mechanical properties of dual-phase steels.pdf

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Lehrstuhl fur Werkstoffkunde und Werkstoffmechanik mit

Materialprufamt fur den Maschinenbau

Technische Universitat Munchen

Mechanical properties of Dual-Phase steels

Prodromos Tsipouridis

Vollstandiger Abdruck der von der Fakultat fur Maschinenwesen

der Technischen Universitat Munchen

zur Erlangung des akademischen Grades eines

Doktor-Ingenieurs (Dr.-Ing.)

genehmigten Dissertation.

Vorsitzender: Univ.-Prof. Dr.-Ing. Horst Baier

Prufer der Dissertation:

1. Univ.-Prof. Dr. mont. habil. Ewald Werner

2. Hon.-Prof. Dr.-Ing, Dr. Eng. (Japan) Hans-Harald Bolt

Die Dissertation wurde am 14.03.2006 bei der Technischen Universitat Munchen

eingereicht und durch die Fakultat fur Maschinenwesen

am 19.06.2006 angenommen.


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Acknowledgements

This study was carried out during my employment as research assistant at the Institute for

Materials Science and Mechanics of Materials of TU-Munich (Lehrstuhl fur Werkstoffkunde

und Werkstoffmechanik).

I am deeply indebted to my direct advisor Prof. Dr. mont. habil. E. A. Werner for his unfailing

support all these years, for his encouragement to proceed with new ideas and for being daily

available for uncountable scientific discussions. Thank you very much!

My special thanks to my co-advisor Prof. Dr.-Ing., Dr. Eng. H.-H. Bolt (Head of group Mate-rials Synthesis and Materials Characterization of Max-Planck-Institut fur Plasmaphysik, IPP

Garching), as well as to the chairman of my PhD examination, Prof. Dr.-Ing. H. Baier (Head

of the Institute for Lightweight Structures, TU-Munich).

My sincere thanks to the project partner voestalpine Stahl Linz GmbH for supplying the

testing material and making possible to conduct the annealing simulations as well as the tensile

and hole expansion tests. Special thanks to Dr. A. Pichler for the support and to Dipl.-Ing.

E. Tragl for the close and continuous collaboration.

I would also like to thank the Christian Doppler Research Association (CDG) for sponsoring

and supporting this study during the years 2002-2005 (as a project/module of the Christian-

Doppler-Laboratory for Modern Multiphase Steels).

I should not forget to thank Dr. G. Triantafyllides (Dep. of Chemical Engineering of Aristotle

University of Thessaloniki) for motivating me to choose the field of steel research.

Last, but not least, i would like to thank my colleagues and the technical staff of the Chair,

without whose help this work would be never completed. But most of all, thank you forcreating a pleasant and friendly working environment and for helping me to become an active

member of our group.

Munich, June 2006 Prodromos Tsipouridis


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Contents

1 Introduction 1

1.1 Low-alloyed dual-phase steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

1.2 Grain refinement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

2 Material production and experimental procedure 152.1 Production of the material . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

2.2 Thermodynamical calculations . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16

2.3 Pre-processing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16

2.4 Cold-Rolling trials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

2.5 Dilatometric investigations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

2.6 Annealing simulations and austenitization kinetics . . . . . . . . . . . . . . . . . 19

2.7 Microstructure investigations. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

2.8 Mechanical testing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

2.8.1 Tensile testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20

2.8.2 Ultramicrohardness and microhardness testing . . . . . . . . . . . . . . . 21

2.8.3 Hole expansion measurements (Stretch flangeability) . . . . . . . . . . . 21

3 Results 23

3.1 Microstructure investigations. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

3.1.1 Recrystallization/austenitization kinetics . . . . . . . . . . . . . . . . . . 23

3.1.2 Dilatometric investigations . . . . . . . . . . . . . . . . . . . . . . . . . . 26

3.1.3 Annealing simulations . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33

3.1.4 Quantitative analysis- Grain size measurements . . . . . . . . . . . . . . 443.2 Mechanical properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50

3.2.1 Tensile testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50

3.2.2 Hardness measurements . . . . . . . . . . . . . . . . . . . . . . . . . . . 65

3.2.3 Hole expansion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 71

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4 Discussion 77

4.1 Grain refinement and microstructure investigations . . . . . . . . . . . . . . . . 77

4.2 Mechanical properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80

5 Summary 102

II


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Chapter 1

Introduction

The increasing automobile market demands for reduced fuel consumption as well as the need

to comply with the international environmental regulations regarding greenhouse gas emissions(GHG), resource reduction and recyclability have motivated and/or even forced the automotive

industry to produce more fuel-efficient vehicles by reducing their weight. In order to provide a

steel-based structural platform that fulfills the auto-makers requirements and takes advantage

of the new high strength steels, a new vehicle architecture based on novel design concepts has

been developed. The application of advanced high strength sheet steels exhibiting both high

strength and excellent formability offers the unique option of combining weight reduction (by

using thinner gauges of sheet material) with improved passive safety, optimized environmental

performance and manufacturing feasibility at affordable cost.

The most common steels used presently in the automobile industry are mild steels, whichare low-carbon steels characterized by a yield strength level of 140 MPa and an excellent deep

draw ability. Despite their forming and cost advantages over high strength steels, the ultimate

strength level of mild steels remains at relatively low levels, so that the crash performance is

mainly dependent on the sheet thickness. By consistent controlling of the alloy chemistry, con-

sidering the presence of interstitial carbon in ferrite, interstitial-free grades (IF) possessing an

ultra-low carbon content were produced. Traditional strengthening mechanisms such as solid

solution hardening (with phosphorous to be the most common hardening element), precipita-

tion hardening and grain refinement by carbides and/or nitrides were employed to increase the

strength of IF-steels, while maintaining their excellent formability. Micro-alloying with vana-

dium, niobium or titanium accompanied with fine carbide precipitation and grain refinement

leads to even higher strength levels and increases the ratio of yield to tensile strength. Bake-

hardening steels (BH) offer a combination of good formability during stamping and provide

an increased yield strength after the paint-baking process. To take advantage of the bake-

hardening effect a certain content of solute carbon and an appropriate aging heat treatment

are required, aiming at the supersaturation of carbon in the ferrite [1]. Due to their sharp

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upper yield point, BH-steels exhibit a high dent resistance, which makes them good candidate

materials for outer body panel applications. For the steel grades mentioned above, widely

designated as conventional steels, a reduced formability is an unavoidable consequence when

selecting steels with higher strength levels.

Figure 1.1: Strength-Formability relationship of thin sheet steels

To overcome this problem new types of high strength steels, the so-called Advanced High

Strength Steels (AHSS), were developed from the sheet steel suppliers in cooperation with the

automakers and design engineers. These grades exhibit higher rates of work hardening than

conventional steels as a result of their low yield strength to tensile strength ratio, show good

press formability and reach higher ultimate tensile strengths, therefore they have the potential

for significantly improved crash performance [2]. AHSS steels are multiphase steels consist

of hard islands of martensite, bainite and/or retained austenite dispersed in a ductile ferritic

matrix, in quantities and combinations sufficient to produce desired mechanical properties.

The multiphase AHSS family includes Dual-Phase (DP, ferritic-martensitic), TRansformation

Induced Plasticity (TRIP) and complex multiphase steels. Ferritic-bainitic steels, also known as

stretch-flangeable (SF) steels, are considered as a subgrade of the DP products. The mechanical

properties of conventional and AHSS thin sheet steels with respect to ductility and ultimate

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tensile strength are shown in Figure 1.1. As can be observed, a partial overlap between the

strength levels of different steel grades is possible.

According to the results of the ULSAB-AVC Program (Ultra Light Steel Auto Body-Advanced

Vehicle Concepts), an automotive body could be constructed by utilizing approximately 85 %

of AHSS, achieving a weight reduction of 25 % compared with a bench-marked average

base model and without any increase of the manufacturing costs. Figure 1.2shows that the

clear majority of autobody components is designed using dual-phase steels [3, 4]. Different

criteria such as formability, weldability, spring-back behavior and of course static and dynamic

properties play a significant role in the material selection, even though for some parts more

than one steel grades fulfill the mechanical property standards and hence are also applicable.

In particular, for a number of components that both DP and TRIP steels are viable candidates,

cost-effective DP grades were preferable.

