Advanced hot rolling strategies for IF and TRIP steels

(1)

Advanced hot rolling strategies

for

(2)

(3)

Advanced hot rolling strategies

for

IF and TRIP steels

Proefschrift

ter verkrijging van de graad van doctor aan de Technische Universiteit Delft,

op gezag van de Rector Magnificus prof. dr. ir. J.T. Fokkema, voorzitter van het College voor Promoties,

in het openbaar te verdedigen op maandag 20 juni om 13:00 uur door

Alexander ELSNER

ingenieur mechatronica HBO Venlo geboren te Düsseldorf, Duitsland.

(4)

Prof. Dr.-Ing. K. Steinhoff

Samenstelling promotiecommissie:

Rector Magnificus voorzitter

Prof. dr. ir. S. van der Zwaag Technische Universiteit Delft, promotor Prof. Dr.-Ing. K. Steinhoff Universität Kassel (Duitsland), promotor Prof. dr. ir. R. Benedictus Technische Universiteit Delft

Prof. Dr.-Ing. W. Bleck Rheinisch-Westfälische Technische Hochschule Aachen (Duitsland) Prof. dr. ir. L. Kestens Technische Universiteit Delft

Prof. Dr.-Ing. D. Raabe Max-Planck-Institut für Eisenforschung (Duitsland) Prof. dr. I.M. Richardson Technische Universiteit Delft

Published and distributed by DUP Science

DUP Science is an imprint of Delft University Press

PO Box 98 2600 MG Delft The Netherlands Telephone: + 31 15 27 85 678 Telefax: + 31 15 27 85 706 E-mail: Info@Library.TUDelft.NL

keywords: IF steel, TRIP steel, ferritic rolling, intercritical rolling, microstructure, mechanical properties, texture development

ISBN 90-407-2591-8

Copyright c 2005 by A. Elsner

All rights reserved. No part of the material protected by this copyright notice may be reproduced or utilized in any form or by any means, electronic or mechanical, including photocopying, recording, or by any information storage and retrieval system, without written permission from the publisher: Delft University Press.

(5)

Financial support

The present work has been supported by the "Verein zur Förderung von Forschungsarbeiten auf dem Gebiet der Walzwerkstechnik in der Hüttenindustrie (VFWH)", Düsseldorf (Germany).

(6)

(7)

Chapter 1 Scope of the Thesis

The production of mild and high strength steel sheets with a thickness of less than 2 − 3 mm usually consists of slab reheating, hot rolling, pickling, cold rolling and final annealing. Various metallurgical changes take place during the whole production process. By controlling the com-plex interaction of these changes the mechanical and technological properties can be tailored in wide range, depending on the chemical composition of the steel and the processing [1]. A major role for the final properties of the cold rolled strip plays the rolling reduction applied during cold rolling and the design of the annealing cycle.

To maintain or even increase their market share, steel producers are forced to reduce the produc-tion cost on the one hand and increase the performances of the products on the other hand. Next to this, the minimisation of the environmental impact is of increasing importance. Sustainable solutions for these challenges can be only achieved by continuous research and development in all stages of the production process. Some possible areas of further research on this are for example formulated in [2, 3].

A likely method to save on production costs is to cut down the rather long production chain of a conventional cold rolled strip, and to substitute certain cold rolled steel grades by hot rolled steel. Certainly this can be hardly achieved for exposed parts, requiring a perfect surface finish, whereas for unexposed parts, e.g. structural components, with lower surface requirements, the use of hot strips might be a cost saving alternative [2]. To facilitate this substitution a couple of prerequisites have to be fulfilled:

• Strip thicknesses below 1.5 mm or even 1 mm

• Improved mechanical properties, comparable to that of a cold rolled strip • Improved surface quality

• Improved dimensional tolerances

(14)

The dimensional tolerances of hot rolled strip could have been markedly improved in the past decades with the introduction of hydraulic roll adjustment and automated gauge control systems. A further improvement, especially in terms of the strip flatness, could have been achieved by the introduction of the CVC technology (Continuous Variable Crown). The CVC technology utilises S-shaped (or bottle shaped) working rolls [4]. By axial shifting of the work rolls the roll gap geometry can be adjusted to produce the desired strip crown. Together with the roll skewing and bending a markedly improvement of the strip flatness can be obtained.

In conventional hot rolling the main objective is to obtain the desired final thickness with a prefe-rably low rolling force, mostly disregarding the required final mechanical properties. These strips do not usually need any supplemental treatments to meet the desired properties. The industrially applied thermomechanical rolling strategies, so called normalising rolling or thermomechani-cal rolling, allow to produce hot rolled strips with improved strength and toughness properties. The increase in strength and toughness are obtained by grain refinement and precipitation harde-ning. The microstructure of these hot rolled steels is obtained directly without a final annealing treatment [5].

In normalising rolling the finish rolling temperature is in the range of the normalising temperatu-re, i.e. well above the Ar3temperature, close above the austenitic recrystallisation temperature. The properties of the steel strip are achieved by the rapid recrystallisation of the deformed auste-nite and the subsequent α/γ transformation. In the thermomechanical rolling strategy the finish rolling temperatures are shifted near to the Ar3 temperature. By the addition of microalloying elements, such as titanium and/or niobium, the non-recrystallisation temperature, Tnr, is raised, so that the strips are finished rolled in a temperature region where no recrystallisation takes place [6]. In this case the α/γ transformation starts from a deformed austenite microstructure, produ-cing a fine ferrite grain size. The small grain size, together with the precipitation hardening of the microalloying elements lead to a high strength and a good toughness of these strips [7]. Ne-vertheless, the improved strength and toughness of these thermomechanical rolled steel grades may be associated with a decreased ductility.

In conclusion, the current industrially applied hot rolling strategies mentioned above are unsuit-able neither for the production of directly applicunsuit-able hot rolled deep-drawing steel grades nor for hot rolled low alloy TRIP steels, as these steels require additional processing steps to obtain the desired mechanical properties. However, in the case of deep-drawing steels, a so called ferritic rollingseems to be a promising hot rolling strategy to produce thin hot strips with desirable deep-drawing properties [8–11]. For the production of hot rolled TRIP steels, an intercritical rolling might be a promising rolling strategy to improve the mechanical properties by a justifiable simple production process [12–15].

1.1 Ferritic Rolling of Deep-Drawing Steels

Due to the low carbon and reduced manganese content of deep-drawing steels, which are the necessary preconditions for obtaining a good deep-drawability, the α/γ transformation is shifted

(15)

1.2 INTERCRITICAL ROLLING OF LOW ALLOY TRIP STEELS 3

to higher temperatures. Therefore, high reheating and finish rolling temperatures are required for the conventional (austenitic) rolling. A reduction of the strip thickness below 1.5 mm involves some difficulties in terms of the temperature control of the strip in the finishing mill. The thin strip cools rapidly in the last stands of the finishing mill and the rolling temperatures can easily drop down into the temperature region of the α/γ transformation. Uncontrolled rolling in the intercritical temperature region of deep-drawing steels is known to deteriorate the mechanical properties [10, 11, 16, 17].

A solution for this problem is the ferritic rolling strategy. In ferritic rolling the finish rolling temperature is shifted intentionally down into the fully ferritic region. This opens the process window to produce thin (ts≤ 2 mm) and even ultra thin (ts≤ 1 mm) hot strip. In addition to the reduced strip thickness it becomes possible to produce hot strips with adequate deep-drawing properties [18].

