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.

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 for

creating a pleasant and friendly working environment and for helping me to become an active

member of our group.

Munich, June 2006 Prodromos Tsipouridis

Contents

1 Introduction 1

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

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

2 Material production and experimental procedure 15

2.1 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 . . . . . . . . . . . . . . 44

3.2 Mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50

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

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

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

I

4 Discussion 77

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

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

5 Summary 102

II

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, which

are low-carbon steels characterized by a yield strength level of 140MPa 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

1

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.

0 200 400 600 800 1000 12000

10

20

30

40

50

60

Elon

gatio

n to

frac

ture

A (%

)

Tensile Strength Rm (MPa)

IF

Mild High Strength IF

BHIS

C-Mn

HSLA DPTRIP

Multiphase Steels

AHSS

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

2

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 1.2 shows 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.

MiscHSLA

IF

CPMART

TRIP

BH

DP

DP BH TRIP MART CP IF HSLA Misc

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

3

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.

4

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 affordable

levels. 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 α + γ 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

[15–19].

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 γ → α 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

5

[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, 21–23].

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

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

than 0.15wt. % carbon (for weldability reasons) and 1.5wt.% 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, 24–28]. 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 addition

of 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 [%]:

log CR [K/s] = − 1.73 (Mn)eq [%] + 3.95 (1.1)

6

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 microstructure

and 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 [29–33]. 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 (Cγ=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 γ/α + γ boundary

(C0=Cγ). 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.

7

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, 34–36]. 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− 423 C− 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)

8

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 of

many 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 galvanizing

line. 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.

9

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 [26–28, 39–44] (usually linear regressions between the constituent phases’ proper-

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

micromechanical models [46–51] (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.

10

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 “ultrafine”

is somewhat vague, it reflects the objective of the Japanese super metal project of producing

a strip having a ferrite grain size of ∼ 1 µm throughout a minimum thickness of 1mm [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 (σy):

σy = σ0 + ky d− 1/2. (1.5)

For a low-alloyed steel, for example, a decrease of the ferrite grain size from 5 µm to 1 µm would

ideally produce an increase of the yield strength by approximately 300MPa. 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 [57–61], grain refinement is pro-

11

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 with dF < 2 µm form on the

surface layers of the strip (reaching a depth of ∼1/4- 1/3 of its thickness). In the core of the

strip forms either a mixture of carbides, bainite and coarse ferrite grains (air-cooling) or tem-

pered martensite (dual-phase steel formation caused by accelerated cooling). The Hot Torsion

method is based on the similar SITR concept, with the difference that the deformation mode

involves torsion at elevated temperatures instead of simple rolling [62, 63]. In each case, a very

precise temperature control is required, thus allowing a relatively narrow thermomechanical

production window. The inhomogeneity of the microstructure in thickness remains one of the

main drawbacks of the procedure. Tensile testing has revealed that UFF steels obtained from

this production route exhibit very high lower yield stress to tensile strength ratios, approaching

0.90-0.95. The complete lack of strain hardening inhibits their application in forming processes.

An alternative suggestion was made by Ueji et al. [64, 65], introducing conventional cold-

rolling and subsequent annealing of a martensitic start microstructure. The method was applied

to plain low carbon steels, undergoing a 50% cold-rolling reduction and then was “warm”-

annealed at 823 K (550℃) to produce a multiphase ultrafine microstructure consisting of UFF

grains and uniformly precipitated fine carbides. Blocks of tempered martensite were occasion-

ally observed. The initial martensitic microstructure and the formation of fine carbides during

warm annealing were identified as the keys to success of the process. Annealing at higher tem-

peratures (700℃) proved to be detrimental for the mechanical performance of the material, by

producing properties similar to a ferritic-pearlitic steel.

The production of fine-grained alloys on the micro-meter scale by conventional thermome-

chanical processing is limited by the achievable equivalent strains, which are typically in the

order of 3-4, if the final product has to exhibit a minimum thickness of 1mm. Several novel

processing routes that allow for larger strains and are based on the concept of accumulat-

ing “unlimited” strain in the material are currently under development on a laboratory scale.

Figure 1.4-a describes the procedure followed in Torsion under Hydrostatic Pressure (or High

Pressure Torsion, HPT). A thin disc is deformed in torsion using the friction provided by the

application of a large hydrostatic pressure. By imposing complex strains in the specimen, very

large equivalent strains in the order of 7 can be induced in the material, resulting in grain sizes

of 2µm produced at room temperature [55, 67].

Equal Channel Angular Pressing (ECAP) can be considered as a rapidly developing tech-

12

(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 [68–71]. According to the method, the

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

angle of 2φ (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 die

angle 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 10min. UFF grains with uniformly distributed blocky martensite

islands of 1µm 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 sample’s thickness and geometry, given that the reduction in thickness

13

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 µm. 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.

14

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 210mm 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 2.40mm. 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 1.00mm

(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 50mm 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.

15

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 α + γ 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 ofthe investigated dual-phase steel (the dashed line indicates the carbon content) and (b) equilibriumfraction 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 5min in

16

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,

DP II and M, 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= 5min, WQ 23%DP II Tan= 800℃, tan= 5min, WQ 40%M Tan= 900℃, tan= 5min, 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 m·min−1 to a final thickness of 1.00 or 0.80mm (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

17

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

18

in the rolling direction (10 x 4.0 x 1.0 or 0.80 mm, depending on the sheet’s 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.

00

300

600

900

tem

pera

ture

(°C

)

time (s)

25 K/s CR: 5-80 K/s

Tan

; tan

00

300

600

900

tem

pera

ture

(°C

)

time (s)

(a) (b)

Figure 2.4: (a) Intercritical annealing, (b) Repeated heating cycles oscillating between the α+γ andγ 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 1.00mm 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 α+γ two-phase and γ single-phase regions for different representative holding

times and then were water-quenched.

19

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 LePera’s 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

LePera’s etching agent could not produce the desirable effects, a 2% Nital etchant (2% nitric

acid 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 LEO1450 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 µm followed by electrolytic thinning

in 5 % perchloric acid in acetic acid. The thin foils were analyzed in a Philips CM20STEM

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 RKM250 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 80mm. 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

20

improve roughness.

2.8.2 Ultramicrohardness and microhardness testing

The ultramicrohardness of the main constituents of the dual-phase microstructure (ferrite

and 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 =189F

d 2 , (2.1)

where d [µm] 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, LeicaAG). The indentation load was set to 150 p (approximately

1471mN), 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 [75–79]. 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 12mm (±0.15mm) hole was punched

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

21

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

any crack (which extends all through the sample’s 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:

η =df − d0

d0

× 100, (2.2)

where η [%] 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

22

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 of Figure 3.1 which involves annealing not

only in the intercritical two-phase α + γ 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 25K/s

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

was investigated.

20 40 60 80 100 120 140 160600

650

700

750

800

850

900

WQ WQWQ

Ann

ealin

g te

mpe

ratu

re (°

C)

Annealing time (s)

750 °C 775 °C 800 °C 825 °C 850 °C

WQ: water-quenching

Heating rate: 20 K/s

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

0 20 40 60 80 100 1200.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

750 °C 775 °C 800 °C 825 °C 850 °C

Mar

tens

ite fr

actio

n

Annealing time (s)

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

23

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 3.2 shows 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 3.3 shows 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 α + γ → γ 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 austenitefraction 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 annealing

temperature 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

24

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 in Figure 3.5, the results are in good agreement with each other.

25

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 functionof 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 microstructure

of 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 25K/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 α + γ 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 α + γ → γ transformation temperature does not actually contradict the results from the

26

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.

650 750 850 95090

100

110

120C

hang

e in

leng

th (µ

m)

Temperature (°C)

T

Figure 3.6: Dilatation curve during heating. Tγ indicates the temperature at which the ferrite toaustenite 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 3.7 demonstrates the influence of holding time on the microstructure of

the grade M, intercritically annealed at 800℃ and quenched with 40K/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

27

(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.

28

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-b and 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 40K/s was

applied 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 80K/s,

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

29

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 10K/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 40K/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 0.80mm, 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 α + γ ←→ γ transformation temperature was to provoke the

continuous formation of new grains, so to impede grain growth. By this repeated nucleation and

preventing 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 α + γ/γ - 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.

30

(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.

31

(a) (b)

(c) (d)

Figure 3.10: Ultrafine ferritic-martensitic microstructures produced from the following labora-tory cold-rolled grades: (a)DP II (0.80 mm), (b)DP I (0.80mm), (c)M (1.00mm) and (d)DP II(1.00mm). 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 and

3µm and with martensite grains generally finer than 1 µm 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

32

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 25K/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-c to 3.11-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

33

(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 30 s and quenched withdifferent cooling rates, etched with LePera.

34

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 3.12 and in reference to Figure 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-c with that in Figure 3.12-d, represent-

ing quenching with 20K/s and 40K/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 40K/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 20K/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 cooling

rates of 5K/s and 80 K/s were applied after annealing in the austenitic region (840℃) for 30 s.

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 5K/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

35

(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.

36

(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 20K/s, etched with LePera.

37

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. LePera’s

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.

38

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

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

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

(g)M, CR= 5K/s (h)M, CR= 80K/sFigure 3.14: Microstructures of the grades R, DP I, DP II and M after annealing at 840℃ for 30 sand quenched with the lower and the higher cooling rates of 5 K/s (left column) and 80 K/s (rightcolumn) respectively, etched with LePera.

39

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 with60K/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 40K/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.

40

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 30 sand 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

41

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℃ (γ- phase field) for 30 s and quenched with 5K/s

and 80K/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 discussed

in Figures 3.14-c and 3.14-d. The microstructure of the samples cooled down with 5K/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 80K/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

42

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

80K/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 carbideprecipitates 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 the

samples cooled down with 5K/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 80K/s (Figures 3.19-b and 3.19-c) show an ultrafine ferritic-

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

43

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 with5K/s and 80K/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.

3.1.4 Quantitative analysis-Grain size measurements

The determination of the quantitative characteristics of the industrially as well as of the labo-

ratory produced materials was necessary in order to study the influence of the pre-processing

route and of the final annealing parameters on the microstructure evolution, regarding volume

fractions and mean grain sizes of the existent phases. To carry out the quantitative evaluation

of such fine grained microstructures, ferrite was considered as the dominant phase while the sum

of all other phases (that is martensite, autotempered martensite, retained austenite, bainite

and cementite) was characterized as a single second phase. The line intercept measurements

were made on light microscopy images (etched with LePera) using a magnification of 1500× or

2000× on purpose. Since grain refinement of ferrite was a main goal of this work, only ferrite

was measured in detail assuming that the rest represents the fraction of the second phase. This

44

0 30 60 9020

25

30

35

40

45

50

55

60

M DP II DP I R

Mar

tens

ite fr

actio

n (%

)

Cooling rate (K/s)

(a)Tan= 800℃

0 30 60 9020

25

30

35

40

45

50

55

60

M DP II DP I R

Mar

tens

ite fr

actio

n (%

)

Cooling rate (K/s)

(b)Tan= 840℃

Figure 3.20: Martensite fraction as a function of cooling rate, for two annealing temperatures.

45

rough classification into ferrite and non-ferrite constituents diminishes the human error sources

stemming from the complexity of phase distinction due to the tint etching effects.

However, for the evaluation of mechanical properties and their correlation with the mi-

crostructural characteristics it is essential to estimate the martensite fraction of the steels.

Taking into account that the fractions of bainite and cementite in the materials investigated

are extremely low (esp. for the laboratory grades) and that retained austenite hardly exists,

and, furthermore, accepting an unavoidable methodical error due to the increased surface area

of the grain boundaries - which becomes significant in ultrafine-grained microstructures, the

second phase fraction provides an acceptable approximation of the martensite fraction (with-

out distinguishing between martensite and autotempered martensite). This roughly means that

the martensite fraction shown in the following diagrams is slightly overestimated, particularly

for cooling rates lower than 10K/s, and for this reason a cautious interpretation of the results

is demanded.

Given that all grades are produced from the same hot-rolled material and subsequently have

identical chemical compositions, the observed differences in martensite fraction result exclu-

sively from the applied heat treatment schedules and their impact on the particular character-

istics of each pre-processed grade. The influence of the cooling rate on the volume fraction of

martensite is shown in Figure 3.20. Irrespective of the annealing temperature, an increase in

cooling rate is generally accompanied by an increase in the martensite fraction for all grades.

The maximum martensite fraction achieved by intercritical annealing was ∼ 48% after cooling

with 80K/s (Figure 3.20-a, M curve). The martensitic grade M seems to react more sensitive

to the changes of cooling rate. Unlikely, the martensite fraction of the grades R, DP I, and

DP II remains relatively stable or slightly increases, ranging between 35 and 40% for cooling

rates higher than 10 K/s.

Austenitic annealing produces dual-phase microstructures with higher martensite fractions

than intercritical annealing (Figure 3.20-b), as predicted from the thermodynamic equilib-

rium. At higher cooling rates two “subgroups” of curves are observed, showing a different

behavior regarding their martensite fraction. Grades M and R possess almost 10% more

martensite than the grades DP I and DP II at the same cooling rates, reaching fractions in

the order of ∼ 55%.

Figure 3.21 shows the dependence of the ferrite grain size on the applied cooling rate.

Increasing the cooling rate results in a decrease of the mean grain size of ferrite (reaching

1.5µm) and leads to the formation of ultrafine-grained microstructures. The measurements

reflect the microstructural observations and, moreover, reveal the refinement degree in

cases that the changes in grain size are beyond the discriminability of the human eye.

Analogous to Figure 3.20-a, grade M responds positively to the changes of cooling rate,

throughout the whole range. The other three grades maintain a mean ferrite grain size of

46

0 30 60 901.0

1.5

2.0

2.5

3.0

3.5

M DP II DP I R

Ferr

ite m

ean

grai

n si

ze (µ

m)

Cooling rate (K/s)

(a)Tan= 800℃

0 30 60 901.0

1.5

2.0

2.5

3.0

3.5

Ferr

ite m

ean

grai

n si

ze (µ

m)

Cooling rate (K/s)

M DP II DP I R

(b)Tan= 840℃

Figure 3.21: Influence of the cooling rate on the mean grain size of ferrite.