Figure 1.2: ULSAB autobody structure steel grade distribution

Due to their special microstructural characteristics, ferritic-martensitic dual-phase (DP) steels

provide an attractive combination of strength and ductility and furthermore exhibit continuous

yielding behavior accompanied with a high work hardening rate. In order to meet the

different design requirements of individual components, various DP grades regarding tensile

strength and formability are produced industrially. This variation of mechanical properties

is mainly achieved by controlling the carbon content of the steel. Addition of other alloying

elements such as manganese, chromium, vanadium, molybdenum and nickel, individually or in

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combination, can also increase hardenability. Another method to increase the tensile strength

of dual-phase steels is to increase the martensite fraction by applying appropriate annealing

schedules, even though this procedure is followed by an expected loss in formability. In the

present study, grain refinement is proposed as an alternative way/solution to improve the me-

chanical properties of a low-alloyed dual-phase steel. To achieve this, severe plastic deformation

was applied to a pre-processed dual-phase steel by means of conventional cold-rolling, followed

by an appropriate final heat treatment with the aim to produce a homogeneous fine-grained

dual-phase ferritic-martensitic microstructure. The impact of grain refinement on the mechani-

cal properties of the dual-phase steel, regarding tensile properties, ultramicrohardness and hole

expansion behavior, is then investigated in this work.

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1.1 Low-alloyed dual-phase steels

Dual-phase (DP) steels were developed in the mid-seventies in order to satisfy the increasing

needs of automotive industry for new high strength steels which combine significant weight

reduction and improved crash performance, while keeping the manufacturing costs at affordablelevels. The high commercial potential of the newly developed alloy has motivated extensive

research in numerous laboratories, resulting in DP-grades having a wide range of chemical

compositions and being produced with various processing routes.

Dual-phase steels are characterized by a microstructure consisting of a fine dispersion of

hard martensite particles in a continuous, soft, ductile ferrite matrix. The term dual-phase,

firstly reported by Hayami and Furukawa [5] and thereafter adopted by the steel research

community, refers to the presence of essentially two phases, ferrite and martensite, in the

microstructure, although small amounts of bainite, pearlite and retained austenite may also be

present. Irrespective of the chemical composition of the alloy, the simplest way to obtain a dual-phase ferritic-martensitic steel is intercritical annealing of a ferritic-pearlitic microstructure in

the + two-phase field, followed by a sufficiently rapid cooling to enable the austenite to

martensite transformation.

Three basic approaches exist for the commercial production of dual-phase steels: (a) the

as-hot-rolled approach, where the dual-phase microstructure is developed during the con-

ventional hot-rolling cycle by careful control of chemistry and processing conditions [6

12], (b) the continuous annealing approach, where hot- or cold-rolled steel strip is un-

coiled and annealed intercritically to produce the desired microstructure [13, 14] and (c)the

batch-annealing, where hot- or cold-rolled material is annealed in the coiled condition

[1519].

Box- or batch-annealing was mainly considered for economical and practical reasons by steel-

makers where continuous facilities were not available. Dual-phase steels could be obtained by

means of batch-annealing in the intercritical region for approximately 3 h to ensure homogene-

ity, followed by very slow cooling. Due to the extremely low cooling rates, much higher alloying

contents were necessary to achieve the desired hardenability (i.e. 2.5 % Mn, 1.5 % Si, 1.0 % Cr).

On the other hand, dual-phase steel production in the hot strip mill demands precise control of

the transformation, because the transformation starts from single phase austenite. The

determination of an accurate CCT- diagram via dilatometry, where the influences of alloying

elements, heat treatment conditions and desired properties are integrated, is of great impor-

tance for the success of the process. The critical point is the correction of the CCT- diagram

to include the presence of strain in austenite (to simulate the real process conditions).

The use of continuous annealing lines (CAL) offers the advantages of high production rates,

better uniformity of the steel properties and, furthermore, the possibility of using lower alloyed

steels (having a lean chemistry). In continuous annealing lines three types of cooling are utilized

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[20]: (a) water-quenching, (b) gas-jet cooling and (c) air-cooling. The use of CAL equipped

with water quenching facilities makes possible an easy and economical production of dual-

phase high-strength steel. The basic heat cycle involves intercritical annealing and subsequent

water quenching to form the ferritic-martensitic microstructure. If necessary (according to

application), a tempering stage follows [13,2123].

From the above short review it becomes clear, that dual-phase steels can be produced by

cooling from the annealing temperature (intercritical or austenitic) by any cooling rate in the

range between batch-annealing and water-quenching. Since every steel producer has different

melting, rolling and cooling facilities, the choice of the alloying elements best suited to the

existing production capabilities is mandatory. Thus, a single widely accepted alloy composition

for each grade of the dual-phase steel family is out of consideration. There are many combi-

nations of alloying elements such as Mn, Si, Cr, Mo and V that can be added to low carbon

(0.1 wt. % C) iron to obtain the desired ferritic-martensitic microstructure. A basis with less

than 0.15 wt. % carbon (for weldability reasons) and 1.5 wt. % Mn is generally acceptable. To

achieve good ductility and toughness, an initial carbon content of about 0.1 wt. % C is ideal, so

that the carbon content and the morphology of martensite can be controlled. For strong and

tough martensite this value is approx. 0.4 wt. % C, depending on the total alloy composition.

For martensite carbon contents higher than 0.4 wt. % C twinned martensite may be formed.

Concerning the other alloying elements, manganese (Mn), chromium (Cr), molybdenum (Mo),

silicon (Si) and vanadium (V) are commonly used to increase the hardenability of austenite

[11,14,19,2428]. The role of Si and Mn is more complex, since they may also contribute to

solid solution strengthening of ferrite and hence to the strength level of the steel. The additionof elements such as Cr, Mo and V that promote carbide formation demands a careful process

control.

For each applied alloy chemistry there exists a critical value of the cooling rate which sets the

lower limit for the formation of a ferritic-martensitic microstructure. Applying cooling rates

lower than this value results in a ferritic-pearlitic microstructure while higher cooling rates are

capable of producing a microstructure consisting of martensite islands in a ferritic matrix. The

overcooling degree is responsible for the martensite fraction in the final product. Equivalently,

the critical cooling rate (which as previously shown is prompted/necessitated by the production

line capabilities) defines a minimum of alloying content for the dual-phase steels.Based on experimental results and accepting a similar alloying behavior of Cr and Mo as Mn

(mainly due to the hardenability effects), Tanaka et al. [14] expressed the influence of alloying

elements on the critical cooling rate in terms of a manganese equivalent, Mneq[%]:

logCR [K/s] = 1.73 (Mn)eq[%] + 3.95 (1.1)

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where

(Mn)eq= (Mn) + 1.3 (Cr) + 2.67(Mo)[%]. (1.2)

Though alloying elements may look versatile via equation eq. 1.1, their effects on microstructureand deformation behavior are not necessarily the same. It is clearly shown that the higher the

amount of alloying elements in the steel the lower is the critical cooling rate necessary to form

a dual-phase microstructure. This explains the reason why batch-annealing processes require

much higher alloying additions. Practically, the critical cooling rate is easily determined with

a series of simple cooling experiments for each individual alloy.

To understand the formation of the ferritic-martensitic microstructure and to be able to in-

terpret the products of the heat treatments it is essential to consider the phase transformations

taking place during heating, intercritical annealing and quenching. The formation of austen-

ite in low carbon C-Mn steels was studied by a number of investigators [ 2933]. In the case

of cold-rolled ferritic-pearlitic steels, the recrystallization of the cold-worked microstructure is

completed already before reaching the annealing temperature, even during the rapid heating

rates applied on most continuous annealing lines. By entering the intercritical two-phase re-

gion, austenite nucleates rapidly at pearlite or in the vicinity of cementite particles and grows

rapidly until the carbides are dissolved. A slower growth of austenite into ferrite is continued

at a rate initially controlled by carbon diffusion in austenite and finally by manganese diffusion

in austenite, until the system reaches the equilibrium state. Practically, due to the very short

annealing times, only carbon redistribution between the phases takes place, because substitu-

tional manganese diffuses much more slowly than interstitial carbon. This effect is referred to

as paraequilibrium.

