Advanced Hot Rolling Strategies for IF and TRIP Steels

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. Dit proefschrift is goedgekeurd door de promotoren: Prof. dr. ir. S. van der Zwaag 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: [email protected] keywords: IF steel, TRIP steel, ferritic rolling, intercritical rolling, microstructure, mechanical properties, texture development

ISBN 90-407-2591-8

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.

Printed in the Netherlands. 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).

Contents

1 Scope of the Thesis 1 1.1 Ferritic Rolling of Deep-Drawing Steels ...... 2 1.2 Intercritical Rolling of Low Alloy TRIP Steels ...... 3 1.3 Outline ...... 4

2 Introduction 7 2.1 Properties of Deep-Drawing Steels ...... 7 2.2 Crystallographic Textures ...... 10 2.2.1 Introduction ...... 10 2.2.2 Orientation Description ...... 10 2.2.3 Macro Texture Measurement and Representation ...... 11 2.3 Conventional Production of Deep-Drawing Steels ...... 14 2.4 Ferritic Rolling of Deep-Drawing Steels ...... 16 2.4.1 Hot Strip Grades ...... 17 2.4.1.1 ULC/ELC Steel ...... 18 2.4.1.2 IF Steels ...... 20 2.4.2 Cold Strip Grades ...... 22 2.4.2.1 ELC/ULC Steels ...... 22 2.4.2.2 IF Steels ...... 23 2.5 Inﬂuence of Solute Carbon ...... 25 2.5.1 Recrystallisation Mechanisms ...... 25 2.5.2 Recrystallisation Texture of Ferritic Rolled Hot Strip ...... 27 2.6 Inﬂuence of Lubrication ...... 31

vii viii CONTENTS

3 Experimental 37 3.1 Materials ...... 37 3.1.1 Chemical Composition ...... 37 3.1.2 Specimen Preparation ...... 38 3.2 Deformation Dilatometry ...... 38 3.2.1 Introduction ...... 38 3.2.2 Continuous Cooling Transformation Diagrams ...... 38 3.3 Hot and Warm Rolling ...... 40 3.3.1 Laboratory Hot Rolling Mill ...... 40 3.3.2 Lubrication of the Roll Gap ...... 41 3.3.2.1 Lubrication System ...... 41 3.3.2.2 Selection of the Lubricant ...... 41 3.3.3 Rolling Schedules ...... 42 3.3.3.1 "Soft" Hot Strip ...... 43 3.3.3.2 "Hard" Hot Strip (annealed) ...... 44 3.3.3.3 "Cold Strip" ...... 44 3.4 Pickling ...... 45 3.4.1 Pre-Tests ...... 45 3.4.2 Laboratory Pickling ...... 46 3.5 Cold Rolling ...... 47 3.6 Annealing ...... 47 3.6.1 Batch Annealing Simulation ...... 47 3.6.2 Continuous Annealing Simulation ...... 47 3.6.2.1 RHESCA Simulator ...... 49 3.6.2.2 Salt Bath ...... 49 3.6.2.3 Comparison of the continuous annealing pre-tests ...... 51 3.7 Mechanical and Technological Testing ...... 52 3.7.1 Tensile Tests ...... 52 3.7.2 Cupping Tests ...... 52 3.7.3 Texture Measurements ...... 53 3.7.4 EBSD Measurements ...... 54 CONTENTS ix

4 Hot Strip Grades – Results and Discussion 55 4.1 Minimum Coiling Temperature for the "Soft" Hot Strip ...... 55 4.2 Inﬂuence of the Ferritic Rolling Reduction ...... 57 4.3 Mechanical Properties ...... 59 4.3.1 "Soft" Hot Strip ...... 59 4.3.2 "Hard" Hot Strip ...... 61 4.3.2.1 Conventional Batch Annealing ...... 61 4.3.2.2 "Direct" Batch Annealing ...... 62 4.4 Texture Development ...... 63 4.4.1 "Soft" Hot Strip ...... 63 4.4.2 "Hard" Hot Strip ...... 65 4.4.2.1 Conventional Batch Annealing ...... 65 4.4.2.2 "Direct" Batch Annealing ...... 65 4.5 Discussion ...... 65 4.5.1 "Soft" Hot Strip ...... 65 4.5.2 "Hard" Hot Strip ...... 68 4.6 Industrial Implications ...... 69 4.6.1 "Soft" Hot Strip ...... 69 4.6.2 "Hard" Hot Strip ...... 72 4.6.3 Financial Aspects ...... 73 4.6.4 Conclusion ...... 73

5 Cold Strip Grades – Results and Discussion 75 5.1 Microstructural Development ...... 76 5.1.1 Initial Hot Strip "Soft" ...... 76 5.1.2 Initial Hot Strip "Hard" ...... 78 5.2 Mechanical Properties ...... 78 5.2.1 Initial Hot Strip "Soft" ...... 81 x CONTENTS

5.2.2 Initial Hot Strip "Hard" ...... 82 5.3 Texture Development ...... 82 5.3.1 Initial Hot Strip "Soft" ...... 82 5.3.2 Initial Hot Strip "Hard" ...... 86 5.4 Cupping Tests ...... 86 5.4.1 Initial Hot Strip "Soft" ...... 86 5.4.2 Initial Hot Strip "Hard" ...... 88 5.4.3 Calculated r-Value Distribution and Relative Earing ...... 88 5.5 Discussion ...... 90 5.5.1 Microstructural Development ...... 90 5.5.2 Mechanical Properties ...... 91 5.5.3 Texture Development ...... 92 5.5.4 Cupping Tests ...... 93 5.6 Industrial Implications ...... 93

6 Towards a new Hot Rolling Strategy for Low Alloy TRIP Steels 95 6.1 Processing Routes ...... 96 6.1.1 Cold rolling and annealing ...... 96 6.1.2 Hot rolling ...... 97 6.1.3 Intercritical Rolling ...... 99 6.1.4 The Hypothesis ...... 99 6.2 Experimental ...... 100 6.2.1 Material ...... 100 6.2.2 Flat Compression Tests ...... 100 6.2.3 Deformation Dilatometry ...... 102 6.2.4 Tensile Tests ...... 104 6.2.5 Determination of the Fraction Retained Austenite ...... 104 6.3 Results and Discussion ...... 106 CONTENTS xi

6.3.1 Deformation Dilatometry Tests ...... 106 6.3.2 Flat Compression Tests ...... 108 6.3.2.1 Inﬂuence of the Coiling Temperature ...... 108 6.3.2.2 Inﬂuence of the Intercritical Deformation Temperature . . . . . 109 6.3.3 Mechanical Properties ...... 113 6.3.4 Conclusion ...... 113

Summary 115

Samenvatting 121

References 127

List of Frequently used Symbols 141

List of Publications 143

Acknowledgements 145

Curriculum Vitae 147

Chapter 1

Scope of the Thesis

The production of mild and high strength steel sheets with a thickness of less than 2 − 3mm 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 complex 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 production 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 1.5mm or even 1mm

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

• Improved surface quality

• Improved dimensional tolerances

1 2 CHAPTER 1 SCOPE OF THE THESIS

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 thermomechanical 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 hardening. 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 temperature, i.e. well above the Ar3 temperature, close above the austenitic recrystallisation temperature. The properties of the steel strip are achieved by the rapid recrystallisation of the deformed austenite 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, producing 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 applicable 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 rolling seems 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 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 1.5mm 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 ≤ 2mm) and even ultra thin (ts ≤ 1mm) 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 {111} 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 cooling 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. 4 CHAPTER 1 SCOPE OF THE THESIS

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 operations. 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 processing 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 testing 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. 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.

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 ≤ 2mm 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 acceptable 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 performance 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

7 8 CHAPTER 2 INTRODUCTION

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

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

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 conventionally 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 ◦ − r ◦ + r ◦ ∆r = 0 45 90 . (2.2) 2

In equation 2.2 the indices designate the angle of the direction, with respect to the rolling direction, 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, rm, deﬁned in equation 2.3, yields an optimal deep-drawing performance.

r ◦ + 2r ◦ + r ◦ r = 0 45 90 (2.3) m 4

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 distributions 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 {111} 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 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 {111} texture. Steel Grades DC 04 and DC 05 are aluminium-killed low and extra low carbon steels (typically 0.05mass% 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 Ti or > 1 in [atom%] (2.5) C + N C + N 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 parallel 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]. 10 CHAPTER 2 INTRODUCTION

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 tensile straining of a single 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 geometry, C = {c1,c2,c3} = {RD,TD,ND} and the crystal coordinate system parallel to the crystal unit cell, S = {s1,s2,s3} = {[100],[010],[001]}. RD denotes the rolling direction, TD the transverse 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 (hkl)[uvw], or {hkl}huvwi for non speciﬁc indices [27]. {hkl} 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. huvwi 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]. 2.2 CRYSTALLOGRAPHIC TEXTURES 11

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 measurement 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 = 2dhkl sinΘ (2.6) a dhkl = √ (2.7) h2 + k2 + l2

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) 12 CHAPTER 2 INTRODUCTION

a)a-fibre <110>||RD b) j1 0°45° 90° 0 153045607590 j 0 2 {001}<110> 90° 15

30 F {112}<110> 45 F {111}<110> {111}<112> {554}<225> 60 {111}<112> {111}<110> 90° {332}<112> 75 {110}<110> {011}<100> 90

j2 = 45° j1 g-fibre <111>||ND

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. 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 – h110i||RD: {100}h110i to {110}h110i The α-fibre includes all orientations with their h110i 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 component {001}h110i (Φ = 0◦), the inverse brass component {112}h110i (Φ = 35◦) and {111}h110i (Φ = 54.7◦). At Φ = 54.7◦ the α-fibre intersects the γ-fibre.

2. γ-fibre – h111i||ND: {111}h110i to {111}h112i The γ-fibre contains all orientations with their {111} plane parallel to the sheet plane. The γ-fibre is also called {111}-fibre. The γ-fibre contains the important texture components 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 often observed that the maxima in a recrystallisation do not lie perfectly along the γ-fibre at {332}h113i or {554}h225i. 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 inhomogeneous 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 {001}h110i (Φ = 0◦), the copper orientation {112}h111i (Φ = 35◦), the intersection point with the γ-fibre {111}h112i (Φ = 54.7◦) and the Goss orientation {011}h100i (Φ = 90◦). The Goss component ◦ ◦ ◦ {011}h100i 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. 14 CHAPTER 2 INTRODUCTION

roughing

finishing roughing Ar3 ag+ Ar ag+ 1

Temperature CMn finishing

coiling a)coiling b) Time

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 region. 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 Ar3 and Ar1 temperatures 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 ≥ 2mm, 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 2.3 CONVENTIONAL PRODUCTION OF DEEP-DRAWING STEELS 15

Figure 2.5: Inﬂuence of the cold reduction on rm for 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 formation 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 ≈ 150MPa, Rm ≈ 280MPa 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 requires 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 h111i||ND. The intensity of the {111} 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. 16 CHAPTER 2 INTRODUCTION

2.4 Ferritic Rolling of Deep-Drawing Steels

The production of thin gauge hot strip (t f < 2mm) by conventional (austenitic) hot rolling involves 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 finishing 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 improvements 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 industrially 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. 2.4 FERRITIC ROLLING OF DEEP-DRAWING STEELS 17

260 ag+ 240

220 e = 0.6 200 e = 0.4 180

160

flow stress [MPa]

140 e = 0.4 120 600 650 700 750 800 850 900 950 1000 Temperature [°C]

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 (h111i||ND) and a partial α-fibre (h110i||RD). Hence, using an adapted post rolling processing it is possible to produce a final thin hot strip with a desirable {111} 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 18 CHAPTER 2 INTRODUCTION

a) "Soft” hot strip b) “Hard” hot strip (annealed) c) “Cold” strip (annealed)

a + g a + g a + g

a annealing a annealing a finishing

Temperature

coiling

Time Time Time

Figure 2.7: Variants of ferritic rolled products

for the production of deep-drawable thin hot strips in a thickness range of t f ≈ 1−1.5mm. 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 t f ≤ 1mm [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. 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 εFR is 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 temperature 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 intensity 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 recrystallisation 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 (180MPa versus 230MPa) and tensile strength (370MPa versus 300MPa) 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]. 20 CHAPTER 2 INTRODUCTION

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 − 40µm in transverse and about 50 − 80µm parallel to the rolling direction [8]. "Hard" Hot Strip For the case of lower finish rolling and, hence, coiling temperatures, the strip does not recrystallise 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 samples 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} texture. The grains with γ-fibre orientation were completely consumed in favour of the α-fibre orientations. The annealing textures exhibit a maximum at {100}h110i 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 {111} 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 2.4 FERRITIC ROLLING OF DEEP-DRAWING STEELS 21

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 50m and a rolling speed of 12 m/s and a strip thickness of 2mm 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 {332}h113i. 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 90MPa, a ultimate tensile strength of 275MPa 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 texture of the strip consisted for both coiling temperatures a strong rolling texture with an intense 22 CHAPTER 2 INTRODUCTION

γ-ﬁbre and and a maximum at {112}h110i along the α-ﬁbre [8]. During the subsequent annealing treatment – either continuous or discontinuous – the rolling 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 − 290MPa;Rp0.2 ≈ 90 − 100MPa;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 transferred 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 − 60 K/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 conventionally 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 2.4 FERRITIC ROLLING OF DEEP-DRAWING STEELS 23

{100}h110i along the α-fibre [8]. The microstructure was fully recrystallised with polygonal ferrite grains. With increasing cold rolling reduction the texture components near {100}h110i 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 {111} 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 mechanical 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 The "soft" hot strip produced of a IF steel exhibits, presupposing optimised ferritic rolling pa- 24 CHAPTER 2 INTRODUCTION

rameters, a strong γ-fibre texture with a maximum intensity near {332}h113i. The initial microstructure 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 {100}h110i and {112}h110i. For 80% reduction the γ-fibre was more homogeneous and the typical rolling texture components along the α-fibre formed, with a strong maximum near {112}h110i[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 {111} 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 subsequent cold rolling of these strips, the γ-fibre was not intensified to a large extent. However, the rolling texture component near {112}h110i 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 {554}h225i developed for a cold reduction of 80% [57]. The mean r-values calculated in [8] together with the 2.5 INFLUENCE OF SOLUTE CARBON 25

mechanical propreties were already for a cold reduction of 40% comparable to that of a conventionally 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 {001}h110i 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

2γ 4γ r = = , (2.8) c p ρGb2 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 dislocation distribution or by large local sub grains, which are commonly formed during recovery phase. 26 CHAPTER 2 INTRODUCTION

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 composed 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 {111}h110i and {110}h110i orientations. The fraction of {111} grains is larger than that of {110} grains, so that the ﬁnal texture will be dominated by {111} grains.

Nucleation by subgrain boundary relaxation: This model proposes that preferentially elongated 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 nucleus forms by bulging of existing grain boundaries of the deformed microstructure. This mechanism allows subgrains with a low stored energy, e.g. {001}h110i or {112}h110i subgrains, to grow into areas of high stored energy, promoting α-ﬁbre texture components. 2.5 INFLUENCE OF SOLUTE CARBON 27

The most important nucleation mechanism for the recrystallisation of deep-drawing steels are subgrain coalescence SGC and strain induced boundary migration SIBM [26]. Selective Growth The selective growth theory proposes that the resulting texture is only inﬂuenced by different growth rates of the randomly oriented nuclei into the deformed matrix. The development of a pronounced {111}h112i is explained by a favourable growth relation (35◦ rotation about the h110i direction) to the {112}h110i orientation [44, 61].

2.5.2 Recrystallisation Texture of Ferritic Rolled Hot Strip

In [51] a low carbon (LC), a low Manganese (LMn) and an IF steel were rolled at temperatures between 70 ◦C and 700 ◦C and subsequently annealed at 700 ◦C. The chemical compositions of the steels are given in Table 2.4.

C P S Mn Al Ti LC 0.014 0.01 0.01 0.22 0.0001 - LMn 0.016 0.003 0.002 0.009 0.041 - IF 0.005 0.01 0.01 0.13 0.042 0.084

Table 2.4: Chemical compositions in mass% of the steels used in [51, 52]

The microstructure of cold (70 ◦C) and warm (700 ◦C) rolled deep-drawing steels, observed in a light optical microscope (LOM) or scanning electron microscope (SEM), consists of smooth and banded grains. The banded grains contain deformation inhomogeneities, the so-called in-grain shear bands. Examples of the appearance of grains containing this kind of banded structure can be found for example in [51, 62]. The same microstructural feature has also been named, shear band, creased, "fish bone" or microband structure by other authors [51, 59, 63]. In-grain shear bands carry localised deformation accompanied by a large crystal fragmentation across the band [60]. In-grain shear bands predominantly occur in γ-fibre grains. In [60] 85% of the measured γ-fibre orientations (±15 ◦ along h111i||ND) in a cold rolled IF were associated with in-grain shear bands. The microstructure of the LC samples [51] rolled at temperatures below approximately 400 ◦C contained some grains with a well developed in-grain shear bands, whereas those rolled at higher temperatures within the ferritic rolling temperature region only few bands were observed. These bands were only thin and short, indicating that the flow along them was retarded. In contrast to the LC steel, the IF steel exhibited a relative high amount of banded grains independently of the rolling temperature [45, 64, 65]. Apart from the rolling temperature, also the rolling strain influences the density of in-grain shear bands. A higher rolling strain intensifies the in-grain shear bands [66, 67]. 28 CHAPTER 2 INTRODUCTION

0.25 Low Carbon Steel 0.20 IF Steel Pure Iron 0.15 Mild Steel 0.10 Austenite

0.05 Low Carbon Steel

0.00 IF Steel

strain rate sensitivity m [-]

-0.05 0 0.2 0.4 0.6 0.8 1

homologous deformation temperature Tdef /T m0 [-]

Figure 2.8: Strain rate sensitivity as a function of the homologous deformation temperature [51]

The different behaviour of the LC and the IF steel in terms of the shear band density and temperature dependency can be attributed to a different strain rate sensitivity, m , and hence to the dynamic strain ageing (DSA) behaviour of the steels [53, 60, 62, 65, 66, 68–70]. The strain rate sensitivity is deﬁned as

d lnσ m = . (2.9) d lnε˙

The DSA is in turn caused by solute carbon (and/or Nitrogen), in a temperature region allowing the interstitial atoms to diffuse to and pin the moving dislocations during straining. The DSA peak is often observed in tensile tests (ε˙ ≈ 10−2 s−1) in a temperature region of 200−350 ◦C. For the case of ferritic rolling conditions with higher strain rates, e.g. ε˙ ≈ 102 s−1, the temperature of the DSA peak has to be higher to compensate for the higher strain rate [51].

From Fig. 2.8 it can be seen that the LC steel exhibits at approximately Tde f/Tm ≈ 0.25 or Tde f ≈ 300 ◦C a negative strain rate sensitivity, promoting the localisation of ﬂow. In this temperature region a high amount of in-grain shear bands is observed. For higher rolling temperatures, like in ◦ ferritic rolling, of about 500 − 700 C, Tde f/Tm ≈ 0.5 − 0.6, the LC steel exhibits a rather high m- value. This provides a more homogeneous distribution of the strain and explains the retardation of the formation of the in-grain shear bands in the ferritic rolled LC and LMn steels. The IF steel shows a low strain rate sensitivity, m ≈ 0, for temperatures up to the ferritic rolling temperature region. The low m-value indicates the sufﬁcient high density of in-grain shear bands. The rolling and annealing textures for the different ferritic rolling temperatures are shown in Fig. 2.9 for the LC and IF steel. Similar texture results of ferritic rolled strips were also presented in [8, 26, 57] for rolling temperatures between 700 ◦C and 760 ◦C. The rolling textures of the warm rolled steels, either ELC or IF, are similar to that of cold rolled products [59, 62, 71] 2.5 INFLUENCE OF SOLUTE CARBON 29

j j 0°1 90° 0°1 90° 0° 0°

F F

90° 90°

a) rolling texture b) recrystallisation texture

Figure 2.9: a) Rolling texture and b) recrystallisation texture of warm rolled LC and IF steels for different ferritic rolling temperatures (intensity levels 2,3,4,5,...) [51, 52]

and are characterised by a partial α-ﬁbre and a complete γ-ﬁbre. The rolling textures of the LMn steel were however similar to that of the LC steel. From Fig. 2.9 a) it is visible that the rolling texture for the LC steel is sharper at 700 ◦C than that at 70 ◦C. With decreasing rolling temperature the texture intensity of the LC steel slightly decreased. This can be attributed to the increasing amount of in-grain shear bands with decreasing rolling temperatures [51]. The crystal fragmentation caused by in-grain shear bands serves to weaken the deformation texture [72]. For the case of the IF steel the texture is relatively insensitive to the rolling temperature, which correlates to the approximately constant amount of in-grain shear bands.

The influence of the rolling temperature on the recrystallisation texture is evident from Fig. 2.9 b). The influence of the rolling temperature is greater in the LC than in the IF steel. The recrystallisation texture of the sample rolled at 700 ◦C is composed of a partial α-fibre with a strong cube component near {001}h110i. The γ-fibre orientations observed in the rolling texture have disappeared. For lower rolling temperatures of 300 ◦C the texture changed to a γ-fibre and an intense Goss component ({011}h100i). A further reduction of the rolling temperature to 70 ◦C intensified the γ-fibre and reduced the Goss component. The development of the different recrystallisation textures is also explained in terms of in-grain shear bands [66, 73, 74]. 30 CHAPTER 2 INTRODUCTION

In [60, 63, 75] it has been shown that the in-grain shear bands, accompanied by a high crystal fragmentation, act as additional nucleation sites. The dominant nucleation mechanism is in this case the SGC mechanism, which is known to favour the formation of γ-fibre texture. For the case of LC steels, with a high strain rate sensitivity in the temperature region of ferritic rolling, the formation of the in-grain shear bands is retarded. The absence of the in-grain shear bands and consequently a more homogeneous strain distribution gives rise to the SIBM nucleation mechanism, promoting the formation of an α-fibre texture [76]. The same correlation between the density of in-grain shear bands and the intensity of the γ-fibre in the annealing texture is observed in [66] for different cold rolling reductions. It has been found that for lower reductions of approximately 50% only a weak γ-fibre texture developed. For increasing rolling reductions of up to 90% the intensity of the γ-fibre texture markedly increased. This was explained in terms of preffered nucleation of γ-fibre orientations within γ-fibre grains. This is again consistent with the increasing amount of in-grain shear bands in γ-grains for higher rolling reductions.

The electron backscatter diffraction (EBSD) measurements in [60] show that the orientations of in-grain shear bands form a strong γ-ﬁbre texture with a maximum near {554}h225i. For these measurements an IF steel was ferritic rolled at 650 − 700 ◦C. These results might give an explanation for the occurrence of the peak like {554}h225i texture observed in a highly deformed ferritic and cold rolled steel as observed in [57]. A similar development of the {554}h225i texture component with increasing cold reduction has been observed in [71]. In this work IF steel has been cold rolled from 70 − 90%. On the basis of these results it is postulated that the in-grain shear bands are involved in the nucleation of the often observed {554}h225i texture component. This change in the recrystallisation texture (homogeneous to peaked γ-ﬁbre texture) has been explained in [77] in terms of an alteration in the nucleation mechanism from oriented nucleation for lower rolling reductions (75%) to selective growth for higher rolling reductions (95%). In [78] it is stated, that for reductions below 80% the develoment of the recrystallisation texture is almost completely controlled by oriented nucleation.

On the basis of the above results found in literature it can be concluded, that it is not the solute carbon per se, which causes the detrimental recrystallisation texture in ferritic rolled LC steel but the absence of in-grain shear bands. However, the absence of the in-grain shear bands is caused by the occurrence of DSA and the DSA behaviour is in turn depending on the content of solute carbon (or nitrogen). It is hence the effect of the solute carbon during the ferritic rolling and not during annealing, causing the detrimental effect on the recrystallisation texture. This was also shown in [8] by cold rolling and annealing of an austenitically rolled ELC steel with a high and low amount of solute carbon. Whereas the two rolling texture showed no signiﬁcant difference, the annealing texture of the strip with the lower amount of solute carbon showed a much stronger γ-ﬁbre. By a special over-ageing treatment it was proven, that the difference in the annealing texture was not caused by the different amount of solute carbon during annealing. Hence, it must have been caused by the solute carbon during deformation. Since the rolling textures were nearly identical, these results strengthen the in-grain shear band hypothesis.

The hypothesis is also strengthened by the results presented in [69]. In this investigation chromium (Cr) and boron (B) were added to an LC steel. The addition of Cr to the steel modiﬁes the 2.6 INFLUENCE OF LUBRICATION 31

DSA behaviour of the steel by creating a plateau to the DSA peak, thereby shifting it to higher temperatures. This markedly reduced the strain rate sensitivity of the steel in the range of ferritic rolling temperatures. It has been found that in the LC steel alloyed with Cr at a rather high rolling temperature of 780 ◦C still more than 50% of the grains contained in-grain shear bands. Also in [59, 74] it has been shown that the addition of Cr to LC steels enables the production of a {111}-texture by ferritic rolling, without removing the carbon from solution.

2.6 Inﬂuence of Lubrication

From the previous sections it is obvious that a strong {111}-texture is necessary to produce a ferritic rolled hot strip with desirable deep-drawing properties. The texture of the strip is however not constant over the strip thickness, t. This is caused by an inhomogeneous deformation across the strip thickness, which is in turn produced by redundant shear introduced by the frictional forces between the rolls and the work piece [79, 80]. In [81] it has been shown that the intensity and depth of the sheared layers below the surface is depending on the friction in the roll gap and the geometry of the roll gap. This was shown by rolling of a specimen with a constant roll gap geometry and varying lubrication conditions and by the rolling with different roll diameters without lubrication. From Fig. 2.10 a) - c) it can be seen that the marking, originally perpendicular to the strip surface, shows an increasing curvature and strip thickness with increasing friction. The increasing strip thickness is caused by the increasing rolling force, due to the higher friction forces. Similar results have been also presented in [34, 82, 83]. The intensity and depth of the sheared zone is, however, not only depending on the lubricant but also to a great extent on the roll gap geometry. The roll gap geometry is conventionally deﬁned by

l pr(t −t ) d = 0 1 , (2.10) t0 t0

with ld is the contact arc, r the roll radius, t0 the initial strip thickness and t1 the ﬁnal strip thickness. This tendency can be seen in Fig. 2.11, as the magnitude of the sheared zone near the surface increases with an increasing roll diameter.

With increasing ld/t0 ratio, the inhomogeneous deformation across the thickness increases [33], as it can be seen in Fig. 2.11 a) – c). The ld/t0 ratio increased in this case from approximately 1.63 to 5.19. The increasing ld/t0 ratio involves in this case (constant rolling reduction) an enlargement of the contact area between the roll and the strip. Therefore, higher rolling forces are required to obtain the same reduction. From a tribological point of view the inﬂuence of friction is more pronounced with an increasing ld/t0 ratio, due to the larger surface areas in contact. Additionally the lubricant entrainment and pressure distribution is considerably inﬂuenced by the changed roll gap geometry. 32 CHAPTER 2 INTRODUCTION

Figure 2.10: The effect of friction on the through thickness strain distribution; Rolled at 700 ◦C, with the roll gap adjusted for a reduction of 30%, a constant roll diameter and a coefﬁcient of friction of a) µ = 0.1 b) µ = 0.2 c) µ = 0.3 [81]

Figure 2.11: The effect of the roll diameter on the through thickness strain distribution. Rolled at 700 ◦C, with the roll gap adjusted for a reduction of 30% without lubrication and a roll diameter of a) 70mm, b) 250mm and c) 710mm [81] 2.6 INFLUENCE OF LUBRICATION 33

It is reported in literature, that lubrication is more effective under the condition of a higher ld/t0

ratio [76, 84]. The ld/t0 ratio conventionally increases from one finish rolling stand to the other, so that the lubrication in the last finishing stand(s) is more effective. The inhomogeneous deformation across the thickness also influences the distribution of the texture across the thickness. To quantify the through thickness texture gradient, the texture is usually measured at planes with fixed distances, x, to the strip centre. These planes are conventionally identified by the s-parameter:

x s = (2.11) 0.5t with x is distance between the observed plane and the specimen centre and t the strip thickness. A value of s = 0 identifies the centre layer and s = 1 the surface layer. The layers within s = 0 − 0.4 are classified as near centre and those within s = 0.7 − 1 as near surface layers [33]. The near centre layers are predominantly deformed by plane strain compression, whereas the near surface layers are predominantly deformed by shearing [33, 79, 85–87]. The texture of the sheared layers are dominated by the Goss ({011}h100i) component, whereas the plane strain compressed layers exhibit a typical rolling texture consisting of a partial α- fibre and a complete γ-fibre [82, 83, 85, 88]. The maximum intensity of the shear texture is conventionally found at about s = 0.8 [33]. During the subsequent recrystallisation of the IF steel strip, the Goss texture near the surface remains unchanged, whereas the rolling texture in the centre layer is transformed into an desirable {111} texture [85]. Even during further processing, such as cold rolling and annealing, the undesirable Goss texture cannot be removed [82]. The deep-drawing properties of the final ferritic rolled strip are, however, depending on the through thickness texture. Hence, unlubricated rolling in the ferritic temperature region results in a hot strip with a rather poor deep-drawing properties, because of the detrimental Goss texture near the surface [16, 82]. To improve the deep-drawing properties of the strip the shear texture near the surface has to be minimised. The use of a proper lubrication during ferritic rolling is thought to be a solution to suppress the shear texture [81, 82]. In Fig. 2.12 a) it can be seen, that with a decreasing friction coefficient, µ, the r-value increases markedly. Fig. 2.12 b) shows the influence of lubrication on the r-value. It can be seen, that the use of proper lubrication increases the r-value from 1 − 1.2 to approximately 1.8 for rolling temperatures below 700 ◦C. In lubricated ferritic rolling the deformation is much more homogeneous across the thickness, Fig. 2.10 a), so that a desirable rolling texture and, hence, recrystallisation texture develops uniformly through the thickness [16, 82, 86]. A friction coefficient of µ ≤ 0.15 is claimed in the literature to ensure a sufficiently uniform microstructure and texture in the as-rolled and annealed specimens [59, 81, 81, 82]. A variety of lubricants is used, such as:

• glass or salt [80, 84], 34 CHAPTER 2 INTRODUCTION

1.6 with lubrication 2.0 without lubrication 1.4 1.8 Ar3 1.6 1.2 1.4

r - value [-] r - value1.2 [-] 1.0 1.0 0.8 0.8 0.1 0.2 0.3 500 600 700 800 900 a) friction coefficient µ [-] b) rolling temperature,temperature [°C] °C

Figure 2.12: a) Effect of the friction on the r-value of a sheet steel rolled at 700 ◦C, b) Effect of lubrication and the rolling temperature on the r-value (rolling reduction 65%; annealed at 750 ◦C for 5h) [79]

• parafﬁn oil based lubricant [86],

• tallow lubrication [83, 84],

• mineral oil based lubricant [16, 76, 79, 81–83],

• ester based oil [16, 76, 82, 83].