The ferritic rolling strategy allows for the production of two different hot strip grades, a "soft" hot strip and a "hard" hot strip. The "soft" hot strip is rolled at higher temperatures within the ferritic region using a sufficiently high coiling temperature that ensures a complete recrystallisation of strip directly in the coil. The "hard" hot strip is rolled and coiled at lower temperatures in the ferritic region, so that a complete recrystallisation in the coil can not occur. Hence, these strips exhibit a strained microstructure after coiling and a subsequent recrystallisation annealing is required, to obtain the desired deep-drawing properties.

Due to the absence of the phase transformation after the ferritic rolling and an intentional reduced dynamic and static softening during rolling, it is possible to develop an advantageous "as-if cold rolled" texture consisting of a pronounced γ-fibre and a partial α-fibre. Presupposing an adapted post rolling processing it becomes possible to produce a thin hot strip with a final desirable {1 1 1} texture [19]. Moreover, the two ferritic hot strip grades can be used produce "cold" strip. In this case, the ferritically rolled strips are used as starting material for a subsequent cold rolling treatment and the initial microstructure and texture may partially be bequeathed into the cold strip.

1.2 Intercritical Rolling of Low Alloy TRIP Steels

The conventional production of a low alloy TRIP steel employs the finish rolling in the austenitic temperature region. However, in order to obtain the desired multi phase microstructure, mainly consisting of ferrite, bainite, retained austenite and martensite, a special attention has to be drawn on the cooling cycle on the run-out table and the coiling temperature. Conventionally a two step cooling cycle is utilised (water / air / water). After finish rolling the strip is immediately cooled to prevent the recrystallisation process. In the range of the maximum ferrite transformation the coo-ling rate is reduced to form the required fraction of about 60 − 70 % ferrite. A sufficient amount of ferrite is required for the desired carbon enrichment of the remaining austenite. After this short break in the fast cooling, the cooling rate is raised again to aproach the coiling temperature.

(16)

The coiling temperature has to be in the range of 300 − 500◦C for the bainite transformation, required to finally stabilise the retained austenite below the room temperature. Such a cooling cycle requires a highly adjustable cooling section [6].

The processing route employing a so called intercritical rolling, that means the finish rolling in the intercritical α/γ temperature range, might be a promising rolling strategy to avoid the difficult adjustable the cooling cycle on the run-out table during the conventional production. The hypothesis is to promote the ferrite formation by an intensive deformation in the upper intercritical temperature range or closely above this range, in order to obtain a similar or possibly better α/γ phase composition as after the two-step (water/air/water) cooling, but in a much shorter time span. Hence, the cooling cycle could be reduced to a simple fast cooling step down to the desired coiling temperature. The results obtained from the plane strain compression experiments suggest that the proposed rolling schedule might bring about an improved process control, due to the possible simplification of the cooling cycle.

1.3 Outline

The current thesis is structured according to the two hot rolling strategies:

a) Ferritic rolling of deep drawing steels

Chapter 2 gives an introduction to the key properties of mild steels for cold forming opera-tions. Special attention is drawn to the deep drawing properties. After a short introduction of required mechanical and deep drawing properties of mild steels, the conventional pro-cessing routes of deep drawing steels are discussed, focussing on the restrictions of such routes. The ferritic rolling strategy is proposed as a possible solution without the most of these restrictions. Finally, the recent literature in the field of ferritic rolling is reviewed. As a possible ferritic rolled products are "soft" hot strip, "hard" hot strip and "cold" strip presented and discussed.

Chapter 3 concentrates on the experimental procedures. Special attention is drawn to the determination of the transformation behaviour of the steels, the lubrication in the roll gap and the simulation of the annealing cycles. Finally the mechanical and technological tes-ting is described.

The results for the hot strip grades ("soft" and "hard") are presented and discussed in chapter 4. The most crucial parameters for the production of a "soft" hot strip proved to be a sufficient high coiling temperature and the ferritic rolling strain. Presupposing optimum processing parameters mechanical properties comparable to that of the steel grade DC 04 have been reached in the "soft" hot strip. The "hard" hot strip revealed slightly better properties comparable to that of the conventional cold strip grade DC 05. On the basis of the results recommendations for the industrial production of ferritic rolled strips are formulated.

(17)

1.3 OUTLINE 5

In chapter 5 the results for the "cold" strip grades produced of an initially either "soft" or "hard" hot strip are presented and discussed. The different previous microstructure and texture of such two hot strip grades are transferred and partially bequeathed to the "cold" rolled strip, leading to a different texture development during cold rolling and annealing. This results also in a different normal and planar anisotropy of the final strips.

b) Intercritical rolling of TRIP steels

Chapter 6 describes the work on the intercritical rolling of TRIP steels. In section 6.1, the conventional production routes of low alloy TRIP steels are introduced. With their close control of the cooling cycle on the run-out table necessary. The intercritical rolling strategy is proposed as an alternative processing route for hot rolled TRIP steels, with a less complicated cooling cycle. The experimental procedure is described in section 6.2. Finally in section 6.3, the laboratory results are presented and discussed.

(18)

(19)

Chapter 2 Introduction

Ferritic rolling of deep-drawing steels is a relative new rolling strategy, which is on its way to industrial application. Ferritic rolling introduces the possibility of producing thinner hot strips in the range of ≤ 2 mm with desirable deep-drawing properties. Therefore, the optimisation of the process parameters is of great interest to steel producers.

The first section 2.1 gives an overview of the technological requirements of deep-drawing quality (DDQ) steels. Section 2.2 introduces the macro texture measurement by x-ray diffraction and the texture representation conventionally used for bcc steels together with the texture development during the process. One of the most important features of deep-drawing steels is a large intensity of the so-called {111} texture, which is essential for a good deep-drawing quality.

Section 2.4 introduces the ferritic rolling process with the three possible product groups, "soft" and "hard" hot strip and "cold strip" produced of ferritic rolled hot strip and reviews the most relevant literature. The development of the desirable {111} texture is strongly influenced by the amount of solute carbon in the matrix and the amount of redundant shear deformation due to the friction in the roll gap. Sections 2.5 and 2.6 review the literature on the above subjects.

2.1 Properties of Deep-Drawing Steels

Mild steels for cold forming operations are characterised by a low 0.2 % proof strength, Rp0.2, a high uniform elongation, Ag, and a sufficiently high tensile strength, Rm, to provide accepta-ble low forces during forming together with a sufficient strength of the produced component. Furthermore, a high strain hardening coefficient is necessary to ensure good stretch-forming per-formance of the strip [20]. The mechanical properties of commonly used steels grades for cold forming operations are defined in DIN EN 10130 [21]. Deep-drawing steels additionally require a high deep-drawing ration, β, and a low tendency to form ears, in order to reduce the cut-off scrap. These properties require a high normal anisotropy, characterised by a preferred material

(20)

flow within the strip plane. The Lankford value, r, is a measure for the normal anisotropy and can be determined in tensile test. The r-value is defined according to the equation 2.1

r= εwidth εthickness = ln( w0 w) ln(t0 t) . (2.1)

The Lankford value denotes the ratio of the strain in width direction (w0 initial width; w width after straining) and thickness direction (t0 initial thickness; t thickness after straining). Hence, r< 1 characterises a preferred material flow within the thickness direction, r = 1 indicates a perfect isotropic material flow and r > 1 a preferred material flow within the strip plane. It is obvious that deep-drawing steels require a high r-value (r > 1) to reduce the local reduction of the strip thickness during the forming operation, which leads to a geometrical softening. This geometrical softening becomes critical for the deep-drawing process, as soon as the reduced cross-sectional area can not resist the load demanded and hence fails.