Assuming that no manganese redistribution occurs, then a vertical section corresponding to

paraequilibrium conditions can be constructed for constant Mn content. Bearing that remark

in mind, Figure 1.3 demonstrates the production concept/scheme of a dual-phase steel by

intercritical annealing. According to the lever rule, for any given carbon content (C0) the

amount of austenite increases with increasing the intercritical temperature, becoming equal

to 100% at the Ac3 -temperature while, as a consequence, the carbon content of austenite

decreases, reaching its minimum value (C

=C0). In an analogous manner, for any given inter-critical annealing temperature the amount of austenite increases with increasing alloy carbon

content, becoming equal to 100 % at a carbon content corresponding to the/+ boundary

(C0=C). Additional alloying elements may cause some changes in the austenite formation

process and/or even widen or tighten the intercritical temperature field. Such alloying effects

can be roughly estimated by thermodynamical codes (e.g. ThermoCalc) but remain beyond

the scope of this work.

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Figure 1.3: Production of dual-phase steels by intercritical annealing. The equilibrium fractionsof austenite and ferrite as well as their carbon content at the annealing temperature can be easilyestimated by applying the lever rule.

The importance of intercritical annealing becomes apparent, since for a given alloy compo-

sition and for a pre-selected annealing temperature the maximum amount of austenite that

can be (ideally) transformed to martensite as well as its carbon content -i.e. its hardenability-

are already pre-determined. Therefore, high intercritical annealing temperatures result in high

austenite fractions of decreased hardenability while low annealing temperatures result in low

austenite fractions with increased hardenability.

Even though the products of the austenitic transformation are strongly dependent on the

intercritical annealing parameters, the cooling rate is the final decisive step for the production

of dual-phase steels. The combined influence of cooling rate and intercritical annealing on

the formed microstructures was studied individually by many investigators [26,28,3436]. In

each case, a very important parameter that should not be forgotten is the effect of alloying

elements on the stability of austenite, which can be qualitatively measured by means of the

martensite-start temperature (MS). For low carbon steels it was proposed by Andrews [37]

that:

MS [C] = 539 423C 30.4 Mn 17.7 Ni 12.1 Cr 7.5 Mo (1.3)

or similarly by Eldis [38] for dual-phase steel compositions (also in wt. %):

MS [C] = 531 391.2 C 43.3 Mn 21.8 Ni 16.2 Cr. (1.4)

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Quenching with high cooling rates from low intercritical temperatures ensures a complete

transformation of austenite into martensite because of the austenite stability; by applying

relatively low cooling rates the intercritically formed austenite partly transforms into ferrite,

enriching further the remaining austenite with carbon, increasing its hardenability and lower

even more the MS-temperature. In this case, isolated retained austenite particles may be

detected in the microstructure in the form of interlath films. However, the formation of

ferrite-carbide aggregates is usually unavoidable. Rapid quenching from high intercritical

temperatures produces even more complex effects: the decreased austenite hardenability

and the absence of time for the necessary diffusions during cooling, both reflected in a high

MS-temperature, result in the formation of autotempered martensite.

Dual-phase steels exhibit a number of superior mechanical properties, such as continuous

yielding behavior, low 0.2 % offset yield strength, high work hardening rate, high tensile strength

and remarkably high uniform and tensile elongations. The mechanical properties of dual-phase

steels arise from structural features, that is the fine dispersion of hard martensite particles

in a ductile ferritic matrix and all the related phenomena that accompany this coexistence.

The yielding and the work hardening behavior have been interpreted in terms of the high

dislocation densities and residual stresses arising in ferrite, as a consequence of the volume

expansion associated with the austenite to martensite transformation. The strength of dual-

phase steels was found to be dependent primarily upon the volume fraction and the carbon

content of martensite; solid solution strengthening of ferrite may also contribute to strength.

The excellent ductility reported for most of the dual-phase steels is the combined result ofmany factors. Among them are the volume fraction and the carbon content of martensite, the

ductility of martensite, topological parameters such as the martensite grain distribution in the

ferritic matrix, the alloy content of ferrite, the dislocation density in ferrite, the presence of

carbides and/or retained austenite. Additionally, lattice imaging from Koo and Thomas [24]

has revealed a good coherency at the ferrite/martensite interface, which prevents decohesive

interface failure during deformation and thus enables the full toughness of ferrite to be realized.

Tempering may be applied as part of the process in some continuous annealing lines, after

water-quenching from the intercritical temperature, to regulate the properties of the dual-phase

steel. It may also be an unavoidable side-effect of an operation, e.g. in a hot dip galvanizingline. Finally, tempering may be useful in some production practices such as bake hardening

(paint baking). The change in yield strength upon tempering is complex because of the relief

of residual stresses, carbon segregation to dislocations and the return of discontinuous yielding.

After tempering at low temperatures the yield strength increases but discontinuous yielding

returns only to the steels containing low volume fractions of martensite. When tempering at

high temperatures the yield strength decreases but discontinuous yielding appears in all steels.

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The tensile strength decreases, while post-uniform and uniform elongations increase due to the

change in hardness of martensite.

There has been a long discussion on the construction of a model correlating the mechanical

properties of dual-phase steels with their microstructural characteristics. Simple empirical rules

of mixtures [2628,3944] (usually linear regressions between the constituent phases proper-

ties, based on Mileikos theory of composites [45]) as well as more complicated/sophisticated

micromechanical models [4651] (introducing the importance of the effective in-situ prop-

erties and topological microstructural parameters) have been developed, most (if not all)

of them based on individual sets of experimental data. In some cases, there was a good

agreement between the predicted properties and the experimental results -even for martensite

fractions in the order of 80 %, in other cases some fair and well-established modifications have

to be made. Since each study refers to a specific alloy composition and a different kind of

heat treatment, a comparison between the obtained results contributes to the disagreements

reflected in literature over the past 25 years.

In order to meet the different design requirements of individual automobile-body components

for strength, crashworthiness, energy absorption, part complexity and dent resistance, a variety

of dual-phase grades exhibiting different strength and ductility levels is currently industrially

produced. Despite the numerous studies on the relationship between the mechanical properties

and the microstructural characteristics of dual-phase steels over the last decades, the chal-

lenge of increasing their formability at a constant strength level (or equivalently increasing the

strength while maintaining a high ductility) remains still unanswered.The statement of many researchers that for the improvement of properties of dual-phase

steels the ferrite should be fine-grained, free of ultrafine carbide precipitates and strength-

ened solely by alloying additions which have minimum effects on ductility is still not fully

explained/affirmed.

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1.2 Grain refinement

The formation of ultrafine-grained ferritic microstructures in low-alloyed and plain carbon steels

has been intensively investigated during the last decade, since grain refinement is expected to

have a beneficial effect on the mechanical properties of the steel. Though the term ultrafineis somewhat vague, it reflects the objective of the Japanese super metal project of producing

a strip having a ferrite grain size of 1 m throughout a minimum thickness of 1 mm [52].

According to the Hall-Petch relation (which is applicable to a variety of polycrystalline single-

and dual-phase metals), a decrease in grain size (d) results in an increase in yield strength (y):

y = 0+ kyd1/2. (1.5)

For a low-alloyed steel, for example, a decrease of the ferrite grain size from 5 m to 1 m wouldideally produce an increase of the yield strength by approximately 300 MPa. Additionally, it

has been reported that grain refinement improves the fatigue resistance of steels [ 53] and can

lead to superplastic behavior at high temperatures and appropriate strain rates [54].

The currently available techniques to obtain ultrafine grains are rapid solidification directly

from the melt, vapor deposition, cryogenic metal-forming, mechanical alloying and severe

plastic straining. Very small grains (with sizes in the nano-scale range) may be produced under

extreme conditions, often leading to impressive physical and mechanical properties. However,

due to the the limited production quantities and the very small grains individually formed

(powder-like), further consolidation and processing is required to produce a bulk material suit-able for structural use. Concerning the refinement methods, there exists a permanent conflict

between the achievable grain size in a material, the amount or the dimensions of the material

that can be processed in this way and -the most important- the cost of processing. In the

case of steel applications, an optimum compromise between these factors would be required [ 55].

Grain refinement in steels can be realized by microalloying. The addition of elements such as

Al, B, P, Sm or Ti in the microstructure can suppress the grain growth of ferrite, utilizing the

pinning effect of secondary phase particles and/or the dragging effect induced by solute atoms

[56]. However, the idea of controlling the grain size and hence the mechanical properties by

thermomechanical processing instead of the classical way of alloying is far more attractive, as

this would result in the production of steels with simpler chemistries and improved recyclability

and would lead to economic benefits as well.