Glass yields the best reduction of the friction in hot rolling, because it is the only lubricant, which offers a sufficiently high viscosity and pressure resistance to produce hydrodynamic lubrication [84]. Hydrodynamic lubrication provides a complete separation of the rolls and the strip, minimising the friction force and shear deformation during rolling. However, the difficulty of removing the glass after the process and the possible abrasiveness of the solidified glass are important reasons, why this lubricant is not industrially used in strip rolling [80, 84]. The lubricants widely used in ferritic rolling are mineral based or ester based oils. The effectiveness of these lubricants is, however, strongly depending on the rolling temperature and the chemical composition of the oils. The results in [76, 83] show that an increasing amount of esters decreases the amount of shear deformation. Only for rolling temperatures below 500 ◦C pure mineral oils yielded a sufficient reduction of the shear texture [76]. Ester oils are more stable at higher temperatures, compared to mineral oils, which become usually unstable above approximately 500 ◦C. Unfortunately it cannot be generalised, that ester based oils provide superior results, because the tribological system is also sensitive to the different additives, the steel grade and to a great extent to the amount and nature of the scale layer. Additionally to the reduction of the undesirable effect of the shear texture near the surface, the use of hot rolling lubrication brings about a couple of other advantages [89–93]: 2.6 INFLUENCE OF LUBRICATION 35

• Reduction of the roll wear, ranging from typical 15% up to 100%

• Reduction of the rolling load by 10 − 20%

• Reduction of the power consumption by up to 25%

• Reduction of the temperature losses during rolling, by utilising exothermic additives

• Improved product quality

However, the most important reason for the use of lubrication in ferritic rolling is the minimisation of the undesirable shear texture near the strip surface. From the literature [76, 79, 81– 83, 85, 86, 88] it is evident, that the production of deep-drawable ferritic rolled hot strip necessitates a proper lubrication of the roll gap, at least in the last ﬁnishing stand(s).

Chapter 3

Experimental

3.1 Materials

From the literature on ferritic rolling of deep-drawing steel it is obvious, that interstitial free (IF) steels are the most promising steel grade for the production of thin gauge deep-drawable hot strips by this processing. Therefore, this research was focused on the inﬂuence of the ferritic rolling parameters on the deep-drawing properties of the ﬁnal strip produced of commercial IF steel. The materials used for this study were two industrially produced IF steels, delivered from the production of ThyssenKrupp Stahl in Duisburg. The samples for the laboratory test were taken from extended cropped portions of the strip after rough rolling with a thickness of about 47mm. These pieces were cooled in still air down to ambient temperature.

3.1.1 Chemical Composition

The chemical composition of the IF steels is shown in Table 3.1. The IF-Ti steel is stabilised with titanium, which corresponds to the classical DC 06 composition. The IF-TiNb is stabilised with a titanium and niobium. Both steels are fully stabilised, to ensure an interstitial free matrix.

Steel C Si Mn P S N Al Ti Nb IF-Ti 0.001 0.017 0.130 0.005 0.008 0.004 0.032 0.064 0.002 IF-TiNb 0.002 0.007 0.097 0.010 0.004 0.003 0.042 0.035 0.028

Table 3.1: Chemical compositions in mass% of the IF steels

37 38 CHAPTER 3 EXPERIMENTAL

3.1.2 Specimen Preparation

After the industrial rough rolling, the material was reheated to approximately 1200 ◦C and pre- rolled to nearly the initial strip thicknesses for the laboratory rolling test and cooled down to the room temperature in still air. Subsequently the specimens were cut to 90mm x 80mm (RD x TD) and milled to the desired initial hot strip thicknesses. The initial hot strip thicknesses were chosen such that the ﬁnal strip thickness was 1 and 2mm for all ferritic rolled hot strip grades and 0.5mm for all cold strip grades. The initial strip thickness for the hot strip grades varied from 3 to 11mm depending on the desired rolling reduction and ﬁnal strip thickness. For the cold strip grades, the initial strip thickness ranged from 4 to 16.6mm, depending the desired cold reduction. As a reference, strips with an initial thickness of 6.7,13.5 and 27mm were machined for conventional austenitic rolling and subsequent cold rolling and annealing. For the deformation dilatometry tests, cylinders with a diameter of 5mm and a height of 10mm were machined parallel to the normal direction.

3.2 Deformation Dilatometry

In section 2.4 it has been shown, that the rolling in the α/γ-transformation region, deteriorates the deep-drawing properties. Therefore, detailed knowledge of the transformation temperatures of the steels for the different cooling rates of strip is essential. The transformation temperatures were determined using the deformation dilatometry.

3.2.1 Introduction

For the determination of the transformation temperature, a Bähr 805A/D deformation dilatometer was used. Using the deformation mode, cylindrical specimens with a typical diameter of 5mm and a length of 10mm are heated in an induction coil. During heating and austenitisation the measurement chamber of the dilatometer is evacuated. Subsequently the specimen can be compressed to a maximum strain of 1.2, with a maximum strain rate of 10s−1 [94]. To generate continuous cooling transformation diagrams, the specimen can be cooled with gas, to simulate the cooling of the real specimen. The temperature is continuously measured with two thermocouples (S-type with a diameter of 0.2 mm), welded to the specimens. One thermocouple is used to control the temperature and the other as a reference. At the same time, the length change of the specimen is recorded. On the basis of this data, the transformation temperatures can be determined using the tangential method, recommended in [95].

3.2.2 Continuous Cooling Transformation Diagrams

The continuous cooling transformation (CCT) diagrams are constructed by plotting a series of cooling curves into a temperature versus time diagram. Then the transformation start temperatu- 3.2 DEFORMATION DILATOMETRY 39

a) b) 940

920

Ar3 Ar 900 3

880 Ar Ar 1 860 1

840

820

800 1020 50 100 200 500 1000 10 20 50100 200 500 1000 c) d) 940

920 Ar 3 Ar3 Temperature [°C] 900

880 Ar1 860 Ar1

840

820

800 1020 50 100 200 500 1000 10 20 50100 200 500 1000 Time [s]

Figure 3.1: Continuous cooling transformation diagrams for the steel IF-Ti a) without and b) with an austenitic deformation of ε = 0.5 at 950 ◦C and for the steel IF-TiNb c) without and d) with an austenitic deformation of ε = 0.5 at 950 ◦C res and transformation ﬁnish temperatures are connected. The α to γ transformation behaviour depends, apart from the chemical composition, on the austenite condition, immediately at the start of the transformation [96–103]. In order to determine the range of cooling rates relevant to the project, the time that a 1 or 2mm thick strip needs to cool from 800 down to 500 ◦C, the so-called t8/5 time, was measured. On this basis, the controlled cooling with t8/5 times of 0.8, 2, 4, 8min was chosen for the dilatometer tests. The tests with the same cooling rate were ◦ performed with and without an austenitic deformation of εAR = 0.5 at 950 C. The CCT diagrams for the steels IF-Ti and IF-TiNb are shown in Fig. 3.1 a) - d) and were constructed according to [104]. A comparison of the CCT diagrams with and without deformation shows, that the Ar3 temperature is relatively independent of the deformation, whereas the Ar1 is raised for both steels due to the deformation. This acceleration of the transformation is usually explained in terms of a higher rate of ferrite formation due to the deformation [98, 99, 101]. In [105] this effect is explained in terms of a reduced undercooling required for nucleationby an increased potency of the nuclei. It can be concluded, that the minimum austenitic rolling temperature is about 900 ◦C and the maximum ferritic rolling temperature is approximately 860 ◦C, in order to avoid rolling in the 40 CHAPTER 3 EXPERIMENTAL

2 1

2 1: Pyrometers

2: Sponges with Lubricant 3 3: Upper Roll

4: Entry Table

4a: Bottom View of 4 4 ´ 5 5: Specimen

Figure 3.2: Laboratory rolling mill of the Max-Planck-Institut für Eisenforschung two-phase region.

3.3 Hot and Warm Rolling

3.3.1 Laboratory Hot Rolling Mill

The hot and warm rolling experiments were conducted on the laboratory hot rolling mill of the Max-Planck-Institut für Eisenforschung, shown in Fig. 3.2. The laboratory rolling mill has a roll-diameter of 300mm and the rolling force is limited to 2MN. The mill is equipped with a hydraulic roll-gap adjustment, which allows for a quick alteration of the roll-gap. A rolling speed of 94 m/min was used for all hot and ferritic rolling tests.

The rolling temperature is monitored with two pyrometers, Fig. 3.2(1), mounted above the entry and exit zone of the mill. The rolling temperature, together with the rolling force, torque and speed are recorded by a data acquisition system. The rolling mill is used in reversing mode, to simulate multi-pass ﬁnish rolling. The minimal reversing time is approximately 4s. 3.3 HOT AND WARM ROLLING 41

3.3.2 Lubrication of the Roll Gap

In section 2.6 it has been shown, that ferritic rolling, without the use of a proper lubrication, only brings about poor deep-drawing properties. Therefore it was chosen to install a lubrication system to the laboratory rolling mill. The lubrication system and the results of the experiments leading to the choice of a proper lubricant are described in the next paragraphs.

3.3.2.1 Lubrication System

To provide a proper lubrication during a series of rolling passes, a continuous lubrication system is needed. The set-up of this lubrication system is shown in Fig. 3.2. Sponges (2), soaked with the desired lubricant are mounted to the mill such that they transfer the lubricant continuously to the rolls. One sponge is mounted above the top roll (3) and two sponges are mounted to the bottom of the entry and exit table (4 & 4a) respectively, to apply the lubricant the bottom roll. After three tests the sponges were removed from the holders, cleaned, immersed in fresh lubricant and reassembled in the holders. This provided a uniform and continuous application of the lubricant to the rolls and so to the roll-gap.

3.3.2.2 Selection of the Lubricant

The effectiveness of lubricants is very sensitive to the rolling parameters, such as the roll gap geometry, the rolling temperature and the nature of the scale, as it has been outlined in section 2.6. Due to the lack of a particular ferritic rolling lubricant, conventional hot rolling lubricants had to be utilised. Four conventional hot rolling lubricants, Table 3.2, were tested at the desired ferritic rolling temperature and a variety of roll gap geometries. On the basis of these preliminary test, the appropriate lubricant for the desired process conditions was chosen.

No. Product* Composition A Kewdol HS 40-A Mixture of mineral oil, grease and additives B Rollub HR 40 Mixture of mineral oil, grease and additives C V 25/99 Mixture of synth. esters, grease and additives D Rollub HR 41 Mixture of mineral oil, grease and additives * Lubricants delivered by HOUGHTON [106]

Table 3.2: Hot rolling lubricants tested in the ferritic rolling temperature region

To test the different lubricants, specimens of the same width of 80mm and initial thickness of 8.5mm were rolled in four passes with a ﬁxed roll gap of 4,2,1 and 0.5mm at a ferritic rolling temperature of 760 ◦C. After each rolling pass the specimens were reheated to approximately 1050 ◦C and allowed to cool down to the rolling temperature. This method allows to compare easily the effectiveness of the different lubricants in the desired process window. A comparison 42 CHAPTER 3 EXPERIMENTAL

700 Rolling Temperature: 760 °C 600

500

400

300 without Lubrication 200 Lubricant A Rolling Force [kN] Lubricant B 100 Lubricant C Lubricant D 0 8.5ð 4 42ð 21ð 1ð 0.5 mm mm mm mm 1. pass 2. pass 3. pass 4. pass

Figure 3.3: Rolling force during lubricated ferritic rolling with the lubricants A – D of the rolling forces, Fig. 3.3, reveals that the lubricant D yielded the lowest rolling force in all four rolling passes. Therefore, this lubricant was chosen for all ferritic rolling test. As it can be seen in Table 3.3, contains lubricant D with approximately 35% the highest amount of the high pressure additives, together with 50% grease. This composition does behave slightly better compared to lubricant C, which consists 95% grease, and only 5% high presure additives. Lubricants A and B contain the highes amount of Minerl oil, wich becomes unstable in the ferritic rolling temperature. Hence, the desirable behaviour of lubricant D can be correlated to the chemical composition of the lubricant.

Constituent Lubricant ABCD Grease 45% 45% 95% 50% Mineral Oil 50% 40% - <35% Sulphur Additive 5% 5% 5% 10% Phosphor Additive - - - 5% Fatty Acids - 10% - <10%

Table 3.3: Additive composition of the lubricants A – D [107]

3.3.3 Rolling Schedules

The rolling schedules employed for the different product groups, already mentioned in section 2.4, are shown in Fig. 3.4 a)-d). A detailed description of the individual rolling schedules will be given in the following paragraphs. 3.3 HOT AND WARM ROLLING 43

a) “Soft” Hot Strip b)”Hard” Hot Strip (annealed)

Taus TAus 123 e at T e at T AR AR AR AR batch annealing ag+ ag+ T direct conv. FRmax e = const. T. FR e 12 FR FRat T FR thold

2h at TC 2h at TC

c) “Cold Strip” (annealed) d) Conventional Hot Strip

TAus TAus

Temperature e e at T AR at TAR AR AR ag+ continuous ag+ continuous annealing annealing e FRat T FR 2h at TC

2h at TC

40 - 80% cold rolling 80% cold rolling

Time

Figure 3.4: Rolling schedules for a) the "soft" hot strip, b) the "hard" hot strip, c) the "cold strip" produced of ferritic rolled hot strip and d) the conventional hot strip produced as reference strip

Independently of the product group three specimens were rolled for each set of parameters.

3.3.3.1 "Soft" Hot Strip

All the specimens for the product group "soft" hot strip, Fig. 3.4 a), were reheated at TAus = ◦ ◦ 1050 C in a muffle type furnace. Subsequently, the specimens were cooled down to TAR = 950 C in still air for an austenitic rolling pass with a total rolling strain of εAR = 0.5. Afterwards, the strip cooled down, through the α/γ-transformation, to the desired ferritic rolling temperature, TFR. To study the influence of the strain distribution in the ferritic temperature region, three different ferritic rolling schedules, 1-3 in Fig. 3.4 a) were designed. The total ferritic rolling strain of εFR = 1.2 was distributed over one, two or three rolling passes, respectively. The final strip thickness was 1 and 2mm for all three rolling schedules.

The distribution of the total strain onto multiple rolling passes requires to raise the temperature of the first rolling pass up to TFRmax, in order to keep the temperature, TFR, of the last rolling pass constant. The finish rolling temperature, TFR, is defined as the temperature just before the last ◦ ◦ rolling pass. The maximum ferritic rolling temperature, TFRmax, was 858 C±10 C, for the case of three rolling passes. 44 CHAPTER 3 EXPERIMENTAL

◦ ◦ The ﬁnish rolling temperature was TFR = 710 C ± 10 C for all three rolling schedules. The ﬁnish rolling temperature was chosen on the basis of the results in [8]. To check whether this optimum temperature was also valid for the present experiments, the coiling temperature, TC, was varied from 610 to 670 ◦C. The coiling process was simulated by an isothermal annealing at the desired coiling temperature for 2h [8]. Finally the strip was pickled and protected against corrosion.

3.3.3.2 "Hard" Hot Strip (annealed)

◦ The "hard" hot strip, Fig. 3.4 b), was also reheated at TAus = 1050 C in a muffle type furnace. ◦ Subsequently the specimens were cooled down to TAR = 950 C in still air for an austenitic rolling pass with a total rolling strain of ε = 0.5. Afterwards the strip cooled down, through the α/γ- transformation, to the desired ferritic rolling temperature, TFR. The strip is rolled in two passes with a total strain of εFR = 1.2. The final strip thickness was 1 or 2mm. The finish rolling temperature for the "hard" hot strip was somewhat lower, when compared to that of the soft hot strip. On the basis of the flow stress data, e.g. Fig. 2.6, the finish rolling temperature was chosen ◦ to be TFR = 660 C, to avoid excessive rolling forces. The strips were "coiled" at an industrially applicable coiling temperature of 550 ◦C (corresponding to the standard length of a run-out table) and at a reduced coiling temperature of 400 ◦C. To simulate the so-called "direct" annealing treatment, the specimens were put directly after the last ferritic rolling pass into a muffle type furnace, which was set to the coiling temperature, cycle 1 in Fig. 3.4 b). After a simulated transportation time of thold = 5, 10 and 20min, required to transfer the coil to the batch annealing furnace, the specimens were subjected to a simulated batch annealing cycle, shown in more detail in section 3.6. For the conventional batch annealing, cycle 2 in Fig. 3.4 b), the strip was reheated — after cooling down to the room temperature — and isothermally annealed for 2h at the desired coiling temperature and finally cooled down to room temperature. The supplemental batch annealing was also carried out in a muffle type furnace, as described in section 3.6. Finally the strips of both product groups were pickled and protected against corrosion.

3.3.3.3 "Cold Strip"

The "cold strip", Fig. 3.4 c) is produced of either ferritic rolled "soft" or "hard" hot strip, coiled at 550 or 400 ◦C, so that the ferritic rolling and coiling conditions were similar to that of the previously in 3.3.3.1 and 3.3.3.2 described schedules. However, the "hard" hot strip was not annealed prior to the cold rolling. The ferritic rolled hot strips were pickled and subsequently cold rolled to a ﬁnal strip thickness of 0.5mm with a total rolling reduction of εCR = 40 to 80%. Finally, the cold rolled strips were subjected to a simulated continuous annealing treatment, described in more detail in section 3.6. 3.4 PICKLING 45

3.4 Pickling

In the industrial production, the hot rolled coils are pickled either discontinuously in push-pull pickling line or continuously in an integrated pickling and cold rolling or pickling and annealing / hot dip galvanizing line. The pickling acid can be either sulphuric or hydrochloric acid with a temperature of approximately 80 ◦C. The main difference between the two acids for the ﬁnal product is the surface appearance of the strip. The sulphuric acid produces a brownish surface, whereas the hydrochloric acid produces a greyish surface. Therefore, the hydrochloric acid is conventionally favoured by the steel producers. During the pickling process the scale layer, together with a maximum of 0.3µm of the surface of the base material is removed. A further removal of the steel surface is called over-pickling and has to be avoided, in order to reduce the material losses during the process and ensure a uniform thickness over the strip length. This can be achieved by a close control of the pickling time, which is deﬁned by the speed of the pickling line and additionally by the addition of pickling inhibitors. The pickling inhibitors protect the base material to a certain extent from the acid. The use of inhibitors becomes more and more popular, because it protects the strip in the pickling line, during a malfunction of the line. For the case of a discontinuous laboratory pickling, with less uniform scale layers, especially because of the different production parameters, the pickling time can be hardly kept constant. Therefore it was decided to use a pickling inhibitor for the laboratory pickling test. To test the effectiveness of the inhibitor for the laboratory test conditions, a couple of pre-test were performed.

3.4.1 Pre-Tests

For the pre-test specimens with a size of 20 x 20mm and a thickness of 1 and 2mm were used. The specimens were cut from the same ferritic rolled sample to obtain a comparable scale layer. The pickling acid was hydrochloric acid (20%) with the addition of 0 or 1200 g/m3 ADACID 328 [108]. The pickling solution was heated to approximately 80 ◦C in a beaker. The specimens were pickled for 10 or 30min. The specimens were weighed before and after pickling to quantify the amount of material removed from the specimens. A pickling time of 10min was in all cases sufﬁcient for a complete removal of the scale. The specimen pickled without the addition of an inhibitor lost about 13.3% of its weight after a pickling time of 10min , whereas the specimen pickled with the addition of 1200 g/m3 inhibitor only lost 3.4%. After pickling for 30min, the specimen pickled without the inhibitor lost about 39.8% and that pickled with the inhibitior only lost 7.8% of its weight. This excessive over- pickling is also visible in Fig. 3.5, showing SEM pictures of the surfaces of different specimens pickled without or with the addition of the inhibitor. The specimens pickled for 10 and 30min without the addition of the inhibitor, Fig. 3.5 a) and b) show deep etching grooves at the grain boundaries and rough grain-surfaces. The specimens pickled with the addtion of the inhibitor, 46 CHAPTER 3 EXPERIMENTAL

a) b)

c) d) 100 µm

Figure 3.5: SEM pictures of the surfaces of pickled specimens; a) and b) pickled for 10 and 30 min in hydrochloric acid, without the addition of an inhibitor and c) and d) pickled for 10 and 30 min in hydrochloric acid, with the addition of (1200 g/m3) inhibitor (All pictures have the same scale) shown in Fig. 3.5 c) and d), exhibit smooth grain-surfaces and no etching grooves after pickling for 10min. After pickling for 30min there are still no etching grooves visible and only a slight roughening of the grain-surfaces occured. The visible surface topography is still below the grain scale, unlike that of the specimens shown in Fig. 3.5 a) and b).

3.4.2 Laboratory Pickling

The strips for a subsequent cold rolling had length of up to 450mm and a width of 80mm, so that the pickling was executed in the chemical laboratory of Corus Special Strip in Düsseldorf, which has the equipment to simulate pickling in a larger laboratory scale. The ferritic rolled strips were pickled in an industrially used pickling solution with the addition of an inhibitor. The pickling solution was heated using a zirconium heating rod to about 80 ◦C and stirred to provide a homogeneous temperature distribution in the pickling bath. After the complete removal of the scale, the specimens were rinsed with clear water and subse- 3.5 COLD ROLLING 47

quently with a mixture of water and a passivating additive Neutracid [108]. Finally the specimens were ﬂushed with desalted water and dried with a heat-gun.

3.5 Cold Rolling

The subsequent cold rolling of the specimens was performed on a laboratory rolling mill of the Max-Planck-Institut für Eisenforschung. This rolling mill has a maximum rolling load of 0.6MN and is equipped with cold rolls with a diameter 150mm. The rolling speed was about 20 m/min. The specimens were rolled in 4 to 8 passes to the ﬁnal strip thickness of 0.5mm, according to the desired cold reduction. The specimens were brushed with a cold rolling lubricant, Houghton Plattieröl [106], prior each rolling pass.

3.6 Annealing

The "hard" hot strip and the "cold strip", both require a supplemental recrystallisation annealing, to provide the desired deep-drawability. This recrystallisation annealing treatment can be carried out discontinuously in an batch annealing furnace, or continuously in a continuous annealing furnace. In a batch annealing furnace a couple of coils are heated slowly to the desired recrystallisation temperature and subsequently cooled. In the continuous annealing process, the strip is uncoiled and pulled through a furnace, where it is rapidly heated to the desired recrystallisation temperature, held isothermally and ﬁnally accelerated cooled.

3.6.1 Batch Annealing Simulation

The batch annealing was simulated in a mufﬂe type furnace. The annealing cycle is shown in ◦ ◦ ◦ Fig. 3.6 a). The specimens are heated with 60 C/h to 600 C and subsequently heated with 8 C/h ◦ ◦ ◦ to 700 C. Afterwards the specimens are cooled with 18 C/h to 200 C and ﬁnally cooled in still air to the room temperature. The overall annealing time adds up to about 50h. The same simulated batch annealing cycle has been used in [8]

3.6.2 Continuous Annealing Simulation

Whereas the batch annealing cycle can be easily simulated in a mufﬂe type furnace, the simulation of the continuous annealing cycle is more complicated. The high heating rates, of up to ◦ 5 C/s, can be simulated using an infrared radiation furnace (IR-furnace), or a salt bath. Galvani- sing simulators, as the RHESCA simulator of the Institut für Eisenhüttenkunde (RWTH Aachen) consist of an IR-furnace, shown in Fig. 3.7 a). This simulator has been employed for a couple 48 CHAPTER 3 EXPERIMENTAL

700°C 810°C, 35s 8 °C/h 600°C

18 °C/h 4.5 °C/s 60 °C/h 200°C Air cool

Temperature 13.5 °C/s 135 °C/s Air cool 1 2 >50h » 230s&80s a)Time b) Time

Figure 3.6: The annealing cycles for a) simulated batch annealing [8] and b) simulated continuous annealing (1: RHESCA simulator; 2: Salt bath for a strip thickness of 0.5mm)

Specimen IR-furnace

Specimens

Salt bath

a) b)

Figure 3.7: a) The IR-furnace of the RHESCA simulator at the Institut für Eisenhüttenforschung (RWTH Aachen) [109] b) Salt bath at the Max-Planck-Institut für Eisenforschung of tests. The IR-furnace has the advantage, that even more complicated multi step annealing cycles can be simulated, as both the heating and cooling rate can be controlled. In a salt bath, shown in Fig. 3.7 b) the heating rate cannot be controlled and is depending on the annealing temperature (salt bath temperature) and the thickness of the specimens. On the other hand, in the salt bath a couple of specimens can be annealed at the same time, which makes this procedure much more productive than the RHESCA simulator, only capable of processing one specimen at a time. Considering the large amount of "cold strip" specimens in this work, the latter aspect was decisive for the choice of the annealing simulation in a salt bath.

However, from the literature (e.g. [110]) it is known that the texture development during the 3.6 ANNEALING 49

annealing is not only depending on the annealing temperature but also to the heating rate. There- fore, it was decided to anneal comparable specimens according to an industrial annealing cycle in the RHESCA simulator and in the salt bath, to compare the resulting textures and deep-drawing properties.

3.6.2.1 RHESCA Simulator

The RHESCA simulator provides a "realistic" laboratory simulation of the annealing and hot dip coating processes. The infrared radiation furnace of the simulator allows for the simulation of continuous annealing cycles, with controlled heating and cooling rates. A more detailed description the RHESCA simulator can be found in [109]. Samples with a size of 24mm x 120mm (width x length) were cut along RD from a cold rolled "hard" hot strip with a thickness of 0.5mm. Two thermocouples were welded to the specimens, one to control the annealing cycle and the other as a reference. The annealing cycle was chosen in accordance with industrially applied cycles and is schematically shown in Fig. 3.6 b) (cycle 1). The specimens were hung into the IR-furnace, as it is shown in Fig. 3.7 a) and heated with ◦ ◦ approximately 4.8 C/s to 810 C and held isothermally for 35s at this temperature. Subsequently ◦ the strip was accelerated cooled with about 135 C/s to room temperature. From the annealed specimens samples for the determination of the texture and tensile test specimens were machined. The mechanical properties measured in the tensile tests are given in Table 3.4. The specimens exhibit a desirably high r-value of r0 = 2.65 and the mechanical properties fulﬁl with Rp0.2 = 93MPa and Rm = 274MPa the requirements for steel grade DC 06.

Rp0.2 [MPa] Rm [MPa] Ag [%] A50 [%] r0 [-] n0 [-] Mean 93 274 29 45 2.65 0.286 StdDev. 6 7 1 3 0.17 0.007

Table 3.4: Mechanical properties of the cold rolled "hard" hot strip annealed at 810 ◦C for 35 s in the RHESCA simulator (Mean = mean values, StdDev. = standard deviations)

The texture of the measured samples exhibited a well developed γ-ﬁbre texture with a maximum intensity of f (g) = 11 near {554}h225i. An representative example of the measured texture is shown in Fig. 3.8 a). This desirable γ-ﬁbre texture correlates very well with the desirable deep-drawing properties, measured in the tensile tests. On the basis of the texture measurements the mean r-value and the planar anisotropy were calculated to be rm = 2.07 ± 0.10 and ∆r = 0.71 ± 0.21. The r-value distribution showed the classical v-shape.

3.6.2.2 Salt Bath

In contrast to the RHESCA simulator, the heating and cooling rates in the salt bath are depending on the thickness of the specimens and the salt bath temperature. The heating rates for the 0.5, 1 50 CHAPTER 3 EXPERIMENTAL

j1 j1 0 1530456075900 153045607590 0

15 11 10 30

75 1 1 90 a) b)

Figure 3.8: Texture of the samples annealed a) in the RHESCA simulator and b) in the salt bath measured at the central plane (distance from the centre s = 0) (intensity levels 1, 2, 3, ...) and 2mm thick specimens were determined(24mm x 120mm) by inserting a thermocouple into the specimens and subsequently immersing them into the salt bath. The salt bath was set to the desired annealing temperature of 810 ◦C. The resulting heating curves are depicted in Fig. 3.9. It was ﬁxed, that the annealing temperature was reached after 10, 20 and 40s for the 0.5, 1 and 2mm thick specimens, respectively. After the specimens were held isothermally for 35s, followed by air cooling to the room temperature. The annealing cycle for a strip thickness of 0.5mm is shown schematically in Fig. 3.6 b) (cycle 2). From the annealed specimens samples for the determination of the texture and tensile test specimens were machined, in order to compare the reslts with those obtained when utilising the RHESCA simulator. The mechanical properties of the strips annealed in the salt bath furnace are listed in Table 3.5.

Rp0.2 [MPa] Rm [MPa] Ag [%] A50 [%] r0 [-] n0 [-] Mean 90 287 28 41 2.27 0.271 StdDev. 4 12 1 2 0.17 0.006

Table 3.5: Mechanical properties of the cold rolled "hard" hot strip annealed at 810 ◦C for 35 s in the salt bath (Mean = mean values, StdDev. = standard deviations)

The specimens exhibit a desirably high r-value of r0 = 2.27 and the mechanical properties fulﬁl with Rp0.2 = 90MPa and Rm = 287MPa the requirements for steel grade DC 06. The texture exhibited a well developed γ-ﬁbre with a maximum intensity of f (g) = 10 near {554}h225i. An representative example of the measured texture is shown in Fig. 3.8 b). On the basis of the texture measurements the mean r-value and the planar anisotropy were calculated to be rm = 1,95±0.14 and ∆r = 0.7 ± 0.28. The r-value distribution showed the classical v-shape. 3.6 ANNEALING 51

1000

800

600

400 Strip thickness: 0.5 mm Temperature [°C] 200 1mm 2mm 0 0102030 40 50 Time [s]

Figure 3.9: Heating rates of specimens with a thickness of 0.5, 1 and 2mm in a salt bath with a nominal temperature of 810 ◦C

3.6.2.3 Comparison of the continuous annealing pre-tests

The mechanical properties of the pre-test samples annealed in the RHESCA simulator and those annealed in the salt bath exhibit similar strength properties with approximately Rp0.2 = 90MPa and Rm = 280MPa. The mechanical properties fulfil the requirements of steel grade DC 06. The deep-drawing properties also meet the requirements of the DC 06 grade, but a slight difference between the specimens annealed in the simulator and the salt bath is visible. The specimens annealed in the simulator exhibit with r0 = 2.65 an approximately 14% higher r-value in the longitudinal direction than the specimens annealed in the salt bath (r0 = 2.27). The texture measured in the centre layer (s = 0) of the specimens, shown in Fig. 3.8 a) and b), exhibit a similar shape of the γ-fibre, with a maximum intensity near {554}h225i. Only the maximum intensity is slightly different. The difference in the texture intensity can be explained in terms of the different heating rates in the RHESCA simulator and the salt bath. The deviation of about 14% observed for the two different heting rates correlates to the deviation measured in [110]. In this work a deviation of 13% was found for heating rates, comparable to those used in the present experiments. If one assumes that γ-fibre grains nucleate first, this effect could be attributed to the time available for the growth of a sufficient amount of γ-fibre nuclei during heating, before signifficant nucleation and growth can occure in the other orientations [39]. The r-value distributions calculated from the texture, exhibited for both simulations the classical v- shape. Even though the deep-drawing properties measured in the tensile test and the texture intensities were slightly worse for the salt bath annealing, the correspondence was still fairly good. Therefore, it was decided to use the salt bath for the simulation of the continous anealing of all cold rolled specimens, considering also the effect, that using the RHESCA simulator , the final results may bee even better. 52 CHAPTER 3 EXPERIMENTAL

3.7 Mechanical and Technological Testing

In order to appraise the mechanical properties, obtained by the ferritic rolling strategies, it was necessary to subject the strips to extensive mechanical testing. The mechanical and deep-drawing properties were primarily determined by ﬂat tensile test. Additionally the standard cupping test was employed, in order to get a better measure of the earing behaviour of the different cold strips.

3.7.1 Tensile Tests

Tensile specimens were cut from the strips, with 0, 45, and 90 ◦ to the rolling direction. To avoid excessive hardening or softening due to the mechanical machining of the specimens, the specimens were cut by spark erosion in a water bath. The dimensions of the tensile specimens are given in Fig. 3.10.

min. 80 50

20 R6

Figure 3.10: Dimensions of the tensile test specimens, spark eroded from the strips

The shortest length of the specimens was deﬁned by the strip width, so that the specimens taken transverse to the rolling direction were about 80mm long. The specimens for the other directions were about 100mm long. The thickness of the specimens was of course equivalent to the strip thickness. For each direction 2 specimens were machined from each strip. The tensile tests were executed on a Zwick 1445 tensile test machine, equipped with a device to measure the strain in length and width direction, in order to calculate the r-value. The test specimens and test equipment were chosen according to DIN EN 10002-1 [111].