The r-value is directly correlated to the deep-drawing ration, β [8, 20, 22]. The r-value is conven-tionally not constant over the various directions within the strip plane, leading to the so-called planar anisotropy, ∆r. This planar anisotropy is responsible for the unfavourable formation of ears during deep-drawing. The planar anisotropy is defined by

∆r = r0

◦− r₄₅◦+ r₉₀◦

2 . (2.2)

In equation 2.2 the indices designate the angle of the direction, with respect to the rolling direc-tion, under which the r-value is measured. A planar anisotropy of r > 0 leads to the formation of ears, in rolling direction and transverse to the rolling direction, whereas r < 0 leads to ears in 45◦to the rolling direction. A low planar anisotropy, ∆r ≈ 0, together with a high mean r-value, r_m, defined in equation 2.3, yields an optimal deep-drawing performance.

r_m= r0◦+ 2r45◦+ r90◦

4 (2.3)

The ∆r-value is only relevant for r-value distributions consisting of a local minimum (v-shape) or maximum (inverse v-shape) at about 45◦ to the rolling direction. In some cases r-value dis-tributions are observed which do not exhibit the classical v-shape. In these cases the ∆r-value yields ambiguous values. Therefore, in [8] a ∆rmax-value is introduced:

∆rmax= max[r0◦, r₄₅◦, r₉₀◦] − min[r₀◦, r₄₅◦, r₉₀◦]. (2.4)

To meet the above mentioned requirements an adapted chemical composition, on the one hand, and a pronounced {1 1 1} texture, on the other hand, are crucial [8]. The chemical composition of such steels is characterised by a low carbon and overall low alloying content, as shown in

(21)

2.1 PROPERTIES OF DEEP-DRAWING STEELS 9

C P S Mn Ti

DC 06 0.02 0.020 0.020 0.25 0.3

DC 05 0.06 0.025 0.025 0.35

-DC 04 0.08 0.030 0.030 0.40

-Table 2.1: Max. chemical compositions in mass% of deep-drawing steel grades according to DIN EN 10130 [21]

Figure 2.1: Relationship of the mean r-value and the {111} texture represented by the ratio of the intensities I(111)/I(100) [25]

Table 2.1. The lean composition ensures a ductile ferrite matrix and promotes the formation of the desirable {1 1 1} texture.

Steel Grades DC 04 and DC 05 are aluminium-killed low and extra low carbon steels (typically 0.05 mass% Al), whereas DC 06 is microalloyed with titanium (Ti) and/or niobium (Nb). Due to over-stoichiometric microalloying with Ti and/or Nb, the carbon and nitrogen precipitate as car-bides and nitrides [23]. This provides a highly ductile interstitial free ferrite matrix. Therefore, these steels are also called IF steels. A minimum ratio of

Nb C+ N or

Ti

C+ N > 1 in [atom%] (2.5)

is the required to produce an interstitial free steel [23]. Usually a higher ratio of about Ti/(C + N) ≈ 4 is necessary to guarantee the stoichiometric fixation of the carbon and nitrogen [20]. To provide the required deep-drawing properties, a pronounced {111} texture is required. This texture is characterised by a predominant amount of grains oriented with their {111} plane pa-rallel to the strip surface [24]. The r-value of the strip is directly correlated to the intensity of the {111} texture, as shown in Fig. 2.1. The higher the volume fraction of {111} oriented grains, the higher is the r-value [8, 25].

(22)

2.2 Crystallographic Textures

2.2.1 Introduction

The mechanical properties of a single crystal are anisotropic. Assuming a random distribution of the crystallites (grains) in a polycrystalline material its properties would be perfectly isotropic. Engineering steel are however mostly polycrystalline and the orientations of their grains is rarely randomly distributed. As soon as the orientation distribution exhibits one or more predominant orientations a material is textured. The mechanical properties of the textured material are hence anisotropic. The intensity of this anisotropy depends on the specific property, the crystal structure and the texture nature and intensity.

The development of a texture can be driven by a variety of processes, e.g. (directed) solidifi-cation, phase transformation, deformation and/or recrystallisation. The plastic deformation of a polycrystalline material leads to the formation of predominant crystal orientations. During tensi-le straining of a singtensi-le crystal the slip direction rotates until it reaches the direction of tension. In compression, the slip direction rotates towards the plane of compression [26]. In the deformation of a polycrystalline material the individual grains have to accommodate complex stress states to maintain connectivity across grain boundaries. This deformation leads to the development of a texture. The nature and intensity of this texture is determined by the mode of deformation and the crystal symmetry [26]. During recrystallisation of such deformed structures new grains nucleate and grow into the deformed matrix. The newly formed grains often possess a specific orientation relationship with their parent structure.

2.2.2 Orientation Description

To describe the orientation of crystal within a specimen two Cartesian coordinate systems, the specimen coordinate system and the crystal coordinate system are necessary as a reference frame. The specimen coordinate system is preferentially chosen in accordance with the process geome-try, C = {c1, c2, c3} = {RD, T D, ND} and the crystal coordinate system parallel to the crystal unit cell, S = {s1, s2, s3} = {[100], [010], [001]}. RD denotes the rolling direction, T D the trans-verse and ND the normal direction of the strip. On this basis the orientation of a crystal (grain) can be described by a rotation matrix, g, representing the rotations necessary to transform the specimen coordinate system into the crystal coordinate system. A more crystallographic and a more demonstrative approach is the notation with Miller indices (h k l)[u v w], or {h k l}hu v wi for non specific indices [27]. {h k l} denotes the crystallographic plane of the grain, which is parallel to the specimen surface, i.e. the normal of this plane is parallel to ND. hu v wi represents the crystallographic direction, which is parallel to RD. A more detailed description of the coordinate transformation and the notation with Miller indices can be found for example in [27–30].

(23)

2.2 CRYSTALLOGRAPHIC TEXTURES 11

a)

b)

c)

Figure 2.2: a) & b) Rotation of a sample, necessary to satisfy Bragg’s law; Diffraction in a texture goniometer with Euler cradle (reflection geometry) [27]

2.2.3 Macro Texture Measurement and Representation

A standard macro texture measurement technique is the x-ray diffraction method. This measure-ment technique is based on Bragg’s law, equation 2.6. Each set of lattice planes with the lattice plane spacing, dhkl, fulfils the Bragg equation for a given wave length, λ, as the incident mono-chromatic x-ray beam and the detector are set to the corresponding angle, 2Θ, and the normal to the lattice plane is the bisector of the angle 2Θ. The lattice plane spacing, dhkl (Eq. 2.7), is depending on the observed crystallographic plane, (hkl), and the lattice parameter, a.

λ = 2dhklsin Θ (2.6)

dhkl =

a √

h2_{+ k}2_{+ l}2 (2.7)

To measure an unknown crystal orientation the sample is tilted and rotated until Bragg’s law is satisfied, Fig. 2.2 a) & b). In the case of a polycrystalline material the intensity measured at a given angle 2Θ is a direct measure for the volume fraction of grains with the corresponding orientation. Once this measurement has been executed for all possible tilt and rotation angles, the results can be readily presented in the form of a pole figure. This pole figure directly reflects the texture of the material [27]. The tilt and rotation angles, α and β, are directly correlated to the pole figure angles (α radial, β azimuthal).

