we gave sections of theODFs for the primary-recrystallization texture of theFe–3% Si steel after different schemes of rolling,namely, single rolling to 70 and 90%, and double roll-ing with the reductions by 70 and 50% during the firstand second rolling, respectively. The textures observeddiffer substantially. The presence of a weak cube-on-edge component is noticeable only after double rolling.However, the sharpest cube-on-edge component aftersecondary recrystallization is obtained in the samplesubjected to a 90% reduction, in which there are presentin the greatest quantity

{111}

⟨

112

⟩

and

{113}

⟨

361

⟩

components that are favorable for the development ofthe cube on-edge orientation.

For the isotropic electrical-sheet steel, it is impor-tant to obtain a greater fraction of the cube-on-edgecomponent and a less fraction of the octahedral orien-tation after the stage of normal grain growth. It is so farunclear how the component composition of the textureafter the grain-growth stage will be affected by thepresence of a sharp octahedral component immediatelyafter primary recrystallization.

The steels No. 1 and No. 2 under consideration differquite strongly in chemical composition (see Table 1).First, they refer to different grades from the viewpoint ofsilicon content; in addition, steel No. 1 contains a consid-erably greater amount of impurities. This affects theirgrain size after primary recrystallization and normalgrain growth. Figure 2 displays the dependences of thegrain size on the temperature of annealing in steel No. 1rolled to a thickness of 0.48 mm (

ε

= 80%) and in steelNo. 2 rolled to 0.35 mm (

ε

= 84%). It is seen that at anyannealing temperature steel No. 1 with the greateramount of impurities possesses a smaller grain sizedespite the fact that it was subjected to a smaller degreeof deformation. The smaller grain size in steel No. 1 canbe connected not only with the retarding influence ofimpurities on the motion of grain boundaries. Taking intoaccount that the largest difference in the grain size wasobserved immediately after the completion of primaryrecrystallization at

800°ë

, it can be supposed that this is

458


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GERVAS’EVA, ZIMIN

caused by the greater amount of nucleation sites in thesteel with the smaller grain in connection with the differ-ent initial structure. The equiaxed grains in the near-sur-face zone of steel No. 1 during cold rolling are dividedinto deformation bands, transient bands, and shearbands, in which then there nucleated primarily recrystal-lized grains. On the histograms of the grain-size distribu-tion after annealing at 800 and

950°ë

for steels No. 1 andNo. 2, it is seen that in steel No. 2 the grain growth

occurs more uniformly, whereas in steel No. 1 along withthe retention of a large quantity of small grains, thereappears a significant fraction of larger grains.

The texture of cold rolling and primary recrystalli-zation of steels after different treatments was studiedwith the aid of ODFs. It is known that the most fre-quently encountered orientations in the texture ofdeformation and recrystallization of bcc metals arepresent in the ODF section at

ϕ

2

= 45°

(see Fig. 3).Among these characteristic orientations, it is possibleto distinguish a limited axial component

⟨

111

⟩

, in whichthe normals to the grain planes {111} are parallel to thedirection of the normal to the rolling plane (ND), and alimited axial component

⟨

110

⟩

, in which the corre-sponding directions in the grains are parallel to rollingdirection (RD) in the sheet. These components aresometimes referred to as the axial components

γ

and

α

,respectively. It is natural that useful orientations in thesheet of isotropic steel after a finishing treatment wouldbe orientations with the planes {100} parallel to therolling plane (the upper side of the square of the sectionof the space of the Euler angles, the axial component

η

)and with the planes {110} parallel to the rolling plane(the lower side of the square of the section of the spaceof the Euler angles), since in the grains with such orien-tations the most frequently encountered easy axis is

⟨

100

⟩

. The planar orientation {111} typical of therecrystallization texture of the iron–silicon alloy will bemost unfavorable.

The steel with 3% silicon is usually produced with athickness of 0.5 and 0.35 mm, and the steel with 1.5%Si, with a thickness of 0.65 or 0.5 mm. Figure 4 dis-plays the textures of rolling and recrystallization of dif-ferent steels. Figure 4a depicts an ODF section for thesample of steel No. 1 rolled to 0.48 mm (with a degreeof deformation

ε

= 80%). After such a treatment, the

RD

(a)

(b)

RD

Fig. 1.

Texture of the hot-rolled strip of (a) steel No. 1 and(b) No. 2 at a distance equal to 8% of the strip thicknessfrom the surface. The pole figures {110}.

Gra

in s

ize,

μ

m

50

45

40

35

30

25

20

15

10

5

0750 800 850 900 950 1000

Annealing temperature,

°

C

Fig. 2.