3.7.2 Cupping Tests

The cupping test were performed on an Erichsen 142/20 sheet metal testing machine of the Institut für Umformtechnik at the University Dortmund. The geometry of the cupping test tool is shown in Fig. 3.11. The specimens were cleaned with alcohol before the test and subsequently drawn without lubrication. 3.7 MECHANICAL AND TECHNOLOGICAL TESTING 53

c = 1.1t

rd = 3.5 Die Sheet metal blank t

r = 4.5 p Blankholder

v = 22 mm/min

D=33mmp

D=60mmb

Figure 3.11: Geometry of the cupping test tool (c: Clearance; Db: Blank diameter; Dp: Punch diameter; rp: Punch radius; rd: Die radius)

t/2+0.1 Mech. grinding

Mech. + Chem. t finishing ND RD

t/2

Figure 3.12: Preparation of the specimens for the macro texture measurements (t: strip thickness)

3.7.3 Texture Measurements

The macro texture measurements presented in this thesis were measured with the Siemens D500 texture goniometer of the Max-Planck-Institut für Eisenforschung. The ODFs were calculated by a series expansion method, proposed by Bunge [31], from the three incomplete pole figures {110}, {200} and {211}. The specimens were cut with a water cooled cutting machine with a size of 14 x 24mm (TD x RD). As the texture was measured in all cases in the centre layer s = 0 (Equation 2.11) the specimens were mechanically ground and finally mechanically and chemically finished, as indicated in Fig. 3.12. The finishing was executed mechanically with water cooled sand paper (grain size 1000) and chemically with a solution of H2O2 and HF(10:1). 54 CHAPTER 3 EXPERIMENTAL

3.7.4 EBSD Measurements

The EBSD pictures were taken within the ND/RD plane. The deformed specimens were measured on the Jeol ﬁeld emission scanning electron microscope (FE-SEM) with a lateral resolution of 0.1µm and the recrystallised ones were measured with a lateral resolution of 2µm. Chapter 4

Hot Strip Grades – Results and Discussion

4.1 Minimum Coiling Temperature for the "Soft" Hot Strip

As it already has been mentioned in section 2.4 a sufficiently high coiling temperature is the crucial parameter for the production of a "soft" hot strip. The coiling temperature has to be high enough in order to obtain full recrystallisation in the coil. Therefore, the lowest possible ferritic rolling temperature is also determined by the necessary coiling temperature. A series of specimens was rolled according to rolling schedule 2 in Fig. 3.4 a). The ferritic rolling temperature was determined to be 710 ◦C. From the literature it is known that higher rolling temperatures lead to less favourable rolling and recrystallisation textures. Rolling temperatures above 760 ◦C lead to undesirable coiling textures, mainly consisting of the {001}h110i(rotated cube) component [8, 26, 53]. The influence of the rolling temperature on the final texture is described in more detail in section 2.4.1. The coiling temperature was varied between 610 and 670 ◦C to determine the minimum acceptable coiling temperature. The coiling of the strip was simulated, as mentioned in section 3.3.3.1, by isothermal annealing for two hours at the desired coiling temperature. The fraction recrystallised was determined by optical microscopy.

Fig. 4.1 a) shows the recrystallised fraction, Xrex, for the different simulated coiling temperatures. It can be concluded that the coiling temperature of 670 ◦C is just sufﬁcient to obtain full recrystallisation in the IF-Ti steel, within the available time. In contrast, the same coiling temperature lead only to approximately 5% recrystallised fraction in the IF-TiNb steel. The minimum coiling temperature of 670 ◦C determined for the steel IF-Ti is in good agreement with the temperature reported in [8]. The steel used in the latter work was also Ti microalloyed and had an overall chemical composition similar to that of the steel IF-Ti, Table 4.1. For the complete recrystallisation of the IF-TiNb a considerably higher higher coiling temperature is necessary, as it can be deduced by extrapolation from Fig. 4.1 a). This behaviour can be explained in terms of the additional microalloying with Nb. It is known from [5, 112, 113] that Nb and Ti

55 56 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

1.0 4 IF-Ti 1mm

IF-Ti 2mm 3 Qrex = 710 kJ/mol 0.8 IF-TiNb 2mm 2 1 0.6 ))), - rex 0

rex -1 X0.4 [-] -2

0.2 ln(ln(1/(1-X -3 -4 0.0 -5 600 620 640 660 680 1/943 1/923 1/903 1/883 a)coiling temperature [°C] b) 1/T, K-1

Figure 4.1: a) Fraction recrystallised for a series of coiling temperatures and the two IF steels (2h isothermal annealing at the desired coiling temperature); b) ln(ln(1/(1 − Xrex))) vs. 1/T plot for the steel IF-Ti and especially their precipitates retard the recrystallisation. The retarding effect of Nb is however much greater than that of Ti. In [5] it has been found, that the time for 90% recrystallistaion is three orders of magnitude larger for Nb alloyed HSLA steel compared to Ti alloyed HSLA steel.

Steel C Si Mn P S N Al Ti Nb IF-Ti 0.001 0.017 0.130 0.005 0.008 0.004 0.032 0.064 0.002 IF [8] 0.002 0.007 0.097 0.010 0.004 0.003 0.042 0.038 0.007

Table 4.1: Chemical compositions in mass% of the IF-Ti steel and the IF steel used in [8]

These results show that the optimum coiling temperature to produce a "soft" hot strip cannot be generalised because it is very sensitive to the chemical composition of the steel used. This means that the optimisation has to be performed for individual alloying concepts.

From the data shown in Fig. 4.1 a) a mean activation energy for recrystallisation, Qrex, can be calculated by adapting a suited model for recrystallised fraction, Xrex. An often used and quoted model for this is the so-called JMAK model with an equation of the form

−βtk Xrex = 1 − e , (4.1)

−Q /RT with the thermally activated rate constant β = β0e rex [114]. The exponent k accounts for the nucleation and growth mechanism and the dimensionality [115–117]. In the case of the present results, the annealing treatment was isothermal and the annealing time, t, was constant (t = 3h). Hence equation 4.1 can be rewritten as follows

−Ae−Qrex/RT Xrex = 1 − e , (4.2) 4.2 INFLUENCE OF THE FERRITIC ROLLING REDUCTION 57

k with the constant A = β0t . By reorganising equation 4.2 the following proportionality can be derived

1 Q ln ln ∝ − rex . (4.3) 1 − Xrex RT

The slope of the curve shown in Fig. 4.1 b) is hence proportional to the activation energy. An −1 activation energy of Qrex = 710kJmol has been calculated for the IF-Ti steel. The calculated activation energy for the IF-Ti steel is much higher than the values found in literature. In [53, −1 114] values for Qrex in the range of 300 − 400kJmol are reported for IF steels. A possible explanation for the high activation energy might be an increased amount of ﬁne and disperse precipitates, retarding the recrystallisation. This might be caused by pronounced strain induced precipitation of TiC during the ferritic rolling passes, as still approximately 5 − 40% of the TiC ◦ are in solution in the temperature range of 750 and 850 C [8, 20, 23]. However, the high Qrex value indicates that the process takes place within a narrow temperature domain.

4.2 Inﬂuence of the Ferritic Rolling Reduction

To investigate the influence of the ferritic rolling reduction on the texture development and hence the mechanical properties, samples were rolled according to Fig. 3.4 a) in two ferritic rolling passes with a total strain of εFR = 0.6 and 1.2. The other rolling parameter were chosen analogous to section 3.3.3.1, in order to produce a "soft" hot strip. The optimised coiling temperature was determined, as described in the previous section. Fig. 4.2 a) and b) shows the rolling texture of the specimens rolled with a ferritic rolling strain of 0.6 and 1.2, respectively. The strips were water quenched after the last rolling pass to fix the ferritic rolling texture. Both textures exhibit a typical rolling texture, with a complete γ-fibre and a partial α-fibre. The α-fibre exhibits a maximum near {112}h110i. The rolling texture of the strip is weaker for a lower rolling strain, Fig. 4.2 a), than that with a higher one, Fig. 4.2 b). Nevertheless both rolling textures are qualitatively similar. The texture development during coiling of the strips is totally different. During the coiling simulation of the strip rolled with εFR = 0.6 the initial rolling texture is slightly weakened, but the nature of the texture did not change, as can be seen in Fig. 4.2 c). The texture still consist of a complete γ-fibre and a partial α-fibre with a maximum near {112}h110i. However, during the coiling simulation of the strip rolled with εFR = 1.2 a desirable {111} texture developed with a maximum intensity of 10 near {554}h225i.

The microstructure of the coiled sample with the lower rolling strain, εFR, is depicted in Fig. 4.3 a). This specimen exhibits a rather large grain size of approximately 80µm and the grain boundaries show a slight bulging. The grains contain a clearly visible substructure, however all strained grains had disappeared. The sample rolled with the higher rolling strain revealed smaller 58 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

◦ Figure 4.2: Rolling texture of the specimens rolled at TFR = 710 C with a ferritic rolling strain εFR of a) 0.6 and b) 1.2; Coiling texture of the specimens rolled with a ferritic rolling strain of c) 0.6 and d) 1.2 (Coiling ◦ temperature TC = 670 C)

100µm 100µm a) b)

Figure 4.3: Microstructure after simulated coiling of the ferritic rolled samples with a) εFR = 0.6 and b) εFR = 1.2 ◦ ◦ (TFR = 710 C, TC = 670 C) 4.3 MECHANICAL PROPERTIES 59

grains with a lower aspect ratio. The grains have a size of approximately 40µm transverse to RD and 80µm parallel to RD.

The different coiling texture and microstructure for the two different rolling strains, can be explained in terms of the major softening mechanism active during the coiling simulations. The fact that no qualitative texture change could be observed for the strip rolled with a strain of 0.6, Fig. 4.2 a) and c), suggests that the specimen mainly recovered during the coiling simulation. During the coiling simulation of the strip rolled with a strain of 1.2 the initial rolling texture is transformed into a desirable γ-ﬁbre texture, as can be seen in Fig. 4.2 b) and d). In this sample recrystallisation occurred. This different behaviour can be related to the stored energy necessary for a complete recrystallisation. Similar results have been presented in [8].

It can be concluded, that a minimum rolling strain is necessary to produce a desirable texture. A ferritic rolling strain of εFR = 1.2 has been found to be sufﬁcient to lead to a "soft" hot strip with desirable {111} texture, assuming a sufﬁciently high coiling temperature.

4.3 Mechanical Properties

The mechanical properties of the strips were determined by tensile tests in 0◦, 45◦ and 90◦ to the rolling direction. For each set of parameters, three strips were rolled. Subsequently two tensile specimens were cut for each direction. Hence, each data point presented in the following paragraphs is the mean value of six individual measurements. The r-values, ∆r values and mean r-values were calculated according to the equations 2.1 to 2.3.

4.3.1 "Soft" Hot Strip

The "soft" hot strips were rolled according to the rolling schedule shown in Fig. 3.4 a) and conditions speciﬁed in section 3.3.3.1.

The mechanical properties after coiling at 670 ◦C are shown in Fig. 4.4 a) and b). The strips exhibit a 0.2% proof strength of approximately Rp0.2 ≈ 90MPa and a tensile strength of about Rm ≈ 260MPa, independently of the number of rolling passes. An ultimate elongation of A50 ≈ 40% has been measured nearly independent of number of rolling passes. The mechanical properties, summarised in Table 4.2, fulﬁl the requirements of steel grade DC 04, Table 4.3.

The r-values measured in the three principal directions of the strip are shown in Fig. 4.5 a) and b) for the 1 and 2mm thick strips. The r-values r0◦ ≈ 1.8, r45◦ ≈ 1 and r90◦ ≈ 1.6 are at a desirably high level. Moreover these values are relatively insensitive to the design of the laboratory rolling schedule and the ﬁnal hot strip thickness. 60 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

Figure 4.4: a) 0.2% proof strength Rp0.2 and tensile strength Rm and b) ultimate elongation A50 of the 1mm & 2mm ◦ ◦ thick "soft" hot strips (εFR = 1.2 distributed over 1, 2 and 3 passes, TFR = 710 C, TC = 670 C)

◦ ◦ ◦ Figure 4.5: Measured r-values in 0 , 45 and 90 to RD for the a) 1mm and b) 2mm thick "soft" hot strips (εFR = 1.2 ◦ ◦ distributed over 1, 2 and 3 passes, TFR = 710 C, TC = 670 C) 4.3 MECHANICAL PROPERTIES 61

1 mm 2mm rolled in 2 passes rolled in 2 passes

Rp0.2 [MPa] 87 ± 13.9 88 ± 8.3 Rm [MPa] 264 ± 2.9 266 ± 4.4 A50 [%] 44 ± 4.6 48 ± 3.5 rm [-] 1.4 ± 0.26 1.42 ± 0.17 ∆r [-] 0.75 ± 0.34 0.95 ± 0.26 nm [-] 0.28 ± 0.006 0.285 ± 0.007

Table 4.2: Mechanical porperties of the "soft" hot strip

Steel Rp0.2 [MPa] Rm [MPa] A80 [%] r90 n90 Grade max. min. min. min. DC 03 280 270 - 370 34 1.3 - DC 04 240 270 - 350 38 1.6 0.180 DC 05 210 270 - 330 40 1.9 0.200 rm nm DC 06 180 270 - 350 38 1.9 0.220

Table 4.3: Mechanical properties as they are deﬁned in [21]

4.3.2 "Hard" Hot Strip

The "hard" hot strips were rolled at relatively lower ferritic rolling temperatures, in order to be- neﬁt from the higher strain accumulation at lower temperatures. A conventionally used coiling temperature of 550 ◦C as well as a reduced one of 400 ◦C were applied. The "hard" hot strips exhibit a strained microstructure after coiling and, therefore, they require a supplemental annealing treatment to obtain the desired deep-drawing properties. For the annealing treatment a simulated batch annealing process was chosen. Additionally, the so-called direct annealing practice, see sections 2.4.1 and 3.3.3.2, was simulated.

4.3.2.1 Conventional Batch Annealing

For both coiling temperatures and hot strip thicknesses the conventionally batch annealed samples exhibit a 0.2% proof strength of Rp0.2 ≈ 73MPa and a tensile strength of Rm ≈ 270MPa, as can be seen in Fig. 4.6 a). The deep-drawing properties exhibit less scatter than those of the "soft" hot strip. The "hard" hot strip fulﬁls with r90◦ ≈ 1.8, Fig. 4.6 b), the requirements for steel grade DC 05. The mechanical properties are summarised in Table 4.4. 62 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

1 mm 1mm 2 mm 2 mm ◦ ◦ ◦ ◦ TC = 550 C TC = 400 C TC = 550 C TC = 400 C Rp0.2 [MPa] 77 ± 4.3 75 ± 3.4 70 ± 2.3 72 ± 3.7 Rm [MPa] 275 ± 1.9 272 ± 4.1 268 ± 2.0 268 ± 2.0 Ag [%] 27 ± 3.7 26 ± 3.0 28 ± 1.7 28 ± 1.8 A50 [%] 42 ± 2.9 38 ± 3.3 45 ± 2.1 46 ± 3.2 rm [-] 1.59 ± 0.08 1.5 ± 0.08 1.44 ± 0.08 1.47 ± 0.10 ∆r [-] 0.83 ± 0.13 0.84 ± 0.13 0.89 ± 0.15 1.03 ± 0.19 nm [-] 0.273 ± 0.002 0.27 ± 0.007 0.275 ± 0.002 0.272 ± 0.008

Table 4.4: Mechanical properties of the "hard" hot strips after supplemental batch annealing

300 3.0 1mm 2mm TC = 550°C 250 2.5 1mm 2mm TC = 400°C 200 2.0

150 1.5

r-value [-]

m p0.2

R , R100 [Mpa] 1.0

50 0.5

Rm Rp0. 2 Rm Rp0. 2 Rm Rp0. 2 0 0.0 a)0 45 90 b) 0 45 90 Angle to RD [°] Angle to RD [°]

Figure 4.6: a) Mechanical properties and b) r-values of the "hard" hot strips after batch annealing for both coiling temperatures, TC, and strip thicknesses

4.3.2.2 "Direct" Batch Annealing

In the direct annealing practice the (still) hot coil is transferred directly to the batch annealing furnace. This process was simulated according to rolling schedule and annealing cycle 1 shown in Fig. 3.4 b). In order to simulate different transportation times between coiling of the strip and the begin of the annealing treatment the strips were held at the coiling temperature for thold = 5, 10 and 20min. From Table 4.5 it can be seen, that the holding time before the start of the batch annealing simulation has no clear influence on the deep-drawing properties. The same can be asserted to the mechanical properties. The slight difference between the values are within the scatter band. It has to be mentioned, that in the case of the directly annealed hot strips the mechanical properties presented in the figures are the mean value of only two individual measurements. Therefore, the statistics are less reliable for this data. The mechanical properties for the different coiling temperatures and strip thicknesses are summarised in Table 4.6 for the transportation time of 10min. The mechanical properties are similar to those of the conventionally annealed "hard" hot strip and hence fulfil the requirements of steel grade DC 05. 4.4 TEXTURE DEVELOPMENT 63

5min 10min 20min rm [-] 1.5 1.57 1.37 ∆r [-] 0.88 1.07 0.69

Table 4.5: Inﬂuence of the transportation time, thold, on the deep-drawing properties of 1mm thick strips (Coiling ◦ temperature TC = 550 C)

1 mm 1mm 2 mm 2 mm ◦ ◦ ◦ ◦ TC = 550 C TC = 400 C TC = 550 C TC = 400 C Rp0.2 [MPa] 74 74 70 68 Rm [MPa] 273 269 268 264 Ag [%] 28 28 27 28 A50 [%] 42 45 43 44 rm [-] 1.57 1.43 1.61 1.22 ∆r [-] 1.07 0.71 0.87 0.47 nm [-] 0.276 0.275 0.274 0.278

Table 4.6: Mechanichal properties of "hard" hot strips (direct annealing), with the holding time thold = 10min

4.4 Texture Development

The texture of the specimens was measured in all cases at the centre layer, s = 0. The texture measurements at different layers would have been beyond the scope of the work and, therefore, they were not executed. Nevertheless, the deep-drawing properties were measured in the tensile tests, which give an integral view of the texture, including surface effects. On the basis of a desirable {111} texture at the centre layer, together with good r-values measured in the tensile tests, one can assume an acceptably low through-thickness texture gradient.

4.4.1 "Soft" Hot Strip

The rolling texture and coiling texture of the "soft" hot strips rolled in 1, 2 or 3 passes are shown in Fig. 4.7 a) - f) and Fig. 4.8 a) - f). The rolling textures of the 1 and 2mm thick strips reveal, independent of the number of rolling passes, a well developed rolling texture with a complete γ-fibre and a partial α-fibre with a maximum near {112}h110i. The maximum intensity f (g) ≈ 6 − 7 has been measured in all specimens. The coiling texture of the 1mm thick strips reveals a desirable deep-drawing texture with a maximum near {554}h225i. The maximum intensity is approximately 6 to 7 for the specimens rolled in 1 and 2 passes. The texture of the specimen rolled in 3 passes, Fig. 4.7 f) exhibits a somewhat lower intensity of about 5 to 6 and the maximum is less sharp defined. The coiling texture of the 2mm thick specimens also exhibits a desirable γ-fibre texture with a maximum near {554}h225i. The maximum intensity is 6 for all three rolling schedules. 64 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

Figure 4.7: Rolling and coiling texture of the 1mm thick "soft" hot strips rolled in a) + d) 1 pass, b) + e) 2 passes and c) + f) 3 passes (intensity levels 1,2,3,...)

Figure 4.8: Rolling and coiling texture of the 2mm thick "soft" hot strips rolled in a) + d) 1 pass, b) + e) 2 passes and c) + f) 3 passes (intensity levels 1,2,3,...) 4.5 DISCUSSION 65

j1 0 153045607590 0

15 1 1 30

75 9 10 10 11 90 a) b)

Figure 4.9: Annealing texture of the "hard" hot strips coiled at a) 550 and b) 400 ◦C (intensity levels 1,2,3,...)

4.4.2 "Hard" Hot Strip

4.4.2.1 Conventional Batch Annealing

The rolling texture of the hard hot strips is similar to that of the of the "soft" hot strips shown in Fig. 4.7 a) - c). The annealing textures for the two different coiling temperatures of 550 and 400 ◦C are shown in Fig. 4.9 a) and b). Both annealing textures reveal an intense {111} texture with a maximum near {554}h225i of f (g) ≈ 10. There is no inﬂuence of the coiling temperature on the texture results.

4.4.2.2 "Direct" Batch Annealing

The rolling and coiling textures of the directly annealed strips are qualitatively and quantitatively similar to that of the conventionally annealed strip.

4.5 Discussion

4.5.1 "Soft" Hot Strip

The "soft" hot strips fulﬁl the requirements of steel grade DC 04, Fig. 4.5, without the necessity of a supplemental cold rolling and annealing and their deep-drawing properties are even superior to those of conventional hot strips. This improvement is caused by the development of the desirable γ-ﬁbre texture of these strips, shown in Fig. 4.7 and Fig. 4.8. Although desirable deep-drawing properties have been reached, it should be pointed out that a close control of the ferritic rolling process is necessary to produce a deep-drawable "soft" hot 66 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

strip. Besides the chemical composition of the steel (section 2.5) the following criteria are crucial to produce a "soft" hot strip with desirable deep-drawing properties:

• A complete recrystallisation recrystallisation of the steel at a sufﬁciently coiling temperature,

• a sufﬁciently low ﬁnish rolling temperature,

• a sufﬁciently high rolling reduction and

• adequate suppression of surface shear by employing roll gap lubrication

These criteria, except for the roll-gap lubrication, are naturally interconnected and all of them have to be fulfilled in order to produce a "soft" hot strip with desirable deep-drawing properties. The above list demonstrates the increasing requirements to the process and the production line. The proper recrystallisation kinetics for the production of a "soft" hot strip with desirable deep- drawing properties are achievable even at a relatively low recrystallisation temperature (provi- ding a sufficient rolling reduction), so that a low coiling temperature can be used, still ensuring complete recrystallisation (section 4.1). In the present work, the activation energy has been calculated to be 710kJmol−1, indicating that the recrystallisation takes place in a narrow temperature interval. Hence, a lower activation energy of 300 −400kJmol−1, as it is conventionally reported in the literature, would be more advantageous for the process control because the transition from not recrystallised to complete recrystallised would be less sharp, which would markedly increase the process stability. ◦ Crucial for the deep-drawability of the strips is the finish rolling temperature of TFRmin = 710 C of the last rolling pass. Similar desirable temperatures have been reported in [8, 26, 53]. While this temperature is rather strictly prescribed, there is still a lot of freedom in the choice of the temperatures, TFRmax, of the preceding rolling passes. In the most unfavourable case of the present rolling experiments (schedule 3 in Fig. 3.4 a); final hot strip thickness 1mm) was the ◦ maximum ferritic rolling temperature TFRmax = 858 C in the first rolling pass and TFRmin = 710 ◦C in the third pass. The γ-fibre texture and deep-drawing properties, Fig. 4.5 a), Fig. 4.7 c) and f), remained at a desirably high value. In [59, 60, 65, 70] it has been suggested, that the low finish rolling temperatures are required to ensure a complete precipitation of carbon prior the ferritic rolling passes, as the solute carbon hinders the formation of strain inhomogeneities during rolling and hence the development of a desirable {111} texture during annealing. The detrimental influence of the solute carbon on the recrystallisation texture has already been explained in more detail in section 2.5.2. As shown in section 4.2 a sufficiently high accumulated strain is required to obtain a desirable rolling and recrystallisation texture. It can be concluded, that the accumulated strain even in the case of three rolling passes was, despite the higher rolling temperatures during the first rolling passes and the inter-pass time of 4sec, high enough to produce a desirable coiling texture. This 4.5 DISCUSSION 67

100

X, % 40

0 600 650 700750 800 850 900 Temperature, °C

Figure 4.10: Softening, X, between two deformation steps with ε = 0.3 each and an inter-pass time of 3sec [8]

can be correlated to the softening behaviour of the IF steel, Table 4.1, determined in [8]. The softening behaviour of this steel is depicted in Fig. 4.10. From this figure it is apparent, that up to a temperature of about 760 ◦C, less than 20% softening occurres, so that most of the ferritic rolling strain can be retained. In addition to this, the deformation induced precipitation of TiC is likely to occure during the first rolling pass, so that the desirable rolling texture and an increased amount of strain inhomogeneities can already form at higher temperatures. Once the restrictions in terms of the finish rolling temperature and coiling temperature are managed, the texture development and the mechanical properties are relatively insensitive to the temperature interval of the ferritic rolling passes.

On the basis of the laboratory roll gap lubrication system and the laboratory rolling trails, any recommendation for the design of a lubrication system in an industrial scale is hardly possible. Nevertheless, the present results indicate, that as the lubrication of the roll gap has been sufficient, also the design of the rolling schedule does not seem to be so critical. This is indicated by the stability of the texture development, shown in Fig. 4.7 and Fig. 4.8, together with the uniform and beneficial mechanical properties for the three different rolling schedules. This combination of a desirable γ-fibre texture in the centre of the specimen and the good mechanical properties suggests a weak through-thickness texture gradient.

In section 2.6 it has been shown, that the roll gap geometry, characterised by the ld/t0 ratio, has a major effect on the homogeneity of the through-thickness texture gradient. Higher values of

the ld/t0 ratio lead, without lubrication, to a larger texture gradient than lower ones. The design

of the laboratory rolling schedule provided a wide range of ld/t0 ratios, summarised in Table 4.7. Since desirably high r-values were measured independently of the different rolling schedules

and, hence, ld/t0 ratio, it can be concluded that the lubrication was in all cases sufﬁcient, in order

to suppress the shear deformation near the surface. The values of the ld/t0 ratio are comparable to that of a conventional 6- or 7-stand hot rolling mill. 68 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

ld/t0 for pass No. 1 2 3 1 pass – 2mm 3.97 1 pass – 1mm 5.62 2 passes – 2mm 3.19 4.31 2 passes – 1mm 4.51 6.09 3 passes – 2mm 2.73 3.33 4.07 3 passes – 1mm 3.86 4.71 5.76

Table 4.7: Nominal ld/t0 ratios for the rolling schedules and strip thicknesses of the "soft" hot strip

4.5.2 "Hard" Hot Strip

The r-values of the "hard" hot strip are slightly better than those of the "soft" hot strip and fulﬁl with a high value of r90◦ ≈ 1.8 and the mechanical properties summarised in Table 4.4 the requirements of steel grade DC 05 (Table 4.3). Also the rolling and coiling textures are improved. The annealing texture measured in the "hard" hot strip, Fig. 4.9 is more pronounced than the coiling textures of the "soft" hot strip, shown in Fig. 4.7 and Fig. 4.8. The slightly better mechanical properties of the "hard" hot strip are presumably due to these texture differences. This is a result of a larger accumulated strain, which arise due to lower rolling temperatures. Even though the effort to produce a "hard" hot strip is lower than that of a "soft" hot strip, a couple of conditions have to be fulﬁlled:

• Suppression of surface shear by roll gap lubrication,

• a low ﬁnish rolling temperature,

• and a sufﬁciently high rolling reduction.

In the case of the "hard" hot strip a proper lubrication of the roll gap is the most difficult to obtain in a conventional hot rolling mill. The desirably high r-values, together with the homogeneous texture inside of the specimens suggest a low through-thickness texture gradient. The lubrication of the roll gap applied for all rolling schedules and ld/t0 values, Table 4.7, proved to be satisfactory. The lower rolling temperature might additionally improve the impact of the lubricant. In [51, 52] it has been shown, that a decreasing finish rolling temperature improves both, the rolling and annealing texture. In the case of the "hard" hot strip the finish rolling temperature is not limited by the coiling temperature so that it can be further reduced, in order to benefit from a higher accumulated strain. The finish rolling temperature is however limited by the maximum allowable rolling force of the hot strip mill. In [8] a marked improvement of the deep-drawing properties has been observed by reducing the coiling temperature from 550 to 400 ◦C. In the present work the r-values shown in Fig. 4.6 b) 4.6 INDUSTRIAL IMPLICATIONS 69

and the texture measurements shown in Fig. 4.9 a) and b) do not conﬁrm the reported positive effect of the reduced coiling temperature. There is probably no distinct difference in the amount of recovery occurring within the two hours of the coiling simulation at the two temperatures. The direct annealing practice, suggested in [8, 16, 56], did not bring about an improvement of the deep-drawing properties, as it can be deduced from a comparison of the properties listed in Table 4.6 to those in Table 4.4. In the literature the positive effect of the direct annealing has been explained in terms of the reduced holding time at medium temperatures favouring recovery processes [55]. The absence of this positive effect in the present work correlates to the results of the conventionally annealed "hard" hot strip, which were found to be independent of the two coiling temperatures, indicating a weak impact of recovery processes in this temperature range.

4.6 Industrial Implications

4.6.1 "Soft" Hot Strip

The "soft" hot strip represents a comparatively cheap product, due to the fact that it exhibits (directly after coiling, without an additional annealing) acceptable high r-values and favourable mechanical properties, so that it meets the requirements of steel grade DC 04, listed in Table 4.3. However, the process control for this grade would be rather difficult and demanding on industrial installations. A prerequisite to produce a deep-drawable "soft" hot strip is the correct chemical composition of the steel. Based on [8, 53, 60, 65, 74] it is concluded, that the deteriorative effect of solute carbon has to be minimised by tying it up with microalloying elements such as Ti and/or Nb or by the addition of Cr in order to shift the dynamic strain ageing (DSA) peak to higher temperatures [69]. The influence of carbon on the recrystallisation texture and, hence, the deep-drawing properties is explained more detailed in section 2.5. Next to this a close knowledge of the recrystallisation kinetics is necessary, in order to determine the minimum allowable coiling temperature to ensure a full recrystallisation in the coil. The finish rolling temperature has to be as low as possible with respect to the coiling temperature to produce a desirable rolling texture. This combination is, however, not similar for all IF steels, as has been shown in section 4.1, because of different microalloying concepts. For the IF-Ti steel ◦ the optimum combination of rolling and coiling temperature has been found to be TFR = 710 C ◦ and TC = 670 C. The low temperature difference between the finish rolling and coiling temperature is the most critical restriction in the production of a "soft" hot strip. Fig. 4.11 shows the cooling behaviour of 1 and 2mm thick strip on the run-out table for a finish rolling temperature of 800, 750 and 700 ◦C. From this figure it can be derived, that a very short run-out table is required to allow for the production of a "soft" hot strip. The maximum run-out table length is approximately 25 to 70 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

50m assuming a finish rolling temperature of 700 to 750 ◦C, a coiling temperature of 670 ◦C, a strip thickness of 1mm and a finish rolling speed of 12.5 m/s. The run-out table of a conventional hot rolling mill is, however, up to 150m long. It is hence impossible to produce a "soft" hot strip on a conventional hot rolling mill, without making special precautions. There are at least two different hot strip mill layouts/modifications possible, in order to produce a "soft" hot strip. The first one is to install an additional coiler closely behind the last rolling stand, as shown schematically in Fig. 4.12 a), to minimise the heat losses on the run out-table and to stay within the process window indicated in Fig. 4.11 a). In order to maintain the overall throughput of the hot rolling mill, an additional cooling section should be installed, as depicted in Fig. 4.12 b). This cooling section can markedly reduce the cooling time necessary between rough rolling in the austenitic temperature region and finish rolling in the ferritic temperature region. Modern hot rolling mills, equipped with an ultra fast cooling section (UFC) are already equipped with a short run-out table of a length down to 10m, so that they would only require the installation of an additional cooling unit between rough and finish rolling. Moreover, an effective roll gap lubrication system has to be installed, at least in the last finishing stands, in order to ensure a sufficient suppression of the formation of the shear texture near the surface. The second possibility is to utilise the so called pony mill concept presented in [118] and shown schematically in Fig. 4.12 c).A pony mill is a powerful single stand hot rolling mill designed for high warm rolling reductions, which is placed next to a conventional hot rolling mill to extend the flexibility and possible product range. The slabs are rough rolled and finish rolled to an intermediate thickness on the conventional hot rolling mill and coiled. The hot coil is subsequently transferred to the pony mill and finish rolledín the ferritic temperature range. Conventional steel grades, such as the IF steel grades should be finish rolled in a single rolling pass due to the low profit contribution of these products. In the case of products with a higher profit contribution the pony mill can be also operated in reversing mode. The pony mill concept is similar to the Steckel mill concept, however, a step later in the production process. The pony mill, Fig. 4.12 c), may consist of a single high reduction stand equipped with roll gap lubrication, two coilers and two induction heating devices, one of each at the entry and exit side of the mill. The induction heating devices can be used to (re)heat the strip before and/or after a rolling pass, which allows for the design of very flexible rolling schedules with a coiling temperature higher than that of the finish rolling and/or defined reheating of the strip before the rolling pass. The pony mill allows for the production of a "soft" hot strip, without diminishing the production throughput of the conventional hot rolling mill, because the "soft" hot strip is only rolled to an intermediate thickness on the conventional hot rolling mill. The strip is rolled in the fully austenitic temperature region and conventionally coiled. Subsequently the strip is finish rolled in the pony mill to the desired thickness in ferritic rolling temperature region. Due to the compact design of the pony mill the narrow process window, depicted in Fig. 4.11, can be met. Moreover, the induction heating devices allow to reheat the strip before the ferritic rolling pass or even to 4.6 INDUSTRIAL IMPLICATIONS 71

800 1 mm; 12.5 m/s 2 mm; 12.5 m/s 750

700 process window

650

600

550 finish rolling temperature, °C a) b) 500 0 25 50 75 100 125 150 0 25 50 75 100 125 150 run-out table length, m run-out table length, m

Figure 4.11: Calculated cooling of a strip with a) 1 and b) 2mm thickness for different ﬁnish rolling temperatures, assuming air cooling and an exit rolling speed of 12.5 m/s [8]

F1 F7 Hot Rolling Mill Exit

Run-Out Table Additional Cooling b) Unit Additional Coiler a) Coiler Transfer of the Hot Coil

Pony Mill

Induction Induction Heating Heating Coiler Coiler Single High c) Reduction Stand

Figure 4.12: Concepts to produce a ferritic rolled "soft" hot strip with desirable deep-drawing properties, by a) installing an additional coiler or by b) utilising the pony mill concept 72 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

reheat the strip after rolling to the required coiling temperature. Hence it becomes possible to roll at a lower temperature than the required coiling temperature, which should bring about a further improvement of the deep-drawing properties, as the texture intensiﬁes with decreasing rolling temperatures [51, 52]. Furthermore, IF steels with slower recrystallisation kinetics, such as the IF-TiNb steel in the present work, could be used to produce a "soft" hot strip because the strip could be reheated after rolling to a higher coiling temperature, ensuring complete recrystallisation.