These texture measurements are usually conducted using a texture goniometer. There are two different measurement geometries possible, transmission and refection geometry. Fig. 2.2 c)

(24)

F j2 j1 0° 90° 90° 90° 45° a-fibre <110>||RD g-fibre <111>||ND F j1 a) b) 0 15 30 45 60 75 90 0 15 30 45 60 75 90 {001}<110> {112}<110> {111}<110> {111}<112> {111}<110> {111}<112> {554}<225> {332}<112> {011}<100> {110}<110> j2= 45°

Figure 2.3: a) 3-dimensional Euler-space with the position of the most important texture fibres, the α-fibre and γ-fibre b) The ϕ2= 45◦cut wshowing some additional texture components important for deep-drawing steels

schematically shows a goniometer in reflection mode, together with the definition of the axes. The rotation axes of the goniometer in the reflection mode have the following correlation to the pole figure angles: β = Φ and α = 90◦− χ. The angle ω is usually kept constant [27]. Because of the limitation of the tilting angle αmax ≈ 60 − 85◦ in the reflection mode, only incomplete pole figures can be measured. The transmission geometry is seldom used because of an elaborate sample preparation.

A pole figure gives only a relative information about the orientation distribution. For the absolute description of the orientations three independent angles are required. The pole figure consists, however, only two rotation angles. A complete description of the orientation distribution can be obtained using the orientation distribution function (ODF), f (g), plotted in a 3-dimensional coordinate system. The ODF is calculated from a set of independent pole figures by a set of series expansion methods [31]. The ODF is a 4-dimensional function, consisting the three rotation angles and the intensity of the corresponding crystallographic orientation.

A frequently used coordinate system to represent the ODF is the so-called Euler space, shown in Fig. 2.3 a) with the positions of the most important bcc texture fibres α and γ. Each point within the Euler space represents the intensity, f (g), of the orientation specified by the three Euler angles. f (g) = 6 means that the measured intensity is six times higher than the intensity of a random orientation distribution. The three Euler angles, ϕ1, Φ and ϕ2(Bunge’s Notation) denote the rotation angles of the specimen coordinate system with respect to the specimen coordinate system [29].

To present the 3-dimensional Euler-space in 2-dimensions, usually equidistant cuts along the ϕ1 axis are depicted with the step size being ∆ϕ1= 5◦. However, this method is hardly suited to compare easily different measurement results. Therefore, the representation is often reduced to so called fibre plots [8, 27, 29, 32]. In a fibre plot the orientation intensity, f (g), is plotted along certain characteristic paths through the orientation space versus the angle which defines this path.

(25)

2.2 CRYSTALLOGRAPHIC TEXTURES 13

The most important fibres for bcc materials are the α-fibre and γ-fibre indicated by the arrows in Fig. 2.3. The rolling and recrystallisation texture are mostly characterised by these two fibres [27]. The important texture fibres for deep-drawing steels are [8, 32, 33]:

1. α-fibre – h1 1 0i||RD: {1 0 0}h1 1 0i to {1 1 0}h1 1 0i

The α-fibre includes all orientations with their h1 1 0i axis parallel to the rolling directi-on (RD). The α-fibre has the coordinates in the Euler space at ϕ1 = 0◦, ϕ2= 45◦ and Φ = 0◦. . . 90◦. Important orientations along the α-fibre are the rotated cube compo-nent {0 0 1}h1 1 0i (Φ = 0◦), the inverse brass component {1 1 2}h1 1 0i (Φ = 35◦) and {1 1 1}h1 1 0i (Φ = 54.7◦). At Φ = 54.7◦the α-fibre intersects the γ-fibre.

2. γ-fibre – h1 1 1i||ND: {1 1 1}h1 1 0i to {1 1 1}h1 1 2i

The γ-fibre contains all orientations with their {1 1 1} plane parallel to the sheet plane. The γ-fibre is also called {1 1 1}-fibre. The γ-fibre contains the important texture com-ponents from cold rolling and recrystallisation. The coordinates in the Euler space are ϕ1= 0◦. . . 90◦, ϕ2= 45◦ and Φ = 54.7◦. Due to the three fold symmetry the intensi-ties are along ϕ1= 0◦. . . 30◦, ϕ1= 60◦. . . 90◦ and ϕ1 = 30◦. . . 60◦ mirrored at Φ = 30◦ identical. Therefore, it is sufficient to consider the orientation between 60◦ and 90◦[33]. The orientations of the γ-fibre are decisive for the deep-drawability of the strip. It is of-ten observed that the maxima in a recrystallisation do not lie perfectly along the γ-fibre at {3 3 2}h1 1 3i or {5 5 4}h2 2 5i. Therefore the fibre plot would only show lower intensities. 3. ε-fibre

The ε-fibre contains important texture components resulting from an shearing due to inho-mogeneous deformation. This is especially the case in unlubricated hot rolling [33–35]. The coordinates in the Euler space are ϕ1= 90◦, ϕ2= 45◦ and Φ = 0◦. . . 90◦. Important orientations an this fibre are the rotated cube orientation {0 0 1}h1 1 0i (Φ = 0◦), the copper orientation {1 1 2}h1 1 1i (Φ = 35◦), the intersection point with the γ-fibre {1 1 1}h1 1 2i (Φ = 54.7◦) and the Goss orientation {0 1 1}h1 0 0i (Φ = 90◦). The Goss component {0 1 1}h1 0 0i a the coordinates ϕ1 = 90◦, ϕ2 = 45◦ and Φ = 90◦ is detrimental for the deep-drawing properties and and has to be avoided.

The other fibres, such as the η, and ζ-fibre are less important for deep-drawing steels. A detailed description can be found for example in [33].

The fibre plots offer a very condensed representation of the texture. Different texture results can be easily compared, by plotting them into one graph. However, as already mentioned it is often the case that the texture maxima do not lie perfectly on the fibres (but still close to them), which might lead to a misinterpretation of the results. A good compromise between plotting the whole ODF and only the fibres plots is the ϕ2= 45◦cut, Fig. 2.3 b). The ϕ2= 45◦ cut contains the α-, γ- and ε-fibre as well as some other important orientations. Points of the same texture intensity are mostly connected with lines.

(26)

Time a g+ a g+ Ar3 Ar1 C Mn roughing finishing finishing coiling coiling roughing Te mperature a) b)

Figure 2.4: Rolling schedule for a) conventional and b) ferritic hot rolling

2.3 Conventional Production of Deep-Drawing Steels

In the conventional production of deep-drawing steel strip the slab is reheated and subsequently rough rolled and finish rolled to its final hot strip thickness in the austenite temperature regi-on. The hot strip is finally cooled to the desired coiling temperature and coiled. Due to the reduced carbon and manganese content of deep-drawing steels the Ar3and Ar1temperatures are noticeably raised, as it is indicated in Fig. 2.4 a).

The hot rolling parameters, such as reheating temperature, rolling reduction during roughing and finishing, the cooling rate and the coiling temperature all influence the mechanical properties of the final hot strip [36]. The main aim during conventional hot rolling is to reduce the thickness of the slab to the desired hot strip thickness, which is usually restricted to ≥ 2 mm, in order to avoid uncontrolled cooling of the strip into the two phase region during finish rolling. Moreover, the grain size and the state of precipitation are important for the properties of the final hot strip. By lowering the reheating temperature and rolling temperature, the austenite and ferrite grain size can be reduced. After pickling the hot strip is cold rolled, annealed and temper rolled. The state of precipitation of the nitrogen in the hot strip is of particular interest for extra low carbon (ELC) steels[7]. ELC steels are conventionally stabilised with aluminium. The required recrystallisation annealing after cold rolling can be done either continuously or discontinuously. The main difference between those two annealing processes are the heating and cooling rate on the one hand and the annealing temperature and time on the other hand [37–39]. Depending on the imposed annealing treatment, the nitrogen is supposed to be precipitated as AlN after hot rolling or still be in solid solution [25, 40, 41].