To achieve steel grain refinement, substantial efforts involving severe plastic deformation

(SPD) by using conventional rolling equipment have been made. According to the dynamic

Strain Induced TRansformation method (SITR), proposed by [5761], grain refinement is pro-

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moted by the continuous intragranular nucleation of numerous small ferrite grains during hot-

or warm-rolling of austenite. The process involves a single-pass rolling of the steel strip at

a temperature just above the Ar3 (i.e. immediately above the temperature at which grain

boundary pro-eutectoid ferrite would begin to form) but below the Ae3, followed by air-cooling

or accelerated quenching depending on the desired microstructure. The required rolling reduc-

tion ranges between 35 and 40 %. Ultrafine ferrite (UFF) grains withdF


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(a) HPT (b) ECAP (c) ARB

Figure 1.4: Schematic illustration of the recently developed methods for severe plastic deformation(HPT and ECAP from [55], ARB from [66]).

nique for obtaining ultrafine-grained microstructures [6871]. According to the method, the

material is pressed through a die, where two channels form an L-shaped configuration with an

angle of 2 (Figure 1.4-b). The process imposes a severe strain on the sample by means of

shearing. Since there is no concomitant change in the cross-sectional dimensions of the samples,

repetitive pressings can produce very high effective strains (in the order of 10), achieving ho-

mogeneous grain refinement with grain sizes on the micro-meter scale. Moreover, the sample is

constrained so also less ductile materials can be processed. The factors influencing the method

are the pressing route by which the sample is rotated during the successive pressings, the dieangle which determines not only the strain per pass but also the geometry of deformation,

the die cross-section geometry, the speed and the temperature associated with pressing. Very

recent investigations revealed that ultrafine-grained ferritic-martensitic dual-phase steels can

be fabricated by ECAP, by applying an effective strain of 4 at 500 and subsequent intercrit-

ical annealing at 730 for 10 min. UFF grains with uniformly distributed blocky martensite

islands of 1 m were produced [72]. Despite the scientific interest and the partial success, the

application of the ECAP method is still restricted in terms of commercial capability while even

the future prospects for steel sheet production are questionable. The potential of up-scaling

ECAP is being currently investigated.Accumulative Roll Bonding (ARB) is also a newly developed technique to realize intense

plastic straining [66, 73, 74]. Following the illustration of Figure 1.4-c, one strip is neatly

placed on top of another strip and the two layers of material are joined together by warm

rolling. The rolled sample is then cut in two halves, which are again stacked together after

an appropriate surface preparation and are roll-bonded. The whole process can be repeated

without any change in the samples thickness and geometry, given that the reduction in thickness

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after each rolling pass is maintained/controlled to 50 %. This process can theoretically introduce

unlimited amounts of strain in the material. Effective strains of 8 have been reached for

aluminium and steels, resulting in grain sizes below 1 m. Critical factors for the success of the

method are surface preparation and cleaning, the deformation temperature and the amount of

induced strain. Although rolling at elevated temperatures is advantageous for joinability, too

high temperature would cause dynamic recrystallization and cancel the effect of accumulated

strain.

All the above proposed techniques are very attractive but seem to encounter several difficulties

in engineering applications. In each case, very high levels of strain or unrealistic inter-pass

times are required, complex processes and special equipment are involved and the production

capacities are still very low to cover the steel market demands. As a combined result of these

factors, the application of the novel refining methods in a continuous industrial process does

not seem possible in the near future.

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Chapter 2

Material production and experimental

procedure

2.1 Production of the material

The dual-phase steel investigated in this work was industrially produced by voestalpine Stahl

Linz. Its chemical composition is given in Table 2.1. Slabs with a thickness of 210 mm were

produced on a continuous casting machine. The slabs were reheated in a pusher-type furnace

to a temperature of 1250 and hot-rolled to a final thickness of 2.40 mm. The finishing

temperature was approximately 900 while the coiling temperature was about 600. Part

of the hot-rolled strip was milled to remove the surface scale and then cold-rolled to 1.00 mm

(cold-reduction of 58 %).

Table 2.1: Chemical composition of the investigated DP-steel.

DP steel C Si Mn Cr+Mo Fe

wt. % 0.1 0.15 1.5 0.8 Bal.

In the main bulk of this work strips from the hot-rolled ferritic-pearlitic material were used,

cut to specimens of 250 mm in length and 50 mm in width in order to satisfy the requirements

of the laboratory annealing and cold-rolling equipment. The industrially cold-rolled strip was

mainly used as reference material (denoted as R in the following) with regard to its response

to the same heat treatment schedules applied to the pre-processed and laboratory cold-rolled

material.

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2.2 Thermodynamical calculations

The phase diagram of the alloy as well as the critical temperatures where phase transformations

occur were initially calculated with the help of the program ThermoCalc, by using the database

TCFE3 and considering ferrite, austenite and cementite as the only phases present in the steel.The intercritical + two-phase temperature range was calculated between 710 and 815.

Additionally, the fraction of austenite at different intercritical annealing temperature steps was

determined, in a first attempt to set the maximum martensite fraction after cooling down from

the annealing temperature (Figures 2.1-a and 2.1-b). It should be taken into account that

all estimated temperatures and phase fractions represent thermodynamic equilibrium.

(a) (b)

Figure 2.1: Thermodynamic calculations: (a) concentration section through the phase diagram of

the investigated dual-phase steel (the dashed line indicates the carbon content) and (b) equilibrium

fraction of austenite during intercritical annealing.

2.3 Pre-processing

The first critical stage of the experimental procedure, described with the term pre-processing,

involves annealing of the as hot-rolled material in a Carbolite Three Zone Tube furnace (TZF)

with an adapted inert gas (Ar) supply to avoid prolonged oxidation. The specimens were

annealed at three different temperatures (760, 800 and 900) with a holding time of 5 min in

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an Ar atmosphere, followed by water-quenching to form a ferritic-martensitic (in the first two

cases) or a completely martensitic microstructure (Figure 2.2). The annealing temperatures

were selected by taking into account the thermodynamical calculations and the holding time

was set long enough to assure thermodynamic equilibrium.

(a) DP I (b) DP II (c) M

Figure 2.2: Micrographs of the pre-processed materials before laboratory cold-rolling, etched withLePera.

In the following chapters, the pre-processed materials-specimens will be denoted as DP I,

DPIIandM, according to the annealing temperature during pre-processing (Table 2.2).

Table 2.2: Denotation of pre-processed material.

Grade Pre-processing conditions Martensite fraction (%)DP I Tan= 760, tan= 5 min, WQ 23 %DPII Tan= 800, tan= 5 min, WQ 40 %M Tan= 900, tan= 5 min, WQ 100 %

2.4 Cold-Rolling trials

The ferritic-martensitic or completely martensitic microstructures produced during pre-

processing were cold-rolled in a Carl-Wezel laboratory mill with a roll diameter of 220 mm

at a roll peripheral speed of 16 mmin1 to a final thickness of 1.00 or 0.80 mm (cold reduction

of 58 % or 67 %, respectively). To avoid cracking of such work hardened microstructures, rolling

was performed in multiple passes.

The microstructures of the as hot-rolled ferritic-pearlitic starting material after industrial

cold-rolling as well as that of the pre-processed material after laboratory cold-rolling are shown

in Figure 2.3. There, the microstructures differ not only qualitatively but also quantitatively

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from each other. Micrograph (a) shows a rather coarse-grained ferritic-pearlitic microstructure,

while in micrographs (b) and (c) dual-phase ferritic-martensitic microstructures are shown,

possessing different martensite fractions due to the initial annealing conditions. The higher

the martensite fraction after pre-processing the finer is the obtained microstructure after

cold-rolling. This is also supported by the last micrograph (d), where a very fine martensitic

microstructure is shown.

(a) (b)

(c) (d)

Figure 2.3: (a) Industrially cold-rolled ferritic-pearlitic microstructure of the as hot-rolled startingmaterial (R), (b) Cold-rolled ferritic-martensitic microstructure, initially annealed at 760 (DP I),

(c) Cold-rolled ferritic-martensitic microstructure, initially annealed at 800 (DP II), (d) Cold-rolledcompletely martensitic microstructure (M), etched with Nital.

2.5 Dilatometric investigations

The influence of various annealing cycles on the microstructural evolution of the cold-rolled

material (with respect to grain refinement) was studied via dilatometry, with the samples cut

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in the rolling direction (10 x 4.0 x 1.0 or 0.80 mm, depending on the sheets thickness). The

dilatometric investigations were conducted on a Bahr dilatometer DIL 805 A/D. The selection

and the appropriate combination of the annealing parameters within a cycle (heating rate, an-

nealing temperature, holding time in the intercritical zone as well as the cooling rate) were made

with regard to the special features of the cold-rolled material, that is the ferritic-martensitic or

the pure martensitic microstructure and the deformation degree. The basic annealing schemes

are presented in Figure 2.4.