Due to the compact design of the pony mill the installation of a roll gap lubrication system is less complicated and cheaper. The costly and complex installation of a roll gap lubrication system to the existing ﬁnish rolling mill becomes unnecessary.

In [119–121] the heavy warm deformation (HWD) process has been proposed and simulated in the laboratory. The HWD process suited for the pony mill allows for the production of improved ferritic-pearlitic steels. In the simulated HWD process the strip is rough rolled and finish rolled in the austenitic temperature region, followed by cooling down to approximately 500 ◦C and subsequently reheated and finish rolled in the upper ferrite region. Such intermediate reheating and ferritic finish rolling can be easily implemented using the pony mill concept.

4.6.2 "Hard" Hot Strip

Due to the fact that the "hard" hot strip is coiled at conventionally used coiling temperatures, the restrictions to the normal production process are much smaller compared to those for the "soft" hot strip, so that it can be produced on a conventional hot rolling mill. The largest restrictions in the case of the "hard" hot strip are the effective cooling system between a conventional roughing and low temperature finishing, the maximum allowable rolling force of the rolling mill and the effective lubrication of the roll gap. The maximum allowable rolling force determines the minimum finish rolling temperature. In the current experiments the finish rolling temperature was chosen to be 660 ◦C. From the flow stress versus temperature curve presented in Fig. 2.6, on page 17, it can be expected, that the rolling force at 660 ◦C is only slightly higher than at 960 ◦C. The coiling temperature of 400 or 550 ◦C can easily be achieved with a conventional cooling section. Hence, ferritic rolled "hard" hot strip can be produced on a conventional hot rolling mill, provided an effective roll gap lubrication in the finish rolling train is imposed. The more efficacious the cooling between roughing and finishing, the higher is the throughput.

The hot strip requires, however, an supplemental annealing treatment to adjust the desired deep- drawing properties. Due to the improved deep-drawing properties and surface quality the "hard" hot strip is suitable for the hot dip galvanising process. Theses strips are particularly suitable for the direct hot strip galvanising process in coupled pickling and hot dip galvanising lines, such as they are already in operation at several of steel producers. These coupled lines are capable of in-line pickling and hot dip galvanising of hot strip and, hence, represent a cheap processing route. 4.6 INDUSTRIAL IMPLICATIONS 73

4.6.3 Financial Aspects

A general cost analysis, performed in [93], predicts a benefit of approximately 25US$ per tonne for a ferritic rolled "soft" hot strip with a thickness of 1mm compared to a similar cold strip. The individual costs are summarised in Table 4.8. In this analysis it is assumed that the mill productivity and hence the fixed costs are unchanged. From the data given in Table 4.8 it can be estimated, that the benefit for the "hard" hot strip would be only 7US$ per tonne, due to the supplemental batch annealing, required to produce the desired mechanical properties.

Parameter Conventional "Soft" Hot Strip Route Route Hot Strip Thickness 2.0 1.0 Costs of HR Coil [US$/t] X X + 15∗1 Final Product Thickness [mm] 1.0 (skin passed) 1.0 (skin passed) Pickling Costs [US$/t] 10 15∗2 Cold Rolling Costs [US$/t] 27 - Batch Annealing Costs [US$/t] 18 - Skin Pass Rolling Costs 7 7 Total Costs 62 + X 37 + X ∗1 added due to the necessity of roll lubrication, yield losses and power consumption ∗2 added due to a possibly reduced pickling speed

Table 4.8: Cost comparison of a conventionally produced cold strip and ferritic rolled "soft" hot strip [93]

This cost comparison suggests that the required investment, necessary for the production of a ferritic rolled thin gauge hot strip, might pay back. However, cost effective analyses are strongly influenced by the fixed costs, mill configuration and market demands, so that the result can markedly differ from mill to mill.

4.6.4 Conclusion

It has been shown, that it is possible to produce a hot strip with improved deep-drawing properties, comparable to that of the cold strip grades DC 04 and DC 05. However, both processing routes require additional investments. For the production of a "soft" or "hard" hot strip, using a conventional hot strip mill, at least the following installations are necessary:

• a roll gap lubrication system and

• a pre-ﬁnish rolling cooling unit.

Using a conventional hot rolling mill the production of the "soft" hot strip, additionally requires the installation of a close coupled coiler. 74 CHAPTER 4 HOT STRIP GRADES – RESULTS AND DISCUSSION

Alternatively, the pony mill concept can be utilised, making the last mentioned investment unnecessary. Due to the uncoupling of the "final" ferritic finish rolling and the "intermediate" conventional hot rolling of the strip is the overall throughput of the hot rolling mill not diminis- hed and, additionally, the production flexibility increases. The compact layout of the pony mill together with the possibility of the pre- and post rolling heating of the strip, even new or more complex thermomechanical rolled products might be produced, which brings about a required product diversification. Chapter 5

Cold Strip Grades – Results and Discussion

The "cold" strip grades described in this work are produced from ferritic rolled hot strips. The aim is to bequeath the microstructure and texture of the initial hot trips to the subsequent cold rolling process, in order to improve the ﬁnal deep-drawing properties or reduce the necessary cold reduction. The initial strips can be either "soft" or "hard" hot strips.

The "soft" hot strip was rolled according to rolling schedule 1 in Fig. 3.4 a). The "hard" hot strip was rolled according to rolling schedule 2 in Fig. 3.4 b), however, without the additional batch annealing treatment. The hot rolling conditions are described in section 3.3.3. The ferritic rolling strains and temperatures were chosen according to the optimised parameters determined in the previous chapter. The ferritic rolled strips were subsequently cold rolled to the ﬁnal thickness of 0.5mm with a cold reduction of εCR ≈ 40 − 80% (section 3.5) and subjected to a simulated continuous annealing (section 3.6). Conventional strips were produced as a reference according to the rolling schedule shown in Fig. 3.4 b). The cold rolling reduction for these strips was εCR ≈ 80%. The simulated continuous annealing process parameters were identical for all products.

The initial microstructure of the "soft" hot strip reveals a fully recrystallised microstructure with a mean grain size of approximately 40 − 50µm. The grains are slightly elongated along RD, Fig. 5.1 a). The "hard" hot strip reveals a strained microstructure, as can be seen in Fig. 5.1 b). The austenitically rolled strip, shown in Fig. 5.1 c), exhibits a polygonal microstructure with a grain size of approximately 30−40µm. In terms of the microstructure the"soft" and the conventional hot strips are similar.

However the initial textures of the hot strips are different. The "soft" hot strip exhibits a desirable {111} texture with a maximum near {554}h225i, Fig. 5.2 a). The "hard" hot strip exhibits a clear rolling texture, Fig. 5.2 b) with a complete γ-fibre and a partial α-fibre with a maximum near {112}h110i. Fig. 5.2 c) shows the initial texture of the conventional hot strip. The conventionally rolled hot strip reveals a much weaker texture. The texture consists of a weak complete γ-fibre and partial α-fibre with maximum near {112}h110i.

75 76 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

TD RD

a) 100 µm b) 100 µm c) 100 µm

Figure 5.1: The initial microstructure prior cold rolling of the a) "soft", b) "hard" and c) conventional hot strip

j1 0 153045607590 0

15 1 1 1 30

75 8 9 6 3 90 a) “Soft” Hot Strip b) “Hard” Hot Strip c) Conventional Hot Strip

Figure 5.2: Coiling texture prior cold rolling for the different initial hot strips (Intensity levels 1, 2, 3, ...)

5.1 Microstructural Development

5.1.1 Initial Hot Strip "Soft"

The microstructural development of the "soft" hot strip during the subsequent cold rolling is shown in Fig. 5.3 a) - d) for a cold reduction, εCR, of approximately 55 and 82% respecively. At a reduction of 55% the microstructure is strained and elongated along the rolling direction, Fig. 5.3 a) and b). The microstructure consits of relatively thick smooth and creased layers, with a thickness of up to 60µm. For a cold reduction 82% the microstructure does not change strongly, as can be seen in Fig. 5.3 c) and d), however, it appears a little coarser. The creased layers carry strain inhomogeneities, already described in further detail in section 2.5.2. These strain inhomogeneoities are mostly referred to as in-grain shear bands [72]. 5.1 MICROSTRUCTURAL DEVELOPMENT 77

55%

e»

t”

a) b)

Initial hot strip: “sof

82%

e»

c) d)

56%

e»

e) f)

Initial hot strip: “hard”

85%

e»

g) h)

Figure 5.3: Microstructure of the cold rolled strips with different cold reductions; initial hot strip "soft" a) - d) and "hard" hot strip e) - h) 78 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

5.1.2 Initial Hot Strip "Hard"

The "hard" hot strip reveals a highly strained microstructure, Fig. 5.3 e) and f), also with a layered structure, consisting of smooth and creased areas. The creased layers also contain in-grain shear bands. The thickness of the layers varies from approximately 20 to 50µm. The strip rolled with εCR = 82% also exhibits a layered structure, Fig. 5.3 g) and h), however, the thickness of these layers is only about 5 to 10µm. The "hard" hot strip exhibits for both cold reductions a much more strained microstructure after cold rolling, compared to the "soft" hot strip. The microstructure of the conventionally rolled strip, shown in Fig. 5.4 a) and b), also reveals a strained microstructure, with smooth and creased layers. This microstructure is comparable to that of the "hard" hot strip rolled with a cold reduction of 55%, depicted in Fig. 5.3 e) and f). After the supplemental continuous annealing treatment all strips exhibit, independently of the initial hot strip and cold reduction, a recrystallised microstructure with polygonal grains. The grain sizes of the different strips are summarised in Table 5.1. The "soft" hot strip reveals an ASTM grain size of 8 independently of the cold reduction. For the "hard" hot strip an ASTM grain size of 8 − 9 was measured for a cold reduction of εCR = 56% and 9 for the highest cold reduction of 85%. The conventional hot strip also exhibits an ASTM grain size of 9.

ASTM Grain size (OM) Grain size [µm] (EBSD) εCR ≈ 55% εCR ≈ 80% εCR ≈ 55% εCR ≈ 80% "Soft" Hot Strip 8 8 30 20 "Hard" Hot Strip 8-9 9 20 20 Conv. Hot Strip 9 20

Table 5.1: Average recrystallised grain size of the different cold rolled strips

The images of the EBSD measurements of the recrystallised specimens are shown in Fig. 5.6 a) and b) , Fig. 5.7 a) and b) and Fig. 5.8 a). All specimens exhibit a high image quality and more than 90% the grain boundaries are high angle grain boundaries with a misorientation angle ≥ 10 ◦. The mean grain sizes, calculated from these measurements are summarised in Table 5.1. An avearge grain size of approximately 30µm has been determined for the cold strip produced of a "soft" hot strip, Fig. 5.6 c) and d). The cold strip produced of a "hard" hot strip revealed with 20µm a slightly ﬁner grain size, Fig. 5.7 c) and d). The conventionally produced cold strip also exhibits a mean grain size of 20µm, Fig. 5.8 b).

5.2 Mechanical Properties

As it has already been described in section 3.7.1, the mechanical and technological properties were determined in tensile tests in 0, 45 and 90 ◦ to the rolling direction. From each strip two specimens for each of these directions was tested. The deep-drawing properties were calculated 5.2 MECHANICAL PROPERTIES 79

RD TD

77%

e»

Conventional hot strip a) b)

Figure 5.4: Mircrostructure of the cold rolled conventional hot strip

TD e»82% e»CR 55% CR RD

t”

Initial hot strip: “sof a) b)

e»85% e»CR 56% CR

Initial hot strip: “hard” c) d)

e»CR 77%

Conventional hot strip e)

Figure 5.5: Microstructure of the cold rolled and annealed "soft", "hard" and conventional hot strip 80 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

a) b) eCR = 56% eCR = 82%

0.16 0.12 0.14 0.10 0.12 0.10 0.08 0.08 0.06 0.06 0.04

Area Fraction 0.04 Area Fraction 0.02 0.02 0.00 0.00 1 10 100 1 10 100 c) Grain Size (Diameter), µm d) Grain Size (Diameter), µm

Figure 5.6: EBSD pictures of the cold strip produced of a "soft" hot strip for a cold reduction of a) 56% and b) 82%; c) and d) show the corresponding grain size distributions

a) b) eCR = 62% eCR = 85%

0.12 0.12 0.10 0.10 0.08 0.08 0.06 0.06 0.04 0.04

Area Fraction

Area Fraction 0.02 0.02 0.00 0.00 c) 110 100 d) 1 10 100 Grain Size (Diameter), µm Grain Size (Diameter), µm

Figure 5.7: EBSD pictures of the cold strip produced of a "hard" hot strip for a cold reduction of a) 62% and b) 85%; c) and d) show the corresponding grain size distributions 5.2 MECHANICAL PROPERTIES 81

a) eCR = 82%

0.14 0.12 0.10 0.08 0.06

Area Fraction 0.04 0.02 0.00 1 10 100 b) Grain Size (Diameter), µm

Figure 5.8: a) EBSD pictures of the cold strip produced of a conventional hot strip for a cold reduction of 82%; b) shows the cosseponding grain size distribution according to the equations 2.1 to 2.4. In Figs. 5.9 – 5.11 circles designate measurements at 0◦ to RD, squares measurements at 90◦ to RD and triangle measurements at 45◦ to RD. The open symbols represent the properties of a conventionally produced cold strip. Additionally, these values are highlighted with grey boxes, to enable an easy visual comparison of the properties of the new grades with those of the conventional cold strip grade.

5.2.1 Initial Hot Strip "Soft"

The 0.2% proof strength, Rp0.2, and the ultimate tensile strength, Rm, are plotted in Fig. 5.9 a) and b). Both properties are insensitive to the total cold reduction in the range tested. A mean value of Rp0.2 ≈ 84MPa and Rm ≈ 285MPa have been determined.

The total elongation and the strain hardening coefﬁcient are also insensitive to the total cold reduction, as it can be seen in Fig. 5.10 a) and b). A total elongation of A50 ≈ 42% and a strain hardening coefﬁcient of nm ≈ 0.264 have been measured, irrespective the degree of the cold reduction.

Fig. 5.11 a) shows the r-values versus the cold reduction for the "soft" hot strip. The r-values in 0◦ and 90◦ to RD rapidly increase up to approximately 3.4 and 3.0 for a cold reduction of ◦ εCR > 55%. The r-value in 45 to RD steadily inceases with increasing cold reduction from approximately 1.4 up to 2.0 for a cold reduction of 80%. This behaviour is reﬂected in the Fig. 5.11 b), as the mean r-value, rm, steadily increases up to 2.6, whereas the planar anisotropy, ∆r, exhibits a local maximum of 1.5 at about 60% and drops to aproximately 1.1 for a cold reduction of 80%. At a cold reduction of 60 − 65% the deep-drawability becomes equal to that of a conventionally produced cold strip, however, with a slightly higher planar anisotropy. 82 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

5.2.2 Initial Hot Strip "Hard"

The hard hot strip has been coiled at two different coiling temperatures, 550 and 400 ◦C. The 0.2% proof strength, Rp0.2, and ultimate tensile strength, Rm, for the two different inititial hot strips is plotted versus the cold reduction in Fig. 5.9 c) - f). These properties were found to be independently of the cold reduction and the coiling temperature of the initial hot strip. A mean value of Rp0.2 ≈ 90MPa and Rm ≈ 294MPa has been calculated. Also the total elongation, A50, and the strain hardening coefﬁecient, n, were measured with A50 ≈ 38% and n ≈ 0.265 independently of the coiling temperature of the initial hot strip and the cold reduction. The r-values measured for the cold strip produced of a "hard" hot strip are shown in Fig. 5.11 c) and e). The r-values in 0◦ and 90◦ to RD increase with increasing cold reduction up to 2.2 (0◦) and 2.4 (90◦) for a cold reduction of approximately 60% and then slightly decreases to 1.8 and 2.2, respectively. The r-value in 45◦ to RD steadily increases with increasing cold reduction and becomes equal or greater than the r-value in 0◦ to RD for cold reductions greater than 75%. This ◦ ◦ behavior leads to the quite high rm-values, despite the somewhat lower values in 0 and 90 to RD. The planar anisotropy achieves at εCR ≈ 80% desirable low values of ∆r ≈ 0, Fig. 5.11 d) and f). Because the conventional ∆r-value is only applicable for a v-shaped r-value distribition, wich is here not the case, this value may be misleading. Therefore the ∆rmax-value, equation 2.4, is introduced in Fig. 5.11 d) and f). The ∆rmax-value is deﬁned as the difference between the highest and lowest value of the r-value distribution. The unusual low ∆rmax-values in Fig. 5.11 d) and f) indicate that the strips will form ears to a much lower extent compared to a conventional cold strip.

5.3 Texture Development

To investigate the development of the rolling and annealing texture of the different initial hot strips, texture specimens were taken before and after the supplemental continuous annealing simulation. The texture was measured in all cases in the thickness centre plane.

5.3.1 Initial Hot Strip "Soft"

Fig. 5.12 a) shows the development of the cold rolling and annealing texture with increasing cold reduction, εCR, for the initially "soft" strip. The rolling texture exhibits a weak γ-fibre texture for a low cold reduction of 35%. Up to a cold reduction of 68% the texture slightly intensifies and the maximum near {554}h225i develops with an intensity of f (g) = 5. For higher cold reductions, additionally the typical cold rolling component {112}h110i develops with an intensity of f (g) = 6. Even for the low cold reduction of 35% a very homogeneous γ-fibre texture developed due to the recrystallisation during anealing with a maximum intensity of f (g) = 7. Up to a cold 5.3 TEXTURE DEVELOPMENT 83

Initial Hot Strip: “Soft” 120 320 0° 310 110 90° 45° Conv. Strip 300

100 290

p0.2 90 m 280

R.MPa

R.MPa 270 80 260 a) b) 70 250

Initial Hot Strip: “Hard” TC = 550°C 120 320

310 110 300

100 290

p0.2 90 m 280

R.MPa

R.MPa 270 80 260 c) d) 70 250

Initial Hot Strip: “Hard” TC = 400°C 120 320

310 110 300

100 290

p0.2 90 m 280

R.MPa

R.MPa 270 80 260 e) f) 70 250 10 20 30 40 50 60 70 80 90 10 20 30 40 50 60 70 80 90

eCR.% eCR.%

Figure 5.9: The 0.2% proof strength, Rp0.2, and the ultimate tensile strength, Rm, of a) - b) the "soft", c) - d) the ◦ ◦ "hard" hot strip coiled at TC = 550 C and e) - f) the "hard" hot strip coiled at TC = 400 C (The open symbols represent the properties of a conventionally produced cold strip) 84 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

Initial Hot Strip: “Soft” 55 0.30 0° 50 90° 0.28 45° 45 0.26 40

n, -

50 0.24 A,% 35

0.22 30 Conv. Strip a) b) 25 0.20

Initial Hot Strip: “Hard” TC = 550°C 55 0.30

50 0.28

45 0.26 40

n, -

50 0.24 A,% 35

0.22 30 c) d) 25 0.20

Initial Hot Strip: “Hard” TC = 400°C 55 0.30

50 0.28

45 0.26 40

n, -

50 0.24 A,% 35

0.22 30 e) f) 25 0.20 10 20 30 40 50 60 70 80 90 10 20 30 40 50 60 70 80 90

eCR,% eCR,%

Figure 5.10: The ultimate elongation, A50, and the strain hardening coefﬁcient, n, of a) - b) the "soft", c) - d) the ◦ ◦ "hard" hot strip coiled at TC = 550 C and e) - f) the "hard" hot strip coiled at TC = 400 C (The open symbols represent the properties of a conventionally produced cold strip) 5.3 TEXTURE DEVELOPMENT 85

Initial Hot Strip: “Soft” 4.0 3.0 2.5 r r 3.5 0° m r90° 2.5 Dr 2.0

3.0 r45° Drmax 2.0 1.5 2.5

2.0 1.5 1.0 r [-]

r [-]

r [-] 1.5 Conv. Strip 1.0 0.5 1.0 Conv. Strip 0.5 0.0 0.5 a) b) 0.0 0.0 -0.5

Initial Hot Strip: “Hard” TC = 550°C 4.0 3.0 2.5

3.5 2.5 2.0 3.0 2.0 1.5 2.5

2.0 1.5 1.0 max

r [-]

r, r [-]

1.5 DD 1.0 0.5 1.0 0.5 0.0 0.5 c) d) 0.0 0.0 -0.5

Initial Hot Strip: “Hard” TC = 400°C 4.0 3.0 2.5 3.5 2.5 2.0 3.0 2.0 1.5 2.5

2.0 1.5 1.0

r [-]

max 1.5 1.0 0.5

r, r [-]

1.0 DD 0.5 0.0 0.5 e) f) 0.0 0.0 -0.5 10 20 30 40 50 60 70 80 90 10 20 30 40 50 60 70 80 90

eCR[%] eCR[%]

Figure 5.11: The values of ,rm,∆r and ∆rmax plotted versus the cold reduction for the different initial hot strips (The open symbols represent the properties of a conventionally produced cold strip) 86 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

reduction of 68% the γ-ﬁbre texture lacks the typical maximum near {554}h225i. As soon as the {112}h110i component becomes clearly visible in the rolling texture, also the {554}h225i becomes more pronounced with a maximum intensity of f (g) = 14 in the annealing texture.

5.3.2 Initial Hot Strip "Hard"

The texture development of the initially "hard" strip is shown in Fig. 5.12 b). The cold rolling texture shows a relative intensive typical rolling texture even for a low cold reduction of 34% with maximum intesity of f (g) = 6 near {112}h110i. The γ-fibre is weakened for increasing cold reductions, whereas the {112}h110i component is strengthened. For cold reduction greater than approximatly 65% the texture is dominated by the {112}h110icomponent with an intensity of up to f (g) = 12. Even for the lower cold reductions of 34% a typical {111} texture develops due to recrystallisation during annealing with a maximum intensity of f (g) = 8 near {554}h225i. The {554}h225icomponent is intensified with increasing cold reduction, so that the texture for a cold reduction of 85% is dominated by the {554}h225i texture component with an intensity of up to 19. The texture development of the cold strip produced of the "hard" hot strip coiled at 400 ◦C is similar to that of the strip coiled at 550 ◦C, which corresponds to the results of the mechanical testing, section 5.2.2. The cold rolling and annealing texture of the conventionally produced cold strip is shown in Fig. 5.13a) and b), respectively. The rolling texture consists of a complete γ-fibre and a partial α-fibre with a maximum intensity of f (g) = 6 near {112}h110i. During the subsequent recrystallisation annealing the rolling texture is transformed into a desirable {111} texture with a maximum of f (g) = 11 near {554}h225i.

5.4 Cupping Tests

In order to get a better measure of the earing behaviour of the different cold strips the standard cupping test, as described in section 3.7.2, was used. The cups were drawn from the cold strips produced of a "soft" and a "hard" hot strip rolled with a total rolling reduction of εCR ≈ 60% and 80%, respectively. Additionally, cups were drawn from the conventionally produced cold strip, as a reference.

5.4.1 Initial Hot Strip "Soft"

The cups drawn from the initially "soft" hot strips exhibit for a cold reduction of 60% and 80% a classical 4-ear shape. Fig. 5.14 a) shows the cup drawn from an initial "soft" hot strip with a cold 5.4 CUPPING TESTS 87

j j 1 [°] 1 [°] 0 153045607590 0 153045607590 0 0

15 15 1 6 1 1 30 30 3 4 1 45 45

F [°] F [°] 60 60

75 75 e = 35% 7 e = 34% 7 8 90 CR 90 CR 1

1 7 4 1 1

e e 12 13 CR = 56% 8 CR = 55%

1 8 1 1 9 1 15

e e 16 CR = 68% 5 10 CR = 65%

1 1

5 6 12 1 1

10 e 11 e 4 CR = 78% CR = 79% 16 15

5 6 1 1 1

12 e 11 14 e 6 CR = 82% 13 CR = 85% 19 18 a)Cold Rolling Texture Annealing Texture b) Cold Rolling Texture Annealing Texture

Figure 5.12: Development of the cold rolling and annealing texture with increasing cold reduction, εCR, for different initial hot strips; a) "Soft" and b) "Hard" Hot Strip (Intensity levels 1, 2, 3, ...) 88 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

j1 0 153045607590 0

6 15 10 1 1 30

75 11 90 a) b)

Figure 5.13: a) Cold rolling texture and b) annealing texture of the conventionally produced cold strip,with εCR ≈ 80% (Intensity levels 1, 2, 3, ...) reduction of 82%. From the measured shape, shown in Fig. 5.15 (dotted line) it can be seen that the maxima of the shape are located parallel and transverse to the rolling direction. The maximum height of the ears is approximately 1.7mm. The shape is qualitatively and quantitatively comparable to that of the conventionally produced cold strip, Fig. 5.14 c).

5.4.2 Initial Hot Strip "Hard"

For a cold reduction of approximately 60% the initially "hard" hot strip reveals 4 ears, similar to that of the conventional hot strip. In contrast to that, a strip rolled with a cold reduction of about 80% yields a cup with 6 ears, as it can be seen in Fig. 5.14 b). The ears are located in approximately 0◦ and in 60◦ to RD, Fig. 5.15, and the maximum ear height is with 1.2mm markedly smaller than that of the conventional cold strip.

5.4.3 Calculated r-Value Distribution and Relative Earing

Additionally, the r-value distribution and the relative earing behaviour was calculated from the texture data. The software program "ANIS-MPI" [122] was used for the calculation. The program utilises the Taylor model and assumes pencil glide. The r-value distribution and the relative earing has been calculated for the different cold strip grades rolled with a cold reduction of approximately 80%. Fig. 5.16 a) shows the r-value distribution for the initially "soft", "hard" and conventional hot strip. Both, the cold strip produced from the "soft" hot strip and the conventionational hot strip exhibit the classical v-shaped r-value distribution with the maximum r-values in 0◦ and 90◦ to RD. The initial "soft" hot strip reveals 5.4 CUPPING TESTS 89

a) 20 mm b) c)

Figure 5.14: Cups drawn from cold strips produced from a) a "soft" hot strip (εCR = 82%), b) a "hard" hot strip (εCR = 85%) and c) a conventional hot strip (εCR = 82%)

2.00 “Soft” Conv. 1.75

1.50 “Hard” 1.25

1.00

0.75

Ear height [mm] 0.50

0.25 0.00 0 45 90 135 180 225 270 315 360 Angle to RD [°]

Figure 5.15: Measured shape of the cups drawn from cold strips produced from a) a "soft" hot strip (εCR = 82%), b) a "hard" hot strip (εCR = 85%) and c) a conventional hot strip (εCR = 82%)

3.0 1.15 Initial Hot Strip: “Soft”e = 78% 2.5 CR 1.12 e “Hard”CR = 79% Conv.e = 77% 2.0 CR 1.09 1.5

r-value [-] 1.06

1.0 Initial Hot Strip: rel. earing, - e “Soft”CR = 78% “Hard”e = 79% 1.03 0.5 CR e Conv.CR = 77% 0 1 0 15 30 45 60 75 90 0 15 30 45 60 75 90 a) b) Angle to RD [°] Angle to RD [°]

Figure 5.16: Calculated r-value distribution a) and calculated relative earing b) for the different "cold strips" on the basis of the thickness centre texture measurement 90 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

slightly higher r-values than the conventional strip. The calculated shape of the cups exhibits correspondingly 4 ears located at 0◦ and 90◦ to RD. The initial "hard" hot strips exhibits a different r-value distribution with a maximum in 30◦ and 90◦ to RD and the r-value in 45◦ to RD is slightly higher than that at 0◦. The caculated relative earing reveals 6 ears, located at approximately 0◦ and 60◦ to RD. These calculated results are in excellent agreement with the cupping tests.

5.5 Discussion

5.5.1 Microstructural Development

During the cold rolling of the "soft" and the "hard" hot strip a very different microstructure developed. The initially "soft" hot strip revealed a relatively coarse strained microstructure after cold rolling with a reduction of about 50 − 80%, as it can be seen in Fig. 5.3 b) and d). The microstructures look quite similar for the low and high reduction. It is obvious, that the smooth and creased leayers are larger than the initial grain size of the hot strip. One can even get the impression, that the thickness of the layers slightly increased with increasing cold reduction. In [123] it has been proposed that in-grain shear bands can cross grain boundaries, as they are shared with other γ-fibre grains of a similar orientation. In the case of the initially "soft" hot strip, the microstructure is dominated by {111} grains, so that the probability is quit high, that the shear bands can cut through a couple of neighbouring γ-fibre grains that behave as one during the deformation, so that these large elongated grains form coarse layers observed in the microstructure. The development of the deformation microstructure of the initially "hard" hot strip, is shown in Fig. 5.3 f) and h). The initially strained grains are further elongated and start to form a thin layered structure. The microstructure of the strip rolled with a rather low εCR ≈ 50%, Fig. 5.3 f) is comparable to that of the conventionally rolled hot strip, rolled with a much higher cold reduction of about 80%, Fig. 5.4 b). The strained microstructure, observed in the specimen rolled with a total reduction of about 80% is much finer than that of a "soft" hot strip with a similar εCR and the layers are only a few microns thick Fig. 5.3 h). The microstructural development reflects the cumulation of the previous (warm) ferritic rolling strain and the cold rolling strain. After the subsequent annealing treatment the microstructure of the initially "soft" and "hard" hot strips appears quite similar. The strips exhibit a recrystallised microstructure with polygonal grains with a similar grain size, irrespective the initial hot strip condition and the cold rolling reduction. From the EBSD measurement, shown in Fig. 5.6, Fig. 5.7 and Fig. 5.8 the mean grain size was found to be 30µm for the initially "soft" hot strip and 20µm for the initially "hard" as well as the conventional hot strip. The high image quality and the high fraction of high angle grain boundaries indicate that all the specimens are fully recrystallised. I tseems that the recrystallisation equalises the different initial microstructures and produces a similar microstructure in all specimens. 5.5 DISCUSSION 91

5.5.2 Mechanical Properties

The mechanical and technological properties, besides the r-values, of the cold strips were found to be independent of the initial hot strip condition and the total cold reduction, Fig. 5.9 a) - f) and Fig. 5.10 a) - f). The mechanical properties, summarised in Table 5.2, fulﬁl generally the requirements of the steel grade DC 06 (Table 4.3).