At conventional reheating temperatures for mild Al-killed steels of about 1250◦C the aluminium and nitrogen are nearly completely in solid solution. The precipitation of AlN can be controlled by the coiling temperature. For the case of batch annealing the AlN is supposed to precipitate

(27)

2.3 CONVENTIONAL PRODUCTION OF DEEP-DRAWING STEELS 15

Figure 2.5: Influence of the cold reduction on rmfor different steel compositions [25]

during the slow reheating in the batch furnace. This requires low coiling temperatures after hot rolling, to keep the Al in solid solution. During reheating of the cold strip coil in the batch furnace the N preferentially precipitates along the elongated grain boundaries. This leads to the forma-tion of the typical "pancake" microstructure [42, 43]. For the case of a continuous annealing cycle the AlN is required to be precipitated prior to cold rolling, which requires higher coiling temperatures [8, 44]. The rapid heating of the strip in the continuous annealing line retards the AlN precipitation. Thus, in the temperature region of recovery and starting recrystallisation the nitrogen is predominantly segregated at dislocations. This hinders the dislocation movement for recovery, which is necessary to initiate the recrystallisation process. For the same reason the recrystallisation is shifted to higher temperatures [43].

The standard mechanical properties of the conventionally produced hot strip are with Rp0.2 ≈ 150 MPa, Rm ≈ 280 MPa and A50 ≈ 40 %, see [8], already comparable to those required in DIN EN 10130 [21] for steel grade DC 06. In contrast to this the deep-drawing properties of such hot strip are only rather poor. Due to the high rolling temperatures and the subsequent α/γ transformation, the hot strips exhibit a nearly random texture [16]. This leads to low mean Lankford values of about rm≈ 0.85 [8]. Therefore, for deep-drawing applications hot strip requi-res supplemental cold rolling and annealing. During cold rolling a sharp rolling texture develops. By subsequent annealing of the strip, the rolling texture is transformed, due to recrystallisation, to an annealing texture. This annealing texture has to be characterised by a high density of grains with h1 1 1i||ND. The intensity of the {1 1 1} texture and hence the normal and planar anisotropy are strongly dependent on the cold reduction.

The optimum cold reductions referred in the literature vary from about 75 % up to 90 % [8, 32] and depend on the chemical composition of the steel [25]. Fig. 2.5 shows the development of the mean r-value with increasing cold reduction for different alloying.

(28)

2.4 Ferritic Rolling of Deep-Drawing Steels

The production of thin gauge hot strip (tf < 2 mm) by conventional (austenitic) hot rolling in-volves serious difficulties in terms of the high Ar3 temperature of deep-drawing steel grades. Because of the low strip thickness and large surface area, the strip cools more rapidly in the finis-hing mill. Hence, the rolling temperatures can easily drop down into or below the temperature region of the α/γ transformation. Uncontrolled rolling in the intercritical temperature region is known to cause problems for the control of the rolling process and to deteriorate the mechanical properties [10, 11, 16, 17].

A possible solution for this problem offers the ferritic rolling strategy. In ferritic rolling the finish rolling temperature is shifted down into the fully ferritic region. Fig. 2.4 b) shows schematically the temperature versus time profile of the ferritic rolling strategy. A variety of possible impro-vements with respect to production costs and process efficiency can be achieved utilising this approach [8–11, 45]:

• cost reduction and increased throughput, due to reduced reheating temperatures • less scale formation in the reheating furnace, due to the lower reheating temperatures • reduction of work roll wear, due to lower rolling temperatures in the finishing mill • better strip flatness control by rolling of a transformed and homogeneous microstructure • reduced strip surface defects

• reduced coolant consumption on the run out table, due to lower finish rolling temperatures • reduced rolling forces, due to rolling of a softer ferrite

• the total reduction in production costs is estimated to be approximately 25 %

Possible disadvantages of the ferritic rolling strategy might be [9]:

• excessive mill-power requirements and

• excessive rolling loads , both due to undesirable low finish rolling temperatures

However, due to the local minimum of the flow stress below the temperature region of the α/γ transformation, the ferritic rolling strategy can be utilised on conventional hot rolling mills within a certain temperature range [46]. The flow stress and thus the rolling force is from about 840◦C down to 750◦C (ferrite) lower than that at above 950◦C (austenite) for the case of an in-dustrially produced Ti stabilised IF steel, Fig. 2.6. Nevertheless, a further reduction of the rolling temperatures would lead to a drastic increase in flow stress and hence the rolling force.

(29)

2.4 FERRITIC ROLLING OF DEEP-DRAWING STEELS 17 600 650 700 750 800 850 900 950 1000 120 140 160 180 200 220 240 260 Temperature [°C] flow stress [MPa] a g+ e = 0.4 e = 0.6 e = 0.4

Figure 2.6: Flow stress versus temperature plot for an industriall produced IF steel, determined by plane strain compression

Austenitically rolled hot strips usually exhibit a random texture, due to the randomising effect of the post rolling phase transformation. Because of the absence of the phase transformation after ferritic rolling and a reduced impact of dynamic and static softening during the rolling process, it is in principle possible to develop an advantageous rolling texture consisting of a pronounced γ-fibre (h1 1 1i||ND) and a partial α-fibre (h1 1 0i||RD). Hence, using an adapted post rolling processing it is possible to produce a final thin hot strip with a desirable {1 1 1} texture [19]. By utilising the ferritic rolling strategy three different product groups can be produced, two hot strip grades and one cold strip grade, Fig. 2.7. These variants will be discussed in more detail in sections 2.4.1 and 2.4.2 respectively.

2.4.1 Hot Strip Grades

Two different hot strip grades can be produced by the ferritic rolling:

1. "soft" hot strip, Fig. 2.7 a)

2. "hard" hot strip (annealed), Fig. 2.7 b)

The "soft" hot strip is finish rolled at higher temperatures in the ferritic temperature region. A corresponding higher coiling temperature ensures a full recrystallisation of strip in the coil. The "soft" hot strip is of such a quality, that it can be sold and used as a finished product. Except for a removement of the oxide layer by pickling it needs no further processing. This practice allows

(30)

a) "Soft” hot strip

a a + g finishing

coiling

b) “Hard” hot strip (annealed)

a a + g annealing

c) “Cold” strip (annealed)

a a + g annealing Time T e mperature Time Time

Figure 2.7: Variants of ferritic rolled products

for the production of deep-drawable thin hot strips in a thickness range of tf ≈ 1 − 1.5 mm. These strips can be used as a substitute for cold strip for some applications [11, 16].

The "hard" hot strip is rolled at relatively lower temperatures in the ferritic temperature region and therefore exhibits a strained microstructure after coiling, indicating a non recrystallised state. Hence, a subsequent recrystallisation annealing is required to obtain the required formability. The lower finishing temperatures allow for the production of even thinner hot strips with tf ≤ 1 mm [9]. The accumulated strain can be expected to be higher, because of the lower finishing temperatures, suggesting that a stronger rolling texture develops, compared to that of a "soft" hot strip. The post rolling annealing may proceed either discontinuously or continuously.