47

20 25 30 35 40 45 501.0

1.5

2.0

2.5

3.0

3.5

M DP II DP I R

Martensite fraction (%)

Ferr

ite m

ean

grai

n si

ze (

m)

(a)Tan= 800℃

20 25 30 35 40 45 50 55 601.0

1.5

2.0

2.5

3.0

3.5

M DP II DP I R

Martensite fraction (%)

Ferr

ite m

ean

grai

n si

ze (µ

m)

(b)Tan= 840℃

Figure 3.22: Relationship between ferrite grain size and martensite fraction.

48

approximately 2µm. Quenching from the austenite field has a more pronounced impact on the

grain size of ferrite; all grades exhibit grain sizes below 2 µm even for cooling rates lower than

40K/s.

The interdependency/interaction between martensite fraction and ferrite grain size is demon-

strated in Figure 3.22. A reciprocal relationship (approaching a linear form) between them is

noticed irrespective of the grade and the annealing temperature, with the exception of grades

DP I and DP II after intercritical annealing (Figure 3.22-a). The finest microstructures are

obtained in steels with increased martensite fractions.

The higher the cooling rate the shorter is the time available for the austenite-to-ferrite trans-

formation, suppressing in that way the further grain growth of ferrite. It is also presumed that

the increase of martensite fraction in the steel is mainly due to the number of martensite (prior

austenite) grains and not due to their size/volume.

49

3.2 Mechanical properties

3.2.1 Tensile testing

The mechanical (tensile) properties of the investigated dual-phase steel are presented in

Table 3.1, valid for a thickness of 1.0 mm. These values, which stem from the industrially

cold-rolled material laboratory annealed with a Hot-Dip Galvanizing cycle (Tan= 815℃, THDG=

475℃) and skin-pass rolled for roughness improvement, will be used as a reference basis.

Table 3.1: Mechanical properties of the HDG-treated industrially cold-rolled DP-steel (DP600). Thequantities Rp0.2 and Rm stand for the yield and the ultimate tensile strength, while Ag and A denotethe uniform and total elongation, respectively.

Mechanical properties Rp0.2 (MPa) Rm (MPa) Ag (%) A (%)

HDG - DP600 380 660 15 22

To study the mechanical properties of the materials subjected to modified processing (grades

DP I, DP II and M), tensile tests were conducted with specimens of the same thickness as

the industrially cold-rolled material (grade R) so that the results are comparable. Taking into

account the observations/results of annealing simulations (see 3.1.3), the laboratory production

of tensile samples included two annealing temperatures (800 and 840℃), one annealing time

(30 s) and six different cooling rates, that is 5, 10, 20, 40, 60 and 80 K/s. The production of

tensile samples by applying non-conventional heat treatment schedules was also attempted, but

due to the insufficient temperature control and the unavoidable deviations from the annealing

schedule, no ultrafine ferritic-martensitic microstructures could be (re-)produced. For this

reason, such samples were not used in the investigations. The reference material (grade R)

was also tensile tested, after being submitted to the same heat treatment with the laboratory

grades. It should be noticed that all samples were tensile tested in the as-annealed condition.

The range of the applied annealing conditions provides with a wide variety of microstructures,

regarding the characteristics directly controlled by the annealing temperature and the cooling

rate such as the fraction of martensite and the presence of additional phases (e.g. bainite).

Moreover, and according to the results of grain size measurements, the study of the cooling

rate impact indirectly incorporates into the analysis the critical parameter grain size.

The existence of a third phase (bainite) makes the interpretation of the results even more

difficult, because in addition to the martensite fraction and the grain size a third parameter is

introduced, which is different for each grade. The influence of each of these parameters on the

mechanical properties can not be isolated and studied separately. Hence, all results are given in

dependence on the applied cooling rate, reflecting in this way the combined effects of all three

parameters (martensite fraction, grain size, presence of a third phase).

50

0 30 60 90260

300

340

380

420

460

M DP II DP I R

R p0.2 (M

Pa)

Cooling rate (K/s)

(a)Tan= 800℃

0 30 60 90240

280

320

360

400

440

480

520

R p0.2 (M

Pa)

Cooling rate (K/s)

M DP II DP I R

(b)Tan= 840℃

Figure 3.23: Influence of the cooling rate on the yield strength (Rp0.2) of the dual-phase steels.

51

Figure 3.23 shows the influence of the cooling rate on the yield strength of the investigated

grades for both intercritical and austenitic annealing temperatures. An increase of the cooling

rate and subsequently of the martensite fraction results in an increase of the yield strength.

A simple qualitative comparison shows that the yield strength curves follow the form of the

martensite fraction curves in Figure 3.20. In the present case, a second contribution to the

yield strength is added by the grain refinement that occurs at high cooling rates. Since the grain

size of ferrite and the martensite fraction are not independent from each other, the quantitative

contribution of the two effects to the yield strength seems to be rather complicated.

In Figure 3.23 two “groups” of curves can be observed. The grades DP I and DP II exhibit

a remarkably lower yield strength than the grade M. This difference becomes more pronounced

as the cooling rate increases and reaches approximately 80MPa. The behavior of the reference

grade (R) is strongly influenced by the annealing temperature. Intercritical annealing keeps

the yield strength of R in the lower strength group (Figure 3.23-a), while austenitic annealing

increases its yield strength to the higher levels of grade M (Figure 3.23-b).

A similar behavior is observed for the ultimate tensile strength of the grades in Figure 3.24.

The tensile strength increases as the cooling rate and the martensite fraction increase. The

two types of curves according to the above classification are still observed, maintaining also

the same form. In comparison with the tensile strength of the HDG-treated dual-phase steel

(Table 3.1) a significant increase reaching up to 200MPa is achieved. It is noteworthy that

grades DP I and DP II maintain their yield strength at the same level as the reference hot-dip

galvanized material (that is between 360 and 380MPa) for a wide range of cooling rates while

achieving higher tensile strengths by an amount of 140-150MPa.

In Figure 3.25 the uniform and total elongations of all grades are plotted as a function of

the cooling rate. If the increase in strength would exclusively originate from the increase in

martensite fraction, then a decrease in total elongation should be expected. In the case of

intercritical annealing (Figure 3.25-a) this effect is not pronounced for any of the investigated

grades, since no dramatic decrease either of the uniform or of the total elongation occurs as

the cooling rate increases. Grades DP I, DP II and R show rather constant elongation even at

high cooling rates, slightly ranging between 13-14% and 18-19 % for Ag and A, respectively.

Though grade M keeps also its elongation at a constant level, the exhibited values of Ag and A

are lower by approximately 2% compared to DP I, DP II and R.

On the contrary, austenitic annealing seems to have a rather beneficial effect on grade

M and a detrimental effect on grade R. For cooling rates above 20K/s, the decrease in

elongation as the martensite fraction increases is more pronounced for the grade R (in

comparison with grade M), reaching a gap of ∼ 5% in total elongation at a cooling rate of

80K/s. Hence, quenching with 80K/s from an annealing temperature of 840℃ results in a

52

0 30 60 90620

660

700

740

780

820

860

R m (M

Pa)

Cooling rate (K/s)

M DP II DP I R

(a)Tan= 800℃

0 30 60 90620

660

700

740

780

820

860

900

R m (M

Pa)

Cooling rate (K/s)

M DP II DP I R

(b)Tan= 840℃

Figure 3.24: Influence of the cooling rate on the tensile strength (Rm) of the dual-phase steels.

53

0 30 60 9010

12

14

16

18

20

22

24 M DP II DP I R

A g, A (%

)

Cooling rate (K/s)

(a)Tan= 800℃

0 30 60 905

7

9

11

13

15

17

19

21

M DP II DP I R

A g, A (%

)

Cooling rate (K/s)

(b)Tan= 840℃

Figure 3.25: Influence of the cooling rate on the uniform and total elongations (denoted as Ag andA respectively) of the dual-phase steels. The hollow symbols in the diagrams stand for the uniformelongation.

54

dual-phase steel with a tensile strength of 860MPa and a total elongation of 10%, in case the

reference grade (R) is used. If, however, the same heat treatment is applied to the martensitic

grade M, then a dual-phase steel with the same tensile strength level (850MPa) but with a 5 %

higher total elongation is produced. In the same example, it should not be overlooked that the

total elongation exhibited by the grade R is lower than the uniform elongation of the grade M.

The ferritic-martensitic grades DP I and DP II show even higher uniform and total elongations

up to 14 % and 19%, respectively, for the entire range of the applied cooling rates (Figure 3.25-

b). Combining the results of Figures 3.20-b to 3.25-b referring to austenitic annealing, it is

unexpectedly observed that the increase in martensite fraction (from 30 to 45% on average)

accompanied by refinement of the ferrite grains does not affect the formability behavior of the

steels. Consequently, the grades DP I and DP II can provide a wide assortment of dual-phase

steels exhibiting a yield strength between 270 and 370MPa and an ultimate tensile strength

between 660 and 810MPa, while maintaining the same high ductility.

Two quantities that are of great importance for the characterization of dual-phase steels are

the yield to tensile strength (Rp0.2/Rm) and the uniform to total elongation (Ag/A) ratios,

providing a critical relationship between the strength and the ductility properties respectively.

The optimum performance of the material regarding both strength and formability is achieved

at low yield to tensile strength and high uniform to total elongation ratios. Figure 3.26 shows

the impact of the annealing conditions on the yield to tensile strength and on the uniform

to total elongation ratios of the investigated dual-phase grades. Intercritically annealed grades

exhibit a yield to tensile strength ratio ranging between 0.43 and 0.50, increasing with increasing

cooling rate. All the grades show the same behavior, with the only exception being grade M

when quenched with 80K/s. The uniform to total elongation ratio lies above 0.68 and does

not exceed 0.78 for all grades. This ratio is slightly higher by 0.05 for the pre-processed grades

than that of reference grade R (Figure 3.26-a).

On the other hand, austenitic annealing results in the formation of two groups of Rp0.2/Rm

curves. The “lower” group formed by the dual-phase grades DP I and DP II exhibits yield

to tensile strength ratios between 0.40 and 0.50 while the “upper group” (M and R grades)

exhibits ratios between 0.50 and 0.58. In both cases, the higher the cooling rate the higher is

the Rp0.2/Rm value. However, such a tendency cannot describe the behavior of the Ag/A ratio.

The uniform to total elongation ratio ranges between 0.68 and 0.80 and does not show any

systematic dependence on the cooling rate. The grade R is only slightly affected by the cooling

rate and, furthermore, reaches higher values than the laboratory grades. The unexpected drop

in Ag/A of the grade M that occurs between cooling rates of 40 and 60K/s cannot be explained

satisfactorily on the basis of the microstructural analysis.

55

0 30 60 900.3

0.4

0.5

0.6

0.7

0.8

0.9

0.3

0.4

0.5

0.6

0.7

0.8

0.9

Ag /

A

M DP II DP I RR p0

.2 / R m

Cooling rate (K/s)

(a)Tan= 800℃

0 30 60 900.3

0.4

0.5

0.6

0.7

0.8

0.9

0.3

0.4

0.5

0.6

0.7

0.8

0.9

M DP II DP I R

R p0.2 / R m

Cooling rate (K/s)

Ag /

A

(b)Tan= 840℃

Figure 3.26: Dependence of yield to tensile strength (Rp0.2/Rm) and of uniform to total elongation(Ag/A) ratios on the cooling rate. The black dashed arrows point out the corresponding y-axis.

56

Strain hardening

In Figure 3.27 some true stress-logarithmic strain (widely referred as true strain) curves of the

investigated materials are shown.

0.00 0.04 0.08 0.12 0.16 0.200

100

200

300

400

500

600

700

800

900

M DP II DP I M

True

stre

ss (M

Pa)

Logarithmic strain0.00 0.04 0.08 0.12 0.16 0.200

100

200

300

400

500

600

700

800

900

M DP II DP I R

True

stre

ss (M

Pa)

Logarithmic strain

(a)Tan= 800℃, CR= 5 K/s (b) Tan= 840℃, CR= 5K/s

0.00 0.04 0.08 0.12 0.16 0.200

100

200

300

400

500

600

700

800

900

1000

M DP II DP I R

True

stre

ss (M

Pa)

Logarithmic strain0.00 0.04 0.08 0.12 0.16 0.200

100

200

300

400

500

600

700

800

900

1000

M DP II DP I R

True

stre

ss (M

Pa)

Logarithmic strain

(c)Tan= 800℃, CR= 80 K/s (d) Tan= 840℃, CR= 80 K/s

Figure 3.27: Strain hardening curves of the dual-phase steels from selected annealing conditions.Diagrams (a) and (b) represent the lowest cooling rate of 5 K/s while (c) and (d) the highest coolingrate of 80K/s.

Again, the curves are, as expected from the results given earlier in this chapter, separated into

“groups”, following the classification observed in Figures 3.23 and 3.24. All grades exhibit

a continuous yielding behavior and no yield point elongation is observed even for the lowest

cooling rate applied. The sudden change of the curve slope which occurs at logarithmic strains

close to 0.015 is due to a slight increase of the strain rate during the tensile test, though not

affecting the mechanical properties of the material (according to EN).

57

For most of the steels, the region of the true stress (σ) versus logarithmic strain (ε) curve

between the onset of plastic deformation (corresponding to the yield strength in the engineering

stress vs. engineering strain curve) and the onset of necking (corresponding to the ultimate

tensile strength in the engineering stress vs. engineering strain curve) can be approximated by

the Hollomon equation [83]:

σ = K εn. (3.1)

The constant K is a strength coefficient depending on the thermomechanical history of the

steel, while the parameter n is defined as the strain hardening exponent, also known as the

n-value. By taking logarithms of both sides of the eq. 3.1, the Hollomon equation can be

linearized:

log σ = log K + n log ε. (3.2)

If the experimental data satisfy the Hollomon equation and, additionally, the curve is plotted

on logarithmic coordinates, then a linear regression line with slope n can be determined.

The strain hardening behavior can be evaluated by determining the strain hardening expo-

nent of the steel (n-value). Physically, n gives a measure of the ability of the steel to distribute

strain along the gage length of the tensile specimen. For low-carbon steels used to form com-

plex shapes, the value of n is ∼ 0.22. The higher the n-value the more uniform is the strain

distribution and, therefore, the greater is the resistance of the steel to necking and the better

is its formability [84].