(a) (b)

Figure 2.4: (a) Intercritical annealing, (b) Repeated heating cycles oscillating between the+and

phase fields.

2.6 Annealing simulations and austenitization kinetics

All annealing simulations were conducted in the laboratory with the Multi-Purpose Annealing

Simulator (MULTIPAS) at voestalpine Stahl Linz on the cold-rolled pre-processed ferritic-

martensitic and martensitic as well as on the reference material. Special attention was paid

to the intercritical annealing heat treatments due to their applicability on an industrial scale.

Specimens with a thickness of 1.00 mm were preferred for this purpose, because they offer

the possibility of direct comparison with the standard industrially cold-rolled material (R,reference).

Although the phase transformation temperatures were calculated with ThermoCalc and the

recrystallization kinetics was investigated via dilatometry, additional annealing simulations were

conducted on the reference material to determine the austenitization kinetics of the alloy in

non-equilibrium conditions on a larger scale. Specimens were annealed in a temperature range

that covers both + two-phase andsingle-phase regions for different representative holding

times and then were water-quenched.

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2.7 Microstructure investigations

A light microscope (Olympus AX70) was used for the microstructural investigations of the

heat treated samples. To reveal the microstructure of the fine-grained materials the specimens

were conventionally prepared and then etched with LePeras etchant, which is a mixture of 1 %sodium metabisulfite in distilled water and 4 % picric acid in ethanol in a 1:1 volume ratio.

This tint etching technique allows the distinction of phases by coloring, staining ferrite brown

and/or blue, bainite dark brown to black while martensite and retained austenite (hardly any

present in the materials of this work) remain white. Quantitative characteristics of the mi-

crostructure such as phase fractions and mean grain sizes of the constituents were determined

by line intercept measurements, considering ferrite as the dominating matrix phase and charac-

terizing martensite, bainite and/or retained austenite as a second phase. For severely deformed

microstructures after cold-rolling and before heat treatment or for annealed specimens where

LePeras etching agent could not produce the desirable effects, a 2 % Nital etchant (2 % nitricacid in ethanol) turned out to be an acceptable alternative.

In cases that a higher magnification analysis beyond the resolution capacity of light micro-

scope was required, e.g. for the investigation of ultrafine-grained microstructures or for the

identification of any third phase present (like bainite), Scanning Electron Microscopy (SEM)

was employed. The SEM observations were conducted on a LEO 1450 SEM and/or on a TOP-

CON SM-520 Field Emission Gun (FEG)-SEM on demand.

Microstructures of selected specimens representing critical heat treatments were further in-

vestigated by means of Transmission Electron Microscopy (TEM). TEM-specimens (thin foils)

were prepared by gentle mechanical thinning down to 80 m followed by electrolytic thinning

in 5 % perchloric acid in acetic acid. The thin foils were analyzed in a Philips CM20 STEM

transmission electron microscope, applying an accelerating voltage of 200 kV. For a first rough

overview of the microstructure, secondary electron (SE) images were taken from the thin foils.

The detailed analysis of the phases was carried out using bright and dark field techniques, while

the identification of the phases was performed using electron diffraction patterns.

2.8 Mechanical testing

2.8.1 Tensile testing

The mechanical properties of the steels were measured on a Roell-Korthaus RKM 250 testing

machine, according to European standard EN 10 002. All tensile specimens were machined

with their tensile axis parallel to the rolling direction and with a gage length of 80 mm. The

laboratory produced grades (DP I, DP II and M) were tensile tested in the as-annealed condition

while the industrially cold-rolled steel (R) was submitted to skin pass rolling of 0.5-1.0 % to

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improve roughness.

2.8.2 Ultramicrohardness and microhardness testing

The ultramicrohardness of the main constituents of the dual-phase microstructure (ferriteand martensite) was measured with a Kammrath & Weiss ultramicrohardness tester UMHT-3

equipped with a Vickers diamond square pyramid. The device was mounted in a LEO 1450

Scanning Electron Microscope (SEM). To identify the phases, the specimens were etched with

Nital etchant. The applied indentation load (15 mN) as well as other critical measuring param-

eters like indentation time and speed were selected to be the same for all measurements, so that

a comparison of the hardness between different phases and various heat treatments becomes

possible. Furthermore, to eliminate any possible errors derived from the optical measurement

of the indentation diagonals, all indentations and measurements in SEM were performed at the

same magnification (12000). Finally, the hardness was calculated from eq. 2.1:

HV=189 F

d2 , (2.1)

where d [m] stands for the mean length of the indentation diagonals and F [mN] for the

indentation load.

Additionally, the Vickers microhardness of the investigated dual-phase steels was determined

with a Micro-Duromat 4000 E microhardness tester from Reichert-Jung mounted in a light mi-

croscope (Reichert Metaplan 2, Leica AG). The indentation load was set to 150 p (approximately

1471 mN), producing Vickers impressions which are bigger by one order of magnitude compared

to ultramicrohardness measurements.

2.8.3 Hole expansion measurements (Stretch flangeability)

Hole flanging is a process widely applied in thin-sheet forming operations, which employs a

punch for producing structural parts with short necks that are subsequently used for assembly

with other components. During stretch flanging, the deformation mode at the edge of the hole

is a combination of bending and stretching which in some cases causes splitting failure and

therefore cannot be grasped by the conventional uniaxial tensile test.

The hole expansion behavior provides a way to measure the tendency of steels to split as a

hole is expanded under external forces and is characterized by the percentage increase in the

size of the hole at the moment that a crack forms [7579]. A schematic diagram of the hole

expansion equipment is shown in Figure 2.5. Selected dual-phase steel grades were cut to

125mm 125 mm square test pieces and before testing a 12 mm (0.15 mm) hole was punched

into the centre of each sample. The hole expansion test is conducted by expanding the initial

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punched hole using a 50 mm diameter punch; the punch driving is immediately stopped when

any crack (which extends all through the samples thickness) is observed at the edge of the hole.

The final diameter of the hole is measured by averaging two readings taken perpendicularly to

each other. The property is expressed as the ratio of the expanded hole size to the original hole

size, as defined by the following equation:

= df d0

d0 100, (2.2)

where [%] is the hole expansion ratio, df[mm] the average final hole diameter (after rupture),

and d0 [mm] is the initial hole diameter.

Figure 2.5: Schematic illustration of hole expansion testing equipment of voestalpine Stahl Linz

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Chapter 3

Results

3.1 Microstructure investigations

3.1.1 Recrystallization/austenitization kinetics

In order to study the austenitization/recrystallization kinetics regarding cementite dissolution

and austenite formation, specimens from the industrially cold-rolled material (R) were subjected

to heat treatments according to the annealing plan ofFigure 3.1which involves annealing not

only in the intercritical two-phase +ferritic-austenitic field but also in the pure austenitic

region. Five annealing temperatures covering a range of 100 in temperature intervals of 25

for holding times of 0 s, 10 s, and 100 s were applied, while the heating rate was set to 25 K/s

to reproduce the industrial conditions. After water-quenching the microstructure of the steelwas investigated.

Figure 3.1: Annealing plan to determine theaustenitization kinetics of the as cold-rolleddual-phase steel.

Figure 3.2: Influence of annealing tempera-ture and annealing holding time on the frac-tion of martensite.

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The fraction of martensite formed during water-quenching was determined by line intercept

measurements, providing the amount of the former austenite prior to the martensitic trans-

formation. Figure 3.2shows the influence of the annealing temperature for different holding

times on the martensite fraction. Although the martensite fraction increases dramatically in

the first 10 s of annealing, a further increase of the holding time has no significant effect except

for the annealing temperature of 775. Figure 3.3shows the microstructure of the steel (R)

as a function of annealing temperature for a holding time of 100 s. It is remarkable that for

annealing temperatures over 800 (even though the calculated + transformation

temperature is 815) the martensite fraction approaches the maximum value of 1.0.

(a)Tan= 750 (b)Tan= 775 (c)Tan= 800

(d)Tan= 825 (e)Tan= 850

Figure 3.3: Influence of the annealing temperature on the martensite fraction (former austenite

fraction before water-quenching) for a holding time of 100 s (LePera).