The results found in literature for the properties of ferritic rolled hot strips are diverse and not consistent even for similar alloy compositions [8]. This is caused by the fact that the production parameters, such as the rolling and coiling temperatures, strongly influence the results obtained. In addition, a major difference is caused by the chemical composition. Therefore, it is reasonable to group the results by the most common deep-drawing steel grades, ELC/ULC (ultra low carbon) and IF (interstitial free) steel. These two steel grades yield very different mechanical properties, especially in terms of the deep-drawability after ferritic rolling.

2.4.1.1 ULC/ELC Steel

ULC/ELC steels belong to the steel grade DC 04 or DC 05 and are both stabilised with Al. The typical chemical compositions are given in Table 2.1. Independently of the steel grade it can be stated that the texture development and so the r-value are strongly dependent on the ferrite rolling temperature. In [8, 47, 48] plane strain compression tests were performed in a temperature range from 820◦C down to 620◦C using the hot deformation simulator "WUMSI" [49]. The steel used in this investigation was an industrially produced Al killed ELC steel.

(31)

2.4 FERRITIC ROLLING OF DEEP-DRAWING STEELS 19

Rolling Texture

To measure the rolling texture, the samples were water quenched after ferritic rolling to retain the as rolled condition. For high rolling temperatures of TFR≈ 820◦C the texture consisted of a weak γ-fibre and a strong intensity on the α-fibre in the region of {100}h110i. At about 670◦C, the rolling texture changed and a texture comparable to that of a cold rolled specimen developed [8, 50]. Rolling at 620◦C produced a strong rolling texture consisting of a stronger γ-fibre and a partial α-fibre with a maximum near {112}h110i.

Another important process parameter was found to be the total logarithmic strain, εFR, applied in the ferritic temperature region. The results in [8] show an increasing intensity of the rolling texture when εFRis raised from 0.3 to 1.2.

"Soft" Hot Strip

The coiling temperature necessary to produce a recrystallised hot strip was approximately 660 − 680◦C for a reheating temperature of 1200◦C and 610 − 620◦C for a reduce reheating tem-perature of 1000◦C [8]. The largest possible temperature difference of approximately 20 − 30◦C achieved between the necessary rolling temperature to produce a desirable rolling texture and the minimum coiling temperature that would guarantee a full recrystallisation, was too low to be realised in a conventional hot rolling mill.

The coiling textures presented in [8] consisted of a rather strong α-fibre texture with a peak in-tensity at {100}h110i. This texture yields only a poor deep-drawabilty of the strip with rm< 1. These results are consistent with those found in [26] with a similar ELC steel. The tests in [26] were however executed on a laboratory rolling mill. For all possible rolling textures the recrystal-lisation in the coil led only to an undesirable annealing texture. In [51, 52], the laboratory-rolling tests with low carbon steel without the addition Al were performed. The specimens were rolled at temperatures between 70 − 700◦C. Raising the rolling temperature up to 700◦C also led to the formation of an undesirable recrystallisation texture with a maximum along the α-fibre at {100}h110i.

The mechanical properties reported in [8, 10] reveal a lower yield strength (180 MPa versus 230 MPa) and tensile strength (370 MPa versus 300 MPa) in combination with a higher total elongation compared to an austenitically rolled hot strip. Similar results have been found in [9, 11]. It is reported in [9, 10, 42], that the ferritic strips have a very low ageing sensitivity because of

• an incomplete dissolution of the AlN during reheating,

• a faster deformation induced AlN precipitation during ferritic rolling and • a faster AlN precipitation during "high" temperature coiling.

In contrast to the somewhat superior mechanical properties, only poor r-values of about rm≈ 1 and a fairly high negative planar anisotropy were measured for the ferritic rolled strips [8, 9, 11].

(32)

This effect is explained in terms of the undesirable recrystallisation texture consisting a strong α-fibre. This undesirable texture development is explained due to the presence of an unacceptable high amount of solute carbon in the ELC/ULC steel during ferritic rolling [51–53].

The microstructure consisted in all cases polygonal ferrite grains with a somewhat bigger grain size when compared to that of a conventionally rolled hot strip. In [8, 26] the occurrence of a slightly elongated microstructure is described. The mean grain size was about 30 − 40 µm in transverse and about 50 − 80 µm parallel to the rolling direction [8].

"Hard" Hot Strip

For the case of lower finish rolling and, hence, coiling temperatures, the strip does not recrystal-lise in the coil and therefore exhibits a strained microstructure. Due to to the possibly higher accumulated total strain, these strips are expected to have a stronger rolling texture than their "soft" counterparts. Therefore, these strips are expected to yield better deep-drawing properties after recrystallisation. Additionally these strips can be even thinner, because of the wider process window.

In [8] two different coiling temperature were used, 550◦C and 400◦C respectively. All sam-ples exhibited a strained microstructure. To provide the required mechanical and deep-drawing properties the strips had to be annealed. As already stated before, the rolling textures for finish rolling temperatures below 620◦C consisted of an intense γ-fibre and a maximum at {112}h110i along the α-fibre. All specimens consisted this desirable rolling texture before the subsequent recrystallisation annealing.

Nevertheless, the intense rolling texture could not be transformed into a desirable {111} tex-ture. The grains with γ-fibre orientation were completely consumed in favour of the α-fibre orientations. The annealing textures exhibit a maximum at {1 0 0}h1 1 0i on the α-fibre. From these experiments it can be concluded, that the occurrence of a desirable rolling texture is not a guarantee for a desirable {1 1 1} annealing texture, just as the "pancake" microstructure [8, 26]. The resulting deep-drawing properties were, as one could expect from the undesirable texture, also very poor. In [8, 9] the calculated mean r-value was approximately 0.82 − 1.0 for both the batch and continuous annealing treatment. The mechanical properties, measured in the tensile tests were comparable to those found for the "soft" hot strips.

Also for the "hard" hot strip the detrimental effect of solute carbon in the ELC and ULC steel during the recrystallisation was observed. It was found that in contrast to the rolling texture the annealing texture is strongly affected by the amount of solute carbon [8, 51, 52].

2.4.1.2 IF Steels

Rolling Texture

IF steels, grade DC 06, are stabilised with Titanium (Ti) and/or Niobium (Nb). The typical chemical composition is given in Table 2.1. The development of the rolling texture for different

(33)

rolling temperatures is in principle similar to that ELC steels. The work of [8] shows that a typical rolling texture, consisting a strong γ-fibre and a maximum along the α-fibre at {112}h110i, starts to develop at about 760◦C. For higher rolling temperatures a weaker texture with a poor γ-fibre and a maximum along the α- fibre near {100}h110i develops. A similar behaviour has been reported also for IF steels in [19, 26, 51, 52, 54, 55]. Rolling at 710◦C resulted in a very strong rolling texture [8, 47, 48].

"Soft" Hot Strip

The necessary coiling temperature to ensure complete recrystallisation in the coil, was ≥ 670◦C for both the conventional (1200◦C) and the reduced reheating temperature (1050◦C) [8]. Assu-ming a short run-out table of about 50 m and a rolling speed of 12m_/_s _{and a strip thickness of} 2 mm the minimum rolling temperature can be about 710◦C [8]. For rolling temperatures below 760◦C and coiling at 700◦C the recrystallisation textures measured in [8] consisted of a strong γ-fibre with a maximum near {3 3 2}h1 1 3i. The combination of the lowest rolling temperature, TFR= 710◦C, and the lowest allowable coiling temperature, TC = 670◦C, gave the strongest {111} texture. The coiling textures for specimens rolled at temperatures above 760◦C consisted of a weak and undesirable α-fibre texture. Comparable results were found in [26, 54].