The strain hardening behavior of dual-phase steels and the determination of their n-value have

been the subject of numerous investigations. Since the experimental data have shown that n-

value is not constant over the entire range of uniform strain, the empirical strain hardening laws,

such as the Hollomon equation, are unable to describe in detail the work hardening behavior of

the steel. For this reason new methods of analysis have been developed, in which the uniform

elongation stage is divided into several segments and individual n-values are calculated within

specified limits of logarithmic strain. Within the scope of this work and using the experimental

data as digitally recorded by the tensile testing machine, the deformation region/curve was

divided in two or three intervals (depending on the achieved uniform elongation Ag), yielding

segmental n-values at logarithmic strains between 2.0 and 4.0, 4.0 and 6.0 and, where possible,

between 6.0 and 10.0, denoted as n2−4, n4−6 and n6−10, respectively. Although in most cases the

uniform elongation of the materials exceeds 12%, sometimes reaching even 15%, the n-values

at higher logarithmic intervals are not available since the next strain segment automatically

provided by the tensile testing machine is n10−18.

58

0 10 20 30 40 50 60 70 80 900.12

0.14

0.16

0.18

0.20

0.22

0.24

0.26

0.28

0.30where: n

2-4

n4-6

n6-10

M DP II DP I R

n - v

alue

Cooling rate (K/s)

(a)Tan= 800℃

0 10 20 30 40 50 60 70 80 900.12

0.14

0.16

0.18

0.20

0.22

0.24

0.26

0.28

0.30 M DP II DP I R

where: n2-4

n4-6

n6-10

n - v

alue

Cooling rate (K/s)

(b)Tan= 840℃

Figure 3.28: Dependence of the segmental n-values of the investigated grades on the cooling rate.Each deformation segment is represented by an individual curve.

59

Figure 3.28 shows the dependence of the n-values of all grades on the cooling rate for both

annealing temperatures applied. As can be observed, the n-value decreases with increasing

strain, that is n2−4 > n4−6 > n6−10, irrespective of the grade and of the annealing conditions.

Furthermore, intercritical annealing produces grades with higher n-values than austenitic an-

nealing for the same cooling rates. Though not clearly distinguished in the strain hardening

curves of Figure 3.27, the grades DP I and DP II exhibit remarkably higher n-values than the

grades M and R throughout the whole range of cooling rates, reaching values in the order of

0.28 at low cooling rates. This is obviously the reason for their increased uniform elongation

as also of the higher total elongation at fracture (see Figure 3.25).

Cooling rate proves to have a significant influence on the strain hardening exponent. In-

creasing the cooling rate and thus the martensite fraction of the steel results in a decrease of

the corresponding n-values. The form of the n-value vs. cooling rate curve remains the same

irrespective not only of the logarithmic strain segment but also of the grade. The impact is

more pronounced at cooling rates below 40K/s. For higher cooling rates the n-value is roughly

independent of the cooling rate. The data for n6−10 of the grades M and R for Tan= 840℃(Figure 3.28-b) were not available. For this reason the strain hardening behavior of these

materials is graphically represented by using only two n curves (n2−4 and n4−6). A more de-

tailed analysis of the strain hardening behavior of the investigated dual-phase steels is given in

chapter 4.

Microstructure at high strains

The deformed microstructures of the tensile specimens, more pronounced in the neighborhood of

fracture (region where necking takes place), were thoroughly examined by means of light (LM)

and scanning electron microscopy (both SEM and FEG-SEM). Figure 3.29 shows the fracture

profile of the ferritic-martensitic grade DP II, annealed at 840℃ and quenched with 10K/s.

In the region of fracture the number of the voids formed is high, decreasing as the distance

from the fracture location increases. Furthermore, a higher density of voids is present in the

middle of the specimens’ thickness. Generally, the voids are formed at the ferrite-martensite

interface when a certain strain is exceeded. In case that bands of martensite are present in the

specimen, formed commonly in the middle of the sheet where the former pearlitic segregation

bands were located, it is observed that the voids are preferably formed around and along these

martensite bands, either by decohesion between martensite grains in contact or by fracture of

the martensite islands. This phenomenon is observed more often in the tensile specimens of

the industrial grade R, which has not undergone any pre-processing.

The ferrite grains located close to fracture are elongated, with the long axis parallel to the

direction of deformation. The hard martensitic islands do not seem to take part in the de-

formation process. Moving away from the fracture surface, not only the ferrite grains appear

60

less deformed, but also the number of voids decreases (Figure 3.29-b). Nevertheless, in some

cases voids were found even at distances in the order of [mm] away from the fracture location,

suggesting that the deformation mechanism cannot be not strictly determined by the/an ac-

cumulated local strain at the fracture point. Additionally, no significant observation regarding

the size distribution of voids over the region/area of their appearance was made.

(a)DP II, location of fracture (b) DP II, few mm away from fracture

Figure 3.29: Light microscope micrographs of the dual-phase grade DP II, annealed at 840℃ andquenched with 10 K/s. The specimen is etched with LePera (magn. 500× and 1500× for (a) and (b)respectively).

The deformation mechanism as indicated by the fracture behavior is, however, not indepen-

dent of the individual microstructural characteristics of the investigated dual-phase grades.

Figure 3.30 shows the region of fracture of the grade M, annealed at 840℃ and quenched

with two different cooling rates, 5 and 80 K/s. As was already discussed in a former section

of this work, the microstructures produced by applying these extreme cooling rates differ not

only in their martensite fraction but also in their mean ferrite grain size. Figure 3.30-a shows

the fracture profile of the specimen quenched with 5 K/s, possessing a relatively low martensite

fraction. Similar to the fracture area shown in Figure 3.29, a high number of voids is formed,

distributed all over the fracture region and penetrating into the specimen along the tensile axis,

indicating that the fracture process is not simply one of void nucleation and growth. The ferrite

grains carry the plastic deformation alone, while the hard martensite grains remain undeformed.

On the contrary, the fracture profile of the specimen with a high martensite fraction, see

Figure 3.30-b, reveals a different deformation mode. The number of the formed voids is

significantly lower. These few voids are located very close to the fracture surface and, practically,

no voids can be found at distances more than 200 - 300 µm away from the fracture. The grains

of the ferrite within the highly strained necked region are severely elongated. Unlikely to the

61

fracture of Figure 3.30-a, the harder martensite grains appear also severely deformed. This

observation suggests that not only the soft ferritic matrix but also the harder second martensitic

phase takes active part in the deformation process.

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

Figure 3.30: Fracture profile of grade M annealed at 840℃, etched with LePera.

FEG-SEM micrographs corresponding to the fracture samples of grade M (Figure 3.30-a)

are shown in Figure 3.31. The higher magnification (3000×) and the absence of tint-etching

effects enable the observation of deformation “bands” within the ferrite grains and in some cases

the identification of the strain transfer paths along the ferritic matrix. The hard martensitic

islands remain undeformed. The sites where crack initiation and void growth occur are usually

found either at the ferrite-martensite interface or between martensite grains.

Additionally, fracture profiles of a high martensite fraction sample (analogous to the deformed

microstructure of Figure 3.30-b) are shown in Figure 3.32. The grains of martensite, most of

which emphasize a substructure due to autotempering or to their lower carbon content, appear

heavily deformed and elongated in the tensile direction. No cracks or deformation bands are

observed in the ferrite grains. A small number of voids is formed at the ferrite-martensite

interface, located very close to the fracture surface.

A quantitative analysis or even a rough estimation with respect to the formation of voids, their

size and area distribution and their preferential nucleation mechanism (either by decohesion of

the ferrite-martensite interface or by fracture of martensite islands) was not possible. Since the

severely deformed regions and the small size of voids do not allow a reliable quantification of

the microscopic observations, any analysis could lead to misinterpretations.

62

(a)M, CR= 5 K/s (b)M, CR= 5K/s, fracture

Figure 3.31: FEG-SEM micrographs of the fracture profile of grade M annealed at 840℃, etchedwith Nital.

(a)M, CR= 80 K/s (b)M, CR= 80 K/s, fracture

Figure 3.32: FEG-SEM micrographs of the fracture profile of grade M annealed at 840℃, etchedwith Nital. The tensile axis is indicated.

The correlation of the mechanical properties of the specimens of grade M (examined above)

with their fracture behavior appears to be of great interest. Although the higher marten-

site fraction of the steel quenched with 80 K/s results in a remarkably higher ultimate tensile

strength by ∼ 170-180MPa, the uniform and the total elongations of both steels lie on the same

level (10-11 % and 14-15 % for Ag and A, respectively). Moreover, according to Figure 3.28,

the steel with the lower martensite fraction exhibits consistently higher n-values by approx.

0.05 at all deformation stages (n2−4, n4−6). Hence, neither the strain hardening behavior as

expressed by the n-values nor the difference in martensite fraction (which exceeds 25 % for the

63

applied annealing and cooling conditions) can provide a reasonable explanation for the similar

elongation properties of the two steels.

Analogous observations regarding the relationship between the mechanical properties and

the deformation mechanisms were made for the other pre-processed grades (DP I and DP II),

irrespective of the annealing temperature. Again, the unexpected ductility of the dual-phase

steels with high martensite fractions was accompanied by deformation of the martensite grains

in the location of fracture (necking area). This phenomenon motivates further investigations.

Especially the fact, that the second (martensitic) phase undergoes severe deformation during

tensile testing raises critical questions about the hardness of martensite grains.

64

3.2.2 Hardness measurements

Ultramicrohardness (UMH) measurements

In order to find out whether or not there is a relation between the hardness of the phase

constituents and the deformation behavior of the investigated dual-phase steels, the Vickers

ultramicrohardness of ferrite and of martensite was determined with an ultramicrohardness

testing machine (UMHT-3), mounted in a scanning electron microscope. After many trial tests,

the optimum indentation load was found to be 15mN, which was applied to all measurements

irrespective of the phase and the sample. The higher the indentation load the larger the

indentation trace and consequently the more reliable the measurement. In this work, however,

the ultrafine ferrite and martensite grains of the pre-processed steels were a limiting factor in

the load selection, since the indentation trace/impression should definitely be smaller than the

size of the grain. The positioning precision of the testing equipment was approximately 1 µm.

Selected examples of the Vickers indentations are presented in Figures 3.33-3.35. Although

not clearly demonstrated due to the various sizes of the figures, the magnification of all SEM

micrographs is identical (12000×).

(a)DP II, CR= 20 K/s (b) DP II, CR= 10K/s

Figure 3.33: Indentations in coarse ferrite grains: (a) precise indentation, relatively unaffected fromthe neighboring grains, (b) the presence of grain boundaries influences the measurement.

Figure 3.33 shows indentations in ferrite grains. By comparing the two micrographs it can

be noticed that the two diagonals of the pyramid trace are not always equal; a difference in

length up to 5 % between them was often unavoidable. In cases where the presence of grain

boundaries/interphases and/or the existence of a different than the indented phase beneath

the targeted grain strongly influenced the measurement, this difference in length could reach

∼ 10% (Figure 3.33-b). Measurements with higher length deviations were taken as invalid

65

and were discarded.

To enable the martensite indentation, coarse grains with a minimum size of 2.5 µm were

selected. The micrographs of Figure 3.34 show indentation traces in “clear” (unstructured)

martensite grains, taken from the samples quenched with 5 K/s. Even though both measure-

ments are from a first point of view acceptable, the microhardness values calculated from the

mean length of the diagonals differ roughly by a factor of three (904 HV and 297HV for (a)

and (b), respectively). A more careful examination of the Figure 3.34-b shows the existence

of a microcrack (marked with an ellipse) on the left side of the trace. Although the size and

the form of the crack cannot sufficiently explain the calculated hardness, which approaches the

ferrite hardness, this may be an indication of the existence of a softer ferrite “substrate” right

beneath the indented martensite grain.

(a)DP II, CR= 5 K/s (b) DP II, CR= 5 K/s

Figure 3.34: Indentations in unstructured (“clear”) martensite grains.

Increasing the cooling rate from the annealing temperature results in the formation of coarse

martensite blocks showing a substructure in their interior. These grains were identified as

autotempered martensite. Provided that in the samples quenched with cooling rates higher than

40K/s the fraction of tempered martensite becomes significant, the majority of the martensite

indentations in these specimens were conducted in tempered martensite. Impressions of the

Vickers indenter in tempered martensite grains are shown in Figure 3.35. Tempered martensite

appears to be apparently softer than untempered martensite and harder than ferrite.

Figure 3.36 shows the influence of the cooling rate on the microhardness of ferrite and

martensite of the grade DP II, after quenching from the austenitic annealing temperature

(840℃). The cooling rate was preferred over the martensite fraction as the independent vari-

able, because it additionally includes the grain size effects and the martensite morphology. To

determine the microhardness of martensite, a minimum number of twenty measurements was

66

(a)DP II, CR= 60 K/s (b) DP II, CR= 60K/s

Figure 3.35: Indentations in tempered martensite.

necessary. Particularly in the samples where hard martensite islands were present (see Fig-

ure 3.34-a), a distribution of the hardness values within a range of ∼ 100HV was observed.

On the other hand, the microhardness of ferrite remained relatively stable and was determined

as the mean value of ten measurements. In order to minimize the measurement errors, extreme

hardness values that do not make sense (either by being far away from the mean value or by

not corresponding to the indented microstructure) were excluded from the evaluation.

The ultramicrohardness of ferrite remains nearly independent of the cooling rate and, con-

sequently, of the martensite fraction of the specimen. In absolute values it ranges between

230HV and 250HV, apparently higher than the ferrite hardness reported in literature. In or-

der to understand this increase, it should be taken into account that the fine grains of ferrite

do not enable indentations unaffected by the surrounding (incl. also the underlying) F/F and

F/M grain boundaries. Additionally, portion of this increase could be interpreted as the result

of the work-hardened ferritic matrix due to the martensitic transformation, see chapter 4.