Variation of the grain size with an increase in theannealing temperature for (

�

) steel No. 1 (single rolling to0.48 mm) and (

�

) steel No. 2 (single rolling to 0.35 mm).


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steel has a sufficiently sharp texture with typical limitedaxial orientations

γ

and

α

. In the case of the axial com-ponent

γ

, the greatest density is characteristic of the

{111}

⟨

110

⟩

orientations. After the primary recrystalli-zation (Fig. 4b), the texture becomes more diffuse. Aswe noted in our previous work [11], which concernstextural transformations in anisotropic steel, in the lim-its of the axial component

γ

there occurs a redistributionof orientations, and the

{111}

⟨

112

⟩

orientation becomestrongest. Figures 4c and 4e display the rolling texturesof a steel with 3% silicon rolled using one pass (to 84%)and two passes (with reductions of 70 and 46%). Aftersingle rolling, there is present a limited axial texture

⟨

110

⟩

with a strong predominance of cubic orientationtypical of high degrees of deformation. After doublerolling, both axes (

γ

and α) are observed, but, in con-trast to the texture shown in Fig. 4a, in the axial γ com-ponent the strongest orientation is {111}⟨112⟩. As aresult of recrystallization after the two-stage rolling(Fig. 4f) the axial γ component becomes predominant,which, of course, must negatively affect the propertiesof the isotropic steel. The presence of weaker planarcomponents {110} and {100} can be estimated fromthe relative integrated intensities. Table 2 gives suchdata for the planar orientations {110}, {100}, and{111}, as well as the relationship between the amountsof favorable and unfavorable components. It is seen thatafter primary recrystallization (annealing at 800°ë) thegreatest amount of most favorable orientations {110}and {100} and the least amount of unfavorable {111}orientations are observed in steel No. 1, which containsless silicon and was rolled once to 80% (the ratio ofthese orientations is 0.73). In steel No. 2 with a siliconcontent of 3%, after the same treatment the ratio offavorable to unfavorable components is somewhat less(0.65). In the same steel rolled to a greater degree ofdeformation (to a thickness of 0.35 mm), the number oforientations {110} and {100} is still less, and that of{111} orientations is greater. It is important to note thatafter the stage of normal grain growth (annealing at950°ë) the texture in virtually all cases becomes lessperfect; in steel No. 1 the ratio of favorable to unfavor-able components decreases from 0.73 to 0.48; after sin-gle rolling of steel No. 2 to 0.35 mm, from 0.28 to 0.18;after two-stage rolling, it simply remains very low.

From these data, we can conclude that the regime oftwo-stage rolling does not lead to an essential improve-ment in the composition of the texture. When compar-ing steels with the low (No. 1) and high (No. 2) contentsof silicon, we can note that upon annealing at 800°ë inthe first alloy there is observed a more favorable rela-tionship between the components. Even after the roll-ing to the same thickness (0.48 mm) the ratio of favor-able to unfavorable components in steel No. 1 is better.This fact is in accordance with reference data for steelswith different silicon contents. In spite of the highervalues of losses for the magnetization reversal in thesteel with the smaller content of silicon, the values ofthe saturation magnetic induction in it is greater, but it

is known that it is the magnetic induction that to a con-siderable degree is determined by the crystallographictexture.

Influence of the Intermediate PrerecrystallizationAnnealing between the Cold-Rolling Passes on the Tex-ture and Structure of Steels. It was established in [12]that the intermediate prerecrystallization annealingafter rolling to 70% in anisotropic steel rolled using asingle-pass regime with a total reduction of 90% leadsto an increase in the quantity of cube-on-edge and octa-hedral components in the texture of primary recrystalli-zation. In the anisotropic steel this leads to a decreasein scatter and an increase in grain size after secondaryrecrystallization because of the greater rate of growthof cube-on-edge embryos into the favorably orientedmatrix. It would be of interest to trace how such a treat-ment would affect the formation of the structure andtexture in isotropic steels. In this work, such a treatmentof the samples of isotropic steel was performed underthe following conditions:

(1) sample No. 1: rolling to 70%; prerecrystalliza-tion annealing (PA) at 450°ë for 15 min; final rolling toa thickness of 0.48 mm (total degree of deformation79.6%);

(2) sample No. 2: rolling to 70%; PA at 450°ë for15 min; final rolling to 0.35 mm (total degree of defor-mation 84%);

(3) sample No. 2: rolling to 70%; PA at 450°ë for15 min; final rolling to 0.25 mm (total degree of defor-mation 87.6%);

Then, the samples were subjected to annealing toensure primary recrystallization (800°ë) and a moreprolonged annealing to ensure normal grain growth(950°ë). For comparison, similar samples were sub-jected to analogous treatments without the intermediateprerecrystallization annealing (control specimens).