Figure 3.4 shows the microstructure of the steel after water-quenching from the annealingtemperature without a holding stage. For the lower annealing temperature (Figure 3.4-a), the

recrystallization is not completed and undissolved carbides are still detected in the microstruc-

ture. At higher annealing temperatures no carbides could be detected. It should be also noticed

that even after water-quenching from high annealing temperatures (825) a significant amount

of ferrite is present in the microstructure (Figure 3.4-d).

The zero holding time (0 s), which eliminates the possibility of equilibrium, provides valuable

data/information for the case that heat treatments schedules are applied involving flashing

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heating and cooling cycles (see Figure 2.4).

(a)Tan= 750 (b)Tan= 775 (c)Tan= 800

(d)Tan= 825 (e)Tan= 850

Figure 3.4: Influence of the annealing temperature on the martensite fraction (former austenitefraction before water-quenching) for a zero holding time (LePera).

The martensite fraction data are appended to the ThermoCalc diagram in order to compare

the experimental results with the numerical calculation. It was assumed that the martensite

fractions measured for a holding time of 100 s approach equilibrium status and additionally

that the amount of martensite represents the former austenite fraction during annealing. As

can be seen inFigure 3.5, the results are in good agreement with each other.

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Figure 3.5: Experimental data from martensite fraction measurements (square symbols) and equilib-rium fraction of austenite (line) as calculated with the program ThermoCalc, both given as a function

of annealing temperature.

3.1.2 Dilatometric investigations

The laboratory developed materials (DP I, DP II and M) were not submitted to the above water-

quenching process, since their recrystallization and austenitization kinetics are expected from

their design concept to be even faster. Nevertheless, systematic dilatometric investigations were

done in order to study the influence of the critical annealing parameters on the microstructureof these grades, aiming to clarify the recovery-recrystallization-grain growth mechanism and

finally to determine the optimum annealing conditions which lead to grain refinement.

Conventional annealing

Conventional heat treatment schedules following the annealing scheme of Figure 2.4-a were

applied to the laboratory cold-rolled microstructures. A relatively high heating rate of 25 K/s

from room to annealing temperature was applied, in order to simulate the industrial conditions

and additionally to minimize the time available for grain growth after recovery and recrystal-

lization. Three different annealing temperatures (Tan= 750, 800 and 840) were chosen

for the dilatometric investigations. According to the thermodynamical calculations, the first

two temperatures are located in the intercritical + while the third one in the austenite

phase field. However, the ferrite to austenite phase transformation as indicated by the dilata-

tion curves during the heating stage seems to be completed only at temperatures above 840

(see Figure 3.6). This important observation which represents the in-situ measurement of

the + transformation temperature does not actually contradict the results from the

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austenitization experiments, since in the measurement of the dilatation curve no holding time

is taken into account. This simply means that a temperature of 840 refers to intercritical

annealing conditions for a zero holding time and to austenitic annealing in case that a holding

time stage is applied. To avoid any possible misinterpretations in the discussion of the results,

Tan of 840 will be considered as austenitic annealing in the following pages.

Figure 3.6: Dilatation curve during heating. T indicates the temperature at which the ferrite to

austenite transformation is completed.

Holding times in the range between 5 s and 120 s were applied during intercritical annealing.

Recovery and recrystallization of the severely deformed microstructures were completed within

the first 30 s of holding stage for all materials investigated. The cold-rolled martensitic grade

(M) is already recrystallized at an even shorter holding time of 10 s. Microstructural observa-

tions revealed that holding times longer than 30 s are rather detrimental for the formation of an

ultrafine dual-phase microstructure, by leading to an undesirable grain growth (mainly of the

ferrite grains). Figure 3.7demonstrates the influence of holding time on the microstructure of

the grade M, intercritically annealed at 800 and quenched with 40 K/s to room temperature.

The specimens are treated with Nital etchant, which stains martensite brown or dark brown

and ferrite light brown to light cream/beige.Recrystallization is completed even for the shortest annealing time applied. The design and

the production history of the grade M as well as the high cooling rate after annealing eliminate

the possibility of presence of any carbide phases in the final microstructure. Holding time does

not seem to affect the fraction of martensite maintained in the steel after quenching but it is

proved to be a dominating parameter for the achievement of ultrafine-grained microstructures.

A holding time of 30 s provided the optimum solution with respect to grain refinement, as shown

in Figure 3.7-d. Further increase of the annealing time, for example by only 30 s, resulted to

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(a)tan= 5 s (b) tan= 10 s

(c)tan= 20 s (d) tan= 30 s

(e)tan= 50 s (f) tan= 60 s

Figure 3.7: Microstructures of the martensitic grade M after annealing at 800 for the subscribedholding times and quenching with a cooling rate of 40 K/s, etched with Nital.

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rapid grain coarsening of ferrite almost by an order of magnitude as shown in Figure 3.7-f.

Analogous observations regarding the time-dependence of grain refinement at certain annealing

conditions were also made for all laboratory grades investigated.

(a)Tan= 750 (b)Tan= 800 (c)Tan= 840

Figure 3.8: Microstructures of the martensitic grade M after annealing at different annealing tem-peratures for 30 s and quenching with a cooling rate of 40 K/s, etched with Nital.

Based on the experimental determination of an optimum holding time for grain refinement

(30 s in our case), the influence of annealing temperature on the microstructure was addition-

ally studied in the dilatometer before moving to a larger experimental scale. As shown in

Figure 3.8, annealing at higher temperatures leads to higher martensite fractions in the mi-

crostructure, consistent with the increased austenite fraction at these conditions. The difference

in martensite fraction between Figure 3.8-band Figure 3.8-c, referring to annealing temper-

atures of 800 and 840, respectively, is not clearly distinguishable, since for the holding time

applied the austenite fraction in both cases approximates the maximum. It should be noticed

that this temperature dependent increase of the martensite fraction is accompanied by a pro-

nounced alteration of the morphology and subsequently of the etching behavior of martensite

grains. Annealing at lower intercritical temperatures (e.g. at 750, Figure 3.8-a) produces

clear unstructured brown martensite, while higher annealing temperatures (Figures 3.8-b and

3.8-c) lead to the formation of dark-brown structured and generally coarser martensite (the

tints refer to the etching effect of Nital agent).

In all dilatometric investigations described above, a moderate cooling rate of 40 K/s wasapplied in order to select the most favorable conditions/parameters for grain refinement with

respect to annealing temperature and annealing time. Once this became clear, the influence

of cooling rate on the microstructural evolution was thoroughly examined. Greater attention

was paid to the grade M because of its starting microstructure peculiarity in comparison with

the other grades. In order to produce a wider possible range of dual-phase microstructures

depending on the cooling rate parameter, six cooling rates, that is 5, 10, 20, 40, 60 and 80 K/s,

were applied. Higher cooling rates were not achievable along the entire cooling stage, due

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to the restricted cooling capacity that the dilatometer -associated with the cooling gas used

(nitrogen)- could provide.

Figure 3.9 demonstrates the influence of cooling rate on the microstructure of the martensitic

grade M, after annealing at 840 for 30 s. By applying cooling rates below 10 K/s a rather

coarse multi-phase microstructure, containing ferrite, martensite and a third phase indicated

by the black tinted regions located on the ferrite grain boundaries (most probably bainite)

is formed. Grain refinement is achieved only at cooling rates higher than 40 K/s, which is

sufficient to avoid the presence of the third phase. Further increasing of the cooling rate results

also in ultrafine ferritic-martensitic microstructures with increased martensite fractions.

These observations are only qualitatively common to all grades (DP I, DP II and M). More

details about the impact of conventional heat treatments on the microstructure, with regard

to the individual annealing behavior of each laboratory grade investigated, will be described

and discussed in the section of larger scale and more comprehensive annealing simulations (see

3.1.3).

Non-conventional annealing

Non-conventional heat treatment schedules following the concept described in Figure 2.4-b

were applied to the laboratory cold-rolled material, particularly to the thinner grades (thick-

ness of 0.80 mm, corresponding to 67 % cold reduction), where recrystallization kinetics is ex-

pected to be extremely fast due to the higher cold deformation degree. The aim of these

heating-cooling cycles around the + transformation temperature was to provoke the

continuous formation of new grains, so to impede grain growth. By this repeated nucleation andpreventing the system from reaching an equilibrium state by rapidly changing the temperature,

an ultrafine-grained microstructure is finally obtained.