−3 Initial Hot Strip Rp0.2 [MPA] Rm [MPa] Ag [%] A50 [%] nm10 [-] ◦ "Soft" TC = 670 C 84 ± 3 285 ± 5 27 ± 1 41 ± 2 264 ± 3 ◦ "Hard" TC = 550 C 90 ± 3 294 ± 4 27 ± 1 38 ± 2 265 ± 4 ◦ "Hard" TC = 400 C 95 ± 5 294 ± 4 27 ± 1 38 ± 2 265 ± 4 Conventional 86 ± 2 290 ± 2 25 ± 2 38 ± 1 263 ± 2

Table 5.2: Mechanical properties of the different cold strip grades

In contrast to this, the r-values exhibit a strong dependency on the initial hot strip and the total cold reduction. The r-values of the initially "soft" hot strips measured in 0◦ and 90◦ to RD rapidly increase up to approximately 3.4 and 3.0 for cold reductions greater than 55%. The r-value in 45◦ to RD steadily inceases with increasing cold reduction, Fig. 5.11 a). The deep-drawability is after ≥ 60 − 65% comparable to that of a conventionally produced cold strip, however, with a slightly higher planar anisotropy. The development of the r-values can be correlated to the texture development shown in Fig. 5.12 a). The strips exhibit up to a cold reduction of 68% a rather homogeneous γ-fibre texture and the maximum intensity increases from f (g) = 7 to 10. For the higher cold reductions a maximum near {554}h225idevelops with a maximum intensity of f (g) = 14. The high mean r-values can be explained in terms of a very homogeneous and intense γ-fibre texture, which correlates to the findings in [20, 25, 57]. However, it has to be mentioned that these strips exhibit despite the high mean r-value a larger tendency to form ears, which is caused by rather high ∆r-values. On the basis of the present results it can be concluded, that a homogogeneuos {111} texture results in quite high mean r-values, wich are, unfortunately, accompanied with elevated ∆r-values. A high ∆r-value is undesirable because it leads to an increasing amount of scrap during the further production of the deep-drawing part. In the initially "hard" hot strip the r-values parallel and transverse to RD increase with increasing cold reduction up to approximately 60% and then slightly decrease. The r-value in 45◦ to RD steadily increases with increasing cold reduction. For higher cold reductions, εCR ≥ 75%, the r-values in 45◦ are equal or even greater than those in 0◦ to RD, Fig. 5.11 c) and e). This r-value distribution still leads to high mean r-values, however, lower than those of the initially "soft" hot strips. The annealing textures of the initially "hard" hot strips contain, just like the "soft" hot strips, a desirable {111} component. However, is the texture less homogeneously distributed along the γ-fibre and exhibits a peak near {554}h225i. With increasing cold reduction this peak is intensified and reaches a maximum intensity of f (g) = 19 for a cold reduction of 85%. Such a peaked texture leads, hence, to the somewhat lower r-values in 0◦ and 90◦ to RD, but due to the high r-value in 45◦ to RD the mean r-value is for cold reductions ≥ 75% comparable to that of the conventional hot strip Fig. 5.11 d) and f). Moreover, these strips exhibit desirably low ∆r 92 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

and ∆rmax values of approximately 0. Because of such low ∆r a modiﬁed earing symetry with a lower tendency to form ears can be expected, as it will be discussed more deeply in section 5.5.4.

5.5.3 Texture Development

The initial γ-fibre texture of the "soft" hot strip, Fig. 5.2 a), is initially weakened during cold rolling for the lowest reduction of 35%, as it can be seen in Fig. 5.12 a). With increasing cold reduction the maximum near {554}h225i is again intensified. From approximately 78% reduction an additional maximum near {112}h110i develops. The weakening of the initial γ-fibre texture of the "soft" hot strip can be explained by the randomising effect of the deformation, due to crystal fragmentation. In [58] it has been shown that in-grain shear bands frequently lead to a misorientation of up to 15 ◦ along the γ-fibre. The "soft" hot strip exhibits a strong tendency for shear band formation due to its initial γ-fibreThe˙ shear band crystallites partly rotated back to their initial orientation {554}h225i, while others rotate towards {112}h110i at higher rolling reductions. There is no consensus in the literature about the origin of the recrystallisation texture in IF steels. Basically, two main theories of the texture development have been suggested [44]. The oriented nucleation theory supposes that the final texture is determined by the orientation of the nuclei and the orientation of the nuclei depends on the deformed matrix [78]. The oriented growth theory proposes that the nuclei with a large variety of orientations are formed, however those with a certain misorientation with respect to the surrounding matrix will grow faster and dominate the final texture [78]. The development of the homogeneous γ-fibre texture, Fig. 5.12 a), after recrystallisation can be attributed to the oriented nucleation mechanism of γ-fibre grains, as it has been proposed in [78]. The recent results suggest, that {hkl}h111i nuclei predominantly form within the deformed microstructure. As soon as a distinct maximum near {112}h110i becomes visible in the texture, a change in the recrystallisation texture can be observed. The γ-fibre texture becomes less homogeneous and a maximum near {554}h225i develops. This might be explained by the oriented growth of near {554}h225i nuclei at the expense of {112}h110i oriented matrix. The latter assumption is supported by the texture development observed for the initially "hard" hot strip, shown in Fig. 5.12 b). With increasing cold reduction the intensity of the {112}h110i texture component steadily increased and for cold reductions greater than 65% the texture is dominated by this component. In agreement with the assumption mentioned above, the {554}h225i component drastically intensified in the annealing texture. The present results suggest a gradual transition of the recrystallisation mechanism from oriented nucleation to selective growth with increasing cold reduction. The rolling texture of the initially "soft" hot strip is relatively weak and not concentrated into a single component. Therefore, there is no impact of the selective growth on the texture development and hence the oriented nucleation leads a homogeneous γ-fibre texture during recrystallisation. With increasing cold reduction, the texture starts to concentrate on the {112}h110i texture component and, hence, the 5.6 INDUSTRIAL IMPLICATIONS 93 potential role of the oriented growth is increased, leading to a visible peak near {554}h225i in the recrystallisation texture. In the case of the initially "hard" hot strip, the increasing influence of the oriented growth is even more pronounced, as the texture is already concentrated on the {112}h110i component prior cold rolling. A similar trend has been described in [77, 78] on the basis of texture simulations and growth rate calculations.

5.5.4 Cupping Tests

The calculated r-value distribution of the "soft" hot strip and the conventional hot strip, Fig. 5.16 a) matches qualitatively the corresponding experimental r-values, shown in Fig. 5.11 a). Due to the fact, that the calculated earing of the two strips, Fig. 5.16 b), also matches qualitatively the measured shape of the cups shown in Fig. 5.15 (dashed and dotted line), it can be concluded, that measured r-values correspond to the maxima and minima of the calculated distribution.

The calculated r-value distribution of the "hard" hot strip, Fig. 5.16 a), however shows that the measured r-values did not correspond to the calculated maxima and minima. Nevertheless, a qualitatively correspondence can be found, as the r-value in 0◦ to RD is slightly lower than that at 45◦, which is lower than that at 90◦. The calculated shape of the "hard" hot strip, Fig. 5.16 b), qualitatively matches the measured shape, shown in Fig. 5.15 (solid line). Therefore, it can be conclude, that the calculated r-value distribution is qualitatively correct.

The cupping tests, together with the calculated earing of the initially "hard" hot strip rolled with a reduction approximately 80% indicate, that the 6-ear shape can be directly correlated to the measured centre texture. The 6-ear shape per se is not advantageous for the deep-drawing, however in combination with the lower ear height, the cut-off scrap can be markedly reduced.

5.6 Industrial Implications

In terms of cold rolling and annealing equipment there are no additional installations required in order to produce a cold strip from ferritic rolled hot strip. Naturally, the similar requirements as described in 4.6 apply for the production of the initial ferritic rolled hot strip.

Cold strip produced of either a "soft" or a "hard" hot strip can bring about a certain improvement of the deep-drawing properties, compared to a conventional cold strip. The initially "soft" hot strip exhibits mean r-values comparable to that of a conventional cold strip at a cold reduction of only approximately 60 − 65%, however with a slightly higher planar anisotropy. Cups drawn from these strips exhibit 4 ears and the ear height is slightly higher than that of a conventional cold strip. The initially "hard" hot strip reveals a very low planar anisotropy, together with a sufﬁcient high mean r-value for cold reductions ≥ 80%. These strips exhibit in the cupping test 6 ears and a lower ear height compared to the conventional cold strip. 94 CHAPTER 5 COLD STRIP GRADES – RESULTS AND DISCUSSION

Due to the lower cold reduction necessary to produce the required deep-drawing properties in the case of the "soft" hot strip and the improved deep-drawing properties with the low planar anisotropy in the case of the "hard" hot strip, the production of cold strip produced from ferritic hot strip represents an interesting enlargement of the product range. Cost advantages, such as they were reported for the ferritic rolled hot strip grades, are absent in the case of the ferritic rolled cold strip grades. Chapter 6

Towards a new Hot Rolling Strategy for Low Alloy TRIP Steels

Crash resistance and weight reduction are main topics of the automotive industry in the recent years. Weight reduction can be easily realised by reducing the sheet thickness or cross-section of the components. To ensure an equivalent or even improved crash resistance and strength of the components with a reduced thickness, an enhanced material strength is required [124, 125]. An increase in strength is conventionally accompanied by a decrease in ductility. With the introduction of multiphase steels, such as dual phase (DP), transformation induced plasticity (TRIP) and complex phase (CP) steel an increased strength together with a desirably high formability has been obtained.

The increasing interest of the automotive industry in high-strength manifested for example in the Ultra Light Steel Auto Body (ULSAB) and ULSAB AVC (Advanced Vehicle Concept) project [126, 127]. The ULSAB project demonstrated out that a body-in-white (BIW) weight reduction of 20% along with a cost reduction of 10% is possible, by using high-strength steel. Before 1990 the BIW contained less than 40% high-strength steel. 90% of the steel used within the ULSAB project had a minimum ultimate tensile strength (UTS) of Rm ≈ 210MPa, whereas in the more recent ULSAB AVC project, already more than 55% of the steel grades had an UTS of Rm ≥ 500MPa, as is illustrated in Fig. 6.1. The majority of the high-strength steels used in the project was Dual-Phase (DP) steel and approximately 4% of the steel envisaged was TRIP steel.

95 96 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

Mart 1250/1520 Misc. 1% 2% Mart 950/1200 3% CP 700/800 1% BH 210/340 4% TRIP 450/800 BH 260/370 4% DP 700/1000 6% 30% IF 300/420 4%

HSLA 350/400 1%

DP 280/600 7% DP 300/500 DP 500/800 8% 22% DP 350/600 DP 400/700 3% 4%

Figure 6.1: Materials used in the ULSAB-AVC project [126]

6.1 Processing Routes

Low alloy TRIP steels can be produced as hot or cold strip, as shown in Fig. 6.2 a) and b). Possible final products can be cold strip, hot strip and hot dip galvanised hot or cold strip. The principle production processes are explained in more detail in the following sections. A lot of research has been done in the field of cold rolled and annealed low alloy TRIP steels. This work mainly focused on the influence of the chemical composition [130–132], and the design of the supplemental annealing cycle [133, 134] on the microstructure and mechanical properties of low alloy TRIP steels. In contrast to that, only little research has been done in the field of hot rolled low alloy TRIP steels [135, 136]. This chapter aims at expanding the knowledge in this field.

6.1.1 Cold rolling and annealing

Initial hot strips of low alloy TRIP steel to be used for subsequent cold rolling can be soft or a hard hot strips. The soft hot strip exhibits a microstructure consisting mostly of ferrite and pearlite. In this case, a sufficiently high coiling temperature around 700 ◦C is necessary. A lower coiling temperature of about 500 ◦C leads to ferritic-bainitic microstructure. Because bainite is a relatively hard phase, the resulting strip has a higher strength. Consequently, higher rolling forces can be expected, when using a hard hot strip. The final cold strip possibly then exhibits in the latter case a more homogeneous and fine grained microstructure [137, 138]. In the sequence of hot rolling, coiling and subsequent cold rolling only the desired thickness and surface quality is obtained, Fig. 6.2 a). To adjust the desire multiphase microstructure and recrystallise the strip, a supplemental annealing treatment is required. This additional annealing treatment usually consists of two isothermal annealing steps. In the first step the strip is heated into the intercritical temperature region, to recrystallise the 6.1 PROCESSING ROUTES 97

a) b)

hot rolling TM rolling intercritical ag+ ag+ IC rolling annealing coiling coiling

cold rolling

thickness reduction adjustment of the thickness reduction + good surface finish microstructure + adjustment of the microstructure

Final Microstructure Ferrite Retained Austenite Bainite Martensite

Figure 6.2: Comparison of the different processing steps during the production of (a) cold rolled TRIP and (b) hot rolled steel strip

microstructure and to generate the desired volume fraction of austenite, which is usually about 30 − 50% [138]. During the formation of the intercritical austenite, carbon is rejected from the carbide rich phase into the surrounding austenite matrix. The solute carbon is completely incorporated in the austenite because of its poor solubility in ferrite [139]. Subsequently the strip is rapidly cooled down into the temperature region of the bainite transformation. During cooling, the austenite may partly transform into ferrite and the remaining austenite is hence further enriched in carbon. The cooling rate has to be high enough to prevent the formation of pearlite. During the second isothermal annealing step in the bainite temperature region, the residual austenite partly transforms into (carbide free) bainite. The residual austenite is again enriched in carbon and the carbon content is ﬁnally high enough, to bring the Ms temperature below room temperature. Finally the strip is cooled down to room temperature. The annealed strip ﬁnally exhibits the typical TRIP microstructure containing, ferrite, bainite retained austenite and mostly a little martensite, Fig. 6.2 a). The retained austenite content is usually about 5 to 15%.

6.1.2 Hot rolling

The production process of a hot rolled low alloy TRIP steel, Fig. 6.2 b), is much compacter than that involving a supplemental cold rolling and annealing process. The processing route consists of re-heating of the slab into the austenite temperature region, followed by hot rolling, cooling and coiling. The hot rolled strip already exhibits the desired multiphase microstructure after coiling, however, a close control of the process parameters is required, to guarantee the formation of the multiphase microstructure and a desirable surface quality. Special attention has to be given 98 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

to the cooling of the strip on the run-out table, as it is the crucial processing step to control the desired multiphase microstructure. Conventionally a two step cooling cycle (water / air / water) is utilised [6]. After ﬁnish rolling the strip is immediately cooled to prevent the recrystallisation process. In the range of the maximum ferrite transformation the cooling rate is reduced to form the required fraction of ferrite, which is usually about 60 − 70%. Subsequently the strip is again accelerated cooled down to the coiling temperature, which has to be in the temperature region of the bainite formation. The remaining austenite is further enriched in carbon and stabilised below the room temperature. An insufﬁcient temperature control yields undesirable and/or inhomogeneous mechanical properties over the strip length and width [140]. As a subsequent cold rolling and annealing treatment is not required, the use of hot rolled TRIP steels provides a certain potential to reduce the production costs. The usual industrially applied hot rolling strategies to control the microstructure and mechanical properties of low alloyed steels are [6]:

1. Conventional (austenitic) rolling 2. Normalising rolling (Recrystallisation controlled rolling), 3. Thermomechanical rolling (Controlled rolling)

In conventional rolling, the strip is rolled at high temperatures in the austenitic temperature region. Therefore, lower rolling forces are necessary to obtain the desired strip thickness. The final strip often requires an additional heat treatment to meet the demanded mechanical properties. With the introduction of microalloyed steels the normalising rolling and thermomechanical rolling strategy became more important. In normalising rolling, the finish rolling temperature is just above the non-recrystallisation temperature. The grain size is refined by the recrystallisation of the austenite. Additionally small amounts of Ti can be added to retard grain coarsening after each rolling pass. Even though this rolling strategy only yields a moderate improvement of the mechanical properties, it remains an interesting technique, especially for rolling mills with low admissible rolling forces. In thermomechanical rolling, the finish rolling temperature is shifted down into a temperature region, near Ar3. By the addition of microalloying elements, such as titanium and/or niobium, the non- recrystallisation temperature, Tnr, is raised, so that no recrystallisation takes places during finish rolling [6]. Thermomechanical rolling provides a superior grain refinement compared to the normalising rolling. The overall grain refinement possible by thermomechanical rolling depends on the total strain applied in the non-recrystallisation temperature range, Tnr, and the cooling rate after deformation [141]. Thermomechanical rolled low alloy steels provide a desirable combination of high strength and toughness, however with a decreased ductility [142, 143]. However, even though the industrially applied rolling strategies might bring about a certain grain refinement, none of these is suitable for the production of a hot rolled TRIP steel. The conventionally applied rolling strategies are optimised to condition the austenite prior the α/γ transformation. However, for the production of a TRIP steel it is vital to control the cooling of the strip 6.1 PROCESSING ROUTES 99

on the run-out table in order obtain the multi phase microstructure. As the unique properties of TRIP steels are obtained by the multiphase microstructure, is the potential for optimizing the mechanical properties of hot rolled TRIP steel by the conventional thermomechanical rolling strategy is only limited. The intercritical rolling seem to be a promising rolling strategy to produce a hot rolled TRIP steel with a single cooling step on the run-out table [12–15].

6.1.3 Intercritical Rolling

Intercritical rolling is defined as rolling in the temperature range of the α/γ transformation, so that the deformation of the strip is at least partly performed concurrently with the transformation from austenite to ferrite. The plastic deformation can be used to influence or even control the transformation kinetics and hence the microstructure and mechanical properties [144, 145]. An increase in strength and notch toughness for intercritical rolled hot strips has been found for example in [141–143]. This has been explained by a remarkable grain refinement due to dynamic recovery and recrystallisation [146–148]. In [149] an increase in work hardening and total elongation by intercritical deformation of low alloy TRIP steels has been observed. On the other hand, the development of a bimodal grain size distribution has been observed after intercritical rolling C-Mn steels [150], which is known to deteriorate the mechanical properties [143]. An increasing rolling reduction in the intercritical region results in austenite grain refinement, due to dynamic recrystallisation [148]. In [12] an increase in ferrite fraction with increasing rolling strain is observed. This effect is explained in terms of a higher density of nucleation sites in the deformed austenite for the subsequent transformation and an increased potency of the individual nuclei [105]. Additionally the stronger faulting within the deformed austenite grains increases the free energy and the rate of diffusion during solute partitioning [12], accelerating the carbon enrichment in the austenite. Due to the higher density of defects, the nucleation rate of bainite is increased, whereas the growth rate is decreased, as a function of the intercritical rolling reduction [12]. In [14, 15] the formation of deformation induced ferrite has been observed at total deformation ◦ greater than 50% and a deformation temperature of Ar3 + 10 C. A volume fraction of approximately 50% ferrite is reported for a total deformation of 80%. This suggest, that a deformation in the upper intercritical temperature region, near Ar3 could be used to produce a microstructure as intercritically annealed with the desired α/γ phase composition.

6.1.4 The Hypothesis

The hypothesis for the experiments, described in this chapter, is to promote the ferrite formation by a heavy deformation in upper intercritical or low austenitic temperature range, near Ar3, in 100 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

order to obtain a similar microstructure as after intercritical annealing with a desirable α/γ phase composition. Hence, the rather complicated two step cooling cycle on the run-out table could be replaced by simple single step fast cooling of the strip down to the required coiling temperature. Of course, full industrial development of hot rolled TRIP steel is very complex and beyond this thesis. However, to assist the industrial development of intercritical rolled TRIP steels, experiments using the hot deformation simulator (WUMSI) have been performed to study the microstructure development and to get a ﬁrst impression of the mechanical properties to be obtained.

6.2 Experimental

6.2.1 Material

The material used in this study was an industrially produced TRIP steel, taken from the production of ThyssenKrupp Stahl. The samples for the laboratory test were taken from an extended cropped portion of a strip after rough rolling with a thickness of about 50mm. These pieces were cooled in still air down to ambient temperature. The chemical composition is given in Table 6.1 and is typical for an Si-Al based TRIP steel. The presence of Nb allows for enhancement of the properties by HSLA like rolling schedules. After the industrial rough rolling, the material was reheated to approximately 1200 ◦C and pre-rolled to nearly the desired specimen thicknesses for the laboratory hot deformation tests and cooled down to the room temperature in still air. Subsequently the material was cut to length and ﬁnally milled to the desired specimen size of 60mm x 65mm x 33.2mm (RD x ND x TD), suitable for plane strain compression tests using hot deformation simulator WUMSI. For the deformation dilatometry tests, cylinders with a diameter of 5mm and a height of 10mm were machined parallel to the normal direction.

Steel C Si + AL Mn Nb RA-1 0.24 1.8 0.98 0.05

Table 6.1: Chemical compositions in mass% of the TRIP steel

6.2.2 Flat Compression Tests

The intercritical rolling experiments were executed using the hot deformation simulator WUMSI of the Max-Planck-Institut für Eisenforschung. The hot deformation simulator is a computer controlled servohydraulic press, equipped to simulate the deformation sequence of a conventional hot rolling mill [49, 151, 152] by ﬂat compression tests [153, 154]. The maximum upsetting force is 2.5MN and the maximum deformation rate is 150s−1. Additionally the system is equipped with an induction heating unit and different cooling units to simulate more complex temperature 6.2 EXPERIMENTAL 101

Taus Taus Taus Taus

eARat T AR eARat T AR eARat T AR

eICat T IC eICat T IC

ag+ tH tH t t 8/ 5 8/ 5 20 K/s

Temperature A TC B C D salt bath

Time

Figure 6.3: Deformation schedules of the hot deformation simulator tests (A) and the dilatometry tests (B - D)

patterns. The temperature is measured by a NiCr-Ni thermocouple, inserted into the specimen and allows to monitor the temperature until the final cooling. Additionally the punch movement and the compression force is measured. This data can be used to calculate the flow stress, k f [155]. The flat compression specimens are large enough to machine secondary specimens for tensile tests [156].

The specimen is reheated and austenitised in an induction coil. Subsequently the specimen is moved pneumatically into the deformation position. In the deformation position the specimen can be cooled with compressed air to adjust the cooling rate. Depending on the specimen size and air pressure cooling rates up to 1 − 10Ks−1 can be reached. The coiling and coil cooling is simulated by quenching and isothermal annealing for 2h in a salt bath set to the desired coiling temperature, TC.

The thermal-deformation schedule of the ﬂat compression tests is shown in Fig. 6.3 A. The spe- ◦ ◦ cimens were reheated to Taus = 1200 C. Subsequently they were cooled to TAR = 950 C and deformed with εAR = 0.6 for a simulated rough rolling. After further cooling down to the in- ◦ tercritical rolling temperature TIC = 865 − 745 C the specimens are deformed in three hits by εIC = 3x0.2 to the ﬁnal thickness of 10mm. The cooling in all these steps is forced air cooling, which yielded a t8/5 = 128s in the high temperature regime for this specimen geometry. The intercritical deformation temperature, TIC, is the main parameter varied in the experiments presented in this chapter.

To obtain a high cooling rate down to the desired coiling temperature, the coiling simulation (2h at TC) was performed using a salt bath furnace, set to the desired coiling temperature, TC = 400 − 200 ◦C. The cooling cycle from the intercritical deformation temperature down to the coiling temperature is shown in Fig. 6.4 for the highest and lowest coiling temperature. Down to a temperature of 600 ◦C the cooling rates are similar. Below this temperature they start to deviate ◦ ◦ ◦ markedly. At 500 C the cooling rate is 10 K/s for TC = 200 C and only 5 K/s for 400 C. 102 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

900

800 TC = 400 C TC = 200 C

700 20 K/s 600 5 K/s 500 0.3 K/s 400 10 K/s 1.7 K/S 300

Temperature [°C] 4.5 K/S 200 0.2 K/s 100

0 0 50 100 150 200 250 300 350 Time [s]

Figure 6.4: Cooling cycle of a deformed plane strain compression sample in a salt bath with a temperature of ◦ TC = 400 or 200 C; the cooling rates at selected temperatures are indicated

6.2.3 Deformation Dilatometry

For the determination of the transformation temperature, a Bähr 805A/D deformation dilatometer was used, which has already been described in more detail in section 3.2. The thermal- deformation schedules of the dilatometry tests are shown in Fig. 6.3 B - D. The schedule B was used to determine the CCT diagram of the TRIP steel, starting from an undeformed austenitic ◦ state. The specimens were reheated to Taus = 1200 C and annealed 3 minutes for austenitisation. The austenitisation temperature was always 1200 ◦C. Subsequently the specimens were cooled with t8/5 = 3, 7.5, 15, 30, 60, 100s respectively.

The schedule C was used for the determination of the Ar3 temperature of the TRIP steel with ◦ a prior deformation of εAR = 0.4 at TAR = 950 C. The cooling rate was set to be identical to the cooling of the deformed plane strain compression specimen. Schedule D was designed simulate the thermal path of the and to investigate the inﬂuence of a holding time, tH, after the intercritical deformation on the microstructure and the Ms temperature. The specimens were annealed for tH = 4, 8, 16s after the intercritical deformation and then cooled down by 20 K/s to room temperature.

The experimental set-up of the ﬂat compression tests requires a certain holding time between the last intercritical deformation step and the start of the salt bath quenching, because the specimen has to be taken out of the press and brought to the salt bath. The time required for this additional handling is approximately tH ≈ 10s. 6.2 EXPERIMENTAL 103

M10 2

11.5

R2 R2 a)RD b)

Figure 6.5: a) Position of the tensile specimens in the plane strain compression specimen and b) dimensions of the specimens

x [mm] 0 3 6912 15 1.0 16.4

1.8 0.8 13.1 jjLN/ = 0.2 0.4 0.6 1.6 0.6 0.8 1.0 9.9 1.2

1.4 y [mm] y/0.5t [-] 0.4 6.6

1.6 0.2 3.3 Centre of the 0.6 deformation zone 1.2 1.0 0.8 0.0 0.0 0.0 0.2 0.4 0.6 0.8 1.0 x/0.5a [-]

Figure 6.6: Normalised local strain ϕL/ϕN in a quadrant of the deformation zone of the compressed specimen in real coordinates (x,y) and coordinates normalised by the punch width, a and the specimen height, h (lines represent iso strain levels) [156] 104 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

6.2.4 Tensile Tests

The tensile test specimens were taken from the flat compression specimens parallel to the simulated rolling direction (RD). Four tensile specimens were taken from each compressed sample, as is shown in Fig. 6.5 a). In taking the samples for the final mechanical testing, it should be realised, that the strain distribution over the large plane strain compression sample is intrinsi- cally inhomogeneous. The local strain distribution after deformation, ϕL, in the flat compression specimens is a function of the initial specimen thickness, t0, the punch width, a, and the nominal strain, ϕN. An example of the strain distribution is shown in Fig. 6.6. Areas with a nominal strain of ϕL/ϕN = 1 ± 0.2 are shaded in grey.

The distance, xk, of the area with nominal strain along the centre plane (x-axis) of the deformation zone was xk ≈ 11.5mm for the current experiments. The inhomogeneity of the strain distribution has been utilised to investigate the inﬂuence of the total strain on the ﬁnal mechanical properties, by using tensile specimens notched at the centre (ϕL > ϕN) and 11.5mm of centre (ϕL ≈ ϕN), as illustrated in Fig. 6.5 b). The tensile tests were executed using a Zwick tensile test machine, equipped with a device to measure the strain in length direction. The test equipment was chosen according to DIN EN 10002-1 [111].

6.2.5 Determination of the Fraction Retained Austenite

The fraction of retained austenite was determined by image analysis of tint etched samples and XRD. The samples for the image analysis were etched using LePera’s tint etching method [157]. First, the samples were ground and polished down to the use of 1µm diamond paste. After thorough cleaning the specimens were immersed for 10 to 20s in the freshly mixed etchant. The etchant is a 1:1 mixture of two solutions with the compositions given in Table 6.2. To obtain comparable colouring of the specimens it is important to use a freshly mixed etchant for each specimen.

Reagent 1 Reagent 2 1g Na2S2O5 4g picric acid 100ml distilled H2O 100ml ethanol

Table 6.2: Composition of the solutions for the LePera tint etchant [157]

Fig. 6.7 a) shows the microstructure, as revealed by the LePera etching method. Ferrite appears blue, bainite brown and austenite/martensite white [157]. Five micrographs were taken for each specimen around the point N, shown in Fig. 6.6. A typical area used for image analysis is marked in Fig. 6.7 a) by the dashed rectangular. Fig. 6.7 b) shows a resulting picture of the quantitative image analysis. The areas containing austenite/martensite are highlighted in green. On the basis of these pictures the imaging software calculates the fraction of the retained austenite, fRA. 6.2 EXPERIMENTAL 105

a) b)

Figure 6.7: a) Typical micrograph of the microstructure tint etched using the LePera method, ferrite appears blue, bainite brown and austenite / martensite white; the dashed rectangular denotes the area used for the image analysis and b) shows the resulting picture of the quantitative image analysis software in which the austenite / martensite areas are highlighted in green.

For the XRD analysis, the Bruker-AXS x-ray diffractometer of the Max-Planck-Institut für Ei- senforschung was used. The incomplete {111}γ and {110}α pole figures were measured and the net integrate pole figure intensities were used to calculate the fraction retained austenite [158– 160]. The texture of the material was taken into account by using the net integrated intensities of the pole figures [161, 162]. The fraction retained austenite was calculated using the equation

2 2 I111γF110αVγCellH110α fRA = 2 2 2 2 , (6.1) I110αF111γVαCellH111γ + I111γF110αVγCellH110α

in which Fhkl denotes the structure amplitude, VxCell the unit-cell volume and Hhkl deﬁnes the number of all symmetry-equivalent lattice planes [163]. The parameter values used for the calculations are summarised in Table 6.3.

2 Fhkl [-] Hhkl [-] VxCell [Å ] {110}α 36.91 12 23.39 {111}γ 74.70 8 46.66

Table 6.3: Constants used for the XRD analysis [163] 106 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

6.3 Results and Discussion

6.3.1 Deformation Dilatometry Tests

First of all, deformation dilatometry tests were used to determine the CCT diagram of the TRIP steel RA-1, shown in Fig. 6.8 a). The microstructure of the samples cooled with t8/5 ≥ 30s consists of ferrite, pearlite and bainite, Fig. 6.8 b). Cooling with t8/5 ≥ 7.5s leads to microstructure consisting of ferrite, bainite and martensite, Fig. 6.8 c). Faster cooling leads to a pure martensitic microstructure, Fig. 6.8 d). The Ar3 temperature for the ﬂat compression experiments has been determined by a deformation dilatometry test with an austenitic deformation and a subsequent ◦ t8/5 time of 128s, according to Fig. 6.3 B. The Ar3 temperature was found to be 805 C. Fig. 6.9 a) - c) shows the microstructure of the intercritical deformed specimens. It is obvious that the additional holding time promotes the ferrite formation. The microstructure of the specimen with the shortest holding time, Fig. 6.9 a) consists of large martensite islands, transformed from the austenite retained after intercritical deformation and holding. After holding for 8s the fraction ferrite is increased and the martensite islands have become smaller and more homogeneously distributed. The fraction ferrite is once more increased for a holding time of 16s, however, to a smaller extent. The martensite islands are even more homogeneously distributed. The increasing amount of ferrite is also reﬂected by the hardness of the specimens, as shown in Fig. 6.10. The hardness drops from 382HV5 for a holding time of 4s down to 375HV5 for a holding time of 16s.