The mechanical properties reported in [8] meet the requirements of the DIN EN 10130 [21]. A 0.2 % proof strength of approximately 90 MPa, a ultimate tensile strength of 275 MPa and a total elongation, A50, of approximately 50 % were measured. These properties are even superior to those required for steel grade DC 06 [21]. The calculated r-values were within the range of steel grade DC 04 without the necessity of cold rolling and subsequent annealing. For a rolling temperature of 760◦C the calculated r-value was r90◦ ≈ 1.5 and for 710◦C it was r90◦ ≈ 1.8 [8]. The calculated r-value distribution found in the experiments were however different from the classical v-shape. The distribution showed only one maximum at 90◦ to the rolling direction. This might result in a different earing behaviour.

The results in [19] show, that a sufficient high amount of Ti is necessary, to tie up all the carbon during ferritic rolling. Otherwise the solute carbon will deteriorate the r-value. The presence of solute carbon can be proved by the occurrence of a yield point elongation during straining in a tensile test.

"Hard" Hot Strip

For the case of lower finish rolling temperatures it becomes impossible to produce a recrystallised strip. The low coiling temperatures provide a "hard" hot strip with a strained microstructure. These "hard" hot strips require additional batch or continuous annealing. In [8] two different coiling temperatures were utilised, TC = 550◦C and 400◦C. These low coiling temperatures even allowed for finish rolling temperatures of TFR= 660◦C. In [9, 19] similar rolling and coiling temperatures were used.

The rolling texture of the "hard" hot strips is comparable to that of the "soft" hot strip. Due to the lack of recrystallisation in coil, the texture is not changed during coiling. The coiling tex-ture of the strip consisted for both coiling temperatex-tures a strong rolling textex-ture with an intense

(34)

γ-fibre and and a maximum at {112}h110i along the α-fibre [8]. During the subsequent annea-ling treatment – either continuous or discontinuous – the rolannea-ling texture was transformed into a desirable {111} texture. In [8] lower coiling temperatures proved to yield a somewhat stronger {111} texture. This was explained in terms of a reduced impact of recovery processes in the coil [55].

The mechanical properties are also comparable to those found for the "soft" hot strip (Rm ≈ 270 − 290 MPa; Rp0.2 ≈ 90 − 100 MPa; A50 ≈ 40 − 50 %) [8, 19]. The mechanical properties are relatively independent of the coiling temperature, whereas the mean r-value improved for decreasing coiling temperatures. The mean r-value increases form 1.2 to 1.6 when reducing the coiling temperature from 550◦C to 400◦C [8].

In [8, 16, 56] a new processing route, the direct annealing, has been suggested to benefit from the reduced impact of the recovery processes. In this process route the hot coil is directly trans-ferred to a batch annealing furnace and subsequently annealed. This route avoids the recovery process during the slow cooling of the coil to room temperature and during the slow reheating ( ˙T_Coil ≈ 20 − 60K_/_h_{) of the coil to the coiling temperature. Additionally the energy consumption} is noticeably reduced, while the production throughput is increased. The measured r-values after direct annealing of the "hard" hot strip were with rm≈ 2.3 superior to those of the conventio-nally annealed strip [8]. Also for this processing route lower coiling temperatures yielded better deep-drawing properties. The calculated mean r-values and planar anisotropy were comparable to those of conventionally produced cold strip of steel grade DC 06. A detailed study in [26] provides a good and quantitative insight into the role of recovery in texture formation.

2.4.2 Cold Strip Grades

The third product group for the ferritic rolling strategy is the "cold strip" produced of ferritic rolled hot strip, Fig. 2.7 c). The initial hot strip for the subsequent cold rolling and annealing can be either a "soft" hot strip, Fig. 2.7 a), or a "hard" hot strip Fig. 2.7 b). This provides a different initial microstructure and texture prior to the subsequent cold rolling. The idea is to partially bequeath the hot strip texture to the cold strip, in order to improve its final deep-drawing properties [9, 48]. The results regarding the mechanical and deep-drawing properties of such cold strip grades are not only dependent on the initial microstructure and texture, but also to a great extent on the chemical composition of the steel. Therefore, the results found in the literature are again ordered by the two main steel compositions ELC/ULC and IF.

2.4.2.1 ELC/ULC Steels

"Soft" Hot Strip

For the case of a "soft" hot strip as an initial strip for the subsequent cold rolling, the only possible initial texture was an undesirable recrystallisation texture with a maximum intensity at

(35)

{1 0 0}h1 1 0i along the α-fibre [8]. The microstructure was fully recrystallised with polygonal ferrite grains. With increasing cold rolling reduction the texture components near {1 0 0}h1 1 0i were strengthened, while the γ-fibre stayed nearly unchanged. During the subsequent annealing treatment the α-fibre texture was weakened, but no desirable γ-fibre texture developed [8, 48]. Even a cold reduction of 80 % did not lead to a desirabel {1 1 1} texture. It can be concluded, that a "soft" hot strip produced of ELC steel will lead to undesirable deep-drawing properties, independently of the applied cold reduction. Simmilar findigs are reported in [10]. The mecha-nical properties of cold rolled and annealed "soft" hot strip found in the literature are given in Table 2.2.

Initial Hot Strip εCR Rp0.2 Rm A rm Ref.

ELC-Steel [%] [MPa] [MPa] [%] [-]

HFR 73 193 335 35 1.13 [11]

LFR 40 162 277 44 1.73* [8]

AR 73 210 320 35 1.80 [11]

AR 80 209 282 45 2.34* [8]

Table 2.2: Mechanical properties of a cold rolled and annealed ferritic rolled "soft" and "hard" hot strip compared to a conventional cold strip [8] (HFR: "Soft" Hot Strip; LFR: "Hard" Hot Strip; AR: Conventional Hot Strip) * calculated from the texture

"Hard" Hot Strip

For the case of a "hard" hot strip the initial strip for subsequent cold rolling exhibits a strained microstructure and a desirable rolling texture with a strong and uniform γ-fibre and a partial α-fibre. For a cold rolling reduction of 40 % the rolling texture was intensified. A very high peak developed near {112}h110i and the γ-fibre was slightly increased. For a cold rolling reduction of 80 % the γ-fibre was however decreased and an intense α-fibre developed with high intensities between {100}h110i and {112}h110i [8].

During the recrystallisation of the strip with a cold reduction of 40 % a desirable γ-fibre texure developed, with low intensities near {100}h110i. The annealing texture was comparable to that of a conventionally rolled strip. For the strip with a cold rolling reduction of 80 %, exhibiting an undesirable cold rolling texture, the annealing texture was also undesirable. Controversially in [10] the opposite behaviour is described. There an improvment of the deep drawing properties with increasing cold rolling reductions was found. The r-values were rm ≈ 1.25 − 1.35 for a cold reduction of 70 − 75 % and rm≈ 1.4 − 1.6 for a cold reduction of > 80 %. The mechanical properties reported in [8] are comparable to those of a conventionally rolled strip, grade DC 05.

2.4.2.2 IF Steels

"Soft" Hot Strip

(36)

pa-rameters, a strong γ-fibre texture with a maximum intensity near {3 3 2}h1 1 3i. The initial mi-crostructure is recrystallised with somewhat elongated grains in the rolling direction. During cold rolling to 40 and 80 % the γ-fibre was strengthened [8]. For a cold reduction of 40 % the γ-fibre was strengthened, without the development of the typical rolling components along the α-fibre between {1 0 0}h1 1 0i and {1 1 2}h1 1 0i. For 80 % reduction the γ-fibre was more ho-mogeneous and the typical rolling texture components along the α-fibre formed, with a strong maximum near {1 1 2}h1 1 0i[8, 57]. All rolling textures measured in [8] were superior to that of the conventionally rolled hot strip.