The selection of annealing parameters, i.e. the initial heating and final cooling rates, tem-

perature range and heating/cooling rates for the + / - cycling had to be done for each

laboratory grade separately, thereby taking into account its production history and its peculiar

microstructural characteristics. The special case in which only one cycle is applied actually

represents a limiting condition toward conventional annealing with 0 s holding time. It was

observed that more than four cycles have no further beneficial effect towards grain refinement;

on the contrary, this could even lead to grain coarsening.

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(a)CR= 5 K/s (b) CR= 10 K/s

(c)CR= 20 K/s (d) CR= 40 K/s

(e)CR= 60 K/s (f) CR= 80 K/s

Figure 3.9: Microstructures of the martensitic grade M after annealing at 840 for 30 s and quench-ing with different cooling rates, etched with Nital.

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(a) (b)

(c) (d)

Figure 3.10: Ultrafine ferritic-martensitic microstructures produced from the following labora-

tory cold-rolled grades: (a) DPII (0.80 mm), (b) DP I (0.80 mm), (c) M (1.00mm) and (d) DP II

(1.00 mm). The applied heat treatments are given in the attached illustrations. Nital etchant.

Ultrafine homogeneous dual-phase steels produced by the thermal cycling technique are shown

in Figure 3.10. The specimens were etched with Nital: martensite appears dark brown to black

while ferrite remains light colored. Microstructures with a ferrite mean grain size between 2 and3 m and with martensite grains generally finer than 1 m could be obtained. Thinner grades,

corresponding to a higher cold deformation degree, react more sensitive than the thicker grades

to the dramatic changes of annealing conditions, in such a way that one or two flashing cycles

were enough to achieve grain refinement (Figure 3.10-a and 3.10-b). The small number of

heating-cooling cycles is simply translated into a shorter total holding time in the intercritical

or in the pure austenitic region and, finally, results in the formation of fine microstructures

possessing a relatively low martensite fraction. This agrees well with the recrystallization

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experiments described in 3.1.1, according to which for very short (or for zero) holding times

austenitization is not yet completed.

The in-situ variation of annealing parameters during heat treatment was successfully com-

bined with the rapid recovery and recrystallization kinetics of the laboratory cold-rolled pre-

processed grades, producing ultrafine dual-phase microstructures. Nevertheless, up-scaling the

steel production by this method faces serious difficulties, mainly due to the rapid temperature

changes necessary between the annealing segments and the complexity of controlling them with

the currently available industrial equipment.

3.1.3 Annealing simulations

Laboratory annealing simulations were conducted in order to reproduce the ultrafine mi-

crostructures on a larger scale and so being able to investigate mechanical properties. For

the final selection of annealing conditions the data collected from the dilatometric investiga-tions were taken into account.

The influence of the annealing temperature and of the cooling rate on the microstructure

was studied for all laboratory and industrial grades, since these parameters seem to have the

strongest impact on the microstructure evolution. To avoid grain coarsening, the heating rate

as well as the annealing holding time were held constant for all conventional heat treatment

schedules at 25 K/s and 30 s respectively. Two annealing temperatures were selected, the one

in the intercritical region (800, more or less standard in continuous annealing lines) and the

other in the austenitic region (840). Lower annealing temperatures (750) were excluded

from the simulations, because even though they provide with a variety of low martensite-

fraction steels they have a rather detrimental effect on grain refinement. For each annealing

temperature, a set of six cooling rates -the same as in the dilatometric investigations- was

applied, covering a wide range of microstructures regarding the formation of phases, the fraction

and the morphology of martensite and the mean grain size of ferrite.

Figure 3.11 shows the influence of the cooling rate on the microstructure of the laboratory

grade DP II after intercritical annealing at 800 for 30 s. The specimens are etched with

LePera. At cooling rates lower than 10 K/s a third dark/black colored phase (probably bainite)

is present in the microstructure, located on the grain boundaries and at the grain triple points

(Figures 3.11-a and 3.11-b). As the cooling rate increases, the third phase disappears while

the martensite fraction increases and, simultaneously, the ferrite mean grain size decreases

(Figures 3.11-cto3.11-f). Martensite grains remain fine for all cooling rates up to 40 K/s. In

the microstructures quenched with 60 K/s and 80 K/s also coarse martensite grains are found; in

the interior of these coarse grains a brown tinted substructure can be detected (Figures 3.11-e

and 3.11-f).

An analogous investigation for the martensitic grade M (by keeping the same annealing

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(a)CR= 5 K/s (b) CR= 10 K/s

(c)CR= 20 K/s (d) CR= 40 K/s

(e)CR= 60 K/s (f) CR= 80 K/s

Figure 3.11: Microstructures of the grade DP II annealed at 800 for 30s and quenched withdifferent cooling rates, etched with LePera.

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parameters constant) revealed a more pronounced influence of the cooling rate on the formed

microstructures, with respect to both grain refinement and martensite fraction. As shown in

Figure 3.12and in reference toFigure 3.11, a third phase could not be identified even for the

lowest cooling rates applied. Furthermore, coarse structured martensitic grains appear already

after cooling with 10 K/s, increasing in fraction/number with increasing cooling rate and finally

reaching grain sizes of the order of ferrite.

The cooling rate of 40 K/s proves to be a critical point in the direction of grain refine-

ment. Comparing the microstructure in Figure 3.12-cwith that inFigure 3.12-d, represent-

ing quenching with 20 K/s and 40 K/s respectively, the increase in martensite fraction as well

as the change in the martensite morphology - as indicated by its etching behavior - is distin-

guishable. This effect is followed by significant grain refinement of ferrite approximately by

an order of magnitude. Applying cooling rates higher than 40 K/s leads to higher martensite

fractions which are not accompanied with a proportional decrease in the grain size of ferrite.

The influence of the annealing temperature on the microstructure of the laboratory as well

as of the reference material is shown in Figure 3.13. For this purpose, a moderate cooling

rate of 20 K/s was applied in order to minimize the impact of the cooling rate on the formed

microstructures. At this condition, the industrial grade R reacts more sensitive to the annealing

temperature than the laboratory grades DP I and M. Increasing Tan from 800 to 840

results in a finer microstructure of R, possessing a higher martensite fraction (Figures 3.13-a

and 3.13-b). Nevertheless, small amounts of a dark-colored third phase are present in the

microstructure for both annealing temperatures.

On the other hand, the annealing temperature has no significant influence on the microstruc-tural characteristics of the grade DP I. Martensite fraction remains on the same level without

any morphological change while no third phase could be observed.

In the case of grade M, the higher annealing temperature has a major influence on the

martensite morphology by increasing the number of structured martensite grains, although

the total martensite fraction does not increase. A third phase could not be detected in the

microstructure.

A more representative comparison between all the grades investigated, regarding their an-

nealing behavior, is given in Figure 3.14. In order to emphasize the main differences, which

are of great importance for the interpretation of the mechanical properties, two extreme coolingrates of 5 K/s and 80 K/s were applied after annealing in the austenitic region (840) for 30s.

Comparing the microstructures of Figure 3.14 column-wise, it becomes obvious that the

cooling rate has a very strong influence on all materials with respect to grain refinement of

ferrite. This impact is more pronounced for the grades R and M than for grades DP I and DP II.

A cooling rate of 5 K/s is sufficient for a preliminary grain refinement of the dual-phase grades,

as shown in Figures 3.14-c and 3.14-e, partly explained from their already homogeneous

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(a)CR= 5 K/s (b) CR= 10 K/s

(c)CR= 20 K/s (d) CR= 40 K/s

(e)CR= 60 K/s (f) CR= 80 K/s

Figure 3.12: Microstructures of the grade M annealed at 800 for 30 s and quenched with differentcooling rates, etched with LePera.

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(a) R, Tan= 800 (b) R, Tan= 840

(c) DP I, Tan= 800 (d) DP I, Tan= 840

(e) M, Tan= 800 (f)M, Tan= 840

Figure 3.13: Microstructures of the grades R, DP I and M annealed at 800 (left column) and840 (right column) for 30 s after quenching with 20 K/s, etched with LePera.

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ferritic-martensitic starting microstructure. Irrespective of the grade, all microstructures in the

left column (referring to the lowest cooling rate) possess a small fraction of a third phase. For the

grade M this fraction is negligible (Figure 3.14-g). It should be underlined that the industrial

grade R preserves/retains this third phase (bainite and/or pearlite) even after quenching with

80K/s (Figure 3.14-b), a fact that differentiates this grade from the laboratory ones.