1200 t8/ 5 = 100s

t8/ 5 = 100s 1000

t8/ 5 = 15s

800 b) FS t = 15s PS 8/ 5 BS

600 t=3s8/ 5

BF MS 400 Temperature [°C] c) t=3s M 8/ 5 200 F 347 496 448 297 260 240 0 100 101 102 103 a) Time [s] d)

Figure 6.8: a) CCT diagram of the TRIP steel RA-1; b) – c) microstructure of samples cooled with t8/5 = 100,15,3s respectively 6.3 RESULTS AND DISCUSSION 107

a) b) c)

Figure 6.9: Inﬂuence of an additional holding time, tH , of a) 4, b) 8 and c) 16s after intercritical deformation on the ◦ microstructure (TIC = 805 C; εIC = 0.6 T˙ = 20 K/s)

400

390

380

370

360

350

340

M [°C], HV5 [-] 330

320

310 MS HV5 300 02468101214161820

tIR [s]

Figure 6.10: Inﬂuence of an additional holding time, tH , of S) 4, b) 8 and c) 16s after intercritical deformation on ◦ the MS temperature and the hardness (TIC = 805 C; εIC = 0.6 T˙ = 20 K/s) 108 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

The intercritical deformation near the Ar3 temperature accelerates the transformation, as has been already mentioned in section 3.2.2. The transformation preferentially starts in regions with a lo- cally lower Mn concentration. This might explain the rather large martensite islands in the case of the short holding time of 4s. In the case of the longer holding times the austenite islands are split up by the proceeding transformation and the remaining austenite is further enriched in carbon [12]. The carbon enrichment of the austenite results in the reduction of the MS temperature with increasing holding time, tH, as can be seen Fig. 6.10.

On the basis of the Ms temperatures determined in the deformation dilatometry test, the carbon content of the remaining austenite after the intercritical deformation can be estimated using the well established Andrews equation [144, 164]

◦ Ms = 539 C − 423wC − 30.4wMn − 17.7wNi − 12.1wCr − 7.5wMo, (6.2) in which wi denotes the content in mass% of the alloying element i. Susequently, the phase fractions after intercritical defomation and annealing can be estimated using the calculated carbon content of remaining austenite and ferrite. It is assumed for simpliﬁcation, that the carbon diffu- ses completely into the remaining austenite and the maximum solubility of 0.02mass% carbon in the ferrite is reached [144]. The resulting carbon content and the corresponding phase fractions are given in Table 6.4.

tH [s] 4 8 16 %Cγ 0.36 0.39 0.41 fα 0.35 0.40 0.44 fγ 0.65 0.60 0.56

Table 6.4: Carbon content of the remaining austenite and phase fractions after intercritical deformation and annea- ◦ ling (TIC = 805 C, εIC = 0.6, T˙ = 20 K/s)

The carbon content of the remaining austenite slightly increases from 0.36mass% for an intercritical annealing time of 4s up to 0.41mass% for an annealing time of 16s. The corresponding fractions of ferrite are 35% and 44% respectively, which correlates to the microstructural observations in Fig. 6.9 a) - c). It can be concluded, that this thermal-deformation schedule is in principle appropriate to produce a microstructure as after intercritical annealing, utilising a two step cooling cycle.

6.3.2 Flat Compression Tests

6.3.2.1 Inﬂuence of the Coiling Temperature

A coiling temperature of 400 ◦C leads to a microstructure containing ferrite, pearlite, bainite and retained austenite with a mean grain size of 4.5µm, Fig. 6.11 a). The formation of pearlite is 6.3 RESULTS AND DISCUSSION 109 undesirable, as the carbon precipitated is no longer available for the necessary stabilisation of the remaining austenite. This also explains the low fraction of bainite visible in microstructure. The less stable austenite already transformed to pearlite in the beginning of the cooling cycle and the remaining austenite was retained. The pearlite formation is favoured by the quite slow cooling rate in the salt bath between 600 and 500 ◦C at a salt bath temperature of 400 ◦C, Fig. 6.4. A coiling temperature of 300 ◦C is sufﬁcient to prevent pearlite formation and the ﬁnal microstructure consists of ferrite, bainite and retained austenite with a mean grain size of 5.8µm, Fig. 6.11 b). A further reduction of the coiling temperature to 200 ◦C increases the amount of retained austenite and martensite, which cannot be distinguished optically from each other. The ◦ mean grain size is approximately 4.9µm. From the MS temperature of approximately 345 C determined by the dilatometer test, Fig. 6.10, it can be argued, that it is mainly the fraction of martensite which is increased and not the fraction of retained austenite. ◦ On the basis of these results a coiling temperature of TC = 300 C has been chosen for the further experiments, as it represents a compromise to avoid or at least minimise the formation of pearlite on the on hand and martensite on the other hand.

6.3.2.2 Inﬂuence of the Intercritical Deformation Temperature

Fig. 6.12 a) - g) show the micrographs of the samples deformed at a temperature of TIC = 865 − 745 ◦C. The microstructure of the samples is very homogeneous and also the distribution of retained austenite and bainite is similar. The specimens exhibit a mean grain size of approximately 3.9µm. The microstructure of all the specimens is typical for TRIP steels, consisting of ferrite, bainite, retained austenite and some martensite. Basically two main morphologies of the retained austenite can be observed, retained austenite island surrounded by ferrite and mixtures of retained austenite and bainite surrounded by ferrite. The latter one can have either a lath or a washy appearance. In the case of the washy type austenite islands one could assume that these areas partly transformed to bainite during coiling, leading to a ﬁne bainite structure within the austenite grains. This would require that these austenite islands had the lowest carbon concentration after cooling down to the coiling temperature. Fig. 6.12 h) shows a reference sample with the same thermal history as the other samples, only the intercritical deformation was εIC = 0. To obtain the same specimen geometry the total deformation of ε = 2x0.6 was applied at 1100 and 950 ◦C respectively. The microstructure of the reference sample lacks the typical TRIP microstructure and consists only of ferrite and bainite.

Fig. 6.13 shows the fraction retained austenite, fRA, of the samples, determined by quantitative image analysis and XRD analysis. Even though the absolute values determined by the two methods differ, the qualitative trend is similar. The samples with the highest deformation temperatures of 865 and 855 ◦C reveal a fraction retained austenite of approximately 10% determined by the XRD measurements. The other samples exhibit a fraction retained austenite of approximately 8%. 110 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

TC = 400 °C

TC = 300 °C

TC = 200 °C

◦ Figure 6.11: Inﬂuence of the coiling temperature on the microstructure (TIC = 805 C; εIC = 0.6) 6.3 RESULTS AND DISCUSSION 111

TIC = 865 °C TIC = 845 °C

a) b)

TIC = 825 °C TIC = 805 °C

c) d)

TIC = 785 °C TIC = 765 °C

e) f)

TIC = 745 °C eIC= 0; T IC = 805 °C

g) h)

Figure 6.12: Inﬂuence of the intercritical deformation temperature on the microstructure a) - g) εIC = 0.6 h) εIC = 0 112 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS

12 quant. image analysis x-ray analysis

f [%]

0 720 740 760 780 800 820 840 860 880 Temperature [°C]

Figure 6.13: Inﬂuence of the intercritical deformation temperature on the fraction of retained austenite, determined by quantitative image analysis and x-ray phase analysis

The fractions determined by the quantitative image analysis are lower and range between 4 and 6%. A possible explanation for the lower fractions determined by the image analysis compared to the x-ray analysis might be the different morpholigies of the retained austenite islands. The image analysis can easily identify the retained austenite islands surrounded by ferrite. In the case of the lath or washy mixture of retained austenite and bainite it is more difﬁcult to distinguish between bainite and austenite, which might result in the discrepancy of the results. The microstructure obtained by the proposed deformation schedule, shown in Fig. 6.3 C, is comparable to that of a conventionally produced TRIP steel. The fraction retained austenite is with approximately 8 to 10% at a dersirably high value, comparable to that of inducstrially produced TRIP steels [165]. The fractions determind by the quantitative image analysis were approximately 4% lower. This has been atributed to the different morpologies of the retained austenite observed in the microstructure. The retained austenite surrounded by ferrite can be easily distinguished by the software, whereas the lath or whashy mixture of retained austenite and bainite is can be hardly sepperated. The microstructure was found to be independent of the intercritical deformation temperature. An additional intercritical deformation prior cooling and coiling of the strip is required to obtain the multiphase microstructure by the proposed schedule, as the reference sample, produced without the additional intercritical deformation did not reveal the typical TRIP microstructure. This contradicts the results presented in [144]. In this work it has been found that with increasing deformation the amount of ferrite increases, while the amount of bainite decreases. At an intercritical deformation of ε = 0.4 the bainite formation was already completely suppressed and the microstructure lacked bainite. A possible explanation for the different microstructural development might be the different chemical compositions of the steels. In [144] an Al based 6.3 RESULTS AND DISCUSSION 113

1100

Rm, jN Rm, j1. 5N Rp0. 2, jN Rp0. 2, j1. 5N 1000

900

800

p0.2 m

R700 , R [MPa]

600

500 700 750 800 850 900

TIC [°C]

Figure 6.14: Inﬂuence of the intercritical deformation temperature on the tensile strength, Rp0.2, and ultimate tensile strength, Rm

TRIP steel without the addition of Nb was used, whereas the current steel RA-1 is Si-Al based with the addition of Nb.

6.3.3 Mechanical Properties

The tensile strength, Rp0.2 and the ultimate tensile strength, Rm were found to be independent of the intercritical deformation temperature, as can be seen in Fig. 6.14. A mean value of Rp0.2 ≈ 658MPa and Rm ≈ 903MPa has been measured for the sample with a nominal intercritical strain. The strength values increase with increasing strain in the intercritical region, as the tensile specimens notched in the centre relatively to the ﬂat compression sample reveal a tensile strength approximately 60MPa and an ultimate tensile strength 30MPa higher than those deformed with the nominal strain. The tensile strength and ultimate tensile strength of the TRIP steels produced were slightly higher than those of a convention TRIP steel, e.g RA-W 700 with a minimum Rp0.2 ≈ 495MPa and Rm ≈ 700MPa [166]. The slightly higher strength values are caused by the geometry of the tensile specimens, as the tensile strength of a notched specimen is, in the case of a ductile material behaviour, larger than that of a notch free specimen [167].

6.3.4 Conclusion

It can be concluded, that the proposed rolling schedule can be used to produce TRIP steel, without the necessity of a rather complicated two step cooling cycle. Next to this, the process seems to 114 CHAPTER 6 TOWARDS A NEW HOT ROLLING STRATEGY FOR TRIP STEELS be quite stable in terms of microstructure and mechanical properties. The mechanical properties were in terms of the strength values comparable to that of the conventional TRIP steel RA-W 700 [166]. The strain data can not be compared to those of a conventional TRIP steel due to the notched tensile test specimens used for the present tensile tests. The results obtained from the plane strain compression experiments suggest that the proposed rolling schedule might indeed bring about an improved process control, due to the simpliﬁcation of the cooling cycle. However, further experiments are required to investigate the complete mechanical properties. Summary "Advanced hot rolling strategies for IF and TRIP steels"

Steel producers are forced to reduce the production cost on the one hand and increase the performances of the products on the other hand, to maintain or even increase their market share. A likely method to save on production costs is to cut down the rather long production chain for conventional cold rolled strip, and to substitute certain cold rolled steel grades by hot rolled steel. Certainly this can rarely be 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. In chapter 1 the technological and qualitative requirements for a substitution of cold strip by hot strip are summarised. The aim of this thesis is to contribute to the understanding of the microstructure and texture development of deep-drawing steels during ferritic rolling, subsequent cold rolling and annealing on the one hand and to appraise the possible improvement of the mechanical and deep-drawing properties on the other hand. On the basis of these results guidelines for the industrial production of ferritic rolled hot strip should be formulated. Furthermore, a new hot rolling strategy for the production of hot rolled low alloy TRIP steel is proposed and the development of the microstructure and the mechanical properties to be obtained are investigated in a set plane strain compression test using a hot deformation simulator. Chapter 2 gives a general introduction to the properties of deep-drawing steels and the quanti- fication of the deep-drawability using the r, rm, ∆r and ∆rmax values. A high mean r-value, rm, together with a low planar anisotropy, ∆r, yields an optimum deep-drawing performance with a high allowable drawing ratio and a low earing tendency. To obtain the above mentioned properties, an adapted chemical composition and a strong final γ-fibre texture is required. The most important texture components for deep-drawing bcc steels are the α-fibre and γ-fibre. In this context the macro texture measurement and the texture representation using the Euler-space are explained in detail. After a short description of the conventional production process of deep-drawing steels, focussing on the limitations of the conventional hot rolling process, the ferritic rolling strategy is introduced and relevant recent literature is reviewed. In the ferritic rolling strategy the finish rolling temperatures are shifted down into the fully ferritic temperature region. The reduction of the temperature below the α/γ transformation allows for the production of thin hot strips (ts < 2mm) and even ultra thin hot strip (ts < 1mm). Additionally to this it becomes possible to produce such

115 116 SUMMARY "Advanced hot rolling strategies for IF and TRIP steels" thin hot strips with adequate deep-drawing properties. The ferritic rolling strategy allows 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 in the ferritic region and coiled at a sufficiently high coiling temperature to ensure a full recrystallisation of the strip immediately after coiling. The "hard" hot strip is rolled and coiled at lower temperatures in the ferritic region, so that a complete recrystallisation in the coil does 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. Additionally, the ferritic hot strip grades can be used to produce "cold" strip. In this case, the ferritic rolled strips are used as primary hot strips for a subsequent cold rolling and the initial microstructure and texture are imparted to the cold strip. The development of a desirable γ-fibre texture is strongly influenced by the steel chemistry, especially by the amount of carbon in solution. Concerning the rolling process itself, the lubrication in the roll gap is of a large importance. The last two sections of chapter 2 deal with these two topics. The solute carbon generally deteriorates the texture development in deep-drawing steels. On the basis of the literature data it is concluded, that it is not the solute carbon per se which causes the detrimental recrystallisation texture in ferritic rolled LC steel but the absence of so called in-grain shear bands. The absence of the in-grain shear bands is caused by the occurrence of the dynamic strain ageing (DSA) behaviour, which in turn depends on the content of solute carbon. A homogeneous γ-fibre texture through the strip thickness is essential for the production of a strip with excellent deep-drawing properties. However, increasing friction in the roll gap leads to an enhancement of an unfavourable shear texture near the surface, which deteriorates the deep-drawing properties. From the reviewed literature it is evident that the production of deep-drawable ferritic rolled hot strip necessitates proper lubrication of the roll gap, at least in the last finishing stand(s). In chapter 3 the experimental procedure is explained in more detail. Two industrially produced IF steels, stabilised with Ti and TiNb, were used for the rolling experiments. The test materials were pre-rolled and milled to the desired initial hot strip thicknesses. In order to avoid uncontrolled rolling in the α/γ temperature region, deformation dilatometry tests were performed to determine the CCT diagrams of the steels investigated. The minimum austenitic rolling temperature was determined to be in the range of 900 − 910 ◦C and accordingly, the maximum ferritic rolling temperature was chosen to be 860 ◦C. The hot and warm rolling experiments were conducted on a laboratory hot rolling mill in reversing mode. To provide a proper lubrication during rolling, a continuous lubrication system with three sponges was installed to the upper and lower roll. A set of four conventional hot rolling lubricants were tested in the relevant ferritic rolling temperature range. The lubricant yielding the lowest rolling forces was chosen for all hot and warm rolling experiments. It contained approximately 35% high pressure additives, 50% fatty matter and 15% mineral oil. After hot and warm rolling the specimens were pickeled in an industrial pickling solution. The "cold" strip specimens were subsequently cold rolled in 4−8 rolling passes. The "hard" hot strip and the "cold strip", both required a supplemental recrystallisation annealing to provide the desired deep-drawability. The batch annealing was simulated in a muffle type furnace. The continous annealing was simulated using a salt bath furnace. In a series of experiments the texture and mechanical properties were compared after annealing in a salt bath SUMMARY "Advanced hot rolling strategies for IF and TRIP steels" 117

and in a special annealing simulator, the so-called "RHESCA"-simulator.

The results of the "soft" and "hard" hot strip are presented and discussed in chapter 4. A sufficiently high coiling temperature is crucial for the production of a "soft" hot strip, since a complete recrystallisation has to be obtained in the coil. The annealing pre-tests were performed after ferritic rolling at 710 ◦C in order to determine the required coiling temperature. The recrystallised fraction of the specimens was determined by optical microscopy using a point counting method. A finish rolling temperature of 710 ◦C together with a coiling temperature of 670 ◦C has been found to be sufficient to ensure a complete recrystallisation in the case of the IF-Ti steel. It was impossible to obtain a complete recrystallisation in the coil for the IF-TiNb steel. From the recrystallisation data of the IF-Ti steel the mean activation energy has been calculated to −1 be Qrex ≈ 710kJmol , higher than the values reported in the literature. The high Qrex value indicates that the process proceeds in a narrow temperature domain. A ferritic rolling strain of εFR = 1.2 has been found to be sufficient to lead to a "soft" hot strip with a desirable {111} texture. The "soft" hot strips exhibit the 0.2% proof strength of approximately Rp0.2 ≈ 90MPa, the tensile strength of about Rm ≈ 260MPa and the ultimate elongation of A50 ≈ 40%, regardless whether the total strain was distributed over 1, 2 or 3 rolling passes. The r-values r0◦ ≈ 1.8, r45◦ ≈ 1 and r90◦ ≈ 1.6 are at a rather high level and fulfil the requirements of conventional cold strip grade DC 04. The mechanical properties of the "hard" hot strip are comparable to those of the "soft" hot strip. The r-values of the "hard" hot strips were slightly better with r90◦ ≈ 1.8 and thus correspond to the cold strip grade DC 05. The reduction of the coiling temperature from 550 to 400 ◦C and the "direct annealing" of the "hard" hot strip brought about no improvement of the deep-drawing properties. The good deep-drawing properties of the "soft" and "hard" hot strip were directly correlated to the strong γ-fibre texture of the annealed specimens.

On the basis of the results recommendations for the industrial production of "soft" and "hard" hot strips have been formulated. Two different hot strip mill layouts/modifications are mentioned and discussed. The first one is to install an additional coiler closely behind the last rolling stand to minimise the temperature losses on the run out-table in order to meet the required rolling and coiling temperature. To maintain the overall throughput of the hot rolling mill, an additional cooling section should be installed before the finishing train. Next to this an effective roll gap lubrication system is essential, at least in the last rolling stands. The second possibility is the "pony mill" concept. In this case a single high reduction stand, equipped with roll gap lubrication, a coiler and an induction heating device, at the entry and exit side of the mill respectively, is built next to the conventional hot strip mill. The pony mill allows for the production of a "soft" hot strip, without diminishing the production throughput of the conventional hot rolling mill, because the hot strip is only rolled to an intermediate thickness on the conventional hot rolling mill. The hot coil is subsequently transferred to the pony mill and finish rolled in the ferritic temperature region.

The production of the "hard" hot strip does not require the close coupled coiler, because the coiling temperature is comparable to that of a conventional hot strip. However, a roll gap lubrication system and the additional cooling section are still required. Due to the fact, that the "hard" hot strips require a supplemental annealing, they are particularly suited for the direct hot strip 118 SUMMARY "Advanced hot rolling strategies for IF and TRIP steels"

galvanising process in coupled pickling and hot dip galvanising line. In chapter 5 the results of the "cold" strip grade are presented and discussed. The "cold" strip grades are produced from ferritic rolled hot strips. It is the aim to impart the initial microstructure and texture of the ferritic rolled hot strips to the cold strip in order to improve the final deep- drawing properties and/or to reduce the necessary cold reduction. The initial microstructure of the "soft" hot strip exhibits a fully recrystallised microstructure with a mean grain size of approximately 40 − 50µm and a rather strong {111} texture with a maximum near {554}h225i. The "hard" hot strip exhibits a strained microstructure and a typical rolling texture with a complete γ-fibre and a partial α-fibre with a maximum near {112}h110i. At a cold reduction of 55% the microstructure of the initially "soft" hot strip exhibits relatively thick layers, with thicknesses up to 60µm, elongated along the rolling direction. They appear either smooth or creased. After a cold reduction of about 80% the layered structure appears to be a little coarser. The creased layers carry strain inhomogeneities, that may be referred to in-grain shear bands, which promote a favourable texture developement. The initially "hard" hot strip reveals for a cold reduction of approximately 55% also smooth and creased layers containing in-grain shear bands. The thickness of the layers varies from approximately 20 to 50µm. In the strip rolled with about 80% the thickness of these layers is only about 5 − 10µm. The γ-fibre texture of the initially "soft" hot strip is weakened for a low cold reduction of 34%. Up to a cold reduction of 68% the texture slightly intensifies and a maximum near {554}h225i develops with an intensity of f (g) = 5. Additionally, the typical rolling component near {112}h110i develops with an intensity of f (g) = 6, for higher cold reductions. After annealing, even for the low cold reduction of 34% a very homogeneous γ-fibre texture emerges with a maximum intensity of f (g) = 7. Up to a cold reduction of 68% the γ-fibre in the annealing texture lacks the typical maximum near {554}h225i. As soon as the {112}h110i component becomes clearly visible in the rolling texture, the {554}h225i also becomes more pronounced with a maximum intensity of f (g) = 14 in the annealing texture. The rolling texture of the initially "hard" hot strip shows a relatively intense typical cold rolling texture even for a low cold reduction of 35% with maximum intensity of f (g) = 6 near {112}h110i. The γ-fibre weakens with further reduction, whereas the {112}h110i component is strengthened. For cold reductions greater than approximatly 65% the cold rolled texture is dominated by the {112}h110i component with an intensity of f (g) = 12. After annealing, even for the lower cold reductions of 35% a typical γ-fibre texture developed with a maximum intensity of f (g) = 8 near {554}h225i. The {554}h225i component in the annealing texture is intensified with increasing cold reduction achieving intensities up to f (g) = 19 for a cold reduction of 85%. The development of the homogeneous γ-fibre texture after recrystallisation of the initially of the "soft" hot strip has been attributed to the oriented nucleation within deformed γ-fibre grains. With increasing cold reduction, the rolling texture starts to concentrate on the {112}h110i texture component and, hence, the potential role of the oriented growth is increased, leading to a visible peak near {554}h225i in the recrystallisation texture. In the case of the initially "hard" hot strip the volume fraction of SUMMARY "Advanced hot rolling strategies for IF and TRIP steels" 119

the {112}h110i orientation is much higher in the rolling texture, resulting correspondingly in a very sharp {554}h225i recrystallisation texture.

The mechanical properties of the "cold" strips were found to be independent of the initial hot strip condition and of the total cold reduction. With the tensile strength of approximately Rm ≈ 290MPa, the 0.2% proof strength of Rp0.2 ≈ 90MPa and the ultimate strain of A50 ≈ 40% they fulfil the requirements of cold strip grade DC 06. The r-values, on the contrary, exhibit a strong dependence on the kind of initial hot strip and the total cold reduction. The r-values of the initially "soft" hot strip at 0◦ and 90◦ to the rolling direction (RD) rapidly increase up to approximately 3.4 and 3.0 for cold reductions greater than 55%. The r-value at 45◦ to RD steadily inceases with increasing cold reduction. The mean r-value produced by a cold reduction of approximately 60−65% with rm ≈ 2.2−2.5 is comparable to that of a conventionally produced cold strip. However, with a slightly higher planar anisotropy with ∆r ≈ 1.5. The correlation to the texture features suggests, that a homogeneuos γ-fibre texture results in high rm-values but also in a quite high ∆r-values. The r-values at 0◦ and 90◦ to RD of the initially "hard" hot strip increase with increasing cold reduction up to 2.2 and 2.4 for a cold reduction up to approximately 60% and then slightly decrease. The r-value at 45◦ to RD steadily increases with increasing cold reduction. For higher cold reductions ≥ 75% the r-values at 45◦ are equal to or even greater than those at 0◦ to RD. Even though the absolute r-values at 0◦ and 90◦ to RD are lower than that of a conventional cold strip, the mean r-value still fulfils with rm ≥ 1.8 the requirements of cold strip grade DC 06, due to the low ∆r-value of approximately 0 to −0.25. Because of such a marked low ∆r-value a modified earing symmetry with a lower tendency to form ears can be expected. A correlation of the development of the r-values in the initially "hard" hot strip and the texture suggests, that a peaked texture with a very high intensity near {554}h225i can result in a more favourable combination of the deep-drawing properties if compared to a conventional cold strip, with good mean r-values and a very low planar anisotropy. Cupping tests with "cold" strips produced from "soft" and "hard" hot strips rolled with a reduction of approximately 80% were performed in order to get a deeper insight into the deep-drawing behaviour of these strips. The initially "soft" hot strip exhibited a classical 4-ear shape, with slightly higher ears compared to the conventional strip, whereas the initially "hard" hot strip revealed a 6-ear shape with a markedly lower ear height.

The process windows of both ferritically rolled hot strip grades ("soft" and "hard" hot strips) have been determined leading to an excellent formability meeting the DC 04 and DC 05 criteria. The production of these grades requires modiﬁcations of the standard hot rolling mill layout. Using these ferritic rolled hot strips as initial hot strips for the cold rolling process, can bring about a couple of improvements. The initial "soft" hot strip, produced using a reduced cold reduction, exhibits similar rm-values to those of a conventional cold strip, however with slightly higher ∆r-values. The initial "hard" hot strip, produced using a conventional cold reduction, exhibits slightly lower rm-values, however with excellent low ∆r-values. Both "cold" strips fulﬁl the requirements of cold strip grade DC 06.

TRIP steels are of great interest for the automotive industry because of the superior combination of a high strength together with a high ultimate elongation and their superior speciﬁc energy 120 SUMMARY "Advanced hot rolling strategies for IF and TRIP steels" absorption during rapid straining. Chapter 6 introduces the two known processing routes: the route via cold rolling and annealing and the hot strip rolling route. However, the industrially applied hot rolling strategies are not suitable for the optimisation of the process control of the hot strip TRIP steel production. The intercritical rolling strategy, utilising a heavy deformation in the upper intercritical or maybe in the low austenitic temperature range is proposed as a promising route to optimise the process control of hot rolled TRIP steel strip. In a series of deformation dilatometry tests the basic deformation temperatures were determined and a similar α/γ microstructure "as after intercritical annealing" could be produced. The plane strain compression tests showed, that the process is quite stable in terms of the microstructure and a typical TRIP type microstructure could have been produced in all cases. The mechanical properties determined, using small specimens taken from the deformed plane strain compression samples, exhibited strength values comparable to that of a conventional hot rolled TRIP steel RA-W 700. The strain data could not be compared to the conventional TRIP steel due to the notched geometry of specimens used in these tensile tests.

Alexander Elsner 3 May 2005 Samenvatting "Advanced hot rolling strategies for IF and TRIP steels" "Geavanceerde warmwals strategieën voor IF en TRIP stalen"

Staalproducenten zijn gedwongen om aan de ene kant de productiekosten te reduceren en aan de andere kant de prestaties van hun producten te verbeteren, om zo hun marktaandeel te behouden of zelfs te vergroten. Een mogelijke methode om te besparen op productiekosten is het verkorten van de behoorlijk lange productieketen voor conventioneel koudgewalst staalband en door het vervangen van bepaalde kwaliteiten koudgewalst staal door warmgewalst staal. Natuurlijk kan dit zelden worden gerealiseerd voor buitendelen, die een perfecte oppervlakteafwerking vereisen. Voor binnendelen met lagere eisen ten aanzien van de oppervlakteafwerking, zoals delen van een dragende constructie, kan het gebruik van warmgewalst band echter een kostenbesparend alternatief zijn. In hoofdstuk 1 worden de technologische vereisten en de kwaliteitseisen voor een vervanging van koudgewalst band door warmgewalst band opgesomd. Het doel van dit proefschrift is een bijdrage leveren aan het begrijpen van de microstructuur- en textuurontwikkeling van dieptrekstaal tijdens het ferritisch walsen en de aansluitende productie- processen koudwalsen en gloeien. Daarnaast worden de mechanische en dieptrekeigenschappen van het op deze manier geproduceerde dieptrekstaal geëvalueerd. Op basis van deze resultaten worden richtlijnen voor de industriële productie van ferritisch gewalst warmband geformuleerd. Tevens wordt een nieuwe warmwalsstrategie voor de productie van warmgewalst laag gelegerd TRIP-staal gepresenteerd. De ontwikkeling van de microstructuur en mechanische eigenschappen werden door het uitvoeren van stuikproeven bij vlakke vervormingstoestand met een warm deformatie simulator geanalyseerd. Hoofdstuk 2 geeft een algemene introductie in de eigenschappen van dieptrekstalen en het kwan- tiﬁceren van dieptrekbaarheid door gebruikmaking van r-, rm-, ∆r- en ∆rmax-waarden. Een hoge gemiddelde r-waarde, rm, samen met een lage planaire anisotropie, ∆r, leidt tot een optimale dieptrekbaarheid met een hoge toelaatbare dieptrekverhouding en een geringe tendens tot oorvorming. Om de bovenstaande eigenschappen te verkrijgen is een aangepaste chemische samenstelling en een sterke uiteindelijke γ-vezel textuur nodig. De belangrijkste textuurcomponenten voor dieptrek bcc stalen zijn de α-vezel en de γ-vezel. In deze context worden macrotextuurme- tingen en de weergave van textuur in de Euler-ruimte gedetailleerd beschreven. Na een korte beschrijving van de conventionele productiemethoden voor dieptrekstalen met de focus op de beperkingen van het conventionele warmwalsproces, wordt de ferritische walsstra-