Table 3 gives data on the relative integrated intensi-ties of separate planar components, as well as the ratios

î

ϕ1

{001} ⟨110⟩0

⟨100⟩ || RD

η

90°

γ

90°

{113} ⟨110⟩{112} ⟨110⟩

{111} ⟨110⟩

{110} ⟨110⟩ {110} ⟨110⟩

{332} ⟨113⟩{554} ⟨225⟩{111} ⟨112⟩

{112} ⟨111⟩

{001} ⟨110⟩

⟨110⟩ || RD

⟨111⟩ || ND

{111} ⟨112⟩ {111} ⟨110⟩

α

Fig. 3. Positions of ideal orientations in the section of thespace of Euler angles at ϕ2 = 45°.

460

THE PHYSICS OF METALS AND METALLOGRAPHY Vol. 108 No. 5 2009

GERVAS’EVA, ZIMIN

45° 2.0

4.0

6.0

8.0

10.0

12.0

(a) (b)

(c) (d)

(e) (f)

45°

45°45° 4.0

16.0

8.0

20.0

12.0

45°

4.5

1.5

6.0

9.0

7.5

3.0

45° 1.0

4.0

6.0

2.0

5.0

3.0

4.5

1.5

6.0

7.5

3.0

0.8

1.6

2.4

3.2

4.0

4.8

Fig. 4. Texture of (a, c, e) cold rolling and (b, d, f) recrystallization (at 800°ë): (a) steel No. 1, single rolling, ε = 80%; (c) steelNo. 2, single rolling, ε = 84%; (e) steel No. 2, two-stage rolling, ε = 46% during the second cold rolling. ODF section at ϕ2 = 45°.



of favorable to unfavorable texture orientations in thesamples after annealing at different temperatures;Table 4 contains data on the average grain size for thesamples after the above-described treatments (PAs) andfor the control samples (CSs), i.e., without PA.

A common feature of all three samples of differentheats rolled to different degrees of deformation is achange in the average size of the recrystallized grainthat occurs in the samples subjected to prerecrystalliza-tion annealing between the cold-rolling passes. In allthe cases after a treatment that includes a PA the grainsize after primary recrystallization at 800°ë becomessmaller or does not change (in sample No.1). In thecourse of a prolonged annealing at 950°ë, the averagegrain size in the samples that were subjected to a treat-ment including PA grows to greater values than in thecase of control samples.

The amount of the cube-on-edge component afterannealing at 800°ë in samples No. 1 (ε = 79.6%) andNo. 2 (ε = 84%) after treatments that include PA some-what increases; the amount of the octahedral componentdecreases. However, after normal grain growth, just as inthe experiments with a single-pass and two-pass rolling,the texture becomes less favorable; the quantity of thecube-on-edge component decreases and that of the octa-hedral component grows. In this case, the positive influ-ence of PA on the amount of the above components isretained only in the sample No. 2 (ε = 84%).

The structure and orientations of grains after rollingof steel No. 2 to a thickness of 0.25 mm (with a degree ofdeformation of 87.6%) have been investigated by theEBSD method, since after this treatment the grain sizewas maximum after normal grain growth both in the con-trol sample and in the sample whose treatment includedPA. Figures 5 and 6 display the orientation images of the

Table 3. Relative integrated intensity of the texture components after primary recrystallization (800°C) and normal graingrowth (950°C) in the samples subjected to single rolling (SR) and rolling with an intermediate prerecrystallization annealingat 450°C (PA)

Sample Treatment

{110} {200} {222}

800 950 800 950 800 950 800 950

No. 1 0.48 mm 79.6%

SR 0.35 0.11 2.04 2.75 3.80 3.11 0.63 0.92

PA 0.42 0.10 1.82 2.26 3.15 3.45 0.71 0.68

No. 2, 0.35 mm, 84%

SR 0.12 0.03 1.76 1.38 6.15 8.66 0.30 0.16

PA 0.13 0.07 1.78 1.16 5.96 8.45 0.32 0.14

No. 2, 0.25 mm, 87.6%

SR 0.12 0.09 1.66 1.22 6.06 7.47 0.29 0.18

PA 0.09 0.04 1.51 0.76 7.33 12.7 0.22 0.06

110{ } 200{ }+222{ }

-------------------------------------

Table 4. Grain size (μm) in the samples that underwent primary recrystallization (800°C) and normal grain growth (950°C)after single cold rolling (SR) and rolling with an intermediate prerecrystallization annealing (PA)