On the other hand, the ultramicrohardness of martensite grains appears strongly dependent

on the quenching conditions. Increasing the cooling rate results in a well-defined decrease of the

measured hardness, even for the smallest rate steps (e.g. from 5 to 10K/s). As indicated by the

blue dashed line (fitting curve of the martensite hardness data) in Figure 3.36, the decrease

in hardness is stronger in the initial part of the curve, representing cooling rates lower than

40K/s. This drop reflects the pronounced decrease in the martensite carbon content which

accompanies the significant increase in its fraction occurring in this cooling rate interval. At

higher cooling rates not only the martensite fraction increases further but also its morphology

changes, since autotempered martensite is formed (see Figure 3.35). The martensite hardness

keeps on decreasing but with a reduced slope. Quenching with 80 K/s (leading to ∼ 50%

martensite) results in the minimum hardness of 390HV.

67

0 30 60 90200

300

400

500

600

700

800

0.2

0.3

0.4

0.5

0.6

0.7

0.8

k = HV fe

rr / HV m

art

HVMartensite

HV

FerriteV

icke

rs h

ardn

ess (

HV

)

Cooling rate (K/s)

k

Figure 3.36: Vickers ultramicrohardness of ferrite and martensite grains of grade DP II, annealed at840℃. The triangle-curve (green) represents the ratio k of ferrite to martensite hardness at a givencooling rate while the dashed lines represent the fitting curves of the experimental data. Indentationload=15mN.

By extrapolating the fitting curve it turns out that a further increase in cooling rate does

only slightly affect the martensite hardness. If this behavior were solely controlled by the

martensite carbon content, then this assumption would contradict the expected results (that is

a further decrease of the martensite hardness). For this reason, fully martensitic samples were

produced from the same grade by water-quenching from the austenitic annealing temperature.

The hardness of martensite ranged between 390 and 400HV, confirming the prediction of the

extrapolation.

Due to the limitations of the method set by the ultrafine-grained microstructures, the obtained

ultramicrohardness values cannot be used as absolute hardness of the individual phases. For

this reason a new variable k was defined, expressing the ratio of ferrite to martensite hardness:

k = HVF/HVM. (3.3)

The trend of hardness ratio with increasing cooling rate is also plotted in the diagram of

Figure 3.36. The curve of k is actually generated by using the data of the fitting curves of

68

ferrite and martensite hardnesses. The decrease in the hardness of martensite with increasing

cooling rate has a dramatical effect on the hardness ratio, increasing its value from 0.33 at

5K/s to 0.60 at 80K/s. The relative hardness of the phases of the steel given by the hardness

ratio can provide a useful tool for the description of their deformation behavior.

Summarizing, the ultramicrohardness of martensite is between 730HV and 390HV in the

range of the cooling rates applied. The measured values correspond with the microstructural

observations of the fracture behavior, supporting the hypothesis that the deformation of marten-

site in the samples quenched with high cooling rates can be attributed to its substantially lower

hardness.

However, a physically-based interpretation of the drop in martensite hardness in correlation

with the deformation mechanism should take into account the martensite fraction that each

cooling rate represents, the carbon content of this specific martensite and the presence or

absence of tempered martensite as a phase constituent.

Microhardness (MH) measurements

Additionally to the ultramicrohardness measurements, the Vickers microhardness of the inves-

tigated dual-phase steels was determined. The indentation load was 150 p, which converted

to SI units is equal to 1471mN (so actually HV0.15 was measured). Compared to the results

of the ultramicrohardness investigations, the approximately one hundred times higher inden-

tation load produces impressions which are bigger by one order of magnitude. Micrographs of

the impressions in different heat treated samples of the DP II grade are shown in Figure 3.37.

The length of the impression diagonals varied from 30 to 35 µm depending on the heat treat-

ment and, hence, on the microstructural characteristics of the sample (martensite fraction,

grain size). The indented area/volume includes a few tens of plastically deformed ferrite and

martensite grains and thereby a significant amount of grain boundaries.

69

(a)DP II, CR= 10 K/s (b) DP II, CR= 80K/s

Figure 3.37: Vickers impressions in grade DP II, annealed at 840℃ and quenched with differentcooling rates (1000×, Nital). The average diagonal lengths listed correspond to hardnesses of 240 HVand 275 HV for the micrographs (a) and (b), respectively. Indentation load= 150 p.

25 30 35 40 45 50200

210

220

230

240

250

260

270

280

HVDPII

Linear fit

V

icke

rs H

ardn

ess (

HV

)

Martensite fraction (%)

HVDP

= 164.4+2.21 * VMV

Figure 3.38: Dependence of the Vickers microhardness on the martensite fraction of the grade DP II,annealed at 840℃. The dashed line represents a linear fit to the experimental data. Indentationload=150 p.

The results of the microhardness measurements are presented in Figure 3.38. The data are

plotted as a function of the martensite fraction instead of the applied cooling rate. The micro-

70

hardness of the dual-phase steel increases nearly linearly with increasing martensite fraction.

According to this linear relationship, the microhardness of a completely martensitic material

(produced from the same alloy) should be in the order of 385HV. The measurements per-

formed using a fully martensitic material gave a mean microhardness value of ∼ 400HV, very

close to the calculated value. The small deviation is within the acceptable error limits of the

measurement.

3.2.3 Hole expansion

To extend the knowledge about the formability behavior of the dual-phase steels investigated

in this work, the hole expansion property (also known as hole flangeability) of selected spec-

imens was determined using the testing equipment described in section 2.8.3. Due to limited

availability of the pre-processed laboratory material, only the samples of the industrial grade

R were investigated in the entire range of cooling rates and annealing temperatures. A small

number of specimens of the grade M annealed at selected extreme conditions was also tested.

Images of a sample after hole expansion are shown in Figure 3.39. The specimen was

produced from the grade R, annealed at 800℃ and quenched with 60K/s. According to the

method, the punch stops when the first crack formed in the lip of the hole propagates through

the whole thickness of the material. Since the judgement of the “complete” crack is made

visually, a further - though slight - hole expansion may take place during the time that intervenes

between the crack observation and the stop of the punch. It is therefore possible that a second

crack also occurs, as pointed out in Figure 3.39.

It was observed that the crack occurrence is somehow associated to the rolling direction

(RD); in the majority of the examined samples the first tear-crack occurred parallel to the

rolling direction. In case a second crack was additionally formed, this was oriented either

parallel or diagonal to the rolling direction. No cracks in the transverse direction were observed.

To determine the average final diameter of the ruptured hole, the measured perpendicular

diameters were selected to be diagonally oriented to the rolling direction, as illustrated in

Figure 3.39-b. The hole expansion ratio η was then calculated using the equation 2.2, where

df= (d1 + d2)/2. Two specimens were tested for each annealing condition and the average value

of η was calculated.

The results of the hole expansion experiments are presented in Figure 3.40. Independently

of the heat treatment parameters, the experimental data can be classified into two groups

according to the achieved η-value. The hole expansion ratios of grade R are at least two times

higher than the ratios of grade M, for both annealing temperatures and for all cooling rates.

The η-values of R lie between 27 and 35 %, exhibiting a general tendency to decrease as the

cooling rate increases from 5 to 80K/s. This tendency is also present in the samples of grade

M, for which the η-values do not exceed 17%.

71

(a) overview of the specimen (b) “zoom in” on the cracks

Figure 3.39: Specimen of the grade R after the hole expansion test. The formed cracks, the evaluatedperpendicular diameters and the rolling direction are also specified.

The analysis and the interpretation of the hole expansion results was not feasible on the

basis of mechanical (tensile) properties; both the strength and elongation characteristics of

the investigated grades, that is the absolute values as well as the yield and elongation ratios

(Rp0.2/Rm and Ag/A), are lying on comparable levels, with the properties of M to be the most

favorable. In an attempt to explain the remarkable difference in η-values, the microstructures

of selected data points of the diagram of Figure 3.40 are shown in Figure 3.41.

It is of great interest to compare the microstructures between the positions P1↔P4, P2↔P5

and P3↔P6, corresponding to pairs of samples quenched with the same cooling rate (5, 20

and 80K/s for each pair, respectively). The microstructures of the samples P1 and P4 are

quite similar regarding the grain sizes and the martensite fractions. The slightly higher frac-

tion of a third (bainitic) phase in P1 and the more homogeneous microstructure of P4 could

be considered as differences. However, the hole expansion property of P1 reaches 34 % in-

stead of 17% of P4. The same microstructural approach can be made for the positions P3

and P6; for almost identical microstructures, the η-values exhibit a two-to-one relationship

(ηR ' 2 ηM), differing by approx. 16% in absolute values. The last pair involves the samples

P2 and P5, both selected from the grade R. The only observed differences in microstructure are

the more pronounced appearance of martensite bands and the existence of a small amount of

tempered martensite in P2, resulting in a 5 % higher hole expansion ratio. It should be pointed

out that the images of P1 and P2 show the microstructures with the highest η-values achieved.

The microstructural investigations conducted in the light microscope could not provide ade-

quate information for the correlation of the hole expansion property with the microstructural

72

0 30 60 9010

15

20

25

30

35

40

P6

P3800°C

P5

P2

P4

800 °C 840 °C

= (d

f-d0)/

d 0 x 1

00 (%

)

Cooling rate (K/s)

R

M

P1

Figure 3.40: Dependence of the hole expansion ratio η on the cooling rate, for each annealingtemperature applied. The microstructures of the numbered points are presented in Figure3.41.

(a) point P1 (b) point P2 (c) point P3

(d) point P4 (e) point P5 (f) point P6

Figure 3.41: Light microscope images corresponding to the hole expansion samples. The specimensare etched with LePera.

73

characteristics of the samples. The detectable but slight microstructure variations are not suffi-

cient to justify the remarkable difference in hole expansion performance. Further investigations

in FEG-SEM were focused on the study of fracture surface as indicated in Figure 3.42.

Figure 3.42: Overview of the fracture surface of both sides of the crack (magn. 100×). The macro-scopical changes in the specimen’s shape/morphology along the cross-sectional area of the hole in thedirection of the punch can be observed.

Macroscopically, the material exhibits ductile fracture. Since overload is the principal cause

of the fracture, the material fails by a process known as void coalescence. The voids nucleate

at regions of localized strain discontinuity, such as inclusions, grain boundaries, second phase

particles and dislocation pile-ups. As the plastic strain increases, the voids grow, coalesce and

finally form a continuous fracture surface. As a result of this fracture mode, numerous cup-like

depressions are formed, referred as dimples [85, 86]. Figure 3.43 shows the fracture surface

of a hole expansion sample of grade R. The sample was annealed at 800℃ and quenched with

40K/s, exhibiting a hole expansion ratio of 29%. The relatively small size of the dimples reflects

the homogeneous distribution of the voids and the big number of nucleation sites. Large dimples

(corresponding to large voids) were sporadically detected, often associated with the presence of

large non-metallic inclusions. Along with the dimples, the formation of mutually parallel other

microcracks (Figure 3.43-a) was observed, resembling the preferable void nucleation in the

martensite bands. These microcracks propagate along the hole expansion crack in a direction

independent of the loading direction. It was also noticed that the dimples at locations away

from the hole lip appear equiaxed, as if they were formed under conditions of uniaxial plastic

strain in a direction perpendicular to the fracture surface.

As approaching the lip of the hole, the fracture characteristics become consistent with the

complex deformation mode applied during the hole expansion test (Figure 3.43-b). The dim-

74

ples slightly change their shape by becoming shallower, an effect which may be attributed to

the joining of voids by shear along slip bands. On the edge of the crack where the impact of

shearing becomes stronger, the dimples appear elongated with the one end being open, that is

the dimples are not completely surrounded by a rim. Cleavage fracture (transgranular fracture

mode) of coarse martensite grains which undergo very little plastic deformation was sparsely

observed.

(a)microcrack propagation along the crack (b) fracture surface on the lip of the hole

Figure 3.43: FEG-SEM micrographs of the fracture surface of the HE sample. The initials PD standfor the punching direction.

(a) cleavage fracture among the dimples (b) elongated dimples and shearing effects

Figure 3.44: Fracture surface of grade M, focused on the edge of the crack (on the lip of the hole).The arrow is pointing in the punching direction.

The FEG-SEM images of Figure 3.44 present fracture surfaces of the crack, focused on

75

the vicinity of the hole lip. The investigated M grade sample was annealed at 840℃ and

quenched with 80 K/s, exhibiting an η-value of 11.4%. The formation of very small dimples

(approximately two times smaller compared to the dimples shown in Figure 3.43) confirms

the existence of a great number of nucleation sites, representing faithfully the ultrafine-grained

microstructure of the as-annealed material. Cleavage fracture of the martensite grains was

frequently observed in the edge location of the crack. The well-orientated features/substructures

revealed on the cleavage surface (cleavage steps or river patterns) are formed in the direction

of loading. Furthermore, effects such as elongated dimples and shear fracture are observed due

to the complex load state (non-uniaxial loading).

76

Chapter 4

Discussion

4.1 Grain refinement and microstructure investigations

The main concept of the applied grain refinement process involved severe cold-rolling of pre-

processed (i.e. previously heat-treated) dual- or single-phase microstructures, followed by ap-

propriate intercritical or austenitic annealing. The obtained ultrafine microstructures were

qualitatively and quantitatively described in section 3.1. A very wide variety of dual- and

multi-phase microstructures was realized, offering the possibility of a detailed discussion of the

influence of the pre-processing routes and of the annealing conditions on the microstructural

characteristics, such as the formed phases, their phase fractions and grain sizes.

An important step in grain refinement is the subdivision of existing coarse grains in the as

hot-rolled material into finer ones before the final annealing process. Pre-processing before

cold-rolling seems to promote this subdivision. By water-quenching after the initial anneal-

ing procedure, the austenite to martensite transformation, which is associated with a volume

expansion of the transforming grain, introduces a large number of dislocations into the sur-

rounding ferritic matrix. This feature as well as the presence of the hard martensitic islands

enhance grain subdivision during severe plastic deformation [64, 87].

Annealing of the pre-processed cold-worked microstructures offers significant advantages com-

pared to the conventional production method of dual-phase steels by intercritical annealing of

cold-rolled ferritic-pearlitic material. Firstly, the number of the existing sites where nucleation

of new grains can occur (grain boundaries and grain triple points) is considerably higher, in-

creasing in that way the possibility to generate new grains. Secondly, the number of dislocations

introduced by deformation is high, providing the necessary driving force for rapid recovery and

recrystallization. By careful selection and matching of the annealing conditions with the indi-

vidual characteristics of the pre-processed grades, the formation of an ultrafine microstructure

can be achieved.