Additionally, the morphology of martensite is strongly dependent on the cooling rate. Low

cooling rates enhance the formation of white, clear, unstructured martensite grains. LePeras

etchant stains white also grains of retained austenite. However, magnetic volumetric measure-

ments showed that no retained austenite is present in the microstructures investigated. When

a dramatically different/higher cooling rate is applied (microstructures of the right column of

Figure 3.14), part of the martensite grains appear dark brown, which is generally true for the

coarser grains with a substructure in their interior.

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(a) R, CR= 5 K/s (b) R, CR= 80 K/s

(c) DP I, CR= 5 K/s (d) DP I, CR= 80 K/s

(e)DPII, CR= 5 K/s (f) DP II, CR= 80 K/s

(g)M, CR= 5 K/s (h) M, CR= 80 K/s

Figure 3.14: Microstructures of the grades R, DP I, DP II and M after annealing at 840 for 30s

and quenched with the lower and the higher cooling rates of 5 K/s (left column) and 80 K/s (right

column) respectively, etched with LePera.

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Further investigations were carried out by means of electron microscopy (involving SEM,

FEG-SEM and TEM) in order to identify all the phases present in the fine and ultrafine

microstructures. As the resolution scale increases, the microstructure characteristics become

location-dependent and no representative conclusions can be drawn about the overall fractions

and/or about the mean grain sizes of the phases. However, higher magnifications enable a more

detailed qualitative analysis.

To make the distinction between the phases possible, the creation of a topographic instead of

a color contrast is required, so that the grain boundaries between the phases are distinguishable.

For this reason all specimens for scanning electron microscopy were etched with Nital, which

preferentially etches ferrite, bainite and cementite and outlines their grain boundaries while

leaving martensite intact/undissolved.

Micrographs of the industrial grade R taken with a Field Emission Gun-Scanning Electron

Microscope at a magnification of 4000 are shown in Figure 3.15. The microstructures were

produced by annealing at 800 for 30 s and quenching with 60 K/s and were taken at different

locations of the same specimen. SEM micrographs reveal the presence of a third bainitic phase

in the microstructure of the reference grade R (Figure 3.15-b), supporting the observations

made in the light microscope. This third phase is mainly detected at the ferrite-martensite

grain boundaries and its fraction does not exceed 2 %.

(a) (b)

Figure 3.15: Micrographs of the industrial grade R annealed at 800 for 30 s and quenched with60 K/s. The symbols F, M and B marked on the grains stand for ferrite, martensite and bainite,respectively.

Quenching with cooling rates equal to or higher than 40 K/s (in the given example 60 K/s)

results in the formation of coarse martensite blocks, consisted of strong orientated laths even

within the same block (Figure 3.15-a). Since no grain boundaries can be identified between

the laths, these blocks are considered and evaluated as discrete martensite grains.

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Analogous difficulties, but this time regarding the identification/observation of grain bound-

aries between ferrite grains, are presented in Figure 3.15-b. There exist regions that resemble

martensite, which are present in the rim of ferrite grains without any distinct grain boundaries.

Light microscopy does not even allow the identification of these subgrains, irrespective of the

etchant used. To simplify the quantitative part of the evaluation and, moreover, to be on the

safe side in the grain size measurements, these grains are also considered as individual ferrite

grains.

The phenomenon described above becomes more pronounced and hence more critical for

the evaluation of the ultrafine microstructures produced from the laboratory grades, especially

after annealing at high temperatures and quenching with high cooling rates. Figure 3.16

shows FEG-SEM micrographs of the DP I grade, annealed in the austenitic region (840) and

quenched with 60 K/s. Although the existence of martensite subgrains partly surrounding the

coarser ferrite grains is evident, no grain boundaries between them could be detected. The

transition from ferrite to martensite is very smooth and the martensite subgrains appear

free of substructures, making the phase identification more difficult. The only indicative char-

acteristic of the martensite presence is a difference in contrast observed between the phases,

where martensite appears brighter due its higher concentration in carbon and in other alloying

elements (Cr, Mn, etc.). Ultramicrohardness measurements have confirmed these assumptions.

(a) (b)

Figure 3.16: Micrographs of the laboratory ferritic-martensitic grade DP I annealed at 840 for 30sand quenched with 60 K/s, (magn. 2000 and 8000 for (a) and (b) respectively).

Martensite grains exhibiting an unusual morphology are also shown in Figure 3.16. The

substructure in the interior of these grains, indicated by the white arrows (Figure 3.16-a), is

definitely different from the lath martensite structure described in Figure 3.15. Such complex

martensitic morphologies were already observed in the light microscope in the form of brown

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tinted substructures within the coarser white martensite grains (as stained with LePera). Ac-

cording to bibliographic references of related works [80, 81], the emergence of the martensite

substructure can be interpreted as a result of carbide precipitation during tempering - which

in the case investigated refers to autotempering during quenching, since no tempering stage is

applied after annealing.

(a) (b) (c)

Figure 3.17: TEM micrographs of grade DP I demonstrating: (a) ferrite and martensite grains,(b) martensite lath structure and (c) cementite precipitates within a grain of tempered martensite.

To clarify the substructure formation within martensite grains and to confirm the presence

of carbide precipitates, transmission electron microscopy (TEM) investigations were performed

on selected specimens, covering a representative range of grades and annealing conditions.

TEM micrographs of the ferritic-martensitic grade DP I are shown in Figure 3.17. The

investigated samples were annealed at 840 (- phase field) for 30 s and quenched with 5 K/s

and 80 K/s, represented by the images (a) and (b, c) respectively. For a direct comparison,

microstructures of the grade DP I subjected to the same heat treatment were earlier discussedin Figures 3.14-c and 3.14-d. The microstructure of the samples cooled down with 5 K/s

consists mainly of martensite and ferrite. Sometimes, bainite and/or cementite precipitates

could be found close to martensite grains. No tempered martensite could be detected, which is

in agreement with the light microscopy observations.

In the samples quenched with 80 K/s, martensite grains exhibiting a lath microstructure

could be observed (Figure 3.17-b). Additionally, significant amounts of martensite grains

with a substructure in the interior are found (Figure 3.17-c). The substructure is characterized

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by the presence of cementite precipitates, aligned along more than one habit plane variants.

These grains are identified as tempered (or to be more specific as autotempered) martensite

based on the observations of Bramfitt et al. [82], according to which the existence of multiple

carbide habit-variants is indicative of autotempered martensite [82]. In the following, the term

autotempered martensite will be used to describe such phase constituents.

Figure 3.18 provides a more detailed image of autotempered martensite formation, taken

from a sample of the martensitic grade M intercritically annealed at 800 and quenched with

80 K/s. As can been observed, autotempered martensite (marked on the micrograph) is located

in the middle of the grain and exhibits a different morphology from martensite located near

the martensite-ferrite interface. Higher resolution analysis confirms the presence of cementite

precipitates, revealing their multi-directional arrangement as well.

Figure 3.18: High resolution TEM micrographs of grade M, highlighting the formation of carbide

precipitates in autotempered martensite.

Additional TEM micrographs of the grade M are presented in Figure 3.19, corresponding

to dramatically different cooling rates after intercritical annealing. The microstructure of thesamples cooled down with 5 K/s (Figure 3.19-a) consists of fine martensite grains homoge-

neously dispersed in the ferritic matrix, without any tempered martensite present. Despite

the low cooling rate applied, no bainite could be detected. What should be also noticed, is

the presence of relatively coarse cementite precipitates located near martensite grains or in the

ferrite grain boundaries.

The samples quenched with 80 K/s (Figures 3.19-b and 3.19-c) show an ultrafine ferritic-

martensitic microstructure. The presence of autotempered martensite, existing in considerable

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fractions in this grade, was previously discussed in detail. Furthermore, negligible amounts

of retained austenite were identified, trapped in the form of thin plates between the laths of

martensite grains. However, the amount of retained austenite found is too small to be taken

into account in the quantitative analysis of the microstructure.

(a) (b) (c)

Figure 3.19: TEM micrographs of grade M intercritically annealed (Tan= 800) and cooled with

5 K/s and 80 K/s, for (a) and (b,c) respectively. The micrographs show: (a) coarse cementite precip-itates located around a martensite grain, (b) ultrafine ferrite and martensite grains and (c) thin filmsof retained austenite between the martensite laths.

3.

mechanical properties of dual-phase steels.pdf

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