121 122 SAMENVATTING "Advanced hot rolling strategies for IF and TRIP steels" tegie voorgesteld en de relevante literatuur wordt in ogenschouw genomen. Bij de ferritische walsstrategie wordt de uiteindelijke walstemperatuur verlaagd tot in het volledig ferritische temperatuurgebied. De verlaging van de temperatuur tot beneden de α/γ transformatie maakt de productie van dun warmband mogelijk (ts < 2mm) en zelfs die van zeer dun band (ts < 1mm). Daarnaast is het ook nog mogelijk om dun warmband te produceren met geschikte dieptrekeigenschappen. De ferritische walsstrategie maakt de productie van twee verschillende typen warmband mogelijk; een ’zacht’ en een ’hard’ warmband. Het ’zachte’ warmband wordt bij hogere temperaturen in het ferritische temperatuurgebied gewalst en wordt aansluitend bij een voldoende hoge haspeltemperatuur gehaspeld om zo een volledige rekristallisatie direct na het haspelen te bewerkstelligen. Het ’harde’ warmband wordt bij lagere temperaturen in het ferritische temperatuurgebied gewalst en opgehaspeld, zodat een volledige rekristallisatie in de coil uitblijft. Dat betekent dat dit type band een gerekte microstructuur laat zien na het haspelen, zodat een aansluitende gloeibehandeling vereist is om de gewenste dieptrekeigenschappen te verkrijgen. Daarnaast kan ferritisch gewalst band worden gebruikt om ’koud’ gewalst band te produceren. In dat geval wordt het ferritisch gewalste band koudgewalst en de oorspronkelijke microstructuur en textuur wordt aan het koudgewalste band doorgegeven. De ontwikkeling van en wenselijke γ-vezel textuur wordt sterk beïnvloed door de chemische samenstelling van het staal en dan vooral de hoeveelheid opgeloste koolstof. Daarnaast speelt bij het walsproces zelf de smering van de walsspleet een zeer belangrijke rol. De laatste twee delen van hoofdstuk 2 beschrijven deze twee onderwerpen. De hoeveelheid opgeloste koolstof verslechtert de textuurontwikkeling in dieptrekstalen. Op basis van gegevens uit de literatuur kan worden geconcludeerd dat niet het aandeel opgeloste koolstof de rekristallisatietextuur in ferritisch gewalst LC-staal beschadigt, maar dat het ontbreken van zogenaamde ’in-grain shear bands’ daarvoor verantwoordelijk is. De afwezigheid van in-grain shear bands wordt veroorzaakt door het optreden van dynamische rek-veroudering (DSA), die op zijn beurt door de hoeveelheid opgeloste koolstof bepaald wordt. Een homogene γ-vezel textuur over de dikte is essentieel voor de productie van band met excellente dieptrekeigenschappen. Toch leidt een toename van wrijving in de walsspleet tot een toenemende hoeveelheid aan ongewenste schuifkrachten dicht bij het oppervlak, die de dieptrekbaarheid nadelig beïnvloeden. Uit de geraadpleegde literatuur wordt duidelijk dat voor de productie van dieptrekbaar ferritisch gewalst warmband een goede smering van de walsspleet noodzakelijk is; in ieder geval in de laatste rolgestellen van een walsstraat. In hoofdstuk 3 wordt de experimentele werkwijze gedetailleerd beschreven. Twee industrieel geproduceerde IF-stalen, gestabiliseerd met Ti en TiNb, werden gebruikt voor de walsexperimenten. De testmaterialen werden voorgewalst en gefreest tot de gewenste initiële warmbanddikte. Om ongecontroleerd walsen in het α/γ temperatuurgebied te voorkomen werden deformatie- dilatometrie testen uitgevoerd om CCT-diagrammen te bepalen van de onderzochte stalen. De minimale austenitische walstemperatuur ligt rond de 900 − 910 ◦C en overeenkomstig werd de maximale ferritische walstemperatuur vastgelegd op 860 ◦C. De warm- en ferritische walsexperimenten werden uitgevoerd op een laboratorium warmwalserij in de heen en weer modus. Om de juiste smering tijdens het walsen te realiseren werd gebruik gemaakt van een continu smering- systeem bestaande uit drie sponzen, die op de bovenste en onderste walsrol gemonteerd waren. In totaal werden vier conventionele warmwalssmeermiddelen getest voor het van toepassingzijn SAMENVATTING "Geavanceerde warmwals strategieën voor IF en TRIP stalen" 123

in de ferritische walstemperatuurgebied. Het smeermiddel dat resulteerde in de geringste wals- krachten, werd geselecteerd voor alle warm- en ferritische walsexperimenten. Deze bestaat uit 35% hoge druk toevoegingen, 50% vetbestanddelen en 15% minerale olie. Na het warm en ferritisch walsen werden de monsters gebeitst met behulp van een industrieel beitsmiddel. Zo- wel het ’harde’ warmband als ook het ’koude’ warmband hadden beide een extra rekristallisatie gloeibehandeling nodig om de gewenste dieptrekeigenschappen te verkrijgen. De batch gloeibehandeling werd gesimuleerd in een muffle type oven. Het continu gloeiproces daarentegen werd door gebruikmaking van een zoutbadoven gesimuleerd. In een serie van experimenten werd de gloeibehandeling in een zoutbad met een gloeibehandeling in een speciale gloeisimulator, de zogenaamde ’RHESCA’-simulator, vergeleken. Daarbij werden zowel de textuur als ook de mechanische eigenschappen geanalyseerd. De onderzoeksresultaten voor het ’zachte’ en ’harde’ warmband worden in hoofdstuk 4 gepresenteerd en besproken. Een voldoende hoge haspeltemperatuur blijkt cruciaal te zijn voor de productie van een ’zacht’ warmband. Deze temperatuur moet hoog genoeg liggen om een volledige rekristallisatie in de coil mogelijk te maken. De eerste gloeitesten werden uitgevoerd na ferritisch walsen bij 710 ◦C om zo de benodigde haspeltemperatuur te bepalen. Het gerekris- talliseerde gedeelte in de proefstukken werd door middel van optische microscopie vastgesteld op basis van een puntentelmethode. Een ferritische walstemperatuur van 710 ◦C en een haspeltemperatuur van 670 ◦C blijkt voor het IF-Ti staal voldoende om een volledige rekristallisatie te realiseren. Het was daarentegen onmogelijk om een complete rekristallisatie in de coil met het IF-TiNb staal te realiseren. Uit de rekristallisatie gegevens van het IF-Ti staal werd de gemid- −1 delde activeringsenergie berekent op Qrex ≈ 710kJmol ; een hogere waarde dan in de literatuur wordt aangegeven. De hoge Qrex-waarde geeft aan dat het proces plaats vindt in een smal temperatuurgebied. Een ferritische deformatie van εFR = 1,2 blijkt voldoende te zijn om tot een ’zacht’ warmband met een wenselijke {111} textuur te leiden. Het ’zachte’ warmband vertoont en 0,2% rekgrens van Rp0,2 ≈ 90MPa, een treksterkte van ongeveer Rm ≈ 260MPa en een rek na breuk van A50 ≈ 40%, ongeacht of de totale deformatie werd verdeeld over 1, 2 of 3 walsre- ducties. De r-waarden r0◦ ≈ 1,8, r45◦ ≈ 1 en r90◦ ≈ 1,6 zijn van een behoorlijk hoog niveau en vervullen de eisen voor een conventioneel koudband met de kwaliteit DC 04. De mechanische eigenschappen van het ’harde’ koudband zijn vergelijkbaar met die van het ’zachte’ warmband. De r-waarden van het ’harde’ warmband zijn echter net wat beter met r90◦ ≈ 1,8 en correspon- deren dus met koudband in de DC 05 kwaliteit. Een reductie van de haspeltemperatuur van 550 naar 400 ◦C en het ’direct gloeien’ van de ’harde’ warmband brengt nauwelijks verbeteringen ten aanzien van de dieptrekeigenschappen. De goede dieptrekeigenschappen van het ’zachte’ en ’harde’ warmband houden direct verband met de sterkte van het γ-vezel textuur van de gegloeide proefstukken. Op basis van de resultaten zijn aanbevelingen voor de industriële productie van ’zacht’ en ’hard’ warmband geformuleerd. Twee verschillende warmwalsstraat layouts/modificaties worden genoemd en besproken. De eerste is de plaatsing van een extra haspel kort achter het laatste rolgestel om zo het temperatuurverlies op de uitlooptafel te minimaliseren, zodat de benodigde walsen haspeltemperatuur wordt bereikt. Om de algemene doorstroming in de warmwalserij te behouden dient een extra koelsectie te worden geïnstalleerd voor de laatste ’finish’ rolgestellen. De 124 SAMENVATTING "Advanced hot rolling strategies for IF and TRIP steels"

tweede mogelijkheid bestaat uit de installatie van een zogenaamde ’pony mill’. In dit geval wordt een enkel rolgestel dat een grote afname realiseren kan en uitgerust is met walsspleetsmering, een haspel en een inductieve verwarming aan de in- en uitgangszijde van het rolgestel, naast een conventionele warmwalsstraat gebouwd. De pony mill maakt de productie van ’zacht’ warmband mogelijk zonder daarbij de productiecapaciteit van de conventionele warmwalsstraat te beïnvloe- den, omdat het warmband alleen maar tot een bepaalde dikte wordt gewalst in de conventionele walsstraat. De warme coil wordt aansluitend overgebracht naar de ’pony mill’ en gewalst in het ferritische temperatuurgebied. Voor de productie van ’hard’ warmband is geen extra haspel nodig, omdat de haspeltemperatuur vergelijkbaar is met die van conventioneel warmband. Toch zijn ook hier een walsspleet smeersysteem en een extra koelsectie nog steeds een vereiste. Door het feit dat ’hard’ warmband een extra gloeibehandeling nodig heeft is het uitermate geschikt om direct als warmband in een gekoppelde beits- en thermische verzinkingslijn te worden verzinkt. In hoofdstuk 5 worden de resultaten van de ’koudband’-kwaliteiten gepresenteerd en besproken. Het uitgangsmateriaal voor de ’koudband’-kwaliteiten was het ferritisch gewalst warmband. Het is de bedoeling om de initiële microstructuur en textuur over te brengen van het ferritisch gewalste warmband op het koudband om zo de uiteindelijke eigenschappen met betrekking tot de dieptrekbaarheid te verbeteren en/of de benodigde reductie tijdens het koudwalsen te minimaliseren. De oorspronkelijke microstructuur van ’zacht’ warmband vertoont een volledig gerekris- talliseerde microstructuur met een gemiddelde korrelgrootte van 40 − 50µm en een behoorlijk sterke {111} textuur met een maximum rond {554}h225i. Het ’harde’ warmband toont een gerekte microstructuur en een typische walstextuur met een volledige γ-vezel en een onvolledige α-vezel met een maximum intensiteit in de buurt van {112}h110i. Bij een reductie tijdens het koudwalsen van 55% vertoont de microstructuur van het tot dan toe ’zachte’ warmband relatief dikke lagen, met een dikte tot 60µm, uitgerekt in de walsrichting. Ze zien glad of geplooid uit. Na ongeveer een reductie van 80% tijdens het koudwalsen ziet de gelaagde structuur er wat grover uit. De geplooide lagen dragen rek inhomogeniteiten, welke ook wel ’in-grain shear bands’ worden genoemd, die een wenselijke textuurontwikkeling bevorderen. Het van oorsprong ’harde’ warmband laat na ongeveer 55% reductie tijdens het koudwalsen ook gladde en geplooide lagen zien, die ’in-grain shear bands’ bevatten. De dikte van de lagen varieert van ongeveer 20 tot 50µm. In de banden, die met ongeveer 80% zijn gewalst, is de dikte van dit type lagen ongeveer 5 − 10µm. De γ-vezel textuur van het oorspronkelijke ’zachte’ warmband wordt verzwakt na een reductie van 34% tijdens het koudwalsen. Tot een reductie van 68% neemt de intensiteit van de textuur enigszins toe tot een maximum van f (g) = 5 bij {554}h225i. De typische walscomponent in de buurt van {112}h110i ontwikkelt zich tot een intensiteit van f (g) = 6 bij hogere reducties voor het koudwalsen. Na het gloeien ontstaat zelfs bij 34% reductie een zeer homogene γ-vezel texture met een maximale intensiteit van f (g) = 7. Bij 68% reductie ontbreekt voor de γ-vezel het maximum rond {554}h225i in de gloeitextuur. Op het moment dat de {112}h110i component duidelijk zichtbaar wordt in de gewalste textuur, wordt ook de {554}h225i component dominanter met een maximale intensiteit van f (g) = 14 in de gloeitextuur. SAMENVATTING "Geavanceerde warmwals strategieën voor IF en TRIP stalen" 125

De gewalste textuur van het oorspronkelijke ’harde’ warmband toont, zelfs voor een lage reductie van 35% tijdens het koudwalsen, een relatief intense typische koudwalstextuur met een maximale intensiteit van f (g) = 6 in de buurt van {112}h110i. De γ-vezel verzwakt bij verdere reductie tijdens het koudwalsen, waarbij daarentegen de {112}h110i component wordt versterkt. Voor een reductie groter dan ongeveer 65% wordt de koudgewalste textuur gedomineerd door de {112}h110i component met een intensiteit rond de f (g) = 12. Na het gloeien ontwikkelt zich zelfs bij lagere reducties rond 35% een typische γ-vezel textuur met een maximale intensiteit van f (g) = 8 in de buurt van {554}h225i. De {554}h225i component in de gegloeide textuur wordt versterkt door de toenemende reductie tijdens het koudwalsen en bereikt een intensiteit van f (g) = 19 voor een reductie van 85%. De ontwikkeling van de homogene γ-vezel textuur na rekristallisatie van het oorspronkelijke ’zachte’ warmband heeft bijgedragen tot een georiënteerde nucleusvorming binnen de vervormde γ-vezel korrel. Bij toenemende reducties tijdens het koudwalsen begint de walstextuur zich te concen- treren op de {112}h110i component en dus wordt de potentiële rol van de georiënteerde groei vergroot. Dit leidt tot een zichtbare piek in de buurt van {554}h225i in de rekristallisatietextuur. Voor het oorspronkelijke ’harde’ warmband is het volume aandeel van de {112}h110i oriëntatie veel hoger in de gewalste textuur en dat resulteert in een geprononceerde {554}h225i rekristallisatietextuur. De mechanische eigenschappen van het ’koudband’ blijken onafhankelijk van de oorspronkelijke warmbandconditie en de totale reductie tijdens het koudwalsen te zijn. Met een treksterkte van ongeveer Rm ≈ 290MPa, een 0,2% rekgrens van Rp0,2 ≈ 90MPa en een rek na breuk A50 ≈ 40% vervullen ze de eisen van de koudband kwaliteit DC06. Aan de andere kant tonen de r-waarden een sterke afhankelijkheid van de oorspronkelijke warmband variant en de totale reductie tijdens het koudwalsen. De r-waarden van het oorspronkelijke ’zachte’ warmband bij 0◦ en 90◦ ten opzichte van de walsrichting (RD) nemen snel toe tot ongeveer 3,4 en 3,0 bij een reductie tijdens het koudwalsen groter dan 55%. De r-waarde bij 45◦ tot de RD neemt langzaam toe voor een toenemende reductie. De gemiddelde r-waarde behorende bij een reductie van 60 − 65% met rm ≈ 2,2 − 2,5 is vergelijkbaar met die van conventioneel geproduceerd koudband. Er wordt echter een licht hogere planaire anisotropie waarde van ∆r ≈ 1,5 bereikt. De correlatie tussen textuureigenschappen en de dieptrekbaarheid suggereert, dat een homogene γ-vezel textuur re- ◦ ◦ sulteert in hoge rm-waarden, maar ook in hogere ∆r-waarden. De r-waarden voor 0 en 90 ten opzichte van de RD van het oorspronkelijke ’harde’ warmband neemt toe bij toenemende reductie tijdens het koudwalsen tot een waarde van ongeveer 2,2 en 2,4 voor een reductie van rond 60%. Voor hogere reducties ≥ 75% varieert de r-waarde bij 45◦ tussen gelijk of groter dan bij 0◦ ten opzichte van de RD. Zelfs al zijn de r-waarden bij 0◦ en 90◦ absoluut kleiner dan bij conventioneel koudband, toch voldoet de gemiddelde r-waarde nog steeds met rm ≥ 1,8 aan de eisen van koudband met de kwaliteit DC 06, door de lage ∆r-waarde van rond 0 tot −0,25. Door deze lage ∆r-waarde wordt een veranderde oorvorming symmetrie met een lagere tendens tot oorvorming verwacht. Een correlatie tussen de ontwikkeling van de r-waarden in de oorspronkelijk ’harde’ warmband en de textuur doet vermoeden dat een textuur met een erg hoge peak intensiteit bij {554}h225i kan leiden tot een meer wenselijke combinatie van de dieptrekeigenschappen, met een goede gemiddelde r-waarden en een erg lage planaire anisotropie, ten opzichte van een 126 SAMENVATTING "Advanced hot rolling strategies for IF and TRIP steels" conventioneel koudband. Dieptrektesten met ’koudband’ geproduceerd van ’zacht’ en ’hard’ warmband gewalst met een reductie van ongeveer 80% werden uitgevoerd om meer inzicht te krijgen in het dieptrekgedrag van het materiaal. Het oorspronkelijke ’zachte’ warmband toont een klassieke 4-oors vorm, met licht hogere oren in vergelijking tot het conventionele materiaal. Het oorspronkelijk ’harde’ warmband laat echter een 6-oors vorm met een opvallend lagere oorhoogte zien. Het procesvenster dat voor beide kwaliteiten ferritisch gewalste warmband (’zacht’ en ’hard’ warmband) is ontwikkeld, resulteert in een zeer goede vervormbaarheid, vergelijkbaar met de DC 04 en DC 05 kwaliteit. De productie van deze kwaliteiten eist echter modiﬁcaties aan de stan- daard layout van warmwalsstraten. Door gebruik te maken van deze ferritsch gewalste warmban- den als initieel warmband voor het koudwalsproces kunnen enkele voordelen worden behaald. Het oorspronkelijk ’zachte’ warmband, geproduceerd met een verminderde reductie tijdens het koudwalsen vertoont gelijke rm-waarden als die voor conventioneel koudband, echter met licht verhoogde ∆r-waarden. Het oorspronkelijke ’harde’ warmband, geproduceerd door middel van een conventionele reductie tijdens het koudwalsen, toont enigszins lagere rm-waarden, maar wel uitstekende lage ∆r-waarden. Beide typen koudband vervullen de eisen ten aanzien van koudband in de kwaliteit DC 06. TRIP-stalen zijn zeer interessant voor de automobielindustrie, omdat ze een superieure combinatie van hoge sterkte, hoge rek na breuk tonen en daarnaast interessant zijn vanwege de superieure speciﬁeke energie absorptie tijdens snelle rekbelastingen. Hoofdstuk 6 introduceert de twee be- kende productieroutes: de route bestaande uit koudwalsen en gloeien en de warmwalsroute. De toegepaste industriële warmwalsstrategieën zijn echter niet geschikt voor de optimalisatie van de procesbesturing voor TRIP-warmband. De intercritische walsstrategie, die gebruik maakt van hoge deformatie in het boven intercritisch of misschien in het lage austenitische temperatuurgebied wordt voorgesteld als een veelbelovende route ter optimalisatie van het besturingsproces voor warmgewalst TRIP-staalband. In een serie van deformatie-dilatometrie experimenten is de deformatietemperatuur bepaald en kon een vergelijkbare α/γ microstructuur ’als na intercritisch gloeien’ worden geproduceerd. De vlakke rek stuikproeven toonden aan dat het proces behoorlijk stabiel is met betrekking tot de microstructuur. Een typische TRIP-microstructuur kon in alle gevallen worden gerealiseerd. De mechanische eigenschappen zoals vastgesteld aan de hand van proefstukken genomen van de gedeformeerde proefstukken laten waardes vergelijkbaar met die van conventioneel warmgewalst TRIP-staal RA-W 700 zien. De rekwaarden konden niet met de conventionele waarden voor TRIP-staal worden vergeleken vanwege de ingekeepte vorm van de proefstukken die voor deze trekproeven worden gebruikt.

Alexander Elsner 3 mei 2005 References

[1] R.C. Hudd. Processing - Cold Working and Annealing. In R.W. Cahn, P. Haasen, and E.J. Kramer, editors, Materials Science and Technology, volume 7, chapter 6, pages 220–284. VCH, Weinheim, New York, Basel, Cambridge, 1992.

[2] J.C. Charbonnier, H. Herlitz, C.J. Kearton, M. Mirabilec, and R. Steffen. Technology Road Map to Determine the Research Priorities of the European Steel Industry. Technical report, EUROFER, Brussels (Belgium), 1999.

[3] Eurofer. Annual report 2003. http://www.eurofer.org, Januar 2004.

[4] W. Bleck, K. Köhler, L. Meyer, and C. Preisendanz. Warmgewalzte Stähle mit guter Kal- tumformbarkeit für Personekraftwagen und Nutzfahrzeuge. Thyssen Technische Berichte, (1):43–52, 1991.

[5] L. Meyer. Grundlagen der Werkstoffentwicklung durch Verknüpfung thermischer und mechanischer Vorgänge beim Warmumformen. Stahl u. Eisen, 101(7):83–91, 1981.

[6] T. Heller, B. Engl, G. Stich, and G. Thiemann. Thermomechanical Rolling of Hot-Rolled Multiphase Steels. Thermomechanical Processing of Steels, 2:438–445, 2000.

[7] C. Lang, G. Stich, and L. Meyer. Verbesserung der Eigenschaften von Warmbreitband aus weichem unlegiertem Stahl. Stahl u. Eisen, 106(3):122–128, 1986.

[8] A. Tomitz. Optimierung der Prozeßparameter zur Erhöhung der Produktqualität von Tiefziehstählen beim Warmwalzen im Ferritgebiet. PhD thesis, RWTH Aachen, Aachen (Germany), 1999.

[9] J.C. Herman and V. Leroy. Ferritic Hot Rolling and the Potential for Hot Rolled and Cold Rolled Products. In The Future of Flat Rolled Steel Production, pages 1–10, Chicago (USA), June 1995.

[10] Ph. Harlet, F. Beco, P. Cantineaux, D. Bouquegneau, P. Messien, and J.C. Herman. New Soft Steel Grades Produced by Ferritic Rolling at Cockerill Sambre. In G. Tither R. As- fahant, editor, Int. Symposium on Low-Carbon Steels for the 90’s, pages 389–396. The Minerals, Metals and Materials Society, 1993.

127 128 REFERENCES

[11] J.C. Herman, P. Cantineaux, and J.-M. Langlais. Future Prospects of Ferritic Hot-Rolled Thin Strip Steel: A new "Low Cost" approach for steel sheet production. Steel World, pages 48–54, February 1992.

[12] A. Basuki and E. Aernoudt. Inﬂuence of rolling of TRIP steel in the intercritical region on the stability of retained austenite. Journal of Materials Processing Technology, 89-90:37– 43, 1999.

[13] A. Basuki and E. Aernoudt. Effect of deformation in the intercritical area on the grain reﬁnement of retained austenite of 0.4 TRIP Steel. Scripta Materialia, 40(9):1003–1008, 1999.

[14] S.C. Hong and K.S. Lee. Inﬂuence of deformation induced ferrite transformation on grain reﬁnement of dual phase steel. Materials Science and Engineering A, 323:148–159, 2002.

[15] S.C. Hong, S.H. Lim, K.J. Lee, D.H. Shin, and K.S. Lee. Effect of undercooling of Auste- nite on Strain Induced Ferrite Transformation behavior. ISIJ International, 43(3):394–399, 2003.

[16] C.J. Barrett and B. Wilshire. The production of ferritically hot rolled interstitial-free steel on a modern hot strip mill. Journal of Materials Processing Technology, 122:56–62, 2002.

[17] S. Godha, K. Watanabe, and Y. Hashimoto. Effects of the Intercritical Rolling on Structure and Properties of Low Carbon Steel. Transactions ISIJ, 21:6–15, 1981.

[18] A. Elsner and R. Kaspar. Deep-Drawable Steel Strip Produced by Ferritic Rolling. In T. Chandra, J.M. Torralba, and T. Sakai, editors, Thermec 2003 — Int. Conf. on Processing & Manufacturing of Advanced Materials Part 2, pages 1349–1354, Leganès - Madrid (Spain), July 2003. Trans Tech Publications.

[19] A. de Paepe, J.C. Herman, and V. Leroy. Deep drawable ultra low carbon Ti IF steels hot rolled in the ferrite region. Steel Research, 68(11):479–486, 1997.

[20] W. Bleck and W. Müschenborn. Möglichkeiten und Grenzen zur Beinﬂussung der Werk- stoffeigenschaften bei warm- und kaltgewalztem Stahl. Thyssen Technische Berichte, Heft 1:24–32, 1984.

[21] DIN EN 10130. Kaltgewalzte Flacherzeugnisse aus weichen Stählen zum Kaltumformen, 1999.

[22] J. Reisner, W. Schmidt, and M. Meier. Tiefziehen. In K. Lange, editor, Umformtechnik – Handbuch für Industrie und Wissenschaft Band 3: Blechbearbeitung. Springer-Verlag, Berlin, 1990.

[23] L. Meyer. Neuere Entwicklungen bei weichen und hochfesten mikrolegierten Stählen. Z. Metallkde., 81(1):1–7, 1990. REFERENCES 129

[24] L. Meyer. Mikrolegierungselemte im Stahl. Thyssen Technische Berichte, Heft 1:24–32, 1984.

[25] H.L. Steyer. Beeinﬂussungsmöglichkeiten der Tiefziehfähigkeit mikrolegierter Kaltbän- der. In H.-J. Eckstein, editor, Mikrolegieren von Stahl. VEB Verlag für Grundstofﬁndu- strie, Leipzig, 1984.

[26] A. Bodin. Intercritical deformation of low-alloy steels. PhD thesis, TU Delft, Delft (Netherlands), 2002.

[27] V. Randle and O. Engler. Introduction to texture analysis: macrotexture, microtexture & orientation mapping. Gordon and Breach Science Publishers, Netherlands, 2000.

[28] H.K.D.H. Bhadeshia. Worked examples in the Geometry of Crystals. The Institute of Metals — electronically published version, 2. edition, 2001.

[29] G. Gottstein. Physikalische Grundlagen der Materialkunde. Springer-Verlag, Berlin, 2nd edition, 2001.

[30] B.D. Cullity. Elements of X-Ray Diffraction. Addison-Wesley Publishing Company, Inc., Massachusetts (USA), 1956.

[31] H.J. Bunge. Mathematische Methoden der Texturanalyse. Akademie-Verlag, Berlin, 1969.

[32] L. Lahn. Zum Einﬂuss unterschiedlicher Rekristallisationmechanismen auf die Texturaus- bildung in Tiefziehstählen. PhD thesis, RWTH Aachen, Aachen (Germany), 1998.

[33] D. Raabe. Texturen kubisch-raumzentrierter Übergangsmetalle. PhD thesis, RWTH Aachen, Aachen (Germany), 1992.

[34] R. Sebald, B. Beckers, and G. Gottstein. Effect of Friction on Texture Evolution in Ferritic Warm Rolled IF Steel. In W. Bleck, editor, Int. Symposium: Modern LC and ULC Sheet Steels for Cold Forming – Conference Book, volume 2, pages 363–370, Aachen, April 1998. Verlag Mainz.

[35] D. Raabe and K. Lücke. Texture and Microstructure of Hot Rolled Steel. Scripta Metal- lurgica et Materialia, 26:1221–1226, 1992.

[36] W. Bleck, W. Müschenborn, and L. Meyer. Bessere Gleichmäßigkeit der mechanische Eigenschaften van kaltgewalztem Feinblech. Stahl u. Eisen, 109(3):111–117, 1989.

[37] F. Weber and H.-T. Junius. Vergleich der Glühverfahren zur Herstellung von kaltgewalztem Feinst- und Feinblech. Stahl u. Eisen, 11(4):65–74, 1991.

[38] G. Martin, O. Pawelski, and W. Rasp. Glühöfen zum Rekristallisieren von Kaltbändern und ihr Einﬂuß auf Bandkleber. Blech Rohre Proﬁle, 40(6):468–472, 1993. 130 REFERENCES

[39] S. Lanteri and H. Réglé. Inﬂuence of Heating Rate on Recrystallisation and Texture For- mation in Ti-IF Steels. In W. Bleck, editor, Int. Symposium: Modern LC and ULC Sheet Steels for Cold Forming – Conference Book, volume 2, pages 521–532, Aachen, April 1998. Verlag Mainz.

[40] E.F. Walker and R.C. Hudd. Annealing Treatment of Extra Low And Ultra Low Car- bon Steels. In W. Bleck, editor, Int. Symposium: Modern LC and ULC Sheet Steels for Cold Forming – Conference Book, volume 1, pages 211–224, Aachen, April 1998. Verlag Mainz.

[41] P. Messien and V. Leroy. Scavenging additions of boron in low C - low Al steels. Steel Research, 60(7):320–328, 1989.

[42] F.T. Münch. Leistungssteigerung und Einsparung von Wärmeenergie bei der Warmband- herstellung. Stahl u. Eisen, 117(5):97–104, 1997.

[43] C. Lang, W. Bleck, L. Meyer, and C. Straßburger. Zusammenhänge zwischen der Herstel- lung von Warmband und den Eigenschaften von Kaltband. Stahl u. Eisen, 107(3):21–27, 1987.

[44] B. Plutka. Untersuchung zu Vorgängen bei der Bildung von Rekristallisationstexturen in Tiefziehstählen. PhD thesis, RWTH Aachen, Aachen (Germany), 1991.

[45] J.J. Jonas. Effects of Shear Band Formation on the Texture Development in Warm Rolled IF Steels. In T. Chandra, K. Higashi, C. Suryanarayana, and C. Tome, editors, THERMEC 2000 — International Conference on Processing & Manufacturing of Advanced Materials, Las Vegas (USA), December 2000. Elsevier Science – CD Proceedings.

[46] R. Kaspar, A. Streißelberger, and O. Pawelski. Fließspannung im Bereich der (α → γ)- Umwandlung von kohlenstoffarmen Stählen. Arch. Eisenhüttenwes., 54(5):195–198, 1983.

[47] A. Tomitz and R. Kaspar. Ferritic rolling to produce soft deep-drawable thin hot strips. Steel Research, 21(6+7):233–238, 2000.

[48] A. Tomitz and R. Kaspar. Deep-Drawing quality of cold rolled sheet made of ferritically rolled hot strip. Steel Research, 71(12):504–510, 2000.

[49] R. Kaspar, A. Streißelberger, L. Peichl, and O. Pawelski. Fortgeschrittene Technik der Warmumformsimulation. Zeitschrift für Werkstofftechnik, 14(8):272–277, 1983.

[50] M. Ferry, D. Yu, and T. Chandra. Inﬂuence of Hot Deformation Conditions on the Annea- ling Behavior of Cold Rolled Ultra Low Carbon Steel. ISIJ International, 41(8):876–882, 2001.

[51] M.R. Barnett and J.J. Jonas. Inﬂuence of Ferrite Rolling Temperature on Microstructure and Texture in Deformed Low C and IF Steels. ISIJ International, 37(7):697–705, 1997. REFERENCES 131

[52] M.R. Barnett and J.J. Jonas. Inﬂuence of Ferrite Rolling Temperature on Grain Size and Texture in Annealed Low C and IF Steels. ISIJ International, 37(7):706–714, 1997.

[53] M.R. Barnett. Inﬂuence of Warm Rolling Temperature on Ferrite Recrystallization in Low C and IF Steels. PhD thesis, McGill University, Montreal (Canada), 1996.

[54] M.P. Butrón-Guillén and J.J. Jonas. Effect of Finishing Temperature on Hot Band Textures in an IF Steel. ISIJ International, 36(1):68–73, 1996.

[55] R. Kaspar and A. Tomitz. Ferritic Rolling for Production of Deep-Drawable Thin-Gauge Hot Strips of Steel. In T. Chandra, K. Higashi, C. Suryanarayana, and C. Tome, editors, THERMEC 2000 — International Conference on Processing & Manufacturing of Advan- ced Materials, Las Vegas (USA), December 2000. Elsevier Science – CD Proceedings.

[56] M. Degner, P.-J. Mauk, H. Palkowski, and H.-P. Schmitz. Entwicklungen bei der Warm- umformung im Ferritgebiet. Stahl u. Eisen, 122(4):59–62, 2002.

[57] B.J. Duggan, K.T. Lam, and M.Z. Quadir. Increasing desirable recrystallization texture in IF steel by controlled rolling. In T. Chandra, J.M. Torralba, and T. Sakai, editors, Thermec 2003 — Int. Conf. on Processing & Manufacturing of Advanced Materials Part 4, pages 3745–3750, Leganès - Madrid (Spain), July 2003. Trans Tech Publications.

[58] B. Hutchinsion and P. Bate. Mechanism of Texture Development in Deep-Drawable Steels. In H. Takechi, editor, IF Steels 2003 — Int. Forum for the Properties and Applica- tion of IF Steels, pages 337–346, Tokyo (Japan), May 2003. The Iron and Steel Institute of Japan.

[59] M.R. Barnett and J.J. Jonas. Distinct Aspects of the Physical Metallurgy of Warm Rolling. ISIJ International, 39(9):856–873, 1999.

[60] M.R. Barnett. Role of In-grain Shear Bands in the Nucleation of h111i||ND Recrystalli- zation Textures in Warm Rolled Steel. ISIJ International, 38(1):78–85, 1998.

[61] Y.B. Park, D.N. Lee, and G. Gottstein. Evolution of recrystallisation textures from cold rolling textures in interstial free steels. Materials Science and Technology, 13:289–298, April 1997.

[62] D. Liu, A.O. Humphreys, M.R. Toroghinezhad, and J.J. Jonas. The Deformation Mi- crostructure and Recrystallization Behavior of Warm Rolled Steels. ISIJ International, 42(7):751–759, 2002.