The ferrite grain sizes shown in Figures 3.21-3.22 resemble mean values estimated from

77

the grain size measurements. The actual distribution of grain sizes follows the scheme of

Figure 4.1, irrespective of the grade and of the annealing conditions. In the given example the

“response” of all investigated grades (M, DP I, DP II and R) to intercritical annealing followed

by rapid quenching is illustrated. The results are plotted as relative frequency over the range

of ferrite grain sizes present in the microstructure. Each set of bars represents the relative

frequency within a size interval (class) of 0.75 µm, i.e. between 0.0-0.75 µm (class 1), 0.75-

1.5µm (class 2), etc. It can be safely concluded that the finest microstructure is formed by the

martensitic grade M, where the amount/number of ferrite grains which are finer than 2.25 µm

exceeds 80%.

1 2 3 4 5 6 7 8 9 10 11 120

5

10

15

20

25

30

35

40

M DP II DP I R

Rel

ativ

e fr

eque

ncy

(%)

Size class of ferrite grains

Figure 4.1: Ferrite grain size distribution for all grades, annealed at 800℃ and quenched with 80K/s.

Indicative of the size homogeneity of the ultrafine microstructure is the size range of the

formed grains, that is the actual spread of the observed values as defined by the difference

between the coarsest and the finest measured grains (here zero µm is assumed to be the

smallest/finest value, though without physical meaning). The narrower the size range of the

formed grains the more homogeneous is the microstructure. The size range of grade M is lim-

78

ited/restricted to 6µm, with the percentage of grains coarser than 3 µm being less than 10%.

On the contrary, the grain size ranges of the laboratory grades DP II and DP I are slightly wider,

reaching 7.5µm and 8.25µm, respectively, and hence increasing not only the mean ferrite grain

size but also the size inhomogeneity of the material. This effect is even stronger in the reference

grade R, where grains as large as 9 µm are observed.

What should not be overlooked is that an important side-effect of pre-processing is the com-

plete dissolution of pearlite contained in the as hot-rolled steel. The absence of pearlite in the

cold-rolled material minimizes the required annealing time, since no extra time for dissolution of

cementite is necessary and allows the application of short holding times during final annealing.

Additionally, pre-processing turned out to be an effective way to minimize the presence of

martensite bands in the dual-phase microstructure, which are formed in the positions where the

former pearlitic bands were located in the as cold-rolled state. Pearlite banding is closely asso-

ciated with the thermomechanical history of the steel, starting already from the solidification

process and taking its final form in the hot-rolled sheet. The “necessary” but not individually

sufficient conditions that provoke banding are the microsegregation of certain alloying elements

like Mn and Cr, which effectively attract carbon (compared to other elements that reject carbon,

e.g. Si)[88], the slow cooling rates that allow the further carbon enrichment of the manganese-

rich regions by diffusion and, finally, the fineness of austenite grains by providing an adequate

number of nucleation sites. The technological importance of banding is focused on its influence

on the mechanical properties of the sheet, especially in the case of dual-phase steels where the

pearlite bands are replaced/substituted after annealing by martensite bands. Though no sig-

nificant influence of banding on the material’s tensile properties has been reported up to now,

a remarkable reduction of the impact toughness due to increased anisotropy was found [89–91].

According to several experimental and modelling studies [88, 91–94] pearlite banding can

be prevented if one of the above conditions is not any longer fulfilled, with the cooling rate

to be the most effective way to achieve it. Since the aim of this work was to produce ul-

trafine microstructures, the option/solution of austenite coarsening would lead to undesired

coarse-grained martensite. By means of pre-processing, the steel is initially heat-treated in

the intercritical or in the austenitic field so that the pearlite bands are dissolved and trans-

formed into austenite. Though the holding time in the pre-processing stage is relatively short,

a partial redistribution of the alloying elements (mainly of carbon due to its higher diffusion

coefficient) is achieved. This effect is much more pronounced in the case of grade M (Tan=

900℃, austenitic field). Water-quenching (very high cooling rates) from the pre-processing

temperature minimizes the number of martensite bands and also prevents a possible formation

of new pearlite bands. As a result, the starting microstructures of grades DP I, DP II and M

(before and after cold-rolling) possess a significantly lower number of bands in comparison with

the ferritic-pearlitic grade R.

79

4.2 Mechanical properties

Yielding behavior

The influence of grain refinement together with the other microstructural characteristics on

the mechanical properties of the dual-phase steels was studied thoroughly because of its great

technological interest. Special attention was paid to the yielding behavior of the material and

its dependence on the grain size of the matrix phase (ferrite) as well as on the volume fraction

of the second phase (martensite). The influence of the cooling rate on the yield strength (Rp0.2)

was shown in Figure 3.23, where an increase of the yield strength with increasing cooling

rate was observed. In case that this increase is attributed through a simplistic approach only

to one of the both strengthening mechanisms (either grain refinement of ferrite or increase of

the martensite fraction), then two different paths of argumentation can be followed for the

interpretation of the results, as being expressed by Figures 4.2 and 4.3.

1.0 1.5 2.0 2.5 3.0 3.5260

300

340

380

420

460

M DP II DP I R

R p0.2 (M

Pa)

Ferrite mean grain size (µm)1.0 1.5 2.0 2.5 3.0 3.5

240

280

320

360

400

440

480

520

M DP II DP I R

R p0.2 (M

Pa)

Ferrite mean grain size (µm)

(a)Tan= 800℃ (b)Tan= 840℃

Figure 4.2: Yield strength Rp0.2 as a single function of ferrite grain size.

As can be seen in Figure 4.2, the behavior of all investigated grades is similar, irrespective

of the annealing conditions: the yield strength increases by decreasing the mean grain size of

ferrite. This effect, firstly described by Hall and Petch, is well known and has been investigated

in-depth over the last decades. A linear dependence between the yield strength (actually

between the lower yield stress) and the reciprocal square root of grain size was established

from the experimental data, as stated by the Hall-Petch equation (eq. 1.5). The interpretation

of this relationship is based on the dislocation pile-up theory. When dislocations encounter

obstacles their motion is impeded, causing the stress required to continue the deformation

process to increase. Grain boundaries act as such obstacles. As dislocations pile-up at the

80

grain boundary the stress concentration in the vicinity increases, causing eventually the grain

boundary to yield, or new dislocation sources to become active in the neighboring grain.

A very rough comparison between the different grades with regard to their yield strength

“sensitivity” to the grain size could be possible by plotting the yield strengths over the reciprocal

grain size of ferrite (d−1/2F ). The slope of each curve represents a ky-parameter of the Hall-Petch

relationship, referring to the specific grade annealed at a certain temperature (Table 4.1). It

becomes obvious from the range of the obtained ky values that no consistent conclusions can

be drawn. Additionally, the calculated σ0 friction stresses exhibit negative -without physical

meaning- values for most of the curves, indicating the weak points of this analysis.

Table 4.1: Estimation of the experimental constant ky for the applied annealing conditions. The ky

value of 81.33 is obviously overestimated.

Grade M DP II DP I R

ky [Nmm−3/2], Tan= 800℃ 19.68 20.03 81.33 (?) 28.05

ky [Nmm−3/2], Tan= 840℃ 24.31 16.80 26.77 22.45

20 25 30 35 40 45 50260

300

340

380

420

460

M DP II DP I R

R p0.2 (M

Pa)

Martensite fraction (%)20 25 30 35 40 45 50 55 60

240

280

320

360

400

440

480

520

M DP II DP I R

R p0.2 (M

Pa)

Martensite fraction (%)

(a)Tan= 800℃ (b)Tan= 840℃

Figure 4.3: Yield strength Rp0.2 as a single function of martensite fraction.

On the other hand, Figure 4.3 shows the influence of the martensite fraction on the yield

strength of the dual-phase steels without taking into consideration the grain size effect. It is

generally observed that the yield strength of the steel increases almost linearly with increasing

martensite fraction. Based on the grain size measurements, it can be safely assumed that

the increase in the volume fraction of martensite is solely expressed by an increase of the

81

number of martensitic islands finely distributed in the ferritic matrix and is not accompanied

by grain growth of a certain number of martensite grains. The higher number of martensite

grains results in a larger area of the ferrite/martensite phase boundaries, increasing in that

way (proportionally) the possibility for dislocation pile-ups to occur at the ferrite/martensite

interface and, hence, the required stress concentration for these boundaries to yield [48, 95].

According to Fischmeister and Karlsson [48] the initiation of plastic deformation of a coarse-

grained ductile matrix with finely dispersed hard inclusions is governed by the yield strength

of the soft matrix. This statement is valid as long as the elastic properties (Young’s moduli) of

the two constituents are equal, which is roughly the case in dual-phase steels. Therefore, any

variation of the yield strength with increasing cooling rate -i.e. martensite volume fraction-

should be attributed to an alteration of the “in-situ” ferrite yield strength in the two-phase

mixture, assuming that ferrite is always the continuous matrix and martensite remains isolated.

To provide a reasonable explanation for the alteration of ferrite yield strength with vary-

ing martensite fraction, the thermomechanical history of the steel has to be taken into con-

sideration. The austenite to martensite phase transformation taking place upon quenching

from the annealing temperature is accompanied by a remarkable change in the volume of

the transforming region. The specific volume of body centered tetragonal (b.c.t.) martensite

(0.1298 cm3/g) is significantly higher than the specific volume of face centered cubic (f.c.c.)

austenite (0.1265 cm3/g) [49, 96], so that theoretically a volume increase of approx. 2.6 % oc-

curs. In order to accommodate the volume expansion and the shear deformation accompanying

the martensitic transformation, a large number of geometrically necessary dislocations is intro-

duced/produced in the surrounding ferrite grains. The higher the martensite fraction (in terms

of larger number of grains) the higher the number of locations at which ferrite has to accom-

modate the volume and shape effects of the transformation and hence the more homogeneous

the distribution of the mobile dislocations in the microstructure [35].

82

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.82.80

2.84

2.88

2.92

2.96

3.00

3.04

3.08

3.563.583.603.623.64

1.00

1.02

1.04

1.06

1.08

1.7

1.8

c/a

c

a

a

a o

f aus

teni

te (Å

)a

and c

of m

arte

nsite

(Å)

C (wt. %)

Axi

al ra

tio c

/a o

f mar

tens

ite

Figure 4.4: Lattice constants of tetragonal martensiteand austenite in quenched carbon steels (After Hondaand Nishiyama [97]).

However, it should always be kept

in mind that the “intensity”/impact

of the martensitic transformation re-

garding volume expansion, shear de-

formation and changes in crystal

structure is strongly dependent on the

carbon content of austenite prior to

the transformation. A body centered

tetragonal structure is favored in car-

bon steels. As reported by Honda and

Nishiyama [97], the lattice parameters

change as follows: the a axis slightly

shrinks and the c axis markedly grows

with an increase in carbon content

above 0.25 %. The degree of tetrag-

onality, as measured by the c/a ratio,

increases linearly with carbon content

according to the the following empiri-

cal relationship [98]:

c/a = 1.005 + 0.045 (wt. % C). (4.1)

As a consequence of the eq. 4.1, the

volume of the unit cell increases also

roughly linearly with increasing car-

bon content.

In the present study, the wide variety of investigated microstructures was produced from

a low-alloyed dual-phase steel having a fixed chemical composition. The carbon content of

the alloy is Calloy = 0.1wt.% (Table 2.1). If it is assumed that this carbon is completely

partitioned between the ferrite and the austenite in the intercritical region, ferrite is carbon-

free after quenching and no carbides are present in the dual-phase microstructure, then the

final carbon content of the martensite (CM) would be inversely proportional to the martensite

volume fraction (VM):

(wt.% CM) = (wt.% Calloy)/VM. (4.2)

For example, for the lower and the higher martensite fraction of grade DP II annealed at 840℃

83

(Figure 3.20-b), it can be calculated that CM, 0.28=0.357 wt.%C and CM, 0.49=0.204wt.% C

(the numerical indices refer to martensite fractions). Substituting these values into the eq. 4.1

yields (c/a)M, 0.28=1.021, while (c/a)M, 0.49=1.0142. This difference in the axial ratios corre-

sponds to a higher volume of the lower martensite fraction unit cell by ∆V =∼ 0.65-0.70%.

As a result, the impact of the martensitic transformation on the surrounding ferritic matrix

should be definitely stronger for the high carbon content grains (i.e. low martensite fraction mi-

crostructures). However, in a macro-scale approach to the example, the total volume expansion

effect is controlled by the much higher martensite fraction.

Liedl et al. [99, 100] studied the impact of the austenite to martensite phase transforma-

tion on the stress-strain behavior of a dual-phase steel by constructing a micro-mechanical

model employing fully three-dimensional finite element calculations. The simulations revealed

that the ferritic matrix (skeleton) is work-hardened during quenching, due to the local internal

stresses caused by the volume expansion which accompanies the martensitic phase transforma-

tion (Figure 4.5). The extent of this work hardening was found to depend on the martensite

content. The calculations were supported by ultramicrohardness measurements in ferrite grains:

it was clearly determined that the hardness of ferrite is higher close to ferrite/martensite phase

boundaries than close to ferrite/ferrite grain boundaries.

Figure 4.5: Plastic equivalent strain in a center section through a unit cell, after the simulatedcooling process. The dark (red) areas adjacent to the martensite grain correspond to high plasticequivalent strains [99].

Transmission electron microscopy was employed to provide evidence for the microstructural

effects of the martensitic transformation. Figure 4.6 shows high resolution TEM images of the

grades DP I and M in the as-annealed condition. Irrespective of the annealing conditions, it

84

can be observed that the dislocation density in ferrite (white grain) is remarkably higher in the

vicinity of martensite (dark grain) and decreases as moving away from the ferrite/martensite

grain boundary towards the interior of the grain. Analogous TEM observations were reported

by Sherman [101]. According to his measurements, the total average ferrite dislocation density

is a linearly increasing function of the martensite fraction, taking values in the order of 109/cm2.