Temperature, °C

No. 1, 0.48 mm, 79.6% No. 2, 0.35 mm, 84% No. 2, 0.25 mm, 87.6%

SR PA SR PA SR PA

800 8.9 8.9 10.4 8.6 17.8 14.2

950 40.4 47.3 40.7 46.5 53.0 61.7

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GERVAS’EVA, ZIMIN

structure (with the indication of the color of orientationson the inverse pole figure) for the control samples(Figs. 5a and 6a) and samples with a PA (Figs. 5b and 6b)after primary recrystallization (Fig. 5) and normal graingrowth (Fig. 6). From the image of the structure obtainedin the scanning electron microscope, it is also seen thatthe PA between the cold-rolling passes leads to adecrease in the grain size after primary recrystallizationand its increase after normal grain growth. Since thedriving force of normal grain growth is the decrease inthe grain-boundary energy, it is natural that in the case ofthe initially small grain the process occurs more rapidlyand the grains grow to greater sizes. The mechanism ofthe redistribution of orientations during normal grain-growth is not completely clear. In his monograph,Novikov [13] used computer simulation to examine thedifferent variants of the evolution of texture componentsduring normal grain growth. He noted that the texture ofthe matrix is not sharp, even a weak component canobtain advantage as a result of normal grain growth.However, in real metals there exist numerous factors,e.g., the presence of impurity atoms and dispersedphases, which render different retardation effects on theboundaries of different types.

In steel No. 2 rolled to 0.25 mm with the use of a PA,the grain growth strongly increased the fraction of theoctahedral component (see Table 3), but, as was alreadymentioned, virtually all the types of treatment decreasedthe ratio of favorable to unfavorable components. Thefact that during normal grain growth in isotropic steelthe fraction of the octahedral component increaseswhile those of the cube-on-face and cube-on-edge com-

ponents decrease is also noted in some other studies,e.g., in [6]. Possibly, this is connected with the fact thatthe ratio between {111} and {110} orientations is toolarge by the moment of the completion of primaryrecrystallization. As is known, the grain boundariesbetween the grains with these orientations possess highmobility. This explains the development of a single-component cube-on-edge texture during second recrys-tallization in anisotropic steel. The probability for acube-on-edge grain to be surrounded by grains with anoctahedral orientation is considerably greater than forthe contrary case, since the content of the cube-on-edgecomponent in the matrix after secondary recrystalliza-tion is very small. If in such a structure a {111} grainabsorbs a cube-on-edge grain, this will affect the gen-eral texture only a little, but if a cube-on-edge grainbegins increasing at the expense of octahedral grains,the further process will be avalanche-type, since thedriving force of second recrystallization is very large.However, in the case of normal grain growth the situa-tion is different. In was shown in [2] that if the octahe-dral component in the texture of primary recrystalliza-tion substantially predominates over the cube-on-edgecomponent, then during the process of grain growth the{111} grains will absorb the cube-on-edge grains. Inthis case to improve the component composition of thetexture after final annealing it is necessary to attempt tochange the composition of the texture of primaryrecrystallization in such a way as to increase the ratio ofthe cube-on-edge to the octahedral component.

The decrease in grain size after primary recrystalli-zation at 800°ë in the samples that passed a treatment

(a) 100 μm100 μm

RD

PA

(b)001 101

111

Fig. 5. Orientation-image maps of the structure and grain orientations in the samples of steel No. 2 after rolling with a reduction of87.6% and recrystallization at 800°ë: (a) without intermediate annealing; (b) with an intermediate prerecrystallization annealing(a schematic image of orientations of individual grains); the RD is located horizontally.



with a PA can be connected with an increase in thenumber of nuclei capable of growing. As a result ofpolygonization at 450°ë, the cells in the deformedmetal become free of defects, and the misorientationbetween the adjacent subgrains increases, which makesa larger quantity of potential nuclei to become capableof growing. Every and Hartherly (cited in [14]) estab-lished that in low-carbon steel rolled to reductions ofabout 70% the stored deformation energy in regionswith the {111} orientations is greater than in regionswith {110} and {100} orientations. It can be supposedthat upon PA in the temperature range of the occurrenceof polygonization processes the cells with the {111}orientation will be freed of defects more intensely thanothers and will be preserved to higher degrees of defor-mation. This must lead to an increase in the planar com-ponent {111} in the texture of primary recrystalliza-tion. An increase of the number of cells in the deformedmatrix that are relatively free of dislocations and capa-ble of playing the role of nuclei of primary recrystalli-

zation leads to a decrease in grain size. It is interestingthat in steel subjected to PA, colonies of grains with aplanar {111} orientation that are elongated along theRD are observed in the structure after primary recrys-tallization. This is seen in the orientation-image mapsfrom the characteristic colors of the corresponding ori-entations, and also in Fig. 5b showing a schematicimage of grain orientations in the form of bricks. Notethat in this figure the RD is arranged horizontally.