During annealing of the cold rolled strip an unusual strong {1 1 1} texture developed even for a cold reduction of 40 % [48]. In [8] it has been found that for this product group the cold reduction only has a minor effect on the r-value. Cold reductions of 40, 60 and 80 % resulted in similar high r-values of rm≈ 2.71. The mean r-value and planar anisotropy were superior to that of the conventionally rolled strip. The mechanical properties found in literature are summarised in Table 2.3. In a more recent work [57] the development of a strong and very uniform γ-fibre texture was observed. In [57] an IF steel has been warm rolled by εFR= 75 %, followed by intermediate annealing, cold rolling by εFR= 80 % and a final recrystallisation annealing. Although this processing route is different from that used in [8], the microstructure and texture of the resulting hot strip are similar for both processes.

Initial Hot Strip εCR Rp0.2 Rm A rm*

IF [%] [MPa] [MPa] [%] [-]

HFR 40 87 256 38 2.67

LFR1 40 97 275 41 2.31

LFR2 40 97 269 39 2.27

AR 80 94 272 47 2.28

* calculated from the texture

Table 2.3: Mechanical properties of cold rolled and annealed ferritic rolled "soft" and "hard" hot strip compared to conventional cold strip [8] (HFR: "Soft" Hot Strip; LFR1: "Hard" Hot Strip TC= 400; LFR1: "Hard" Hot Strip TC= 580; AR: Conventional Hot Strip)

"Soft" hot strips with an undesirable initial rolling texture, rolled at TFR> 760◦C, yielded after cold rolling and annealing also undesirable recrystallisation textures. Even a cold reduction of 80 % did not improve the texture and deep-drawing properties [48].

"Hard" Hot Strip

The "hard" hot strip exhibits a strained microstructure and a strong rolling texture. During sub-sequent cold rolling of these strips, the γ-fibre was not intensified to a large extent. However, the rolling texture component near {1 1 2}h1 1 0i was conspicuously reinforced [8, 57]. This strong rolling texture reflects the accumulation of the residual warm and cold rolling strain.

During recrystallisation of these strips a very peaked texture with a maximum near {5 5 4}h2 2 5i developed for a cold reduction of 80 % [57]. The mean r-values calculated in [8] together with the

(37)

2.5 INFLUENCE OF SOLUTE CARBON 25

mechanical propreties were already for a cold reduction of 40 % comparable to that of a conven-tionally produced cold strip, Table 2.3. The planar anisotropy was however slightly improved. A further increase in the cold reduction did not improve the r-values. Comparable r-values (DDQ → 1.5 < rm< 2) are also reported in [9].

2.5 Influence of Solute Carbon

The literature review in the previous section has shown that for the "soft" and the "hard" hot strip produced of ELC steels only a poor recrystallisation texture was obtained. The recrystallisation texture was characterised by a strong α-fibre with a maximum at {0 0 1}h1 1 0i independently of the ferritic rolling parameters. In the literature this effect is attributed to an undesirable high content of solute carbon [8, 19, 26, 32, 51, 52, 58–60]. The role of carbon in texture formation is described in this chapter.

2.5.1 Recrystallisation Mechanisms

The primary (discontinuous) recrystallisation is characterised by nucleation and grain growth. During recrystallisation the unstrained nucleus is growing into the deformed matrix, reducing the dislocation density. Both processes, nucleation and growth, determine the recrystallisation texture. These mechanisms, both oriented nucleation and selective growth are proposed as crucial ones [32, 44].

A successful nucleus has to fulfil three different criteria which are also called instability criteria [29]:

Thermodynamic instability: The nucleus requires a critical size, rc, which is caused by the fact that its growth reduces the free enthalpy. The critical radius is defined by

r_c= 2γ

p =

4γ

ρGb2, (2.8)

in which γ denotes the grain boundary energy and p the stored deformation energy. Fur-thermore, is ρ the dislocation density, G the shear modulus and b the Burger’s vector. It is assumed that nuclei with a critical size already exist in the deformed microstructure in the form of cells or subgrains. Recovery processes are however necessary to activate such cells as nuclei.

Mechanical instability: A local non-equilibrium in driving force is necessary to create a grain boundary movement. This requirement is usually fulfilled by an inhomogeneous disloca-tion distribudisloca-tion or by large local sub grains, which are commonly formed during recovery phase.

(38)

Kinetic instability: The boundary of the nucleus must be movable, which is only the case for high angle grain boundaries. The creation of the high angle grain boundary is the hardest part in the formation of a nucleus. There are several mechanisms to form high angle grain boundaries:

• discontinuous subgrain growth

• nucleation on existing grain boundaries • deformation inhomogeneities

• nucleation on larger precipitates • creation of annealing twins.

The difficulty to fulfil all three instability criteria at the same time leads to a strong selection of nucleation sites within the deformed matrix. Especially deformation inhomogeneities and grain boundaries are preferred nucleation sites [32].

Oriented Nucleation

Oriented nucleation means that a direct correlation between the orientation of the nucleus and the nucleation site exists. The following nucleation models are collected from the literature [26, 32, 44]:

Nucleation by subgrain growth: This model proposes that cells coarsen by cell wall movement, similar to the abnormal grain growth. The driving force originates from single dislocations and dislocations from neighbouring cells or sugrains.

Nucleation by subgrain coalescence (SGC): This model proposes that the cell walls are com-posed of dislocations. The cell walls of neighbouring grains can dissolve by dislocation climb and hence grow. This kind of nuclei form within subgrains of high stored energy and favours the formation of {1 1 1}h1 1 0i and {1 1 0}h1 1 0i orientations. The fraction of {1 1 1} grains is larger than that of {1 1 0} grains, so that the final texture will be dominated by {1 1 1} grains.

Nucleation by subgrain boundary relaxation: This model proposes that preferentially elonga-ted cells or subgrains coarsen on the expense of neighbouring subgrains, if a force disequi-librium exists at nodes or boundaries. This is especially the case in the vicinity of micro or shear bands. As a result of this, the cell can grow to its critical size.

Nucleation by strain induced boundary migration (SIBM): This model proposes that a nucle-us forms by bulging of existing grain boundaries of the deformed microstructure. This mechanism allows subgrains with a low stored energy, e.g. {0 0 1}h1 1 0i or {1 1 2}h1 1 0i subgrains, to grow into areas of high stored energy, promoting α-fibre texture components.

Advanced hot rolling strategies for IF and TRIP steels

Advanced hot rolling strategies

for

Advanced hot rolling strategies

for

IF and TRIP steels

Alexander ELSNER

Contents

Chapter 1

Scope of the Thesis

1.1

Ferritic Rolling of Deep-Drawing Steels

1.2

Intercritical Rolling of Low Alloy TRIP Steels

1.3

Outline

Chapter 2

Introduction

2.1

Properties of Deep-Drawing Steels

2.2

Crystallographic Textures

2.2.1

Introduction

2.2.2

Orientation Description

a)

b)

c)

2.2.3

Macro Texture Measurement and Representation

2.3

Conventional Production of Deep-Drawing Steels

2.4

Ferritic Rolling of Deep-Drawing Steels

2.4.1

Hot Strip Grades

2.4.2

Cold Strip Grades

2.5

Influence of Solute Carbon

2.5.1

Recrystallisation Mechanisms