Furthermore, the ferrite dislocation density in grade M (Figure 4.6-b) is much higher than in

grade DP I (Figure 4.6-a).

(a)DP I, Tan=840℃, CR= 5 K/s (b) M, Tan= 800℃, CR= 80K/s

Figure 4.6: Bright field TEM images of the dual-phase grades DP I and M in the as-quenchedcondition.

Each of the two approaches discussed above throws light on the influence of the microstruc-

tural characteristics on the yield strength of the dual-phase steel from a different point of

view. However, these arguments cannot hold independently from each other within the scope

of this study, since an interaction between martensite fraction and ferrite grain size is un-

avoidable (Figure 3.22). Several experimental and theoretical works on the dependence of the

yield strength (or of the flow stress) on the grain size in dual-phase steels have been published

[14, 95, 102, 103], most of them trying to incorporate the effect of martensite fraction. All these

studies showed that a linear relationship between the yield strength and the reciprocal square

root of grain size was satisfied and could be expressed by appropriate modifications of the

85

Hall-Petch relation. Nevertheless, no general agreement on the conclusions of yield strength

dependence on grain size has been reached, mainly because of the different approaches con-

cerning the definition of the controlling grain size. Balliger and Gladman [103], for example,

defined the grain size as the average martensite (second phase) island diameter, Reuben and

Baker [102] used a weighted average of the phase specific grain sizes (by applying a “law of mix-

tures”, d 1/2c = (1 − VM)d

1/2F + VMd

1/2M ), while Chang and Preban [104] have accepted/adopted

the average ferrite grain diameter as equivalent to the mean free path in ferrite. In each case,

the investigators have emphasized the importance of the grain size of one constituent phase

(ferrite or martensite).

It should always be kept in mind that the existence of any kind of interdependencies among

microstructural parameters/variables makes multi-variable quantitative correlations extremely

risky, so that the requirement for a single-variable data analysis becomes mandatory. To enable

a fundamental Hall-Petch-type approach to the net impact of grain refinement on the yield

strength, coarse-grained dual-phase specimens were produced from the industrial as hot-rolled

strip. The coarsening procedure involved annealing in the austenitic region (Tan=1200℃) for

a relatively long holding time to ensure homogenization of the microstructure (tan=10min),

followed by very slow furnace cooling (CR=∼ 1K/min) down to the quenching temperature

(TQ=720℃) and finally oil quenching to room temperature. The quenching temperature was

selected in the low zone of the intercritical region, but high enough to avoid any precipitation

of carbides. The obtained microstructures are shown in Figure 4.7, exhibiting a martensite

fraction of approx. 50-55% and a mean ferrite grain size of dF=12-13µm, six to seven times

coarser than the ultrafine-grained material of analogous martensite fraction (e.g. grade M).

The mechanical properties of the coarse-grained material were determined by tensile testing.

(a) (b)

Figure 4.7: Coarse-grained ferritic-martensitic microstructures (LePera).

86

Figure 4.8 shows the strength dependence of dual-phase microstructures containing 50%

martensite on the grain size of ferrite, at different levels of plastic strain. Since the plotted

data refer to similar martensite fractions (∼ 50-55 %), it was assumed that grain refinement is

the dominant strengthening mechanism. Accepting a priori a linear relationship between the

strength and the reciprocal square root of the mean ferrite grain size (Hall-Petch relation), the

experimental data were fitted by linear regressions. The fitting curves (trendlines) and their

mathematical expressions are also shown in the diagram. At a first qualitative glance, the

approach seems reasonable: the yield strength of the steel increases as the grain size of ferrite

decreases and the estimated σ0 friction stresses take reasonable values.

2 7 12 17 22 27 320

100

200

300

400

500

600

700

800

900

Extrapolation of theor. values to smaller ferrite grain sizes according to published exp. data after:

Hodgson Hickson Han

p=0.2%

u

p=1.0%

Rp1.0= 401 + 10.9* d -1/2F

(MPa)

Rm= 531.3 + 13.4* d -1/2F

(MPa)

Rp0.2= 254.2 + 9.88* d -1/2F

(MPa)

Pure Ferrite (theor.) Linear fit, theor. Fine DP (exp.) Coarse DP (exp.) Linear fit, exp.

R p0.2, R

m (M

Pa)

d -1/2F

(mm-1/2)

Rp0.2= 113.4 + 17.4* d -1/2F

(MPa)

Figure 4.8: Grain size dependence of the strength of a dual-phase steel containing roughly 50%martensite, approached by a Hall-Petch relationship. Literature data referring to pure ferrite enablea representative comparison.

To allow the evaluation of the experimental results, theoretical data for the pure ferrite yield

strength dependence on the grain size are also included in Figure 4.8. The values are estimated

from the empirical relationship of Gladman and Pickering [105, 106], taking into account the

alloy content (weight %) in Mn, Si and free nitrogen (Nf) and the grain size of ferrite (dF in

mm):

σy[MPa] = 53.9 + 32.3 %Mn + 83.2 %Si + 354 Nf + 17.4 d−1/2F . (4.3)

87

For the chemical composition of the investigated steel -and for zero free (soluble) nitrogen-

eq. 4.3 takes the form:

σy[MPa] = 113.4 + 17.4 d−1/2F . (4.4)

Furthermore, the estimated ferrite yield strengths from the Gladman and Pickering equation

were extrapolated to smaller ferrite grain sizes. Experimental data of ultrafine ferritic mi-

crostructures (dF=1.5-2µm) produced from low-alloyed steels verify the validity of this extra-

polation [57, 59, 69, 107, 108].

The most significant observation is that the grain size dependence of the dual-phase steel, as

expressed by the slope of the curve, is remarkably smaller compared to that of the ferritic steel,

taking values ky,DP= 9.88Nmm−3/2 and ky,F= 17.4Nmm−3/2 respectively. Analogous results

were produced by Liu [109] for ultrafine-grained mild steels, where the slope of the Hall-Petch

linear fit expressed in Nmm−3/2 was found equal to 7.48. Priestner et al. [52] have reported a

Hall-Petch slope of 9.75Nmm−3/2 for ultrafine ferrite grain structures in Nb microalloyed low

carbon steels. According to Tanaka [14], who has also observed very low ky values, a possible

explanation could be that in as-quenched dual-phase steels the stress required to unpin the

dislocations is relatively low due to the existence of unlocked dislocation sources.

For a better established comparison, more data points concerning the grain size variation

at a roughly constant martensite fraction should be acquired, especially for lower martensite

fractions. This was not possible within the framework of this study due to the limited flexibility

in simultaneous control of grain size and martensite fraction. Of equal importance becomes

the consideration of martensite grain size; it is evident that the production of coarse ferrite

grains was inevitably accompanied by the formation of coarse martensite grains (Figure 4.7).

Consequently, the impact of the martensitic transformation on ferrite should be different in the

two cases (coarse and fine microstructure) despite the similar martensite fraction.

Although the analysis is focused on the yield strength, interesting observations arise from the

comparison of the “Hall-Petch slope” of the different strength levels (that is Rp0.2, Rp1.0 and

Rm). At higher plastic strains the grain size dependence of strength becomes more pronounced,

so that the slope of the lines gradually increases from 9.88Nmm−3/2 to 13.4Nmm−3/2. This

effect has been observed by other authors, too [14, 48, 50, 110], and was attributed to the

rapid increase in the work hardening of the material taking place in the early stages of the

deformation. The inverse effect takes place in materials exhibiting discontinuous yielding, e.g.

ferritic-pearlitic steels [14].

88

Strain hardening

The key factor that differentiates dual-phase steels from the variety of the ferritic steels (e.g.

mild steels, IF-steels, conventional HSLA, etc.) with respect to the deformation characteristics,

is the absence of yield point elongation, i.e. the ferritic-martensitic steels exhibit continuous

yielding. This feature together with the very high initial strain hardening rates account for the

excellent formability properties of dual-phase steels.

It is widely accepted that the main reason for the continuous yielding behavior as well as

for the high work hardening rate of dual-phase steels is the existence of mobile dislocations,

introduced into the ferritic matrix as a result of the austenite to martensite transformation. The

movement of these mobile dislocations results in the elimination of yield point elongation and in

the observed low yield strength. The interaction of the dislocations with each other and with the

finely dispersed martensite grains results in the high strain hardening exponent. Hahn’s model

[111] for discontinuous yielding predicts that mobile dislocation densities of 102-104/cm2 lead

to discontinuous yielding, while mobile dislocation densities of 106-108/cm2 result in continuous

yielding behavior. Although dislocation densities of 106-108/cm2 are typical in hot-rolled or

annealed steels, the majority of these dislocations are “immobile”, i.e. pinned by interstitial

atoms [36, 111]. As previously referred and cited, the total average ferrite dislocation density

in dual-phase steels is martensite fraction dependent and ranges between 1-1.5×109/cm2, hence

satisfying Hahn’s criterion.

However, an important prerequisite for continuous yielding is that the dislocations produced

in ferrite during the martensitic transformation remain mobile at room temperature. These

dislocations can be pinned by interstitial atoms if a sufficient time at a certain temperature is

available for the necessary diffusion. Since these dislocations are generated below the martensite

start temperature (Ms), the heat treatment steps following the transformation can be critical,

e.g. in the case that an overaging stage is required. In the annealing cycles applied in this work

this effect has no practical importance, because the time that passes between the Ms and room

temperature is too short for any diffusion to happen (even for the higher observed Ms).

According to Speich and Miller [42], the rapid increase in the work hardening of dual-phase

steels at low strains is the outcome of three interacting parameters. Firstly, the residual stresses

generated during the austenite to ferrite transformation upon cooling are relieved already by

small amounts of plastic deformation. Secondly, the dislocation density in ferrite is increased

by generation of both statistically stored and geometrically necessary dislocations [112]. The

statistically stored dislocations are those resulting from simple work hardening of ferrite while

the geometrically necessary dislocations are created by the need to maintain contact between

the two phases during plastic deformation (compatibility). Thirdly, because the plastic in-

compatibility of the two phases cannot be accommodated completely by plastic deformation,

stresses are created within the martensite phase, which are compensated by back stresses in

89

ferrite. These back stresses impede dislocation movement in ferrite.

An accurate description of the plastic behavior of dual-phase steels is of great technologi-

cal interest, since it is essential to understand the forming characteristics of the material. As

described in section 3.2.1, the determination of the n-value is the most common way to charac-

terize the strain hardening behavior of the steels. Several idealized mathematical descriptions

-known as strain hardening laws- have been developed, with the commonly used [113]:

σ = KH εnH (Hollomon, see eq. 3.1), (4.5)

σ = σ0 + KL εnL (Ludwik), (4.6)

σ = KS (ε1 + ε)nS (Swift), (4.7)

where σ and ε stand for the true stress and the logarithmic (or “true”) strain while K, ε1 and σ0

are constants. The n-values are usually determined from a double logarithmic plot of true stress

vs. logarithmic strain by linear regression. Correlations between the n-values and mechanical

properties of dual-phase steels have been reported [35, 41, 114], however no general conclusions

have been reached. From the mathematical point of view all the above equations (eq. 4.5 to

4.7) are well-founded and, furthermore, almost every stress-strain curve can be described under

the appropriate assumptions by nearly every law. Nevertheless, the fact that the n-value can

be strongly strain-dependent can lead to great deviations and misinterpretations, even between

the results of the different approaches.

Ratke and Welch [115] have defined a differential n-value, which is independent of strain

hardening laws and physically expresses the instantaneous slope of the stress-strain curve in its

double logarithmic form:

n(ε) =dlnσ

dlnε| ε,T=constant . (4.8)

An analysis according to this method is shown in Figures 4.9 and 4.10. The differential n-

values for grades DP I and M are calculated from the strain-stress raw data, without making

any prior assumptions about the validity of an arbitrary strain hardening law.

The two curves of Figure 4.9 refer to the same heat treatment schedules, involving austenitic

annealing at Tan=840℃ and quenching with a high cooling rate (CR=80 K/s). Though the

mean ferrite grain size of the two specimens at the specific conditions is similar (dF,DP I=1.59µm

and dF,M=1.57µm), their martensite fractions differ by approximately 8.5% (VM,DP I=45.6%

90

0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0.160.0

0.2

0.4

0.6

0.8

1.0

0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0.160.0

0.2

0.4

0.6

0.8

1.0

DP I MRp0.2 395.7 483.6 MPaRm 809.1 847.5 MPa Ag 12.8 10.6 %A 18.5 14.7 %

n4-6= 0.178

n av,M= 0.

122 M

n4-6= 0.138n2-4= 0.166

n av,M= 0.

145

n =

d(ln

) / d

(ln)

Logarithmic strain ( )

n av,M= 0.

176

n av,D

PI= 0.

152

n6-10= 0.152

DP I

n2-4= 0.212

n av,D

PI= 0.

181

n av,D

PI= 0.

219

(change in strain rateduring the tensile test)

YS

Figure 4.9: Differential n-value of the grades DP I and M, annealed at Tan=840℃ and quenchedwith CR=80 K/s.

and VM,M=54.1%). Both curves follow the same shape: the n-value increases steeply at very

small strains reaching a maximum value, followed by a sharp drop until it reaches a “plateau”

and then a smooth decrease takes place at higher strains. The last portion of the curve is

dependent on the uniform elongation; at strains higher than the necking point, the decrease

of n-value becomes more rapid until it equals zero. The peak observed at logarithmic strains

around 0.015 (higher than the strain where yield strength is recorded) is due to a slight change

in the strain rate, as previously discussed (see 3.2.1). Unfortunately, at very low strains the

low density of the recorded stress-strain data does not allow reliable assessments.

The n-values as demonstrated in Figure 3.28 are fairly reproduced from the curves by es-

timating the mean n-values between the same logarithmic strain intervals. Practically, the

example shows that among two ultrafine-grained ferritic-martensitic microstructures, the one

with the lower martensite fraction exhibits a more pronounced strain hardening for the whole

range of logarithmic strains, which finally leads to higher values of uniform and total elonga-

tions. It is worth mentioning that in the strain interval of interest, that is 0.02 < ε < 0.10, both

91

curves are decreasing with almost equal rates. From one point of view, this is a clear indication

that the initial stages of strain hardening are very crucial for the formability characteristics of

the steel. However, they do not provide any indication of the sudden drop of n-value occurring

at strains beyond necking.