On the other hand, in the process of prerecrystalliza-tion annealing, aging should occur, i.e., the precipita-tion of disperseus particles from the solid solution. Inthe process of the subsequent high-temperature anneal-ing this hampers growth of existing nuclei, which, obvi-ously, increases the temperature of the onset of primaryrecrystallization and results in a finer grain of primaryrecrystallization. That the prerecrystallization anneal-ing slows down processes of recovery and recrystalliza-tion also follows from the electron back-scattering pat-terns of the control sample and the sample subjected to

(a)100 μm

TD

RD

(b)001 101

111

100 μm

Fig. 6. Orientation-image maps of the structure and grain orientations in the samples of steel No. 2 after rolling to a reduction of87.6% and in normal grain growth at 950°ë: (a) without intermediate annealing and (b) with an intermediate prerecrystallizationannealing.

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GERVAS’EVA, ZIMIN

PA, which were annealed after final rolling at a temper-ature of 300°ë to partly remove work hardening. Inthese patterns, separate deformation bands of certainorientations can be distinguished, as well as regions ofthe structure in which the program could not identifythe Kikuchi-line patterns and determine their orienta-tions because of the strongly stressed state of metal. Inthe structure images of the sample whose treatmentincluded PA, the number of such unidentified regions isconsiderably greater, which indicates that the recoveryprocesses in this sample occur at a higher temperature.

Another specific feature of the structure ofdeformed metal subjected to PA is the greater fre-quency of shear bands in the limits of deformationbands. This is seen from the structure-orientation mapsof the samples annealed at 500°ë before the EBSD mea-surements (Fig. 7). From the comparison of Figs. 7a and7b it is seen that the frequency of shear bands in thedeformation bands of cubic and octahedral orientationsis greater in the experimental samples than in control

samples (Fig. 7b). Without additional studies, it is diffi-cult to understand why this occurs, but this explains thedevelopment of finer grain in the material subjected toPA. Figure 5b schematically shows the orientations ofindividual grains in the structure of primary recrystalli-zation after experimental treatment. It is seen that in thecolony of grains with the orientations {111} that iselongated along the RD there are present groups ofgrains that have one or another orientation of two sym-metrical orientations {111}⟨112⟩. It is obvious thatthese grains were formed in the region of the deforma-tion bands with the symmetrical orientations{111}⟨110⟩ from the shear bands with a deviated orien-tation.

The minimum final thickness of sheets produced inthe industry from steel No. 2 with 3% silicon is 0.35 mm.As can be seen from Table 3, the ratio of favorable tounfavorable components after normal grain growth in thecontrol sample and in the sample with a PA in this casediffer by no means strongly. Therefore, in spite of the

(a) 90 μm

TD

RD

(b)

001 101

111

70 μm

Fig. 7. Orientation-image maps of the structure and grain orientations in the samples of steel No. 2 after rolling with a degree ofreduction equal to 87.6%: (a) without intermediate annealing and (b) with an intermediate prerecrystallization annealing. Annealingat 500°ë for stress relieving before EBSD measurements.



absence of a positive effect of the treatment that includesPA on the texture composition, its influence on anincrease in grain size after the final annealing must favor-ably affect steel properties.

Thus, in this work we established that the use of theoperation of prerecrystallization annealing between thecold-rolling passes upon processing of nonorientedelectrical steel does not substantially change the textureof sheet but leads to an increase in grain size after finalannealing, which must have a positive effect on themagnitude of energy losses.

ACKNOWLEDGMENTS

The results on the analysis of structures and orienta-tions based on the Kikuchi patterns were obtained in theCenter of Electron Microscopy, Institute of Metal Phys-ics, Ural Division, Russian Academy of Sciences.

The work was performed according to the plan ofthe Russian Academy of Sciences (themeno. 01.2.00613391) and was supported in part by theRussian Foundation for Basic Research, projectno. 08-02-00327 and by the Integration project of theInstitute of Metal Physics and Institute of Electrophys-ics, Ural Division, Russian Academy of Sciences.

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