[63] D. Vanderschueren, N. Yoshinaga, and K. Koyama. Recrystallisation of Ti IF Steel Investi- gated with Electron Backstattering Pattern (EBSP). ISIJ International, 36(8):1046–1054, 1996. 132 REFERENCES

[64] J.J. Jonas. Effects of dynamic strain aging, rate sensitivity and shearband formation an the annealing behaviours of warm rolled LC and ULC steels. In W. Bleck, editor, Int. Symposium: Modern LC and ULC Sheet Steels for Cold Forming – Conference Book, volume 1, pages 73–83, Aachen, April 1998. Verlag Mainz.

[65] J.J. Jonas. Effects of shear band formation on the texture development in warm rolled IF steels. Journal of Materials Processing Technology, 117:293–299, 2001.

[66] I. Samajdar, B. Verlinden, L. Kestens, and P. van Houtte. Physical Parameters Related to the Developments of Recrystallization Textures in an Ultra Low Carbon Steel. Acta Materialia, 47(1):55–65, 1999.

[67] Y.Y. Tse, G.L. Liu, and B.J. Duggan. Deformation Banding and Nucleation of Recrystal- lisation in IF Steels. Scripta Materialia, 42(1):25–30, 2000.

[68] A.O. Humphreys, C. Galvani, and J.J. Jonas. Effect of Dynamic Strain Aging on Shear Band Formation During Warm Rolling of IF Steels. In T. Chandra, K. Higashi, C. Suryana- rayana, and C. Tome, editors, THERMEC 2000 — International Conference on Processing & Manufacturing of Advanced Materials, Las Vegas (USA), December 2000. Elsevier Science – CD Proceedings.

[69] A.O. Humphreys, D.S. Liu, M.R. Toroghinezhad, and J.J. Jonas. Effect of Chromium and Boron Additions on the Warm Rolling Behavior of Low Carbon Steels. ISIJ International, 42(Supplement):S52–S56, 2002.

[70] M.R. Barnett. Improving the r-value of warm rolled and annealed low carbon steel by adding chromium. Steel Research, 71(8):295–302, 2000.

[71] J. Asensio, G. Romano, V.J. Martinez, J.I. Verdeja, and J.A. Pero-Sanz. Ferritic steels: Optimization of hot rolled textures through cold rolling and annealing. Materials Charac- terization, 47:119–127, 2001.

[72] M.R. Barnett and L. Kestens. Formation of {111}h110i and {111}h112i Textures in Cold Rolled and Annealed IF Sheet Steel. ISIJ International, 39(9):923–929, 1999.

[73] R.D. Doherty, D.A. Hughes, F.J. Humphreys, J.J. Jonas, D. Juul Jensen, M.E. Kassner, W.E. King, T.R. McNelly, H.J. McQeen, and A.D. Rollett. Current issues on recrystalli- sazation: a review. Materials Science and Engineering A, 238:219–274, 1997.

[74] M.R. Barnett. Carbon and the r-Value of Warm Rolled and Annealed Strip. In W. Bleck, editor, Int. Symposium: Modern LC and ULC Sheet Steels for Cold Forming – Conference Book, volume 1, pages 61–72, Aachen, April 1998. Verlag Mainz.

[75] N. Hashimoto, N. Yoshinaga, and T. Senuma. Texture Evolution of IF Steel Due to Re- crystallisation. ISIJ International, 38(6):617–624, 1998. REFERENCES 133

[76] V. Leroy, A. de Paepe, and J.C. Herman. Ferritic Hot Rolling of Thin Gauge Hot Strips: Processing and Properties. In W. Bleck, editor, Int. Symposium: Modern LC and ULC Sheet Steels for Cold Forming – Confernece Book, volume 1, pages 51–60, Aachen, April 1998. Verlag Mainz.

[77] L. Kestens and J.J. Jonas. Modelling Texture Change during Static Recrystallisation of a Cold Rolled and Annealed Ultra Low Carbon Steel Previously Warm Rolled in the Ferrite Region. ISIJ International, 37(8):807–814, 1997.

[78] H. Magnusson, D. Juul Jensen, and B. Hutchinsson. Growth Rates for Different Texture Components During Recrystallization of IF Steels. Scripta Materialia, 44(3):435–441, 2001.

[79] S. Matsuoka, K. Sakata, O. Furukimi, and T. Obara. Development of Super Deep Dra- wable Sheet Steel by Lubricant Hot Rolling in Ferrite Region. In W. Bleck, editor, Int. Symposium: Modern LC and ULC Sheet Steels for Cold Forming – Conference Book, volume 1, pages 85–96, Aachen, April 1998. Verlag Mainz.

[80] J.A. Schey. Metal Deformation Processes: Friction and Lubrication. Marcel Dekker Inc., New York, 1970.

[81] S. Matsuoka, M. Morita, O. Furukimi, and T. Obara. Effect of Lubrication Condition on Recrystallization Texture of Ultra-low C Sheet Steel Hot-rolled in the Ferrite Region. ISIJ International, 38(6):633–639, 1998.

[82] C.J. Barrett. Inﬂuence of Lubrication on through thickness texture of ferritically hot rolled interstitial free steel. Ironmaking and Steelmaking, 26(5):393–397, 1999.

[83] P. Messien and J.C. Herman. r-Value Improvement of Steel Product Using Hot-Rolling Lubrication. In G. Tither R. Asfahant, editor, Int. Symposium on Low-Carbon Steels for the 90’s, pages 383–388. The Minerals, Metals and Materials Society, 1993.

[84] J.A. Schey. Tribology in Metalworking: Friction, Lubrication and Wear. American Soci- ety for Metals, Ohio, 1984.

[85] S.H. Lee and D.N. Lee. Shear rolling and recrystallisation textures of interstial-free steel sheet. Materials Science and Engineering A, 249:84–90, 1998.

[86] T. Sakai, Y. Saito, and K. Kato. Texture Formation in Low Carbon Ti Bearing Steel Sheets by High Speed Rolling in the Ferrite Region. Transactions ISIJ, 28:1036–1024, 1988.

[87] M. Hölscher, D. Raabe, and K. Lücke. Relationship between Rolling Textures and Shear Textures in F.C.C. and B.C.C. Metals. Acta Materialia, 42(3):879–886, 1994.

[88] D. Raabe and K. Lücke. Textures of ferritic stainless steels. Materials Science and Tech- nology, 9:302–312, April 1993. 134 REFERENCES

[89] K.L. Williams. A Survey of the Literare on the Subjekt of Hot Mill Roll Gap Lubrication. Internal Repot No: WL/ME/R/R (WL 49), Welsch Laboratory – British Steel Corporation, Port Talbot, 1979. [90] W.L. Roberts. Tribological Considerations in the Hot Rolling of Low-Carbon Steels. Lub- rication Engineering, 33(11):575–580, 1977. [91] V.B. Ginzburg. Steel – rolling technology: theory and practice. Marcel Dekker Inc., New York, 1989. [92] W.L. Roberts. Flat processing of steel. Marcel Dekker Inc., New York, 1987. [93] S. Ghosh and M.M. Bowness. Production of thin strip through a conventional hot strip mill. Steel World, 5(1):54–59, 2000. [94] Bähr-Thermoanalyse GmbH. Dilatometer. Brochure, 4/97. [95] SEP 1681 2. Ausgabe. STAHL-EISEN-Prüfblätter des VDEh: Richtlinien für die Vorbe- reitung, Durchführung und Auswertung dilatometrischer Umwandlungsuntersuchungen an Eisenlegierungen, June 1998. [96] H. Hougardy. Die Umwandlung der Stähle – Teil 2: Zeit - Temperatur - Austenitisie- rung (ZTA)- und Zeit - Temperatur - Umwandlung (ZTU) Schaubilder, Einführung in die Wärmebehandlung. Verlag Stahleisen, Düsseldorf, 1982. [97] A. Köthe, J. Richter, and G. Backmann. Zur γ − α Umwandlung in ferritisch-perlitischen höherfesten Baustählen unter den Bedingungen des kontrollierten Walzens. Freiberger Forschungshefte B, 244:104–123, 1985. [98] R. Kaspar, W. Kapellner, and O. Pawelski. Umwandlung bei der thermomechanischen Behandlung. HTM, 41(4):177–185, 1986. [99] Y.E. Smith and C.A. Siebert. Continuous Cooling Transformation Kinetics of Thermo- mechanically Worked Low-Carbon Austenite. Metallurgical Transactions, 2:1711–1725, 1971. [100] A. Klumburg, F. Krontheuer, and E. Kaiser. Einﬂuß der Abkühlungsart auf das Umwand- lungsverhalten von Stählen. HTM, 42(2):69–74, 1987. [101] C. Ouchi, T. Sampei, and I. Kozasu. The Effect of Hot Rolling Condition and Chemical Composition on the Onset Temperature of the α − γ Transformation after Hot Rolling. Transactions ISIJ, 22:214–222, 1982. [102] H. Schumann. Austenitumwandlung in Stählen. Neue Hütte, 22(6):296–301, 1977. [103] C.M. Vlad, M. Liphardt, and J. Kügler. Einﬂuß der Warmverformung auf das Umwand- lungsverhalten und das Gefüge verschiedener niedriglegierter Stähle. Steel Research, 57(12):654–657, 1986. REFERENCES 135

[104] SEP 1680. STAHL-EISEN-Prüfblätter des VDEh: Aufstellung von Zeit-Temperatur- Umwandlungsschaubildern, December 1990.

[105] D.N. Hanlon, J. Sietsma, and S. van der Zwaag. The Effect of Plastic Deformation of Austenite on the Kinetics of Subsequent Ferrite Transformation. ISIJ International, 41(9):1028–1036, 2001.

[106] Technische Lieferinformationen. HOUGTHON Deutschland GmbH Werk Aachen, Werk- str. 26, 52076 Aachen-Oberforstbach.

[107] Personal communication with Mr. H. Lennartz (HOUGTHON Deutschland GmbH Werk Aachen, Werkstr. 26, 52076 Aachen-Oberforstbach), 2004.

[108] Technische Lieferinformationen. Keller & Bohacek GmbH & Co. KG, Liliencronstr. 64, 40472 Düsseldorf.

[109] D. Beste. Glüh- und Tauchbeschichtungssimulator am IEHK – PDF Dokument ’Beschrei- bung RHESCA 06-08-02’. Institut für Eisenhüttenkunde der RWTH Aachen, e-mail: [email protected], August 2002.

[110] P. Juntunen, D. Raabe, P. Karjalainen, T. Kopio, and G. Bolle. Optimizing Continuous Annealing of Interstial-Free Steels for Improving Deep Drawability. Metallurgical and Materials Transactions A, 32A:1989–1995, 2001.

[111] DIN EN 10002-1. Metallische Werkstoffe – Zugversuch – Teil 1: Prüfverfahren bei Raum- temperatur, December 2001.

[112] A.J. DeArdo. Physical Metallurgy of the Thermomechanical Processing of Microalloyed Steels. In T. Chandra, K. Higashi, C. Suryanarayana, and C. Tome, editors, THERMEC 2000 — International Conference on Processing & Manufacturing of Advanced Materials, Las Vegas (USA), December 2000. Elsevier Science – CD Proceedings.

[113] Niobium Product Company. High Strength EDDQ Sheet. Niobium Information, (12), 1996.

[114] I. Gutierrez, M.M Petite, J.I. Laburu, J. Zaitegui, W.B. Hutchinson, D. Artimowicz, P.J. Evans, G.J Spurr, K.H.D.H. Badeshia, and N. Chester. Modelling of Microstructural De- velopment During Continuous Annealing Process. Draft Final Report 7210.EC/808, EC- SC, London, 1998.

[115] D.R. Waryoba and P.N. Kalu. Quantiﬁcation of Recrystallisation Kinetics in Heavily Drawn OFHC Copper Wires by Microhardness Technique. In Proceedings of IMECE’03 – International Mechanical Engineering Congress, Washington, D.C. (USA), November 2003. 136 REFERENCES

[116] M. Fontana, B. Arcondo, M.T. Clavagurea-Mora, and N. Clavaguera. Crystallization kinetics driven by two simultaneous modes of crystal growth. Philosophical Magazine B, 80(10):1833–1856, 2000.

[117] M.E. McHenry, F. Johnson, H. Okumura, T. Ohkubo, V.R.V. Ramanan, and D.E. Laugh- lin. The kinetics of nanocrystallization and microstructural observations in FINEMET, NANOPERM and HITPERM nanocomposite magnetic materials. Scripta Materialia, 48:881–887, 2003.

[118] N.J. Champion, P.J. Evans, and R.P. Underhill. Novel Thermo-Mechanical Processing using the Ponymill Process. In Thermomechanical Processing of Steel, pages 625–635, London (UK), May 2000.

[119] L. Storojeva, R. Kaspar, and D. Ponge. Ferritic-Pearlitic Steel with Deformation Induced Spheroidized Cementite. In T. Chandra, J.M. Torralba, and T. Sakai, editors, Thermec 2003 — Int. Conf. on Processing & Manufacturing of Advanced Materials Part 2, pages 1349–1354, Leganès - Madrid (Spain), July 2003. Trans Tech Publications.

[120] L. Storojeva, R. Kaspar, and D. Ponge. Effect of Heavy Warm Deformation on the Mi- crostructure and Mechanical Properties of Medium Carbon Ferritic-Pearlitic Steel. in preperation, 2004.

[121] L. Storojeva, D. Ponge, R. Kaspar, and D. Raabe. Development of Microstructure and Texture of Medium Carbon Steel during Heavy Warm Deformation. in preperation, 2004.

[122] P.S. Bate. Manual for ANIS-MPI. University of Birmingham, 1997.

[123] M.Z. Quadir and B.J. Duggan. Shear band formation in warm rolled IF steel. In T. Chan- dra, J.M. Torralba, and T. Sakai, editors, Thermec 2003 — Int. Conf. on Processing & Ma- nufacturing of Advanced Materials Part 4, pages 3769–3774, Leganès - Madrid (Spain), July 2003. Trans Tech Publications.

[124] B. Engl, E.-J. Drewes, and G. Stich. Eigenschaften und Anwendung von höherfesten Mehrphasenstählen. Stahl und Eisen, 117(10):67–69, 1997.

[125] S. Nomura, Y. Nakazawa, S. Nakai, N. Komatsubara, and Y. Hirose. Development of High-Strength Hot-Rolled Sheet Products for Automobile Application. The Somitomo Search, 7(59):39–46, september 1997.

[126] ULSAB Advanced Vehicle Concept. Technical Transfer Dispatch #6, 26. May 2001.

[127] M. Dukat. Ultra Light Steel Auto Body: Vom Konzept zur realen Rohbaustruktur. In Ent- wicklung im Karosseriebau, number 1398, pages 261–281. VDI Verlag, Hamburg (Ger- many), May 19-20 1998.

[128] Niobium Information 15/97. Dual Phase and TRIP Steels, 1997. REFERENCES 137

[129] P. Jacques, Q. Furnémont, T. Pardoen, and F. Delannay. On the Role of Martensitic Trans- formation on Damage and Cracking Resistance in TRIP-Assisted Multiphase Steels. Acta Materialia, 49:139–152, 2001.

[130] J. Wang and S. van der Zwaag. Theoretical Study of P-Containing Transformation- Induced Plasticity Steel - Part 1: Determination of the Phosphorus Concentration. Z. Metallkunde, 92(12):1299–1305, 2001.

[131] J. Wang and S. van der Zwaag. Theoretical Study of P-Containing Transformation- Induced Plasticity Steel - Part 2: Analysis of the Potential TRIP Effect. Zeitschrift für Metallkunde, 92(12):1306–1311, 2001.

[132] S. Traint, A. Pichler, K. Hauzenberger, P. Staiszny, and E. Werner. Inﬂuence of silicon, aluminum, phosphorous and copper on the phase transformations of low alloyed TRIP- steels. In Int. Conf. on TRIP-Aided High Strength Ferrous Alloys, pages 121–128, Ghent (Belgium), June 19-21 2002.

[133] P. Jacques, X. Cornet, P. Harlet, J. Ladrière, and F. Delannay. Enhancement of the Me- chanical Properties of a Low-Carbon, Low-Silicon Steel by Formation of a Multiphased Microstructure Containing Retained Austenite. Metallurgical and Materials Transactions A, 29A(9):2383–2393, 1998.

[134] P. Jacques, P. Harlet, and F. Delannay. Critical assessment of the phase transformation occurring during the heat treatment of TRIP assisted multiphase steels. In Int. Conf. on TRIP-Aided High Strength Ferrous Alloys, pages 129–134, Ghent (Belgium), June 19-21 2002.

[135] E.V. Pereloma, I.B. Timokhina, and P.D. Hodgson. Transformation behaviour in thermo- mechanically processed C-Mn-Si TRIP steels with and without Nb. Materials Science and Engineering, A273-275:448–452, 1999.

[136] H.J. Koh, S.K. Lee, S.H. Park, S.J. Choi, S.J. Kwon, and N.J. Kim. Effect of hot rolling Conditions on the Microstructure and Mechanical Properties of Fe-C-Mn-Si Mulitphase Steels. Scripta Materialia, 38(5):763–768, 1998.

[137] W. Bleck, R. Kaspar, W. Müschenborn, and F. Rakoski. Grundlagen der integrierten Wä- rmebehandlung. Stahl und Eisen, 117(4):39–46, 1997.

[138] J. Ohlert. Einﬂuss von chemischer Zusammensetzung und Herstellungsverlaufauf die Mikrostruktur und mechanischen EIgenschaften von TRIP Stählen. PhD thesis, RWTH Aachen, Aachen (Germany), 2003.

[139] T. Heller, B. Engl, J. Kruse, and U. Schriever and. Innovative Werkstoffentwicklung als Grundlage für den modernen Automobilbau. In VDI, editor, Entwicklung im Karosserie- bau, number 1398, pages 323–341. VDI, Hamburg (Germany), May 19-20 1998. 138 REFERENCES

[140] A. Bodin, J. Sietsma, and S. van der Zwaag. Texture and Microstructure Development during intercritical Rolling of Low-Carbon Steels. Metallurgical and Materials Transac- tions, 33A:1589–1603, 2002.

[141] S. Gohda, K. Watanabe, Y. Hashimoto, H. Hirayama, and S. Kijima. Effects of Intercritical Rolling at Hot Strip Mill on the Strength and Notch Toughness of Plain Carbon Steel. ISIJ Trans., 21:360–369, 1981.

[142] D.P. Dunne, B. Feng, and T. Chandra. The Effect of Ti and Ti-Nb Additions on X Forma- tion and Restoration during Intercritical Rolling and Holding of C-Mn Structural Steels. ISIJ International, 31(11):1354–1361, 1991.

[143] S. Gohda, K. Watanabe, and Y. Hashimoto. Effects of the Intercritical Rolling on Structure and Properties of Low Carbon Steel. ISIJ Trans., 21:6–15, 1981.

[144] S. Papaefthymiou. Failure mechanisms of multiphase steels. PhD thesis, RWTH Aachen, Aachen, 2004.

[145] H. Luo, L. Zhao, S.O. Kruiver, J. Sietsma, and S. van der ZWAAG. Effect of Intercriti- cal Deformation on Bainite Formation in Al-containing TRIP Steel. ISIJ International, 43(8):1219–1227, 2003.

[146] E.A. Simielli, S. Yue, and J.J. Jonas. Recrystallization Kinetics of Microalloyed Steels deformed in the intercritical Region. Metallurgical Transactions, 23A:597–608, 1992.

[147] L. Chabbi and W. Lehnert. Temperaturkontrolliertes Walzen von Vergütungsstählen im Zweiphasengebiet. Stahl und Eisen, 118(3):63–66, 1998.

[148] L. Chabbi and W. Lehnert. Controlled hot forming of heat treatment steel grades in the intercritical ( ) region. Journal of Materials Processing Technology, 106:13–22, 2000.

[149] S. Godet, P. Harlet, F. Delannay, and P. Jacques. Effect of hot-rolling conditions on the tensile properities of multiphase steels exhibiting a TRIP effect. In Int. Conf. on TRIP- aided high strength ferrous alloys, pages 135–140, Ghent (Belgium), June 19-21 2002.

[150] A. Bodin, J. Sietsma, and S. van der Zwaag. On the nature of the bimodal grain size distribution after intercritical deformation of a carbon-manganese steel. Materials Characte- rization, 47:187–193, 2001.

[151] O. Pawelski, U. Rüdiger, and R. Kaspar. Der Warmumformsimulator des Max-Planck- Institutes für Eisenforschung – ein neues Konzept zur Erforschung schneller Warmum- formvorgänge. Stahl und Eisen, 98(5):181–189, 1978.

[152] A. Streißelberger, L. Peichl, R. Kaspar, and O. Pawelski. Auswertung und metallkundliche Aussagen rechnergesteuerter Warmﬂachstauchversuche. Zeitschrift für Werkstofftechnik, 14(9):305–310, 1983. REFERENCES 139

[153] V. Gopinathan and O. Pawelski. Flat compression – A simulation test for hot rolling. Transactions of the Indian Institute of Metals, 35(5):503–511, 1982. [154] O. Pawelski and V. Gopinathan. Comparison of Meterial flow and deformation resistance of HSLA steel deformed by hot rolling and by flat compression under simulated conditions. Journal of Mechanical Working Technology, (5):267–280, 1981. [155] W. Kapellner, R. Kaspar, and O. Pawelski. Bestimmung von Fließkurven mit dem Warm- flachstauchversuch. steel research, 57:256–261, 1986. [156] J.S. Distl, A. Streißelberger, R. Kaspar, and U. Zeislmair. Messung mechanischer Eigen- schaften an flachgestauchten Proben nach Simulation des TM - (thermomechanischen) Walzens. Materialprüfung, 27(5):131–135, 1985. [157] E. Girault, P. Jacques, P. Harlet, K. Mols, J. van Humbeeck, E. Aernoudt, and F. Delannay. Metallographic Methods for Revealing the Multiphase Microstructure of TRIP-Assisted Steels. Materials Characterization, pages 111–118, 1998. [158] E. Macherauch and H. Wohlfahrt. Prinzipien der quantitativen röntgenographischen Pha- senanlyse (RPA). HTM, 27:230–232, 1972. [159] G. Faninger and U. Hartmann. Physikalische Grundlagen der quantitativen röntgeno- graphischen Phasenanalyse (RPA). HTM, 27:233–244, 1972. [160] U. Hartmann. Quantitative Analyse vielphasiger Gemische mit beliebigem Superpositi- onsgrad der Röntgeninterferenzlinien. HTM, 27:251–260, 1972. [161] U. Wolfstieg. Berücksichtigung der Textureinflüsse bei der quantitativen röntgenographi- schen Phasenanalyse. HTM, 27:245–251, 1972. [162] W. Beumelburg and E. Schreiber. Einfluß von Probenform und Probenvorbereitung bei der quantitativen röntgenographischen Phasenanalyse (RPA). HTM, 27:265–271, 1972. [163] W. Kraus and G. Nolze. PowderCell for Windows – Version 1.0. Federal Institute for Materials Research and Testing (BAM), Unter den Eichen 87, D-12205 Berlin, 1997. [164] C. Capdevila, F.G. Caballero, and C. García de Andrés. Determination of Ms Temperature in Steels: A Bayesian Neural Network Model. ISIJ International, 42(8):894–902, 2002. [165] P. Jacques, A. Petein, and P. Harlet. Improvement of mechanical properties through con- current deformation and transformation: new steels for the 21st century. In Int. Conf. on TRIP-aided high strength ferrous alloys, pages 281–285, Ghent (Belgium), June 19-21 2002. [166] ThyssenKrupp Stahl AG – Werkstoffblatt 1830, July 2003. [167] W. Beiz and K.-H. Küttner, editors. Dubblel Interaktiv – Das elektronische Taschenbuch für den Maschinenbau. Springer-Verlag, 1999.

List of Frequently used Symbols

Symbol Explanation Unit

{hkl} huvwi Miller indices of a crystallograhic plane or orientation - A50 ultimate elongation (uniform length of 50mm) % Ag uniform elongation % ◦ Ar1 γ to α and Fe3C equilibrium temperature during cooling C ◦ Ar3 γ to α equilibrium temperature during cooling C b Burgers vector m d grain size m dhkl lattice plane spacing m, Å ∆r planar anisotropy - ∆rmax maximum planar anisotropy - ε˙ strain rate s−1 εAR austenitic rolling strain - εCR cold rolling reduction % εFR ferritic rolling strain - εIC intercritical rolling strain - εthickness strain in thickness direction of the tensile specimen - εwidth strain in width direction of the tensile specimen - ◦ φ1,φ2,Φ Euler angles fRA fraction retained austenite - G shear modulus Pa γ grain boundary energy J m−2 h,k,l Miller indices - λ wave length m, Å ld arc of contact m m strain rate sensitivity - n, n0◦ ,n45◦ ,n90◦ strain hardening coefﬁcient (indices indicate angle to RD) - nm mean strain hardening coefﬁcient - −1 Qrex activation energy for recrystallisation J mol r, r0◦ , r45◦ , r90◦ Lankford value (indices indicate angle to RD) - rm mean Lankford value -

141 142 LIST OF FREQUENTLY USED SYMBOLS

ρ dislocation density m−2 r roll radius m R gas constant j mol−1 K−1 rc critical radius of a nucleus m rm mean Lankford value - Rm ultimate tensile strength Pa Rp0.2 0.2% proof strength Pa T temperature K t0 initial strip thickness m t1 ﬁnal strip thickness m ◦ t8/5 time required to cool a specimen from 800 to 500 C s ◦ TAR austenitic rolling temperature C ◦ TAus austenitisation temperature C ◦ TC coiling temperature C ◦ TFR ferritic rolling temperature C tH holding time after intercritical deformation s ◦ TIC intercritical rolling temperature C ts,t f strip thickness m w0 initial width of the tensile specimen m w1 width after straining of the tensile specimen m wi content of the alloying element i mass% X softening % Xrex fraction recrystallised - List of Publications

A. Elsner and R. Kaspar. Deep-Drawable Steel Strip Produced by Ferritic Rolling. In T. Chan- dra, J.M. Torralba, and T. Sakai, editors, Thermec 2003 — Int. Conf. on Processing & Ma- nufacturing of Advanced Materials Part 2, pages 1349–1354, Leganès - Madrid (Spain), 7.-11. July 2003. Trans Tech Publications.

A. Elsner, R. Kaspar, D. Ponge, D. Raabe and S. van der Zwaag. Recrystallisation Texture of Cold Rolled and Annealed IF Steel Produced from Ferritic Rolled Hot Strip. In B. Bacroix, J.H. Driver, R. Le Gall, Cl. Maurice, R. Penelle, H. Réglé and L. Tabourot, editors, Proceedings of the 2nd Joint International Conference on Recrystallization and Grain Growth, ReX & GG2, pages 257–262, Annecy (France), 30. August - 3. September 2004. Trans Tech Publications.

A. Elsner and R. Kaspar Bericht Nr. AW 132 Laborwalzversuche zur Optimierung der Pro- zessparameter beim Ferritwalzen tiefziehbarer Stähle. Verein zur Förderung von For- schungsarbeiten auf dem Gebiet der Walzwerkstechnik in der Hüttenindustrie (VFWH) im Stahl-Zentrum, August 2004.

143

Acknowledgements

The research presented in this thesis has been carried out at the Max-Planck-Institut für Eisen- forschung in the department for Microstructure Physics and Metal Forming and Delft University of Technology in the group for Fundamentals of Advanced Materials. First of all I would like to thank my promoters Prof. dr. ir. Sybrand van der Zwaag and Prof. Dr.-Ing. Kurt Steinhoff for their interest in my work and giving me the opportunity to pursue my Ph D degree at the TU Delft. Their professional supervision, continuous encouragement and the many fruitful discussions helped me not to loose the track during the past 4.5 years. Next to this I would like to thank my supervisors from the Max-Planck-Institut Dr.-Ing. Radko Kaspar and Dr.-Ing. Wolfgang Rasp for the idea of the project and the support during my time at the MPI. Especially I would like to thank Dr.-Ing. Radko Kaspar for the many fruitful discussions and helpful suggestions even after his retirement. Furthermore I would like to express my gratitude to Prof. Dr.-Ing. Oskar Pawelski and Prof. Dr.-Ing. Dierk Raabe for giving me the opportunity to work at the MPI. Special thanks go to ir. Johan van Santen for his friendly collaboration and engagement during his diploma-work. His results formed a solid basis for chapter 5. More over I would like to thank him and my former fellow student ing. Albert Hendriks for the translation of the summary from English to Dutch. Of course I also appreciated the time with my former colleagues at the MPI, especially Dr. Chris Wichern, Dipl.-Ing. Cristian Klüber, Dipl.-Ing. Michael Sachtleber, Dipl.-Ing. Hermann Lücken, Dipl.-Ing. Araz Barani and Dipl.-Ing. Sandra Detroy. This thesis would not be what it is without the help and cooperation of the technicians of the group for Microstructure Physics and Metal Forming at the MPI. Particularly I would like to thank Paul Tolksdorf for solving any electric problem and tuning the rolling mill, Herbert Faul for his ﬂexible help and performing the mechanical testing, Frank Schlüter and Gerd Goris for assis- ting the rolling and performing the WUMSI experiments, Heidi Börgershausen for the quick and professional metallography and Michael Adamek for performing the dilatometer experiments. Next to this there were many people and companies outside the MPI and the TU Delft who helped to realise the present thesis. Especially I would like to thank the ThyssenKrupp Stahl AG for delivering the extended crop-pieces for the experiments, Dipl.-Ing. Helmut Lennartz (Houghton Deutschland) for delivering the hot rolling lubricants, Dr. Friedhelm Danowski (Keller & Boha- cek), Dr. Werner Olberding (Institut für Galvno- und Oberﬂächentechnik) for their consultancy

145 146 ACKNOWLEDGEMENTS in the field of pickling, Helmut Kosslers (Corus Special Strip - Hille & Müller) for the opportunity to use his laboratory and equipment for the pickling of my strips and Dipl.-Ing. Daniel Beste (Institut für Eisenhüttenkunde, RWTH Aachen) for performing the annealing experiments using the RHESCA simulator. Moreover I am indebted to the "Verein zur Förderung von Forschungs- arbeiten auf dem Gebiet der Walzwerkstechnik in der Hüttenindustrie (VFWH)" for the financial support of this project. Many thanks go to my parents, parents-in-law, sister and brother for their encouragement and support in every respect. Last, but definitely not least, special thanks go to my wife Daniela and my son Till for their love, encouragement and tolerance, especially during the last months of the completion of this thesis.

Olpe, 3 May 2005 Curriculum Vitae

1975 Geboren, 21 August 1975 te Düsseldorf (Duitsland).

1982 – 1986 Basisschool, Gebrüder-Grimm-Grundschule, Ratingen (Duitsland).

1986 – 1995 VWO, Geschwister-Scholl-Gymnasium, Ratingen (Duitsland).

1995 – 1996 Civil Service Cartitas Pﬂegestation, Ratingen (Duitsland).

1996 – 2000 Studie Mechatronica, Fonty’s Hogeschool, Venlo en Niederrhein University of Applied Science, Krefeld (Duitsland). Afstudeerwerk bij het Max-Planck-Institut für Eisenforschung, Düsseldorf (Duitsland). Onderwerp: "Investigation of the inﬂuence of electromagnetic ﬁelds on the cold deformation process".

2000 – 2003 Werkzaam bij het Max-Planck-Institut für Eisenforschung, Düsseldorf (Duitsland) in de groep Microstructure Physics and Metal Forming o.l.v Dr.-Ing. Radko Kaspar en Dr.-Ing. Wolfgang Rasp.

2001 – 2005 Promotieonderzoek in de groep Fundamentals of Advanced Materials, Technische Universiteit Delft. Onderwerp: "Advanced hot rolling strategies for IF and TRIP steels", o.l.v. Prof. dr. ir. Sybrand van der Zwaag en Prof. Dr.-Ing. habil. Kurt Steinhoff.

2003 – heden Werkzaam bij ThyssenKrupp Stahl AG, Division Industrie

147