0.00 0.02 0.04 0.06 0.08 0.10 0.12 0.14 0.16 0.180.0

0.2

0.4

0.6

0.8

1.0 DP I MRp0.2 320.2 371 MPaRm 718.9 728.7 MPa Ag 14.6 11.5 %A 20.7 15.4 %

n =

d(ln

) / d

(ln)

Logarithmic strain ( )

n av,M= 0.

154

n av,M= 0.

187

n av,M= 0.

218

n4-6= 0.184n2-4= 0.215

n4-6= 0.216

M

n av,D

PI= 0.

181

n6-10= 0.182

DP I

n2-4= 0.251

n av,D

PI= 0.

219

n av,D

PI= 0.

255

(change in strain rateduring the tensile test)

YS

Figure 4.10: Differential n-value of the grades DP I and M, annealed at Tan=840℃ and quenchedwith CR=10 K/s.

Qualitatively, these results are in good agreement with the works of Davies [40] and Hayami

[34] -though here the results refer to much finer microstructures- where higher martensite

volume fractions are accompanied by a decrease in the n-value. On the contrary, Karlsson and

Sundstrom [47] have proposed, that the strain hardening rate increases with increasing volume

fraction of the second phase (here martensite) and/or by increasing the difference in hardness

between the two phases. While the findings of this study do not agree with the first half of

his proposal, the second part turns out to be of great interest, as will be discussed later in

this chapter. Obviously, since the investigations involve only one chemical composition, it is

92

not possible to succeed in fulfilling both Karlsson’s requirements: the formation of either a

high martensitic grade with soft martensite or a low martensitic grade with hard martensite is

feasible.

Additionally, Ashby [112] stated that the compatible deformation of a soft matrix containing

hard, less deformable particles requires the generation of plastic strain gradients (i.e. geometri-

cally necessary dislocations) within the more deformable phase. These plastic strain gradients,

which contribute to work-hardening by impeding the motion of other moving dislocations, are

evidently related to the type and the characteristics of the second phase, e.g. the hardness of

martensite. Speich and Miller [42] verified this statement in dual-phase steels, by investigating

microstructures with different carbon contents of martensite; steels with a high carbon marten-

site (CM=0.4wt.%) exhibited higher work hardening rates than the steels with low carbon

martensite (CM=0.2wt.%). The phenomenon was attributed to the greater plastic incompati-

bility of the harder martensite grains with the ferritic matrix and to the higher residual stresses

associated with their lower transformation temperatures.

On the other hand, Figure 4.10 provides the differential n-values of the same grades as before

(DP I and M) after annealing at Tan=840℃ and quenching with CR=10K/s. This comple-

mentary example involves dual-phase microstructures containing identical martensite fractions

(VM∼= 35.7%) which correspond to different grain sizes (dF,DP I=2.09µm and dF,M=2.54µm).

Here, the finer-grained grade exhibits higher n-values and, consequently, better elongation prop-

erties. Again, the strain hardening behavior at high logarithmic strains cannot be anticipated.

At a constant martensite volume fraction it can be assumed that the in-situ properties of

ferrite and martensite remain constant, so that the hardness ratio of the two phases should also

remain constant. An increase in grain size results in a decrease of the number of ferrite grains

adjacent to martensite and influenced by the martensitic transformation. Hence, the coarse-

grained material would exhibit lower initial dislocation densities distributed inhomogeneously

compared with the fine-grained material, leading to lower strain hardening rates. Analogous

results where reported by Matlock [35].

Deformation behavior at high strains

A metallographic study of the tensile deformation and fracture of the investigated dual-phase

steels has been already presented in the section “Microstructure at high strains”. Light and

electron microscope images of the fracture profiles have revealed two distinct modes of defor-

mation, based on the type (“clear” or autotempered) and the strength-hardness of martensite

present in the microstructure. These characteristics are, however, dependent on the heat treat-

ment schedules and especially on the cooling rate after annealing; low cooling rates produce

hard martensite enriched in carbon while high cooling rates result in low carbon, thus softer,

93

martensite.

The deformation mechanism can be comprehensively described in the strain region between

the initial stages of plastic deformation and the point where necking occurs. According to

Rashid [84], deformation occurs first in the ferrite which is the continuous soft phase. When

reaching its maximum capacity for being strained, ferrite matrix transfers strain across the

ferrite-martensite interface, leading to the deformation of the martensite islands. The nature

(and the properties) of martensite determines the extend of its plastic deformation and defines

the occurrence of the two deformation modes.

As shown in Figures 3.29, 3.30-a and 3.31 for low martensite fraction grades, no pronounced

deformation of the hard martensite grains takes place, even in the necking region. In some cases,

a slight rotation of the martensite grains towards the tensile direction could be identified. On

the contrary (Figures 3.30-b and 3.32), the softer martensite of the high martensite fraction

grades undergoes extensive plastic deformation, occurring mainly in the neck of the specimen.

Irrespective of the deformation type, all grades exhibit macroscopically a dimple-like ductile

fracture. Generally, ductile fracture begins with the formation of voids (mostly after necking),

which nucleate at sites of localized strain discontinuity such as non-metallic inclusions, grain

boundaries, phase boundaries, second phase particles and dislocation pile-ups. With increasing

plastic strain these voids grow further, coalesce into cracks and finally form a continuous fracture

surface.

However, it was found that the void formation process is different in the grades examined and

is affected by the in-situ martensite properties. Lower fractions of harder martensite promote

the formation of a high number of voids, distributed not only close to the fracture but over the

whole necking area. In that case, the ferrite-martensite interfaces act as the preferable void

nucleation sites. If martensite bands are present, then voids are formed around and along these

bands, either by decohesion of martensite grains in contact or by fracture of martensite grains.

Higher fractions of softer martensite result in a limited number of voids, strictly located in the

fracture region. The few voids are formed at the ferrite-martensite interface while no cracking

of martensite was observed.

Since cracking or decohesion of the martensite grains is involved in the void nucleation and

void growth process, the post-uniform elongation (A-Ag) is expected to depend both on the

amount of martensite and on the carbon content of martensite. A number of simultaneous

opposing effects caused by the variations in quantity (fraction) and in quality (composition,

properties) of martensite should be taken into account. This fact is critical for the evaluation

of results in the present work, since only one carbon content is used and therefore CM is

inversely proportional to the martensite fraction. Higher fractions of martensite decrease the

spacing between the voids, so that the plastic strain required to join the voids is lower and,

consequently, should result in a lower ductility. On the contrary, the softer martensite grains

94

yield at lower values of plastic strain (in locations where the martensite yield stress is exceeded)

and so playing an active role in the deformation process, do not easily decohere and do not

crack. In an equivalent manner, martensite grains with higher carbon contents crack or decohere

more easily, so that the ductility should also decrease at high values of CM, which in this study

represents low martensite fractions.

It becomes obvious that a correlation of the microstructural characteristics and the deforma-

tion mechanisms with the elongation properties is not a simple task. An optimization of the

desired microstructures should involve all the possible parameters, esp. the contradicting ones

(martensite fraction, carbon content of martensite, morphology of martensite).

Micro- and ultramicrohardness measurements

Ultramicrohardness measurements have revealed that the in-situ hardness of martensite does

not remain constant over the range of the applied cooling rates. The highest measured value

(HVM=730HV) corresponded to a cooling rate of 5 K/s while the lowest (HVM∼= 390HV) to

a cooling rate of 80 K/s. This variation is intimately connected with the carbon content and

the morphology of martensite. Additional measurements in fully martensitic material of the

same chemical composition showed that an increase of the cooling rate over 80K/s does not

further influence the martensite hardness. The hardness of ferrite remained roughly constant at

240HV, indicating that no severe change regarding the partition of alloying elements in ferrite

takes place during cooling.

The influence of grain refinement on the ferrite hardness was not possible to be incorporated

into the results. It would be expected that as the microstructure becomes finer, the effect of

the increased area of grain boundaries should become greater. However, due to the equipment

limitations, ultrafine ferrite grains (dF < 4µm) could not be indented. To enable the indenta-

tion, ferrite grains coarser than 4 µm were selected. Analogous size-selective sampling was done

for the martensite grains; only grains coarser than 2.5 µm were indented.

To express the microhardness of dual-phase steels as a function of the martensite fraction, a

couple of critical assumptions had to be made. It was primarily assumed that all the grains

of each phase exhibit the same hardness. Even though such a statement could be acceptable

for ferrite grains, its validity in the case of martensite -where a wide scattering of hardness

values is observed due to the changes in its morphology and its carbon content- is questionable.

Moreover, the effect of the grain refinement of ferrite was not taken into account.

There has been an attempt to cast the experimental ultramicrohardness data taken from

specimens of the grade DP II into a linear rule of mixtures:

HVDP = VF HVF + VM HVM, (4.9)

95

where HVDP, HVF and HVM stand for the microhardnesses of dual-phase steel, of ferrite and of

martensite, respectively, while VF and VM express the volume fractions of the phases. The only

independent variables of eq. 4.9 are the volume fraction of martensite VM and the hardness of

ferrite HVF. Once VM is defined, the following two variables are accordingly directly or indirectly

defined: the volume fraction of ferrite equals VF= (1-VM) while the hardness of martensite HVM

depends on its carbon content set by VM. Though HVF remains nearly constant, for the needs

of the example it will be considered as variable. After the necessary substitutions, eq. 4.9 takes

the form:

HVDP = (1− VM) HVF + VM HVM, (4.10)

or equivalently,

HVDP = HVF + VM (HVM −HVF). (4.11)

The values of HVF and HVM were experimentally determined (Figure 3.36). The estimated

microhardnesses of the dual-phase grade DP II are appended to the microhardness diagram

of Figure 3.38, marked with black diamond symbols as demonstrated in Figure 4.11. Obvi-

ously, the dependence between the calculated microhardness and the martensite fraction is not

linear, since eq. 4.11 includes the product of two interdependent parameters (VM HVM). The

estimated results strongly deviate from the measured microhardness values, even though this

deviation becomes smaller with increasing martensite fraction. This tendency indicates that

the assumption of the same hardness for all martensite grains at a certain volume fraction (i.e.

carbon content) is acceptable only at high martensite fractions, where the hardness scattering

is minimized.

A second approach to the estimation of the dual-phase steel microhardness by a rule of

mixtures was attempted, assuming that both HVM and HVF remain constant. In order to

perform a conservative analysis, the minimum hardness of martensite was selected (395HV at

VM=100%) while the hardness of ferrite was set to a mean value of 240HV. Then, a linear

relationship between the steel’s microhardness and the martensite volume fraction is obtained

from eq. 4.11, plotted in Figure 4.11 within a volume fraction range of 20% to 100% marten-

site. The calculated values are consistently overestimating the experimental ones, even for the

smallest hardnesses used. Again, the difference in hardness decreases with increasing marten-

site fraction. When extrapolated to zero martensite fraction the red fitting curve intersects the

y-axis at HVF=240HV, a value which is by ∼ 75 HV higher than the hardness of pure ferrite

given by the extrapolation of the experimental curve (∼ 165HV). On the contrary, there is a

perfect match between the experimental values (both UMH- and MH-measurements) and the

96

20 30 40 50 60 70 80 90 100100

140

180

220

260

300

340

380

420

HVDP

= 120+2.75 * VM

HVDP

= 240+1.55 * VM

HVDP

= 164.4+2.21 * VM

HVDPII, exp. HV100% Mart., exp. Linear fit of HVDPII

mix. rule, HVF= 240 HV mix. rule., HVM= f(VM) mix. rule, HVF= 120 HV

Martensite fraction (%)

Vic

kers

Har

dnes

s (H

V)

Figure 4.11: Correlation of the experimental microhardness and ultramicrohardness values by meansof a rule of mixtures. The measurements refer to DP II grade annealed at 840℃.

calculation of microhardness for the fully martensitic material.

It is quite remarkable that the as-measured microhardness of the dual-phase microstructure

is significantly lower than the microhardness that a rough/simple rule of mixtures predicts,

even though the lowest values obtained from the ultramicrohardness measurements were ap-

plied. It is worth mentioning, for example, that a grade containing 28% martensite exhibited

a microhardness lower than the mean measured ferrite ultramicrohardness.

It has been reported by several authors [116–119] that hardness of metals or ceramics in

nanoindentation or even in ultramicrohardness measurements (low-load Vickers tests, [120])

tends to be overestimated due to a strong size-dependence of the tests. According to this

dependence, known as indentation size effect (ISE), the hardness is observed to increase with

decreasing indentation size (or, equivalently, indent load or indent depth). The effect is already

apparent at indentations shallower than 10 µm and it becomes even more pronounced in the

sub-micron depth regime. Indentation size effect could not be explained by using a conven-

tional plasticity theory-based analysis, according to which the hardness value should remain

97

independent of the applied load. Based on Ashby’s theory of geometrically necessary disloca-

tions, Fleck and Fleck et al. [121–123] have proposed and developed a strain plasticity theory

which includes an internal material length scale. The basic concept was that in the sub-micron

scale the material hardening (and yielding) is controlled by the total density of dislocations,

part of which generates and is proportional to the gradient of plastic strain. Plastic strain

gradients appear either because of the loading geometry or because the material itself is plas-

tically inhomogeneous (e.g. containing non-deforming phases, like dual-phase steels). In order

to accommodate these strain gradients surrounding the indent the generation/production of a

certain density of geometrically necessary dislocations is required, which is inversely propor-

tional to a local length scale (in microhardness tests the length scale is directly related to the

indentation size). Concerning the contribution of the dislocations in the strengthening effects,

the model presumes that the flow stress is proportional to the square root of the total den-

sity of dislocations (statistically stored plus geometrically necessary dislocations, denoted as ρS

and ρG, respectively) and that the density of statistically stored dislocations -which depends

on the uniform strain- remains constant. At large indentation sizes (ρG¿ ρS) the statistically

stored dislocations govern the flow behavior, therefore no size dependence is observed. At

smaller length scales, the contribution of geometrically necessary dislocations to hardening be-

comes more significant. Finally, after entering a critical size/length regime which depends on

the material, the geometrically necessary dislocations outnumber the statistically stored ones

(ρGÀ ρS) and take control of the hardening effect. The validity of this “extended theory” was

experimentally verified by Elmustafa and Stone [124, 125] using a combination of microhardness

and nanoindentation measurements.

The indentation sizes involved in the ultramicrohardness measurements of this study lie in

the length scale where indentation size effects take place. The applied load of 15mN produces

indents with a size of 2-2.5 µm in martensite and 3-3.5µm in ferrite grains, corresponding to

maximum indentation depths of ∼ 0.35µm and 0.5µm, respectively (note that the depth of a

Vickers micro-indentation is about 1/7 of the diagonal length), that is well inside the sub-micron

range. However, in the case of pure martensite the measured microhardness was unexpectedly

very stable for loads (and indent sizes) differing by two orders of magnitude, namely 15mN for

UMH and ∼ 1500mN for MH.

Following the concept of the previous paragraph while keeping in mind the deviations of

the rule of mixtures, it turns out that the measured ferrite hardness may be linked to the

indentation size effect or, generally speaking, to any size effect. As already referred, the plastic

strain gradients during the indentation process generate, for compatibility reasons, a number

of geometrically necessary dislocations. Recalling the discussion made previously on the flow

behavior of dual-phase steels and on the impact of martensitic transformation in the ferritic

matrix, it becomes clear that in the ferritic matrix there exists already before indenting a

98

huge number of geometrically necessary dislocations ready to flow. As the indentation load is

applied on ferrite, these dislocations accelerate the strain hardening of ferrite even further so

that eventually a kind of fictitious microhardness is measured. The additional/collateral effects

of the already work-hardened ferritic skeleton, as well as the existence of grain boundaries in the

vicinity of the indent, on the microhardness lead to significantly overestimated measurements.

To support this argumentation and to clarify the observed discrepancy in the results, inden-

tations in coarse ferrite by applying a higher load had to be conducted. Since the production

of pure ferritic material with the same alloying composition was not feasible, specimens of the

coarse-grained dual-phase steel of Figure 4.7 containing 50 % martensite were used. Addi-

tionally, even coarser ferrite grains from the same dual-phase grade (though with much lower

martensite fraction) were indented to check whether there is an influence of the ferrite grain

size. Figure 4.12 presents the Vickers pyramid impressions in coarse ferrite grains. Four in-

dentations in each micrograph are shown, the two small ones corresponding to a load of 20 p

(196.1mN) and the largest two to a load of 50 p (490.3mN), that is more than an order of

magnitude higher than the ultramicrohardness load. Even though the presence of an ISE effect

was never excluded as possibility, the extent of the size dependence was rather surprising. The

coarse grains of Figure 4.12-a exhibit a hardness of 160HV while the coarser grains in Fig-

ure 4.12-b are even softer with a hardness of only 120HV. These values can be trusted, since

the partitioning of alloying elements does not change. It should be noticed that the micrographs

below refer to different magnifications.

(a)HVF=160 HV (b)HVF=120 HV

Figure 4.12: Microhardness measurements in coarse ferrite grains of the same grade. In this lengthscale, the different loads have only negligible impact on the hardness values (magn. 500× and 200×respectively, Nital). Applied indentation load= 20 or/and 50 p.

As clearly shown, the presence and the density of “ready-to-flow” mobile dislocations play a

catalytic role in the rapid work hardening of ferrite during the indentation. Actually, the ISE

99

effect is coupled with the grain size effect due to the refinement of ferrite. If an indent into

a ferrite grain is assumed, then the possibility of “primary” work hardening -to differentiate

it from the work hardening caused by the geometrically necessary dislocations created by the

plastic strain gradients- due to the existing mobile dislocations in the volume influenced by the

indent increases as the size of the ferrite target grain decreases. It was previously mentioned

that Liedl et al. [100] have experimentally shown that the hardness of ferrite is higher close to

ferrite/martensite phase boundaries than close to ferrite/ferrite grain boundaries. However, the

measured hardness difference was not exceeding 40HV. As the microstructure becomes finer,

the number/amount of grain and phase boundaries influencing the measurement increases and,

hence, a higher microhardness is measured. On the contrary, the absence of grain boundaries in

the vicinity of indent in coarse-grained microstructures leads to measurements that approach the

hardness of pure ferrite. Additionally, a comparison between the example of Figure 4.12-a and

the measurements in fine-grained dual-phase microstructures of the same martensite fractions

verifies the assumption that the finer the distribution (size and topology) of martensite grains

the more pronounced is the impact of martensitic transformation.

The critical question that arises from the above discussion concerns the ferrite hardness value

that somebody could or should use in the simplified rule of mixtures. Apparently, the in-situ

measured HVF,UMH of 240HV has not produced satisfactory results. The first impression was

that the produced deviations were due to the very hard martensite grains. Considering that

a hardness measurement reflects a flow behavior of the material, then the properties of ferrite

-which is the matrix phase- should be more decisive. On the other hand, if a HVF,MH of 160HV

is used by keeping the martensite hardness constant at its lowest value of 395HV, then this

particular rule of mixtures reproduces the linear fit of the experimental results. The reasoning

should be much easier, if no softer ferrite were measured. But since ferrite of the same dual-

phase steel can take even lower values, then why not using these values in the rule of mixtures.

Following this idea, the green curve (joining the crossed-circles) of Figure 4.11 is plotted. In

this case, the predicted hardness values of the dual-phase grade are underestimated, though this

hardness difference could be partly interpreted for the ultrafine microstructures as additional

strengthening attributed to grain boundary effects (analogous to the observations of [126]). It

is representatively displayed that setting up a model to describe the microhardness of a dual-

phase steel can be quite difficult, even for the most simple approach of a linear rule of mixtures.

In order to avoid any misinterpretations, the response of dual-phase steels in microhardness and

ultramicrohardness measurements should be holistically examined within the framework of all

their special microstructural characteristics.

100

Hole expansion behavior

A step forward in the direction of understanding the deformation behavior of dual-phase steels

under more complex loading conditions was attempted by means of hole expansion measure-

ments. The results were presented in section 3.2.3. Unfortunately, the limited availability of

pre-processed material did not allow the determination of the hole expansion property for all

the grades involved in the study.

Based on the results, a couple of general observations can be made irrespective of the grades

investigated. The hole expansion ratio was found to decrease with increasing cooling rate

(Figure 3.40). The effect is more pronounced for the specimens annealed in the austenitic

region, probably because their martensite fraction is higher for the same cooling rate. The

higher the martensite fraction of the dual-phase steel the higher is the number of micro-cracks

formed at the surface of the hole during punching (hole-surface damage) and the higher the

strain hardening of the edges of the hole due to the shear deformation. Furthermore, the

presence of a hard second phase in the vicinity of these micro-cracks causes a significant stress

concentration and accelerates the crack propagation during the hole expansion test [127]. No

conclusions could be drawn about the influence of grain refinement and/or of the martensite

morphology on hole expansion, mainly because of the interdependency of these microstructural

parameters and the lack of a complete series of data/experiments.

If the difference in hardness between the phases is somehow reduced, then the surface damage

of the hole edge becomes less pronounced. Of course, this effect does not refer to the present

investigations, where the decrease in the hardness ratio k occurs due to the reduced carbon con-

tent of the microstructures with a higher martensite fraction. Several authors have studied the

effect of tempering on the hole expansion properties [77, 78, 127, 128]; by softening martensite,

the ferrite/martensite hardness ratio was drastically reduced. The results were encouraging,

however, in most cases an increase of the hole expansion ratio η was accompanied by the return

of discontinuous yielding. Therefore, the application of tempering techniques strongly depends

on the desired properties of the final material. Another way/method proposed for the improve-

ment of hole expansion property was the substitution of one part of martensite with bainite or

retained austenite [75, 76, 127, 129], but such a task was beyond the scope of the present study.

The electron microscopy investigations on the fracture surfaces of the cracks formed during the

hole expansion test provided valuable information about the mechanisms of fracture. However,

since no quantitative fractography was performed, the differentiations in the observed fracture

modes between the different grades and the various annealing conditions can be used only for

a rough interpretation of the results. Since the hole expansion testing methods are currently

under development, and the available up-to-date literature sources are limited, the clarification

of the deformation and fracture mechanisms involved in the test demands further experimental

work.

101

Chapter 5

Summary

In the first section of this work alternative ways of thermomechanical processing were studied,

aiming at the grain refinement of a dual-phase ferritic-martensitic steel by using conventional

industrial facilities/equipment. The main concept involved severe laboratory cold-rolling of

pre-processed microstructures instead of the classical ferritic-pearlitic steel sheet followed by

appropriate annealing. Various classical dual-phase heat treatments (though without an over-

aging stage) were applied in a dilatometer, until the desired microstructural characteristics

were obtained for each laboratory grade. Non-conventional annealing cycles were additionally

performed, aiming to promote the nucleation of new grains and -at the same time- to impede

excessive grain growth.

Based on the results of dilatometric investigations, laboratory annealing simulations were

conducted in order to reproduce the ultrafine dual-phase microstructures on a bigger scale and

to determine their mechanical performance. Although the results of the novel non-conventional

heat treatments were satisfactory, the production of specimens at a bigger pilot scale was not

feasible mainly due to insufficient control of the rapid temperature changes between the anneal-

ing segments with the existent industrial equipment. Therefore, only conventional annealing

cycles were employed, with the objective of studying the influence of the annealing parameters

(annealing temperature and cooling rate) on the microstructure and eventually on the mechan-

ical properties. A wide variety of dual- and multi-phase microstructures was realized, regarding

the phase fractions and the grain sizes. Evidently, the unique response of the pre-processed and

industrial grades to the same heat treatments produced different microstructures. According

to the quantitative microstructural investigations, dual-phase steels with a mean ferrite grain

size between 1.5µm and 3µm and martensite fractions between 25% and 55% were obtained.

The mechanical properties of the grades were determined by means of tensile and hole ex-

pansion tests. Ultimate tensile strengths in the order of 800-850MPa were accompanied by

remarkably high uniform and total elongations, only slightly affected by the martensite frac-

tion. Moreover, yield to tensile strength ratios between 0.4 and 0.5 and very high strain hard-

102

ening rates (n2−4 -values> 0.25) were achievable. Of particular interest were the properties of

the grades DP I and DP II, providing a wide assortment of dual-phase steels exhibiting a yield

strength between 270 and 370MPa and an ultimate tensile strength between 660 and 810MPa,

while maintaining the same level of high ductility. Though grain refinement proved to have

beneficial effects on the tensile properties of the steels, no safe conclusions on its impact on

hole expansion behavior could be drawn.

The corresponding as-annealed and as-tested microstructures as well as the fracture surfaces

and the fracture profiles were characterized by light and scanning electron microscopy. Inten-

sive studies on the deformation behavior of the steels, focused on the neighborhood of fracture,

helped to identify the different failure mechanisms between the grades. Void formation was

more pronounced in the specimens of low martensitic fraction, consisting of hard martensite

particles embedded in the soft ferritic matrix. The sites where void nucleation and crack initi-

ation occur were usually found either at the ferrite-martensite interface or between martensite

grains. On the other hand, the extraordinary high ductility of the dual-phase steels with high

martensite fractions was associated with severe deformation of the martensite grains in the

location of fracture (necking area). The observations were supported by microhardness and

ultramicrohardness measurements, which revealed the changes in the hardness of martensite

grains (softening) depending on its carbon content and its morphology.

The yield and the strain hardening behavior of the dual-phase steels were extensively discussed

on the basis of the unique material characteristics. The main reason for the continuous yielding

behavior as well as for the high work hardening rate of dual-phase steels is the existence of a

large number of geometrically necessary dislocations, introduced into the ferritic matrix due

to the volume expansion and the shear deformation accompanying the austenite to martensite

transformation during cooling. The movement of these mobile dislocations during the initial

stages of deformation results in the elimination of yield point elongation and in the observed

low yield strength. The interaction of the dislocations with each other and with the finely

dispersed martensite grains results in the high strain hardening exponent.

A Hall-Petch approach to the grain refinement impact on the yield strength was attempted,

assuming a linear relationship between the yield strength and the reciprocal square root of the

mean ferrite grain size. The results of the linear regression were adequately interpreted while

the validity of the assumptions was cross-checked with other published works.

103

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Lebenslauf

Personliche Daten

Name, Vorname: TSIPOURIDIS PRODROMOS

Geburtsdatum, Geburtsort: 23.03.1978, Drama (Griechenland)

Staatsangehorigkeit: Griechisch

Adresse: Implerstrasse 4, 81371, Munchen (Deutschland)

E-Mail Adresse: tsipouridis@wkm.mw.tum.de, makis.tsipouridis@googlemail.com

Ausbildung

Seit 2002: Wiss.Mitarbeiter am Lehrstuhl fur Werkstoffkunde und Werk-

stoffmechanik, Fakultat fur Maschinenwesen, TU-Munchen (Pr. Dr.mont.

E. Werner)

• Forschungsgebiet I: Dual- und Multi-Phasen Stahle, Promotion zum

Thema “Mechanical properties of Dual-Phase steels” (June 2006),

• Forschungsgebiet II (Aktuell): “European RFCS-Program: DP-grades

with improved formability”

2001: Wiss.Mitarbeiter am Lehrstuhl fur Technische Chemie II, Fakultat fur

Chemie, TU-Munchen (Pr. Dr. J. Lercher)

2001: Wiss.Mitarbeiter am Lehrstuhl fur Physikalische Chemie, Fakultat fur

Chemie-Ingenieurwesen, Aristoteles-Universitat von Saloniki (GR)

1995-2000: Diplomstudium: Fakultat fur Chemie-Ingenieurwesen, Aristoteles-

Universitat von Saloniki (GR), November 2000

Diplomarbeit zum Thema: “Electrocatalytic Dehydrogenation of SiH4-CH4

mixture in a Proton Conducting Solid Electrolyte Cell”

Stipendien

1998: Stipendium von I.K.Y. (State Scholarships Foundation, GR)

Sprachkenntnisse

Griechisch (Muttersprache), Englisch (fließend), Deutsch (gute Kenntnisse)

Muenchen, October 30, 2008