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8/11/2019 Mechanical properties of Dual-Phase steels.pdf
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Lehrstuhl fur Werkstoffkunde und Werkstoffmechanik mit
Materialprufamt fur den Maschinenbau
Technische Universitat Munchen
Mechanical properties of Dual-Phase steels
Prodromos Tsipouridis
Vollstandiger Abdruck der von der Fakultat fur Maschinenwesen
der Technischen Universitat Munchen
zur Erlangung des akademischen Grades eines
Doktor-Ingenieurs (Dr.-Ing.)
genehmigten Dissertation.
Vorsitzender: Univ.-Prof. Dr.-Ing. Horst Baier
Prufer der Dissertation:
1. Univ.-Prof. Dr. mont. habil. Ewald Werner
2. Hon.-Prof. Dr.-Ing, Dr. Eng. (Japan) Hans-Harald Bolt
Die Dissertation wurde am 14.03.2006 bei der Technischen Universitat Munchen
eingereicht und durch die Fakultat fur Maschinenwesen
am 19.06.2006 angenommen.
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Acknowledgements
This study was carried out during my employment as research assistant at the Institute for
Materials Science and Mechanics of Materials of TU-Munich (Lehrstuhl fur Werkstoffkunde
und Werkstoffmechanik).
I am deeply indebted to my direct advisor Prof. Dr. mont. habil. E. A. Werner for his unfailing
support all these years, for his encouragement to proceed with new ideas and for being daily
available for uncountable scientific discussions. Thank you very much!
My special thanks to my co-advisor Prof. Dr.-Ing., Dr. Eng. H.-H. Bolt (Head of group Mate-rials Synthesis and Materials Characterization of Max-Planck-Institut fur Plasmaphysik, IPP
Garching), as well as to the chairman of my PhD examination, Prof. Dr.-Ing. H. Baier (Head
of the Institute for Lightweight Structures, TU-Munich).
My sincere thanks to the project partner voestalpine Stahl Linz GmbH for supplying the
testing material and making possible to conduct the annealing simulations as well as the tensile
and hole expansion tests. Special thanks to Dr. A. Pichler for the support and to Dipl.-Ing.
E. Tragl for the close and continuous collaboration.
I would also like to thank the Christian Doppler Research Association (CDG) for sponsoring
and supporting this study during the years 2002-2005 (as a project/module of the Christian-
Doppler-Laboratory for Modern Multiphase Steels).
I should not forget to thank Dr. G. Triantafyllides (Dep. of Chemical Engineering of Aristotle
University of Thessaloniki) for motivating me to choose the field of steel research.
Last, but not least, i would like to thank my colleagues and the technical staff of the Chair,
without whose help this work would be never completed. But most of all, thank you forcreating a pleasant and friendly working environment and for helping me to become an active
member of our group.
Munich, June 2006 Prodromos Tsipouridis
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Contents
1 Introduction 1
1.1 Low-alloyed dual-phase steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5
1.2 Grain refinement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11
2 Material production and experimental procedure 152.1 Production of the material . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
2.2 Thermodynamical calculations . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16
2.3 Pre-processing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16
2.4 Cold-Rolling trials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17
2.5 Dilatometric investigations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18
2.6 Annealing simulations and austenitization kinetics . . . . . . . . . . . . . . . . . 19
2.7 Microstructure investigations. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20
2.8 Mechanical testing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20
2.8.1 Tensile testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20
2.8.2 Ultramicrohardness and microhardness testing . . . . . . . . . . . . . . . 21
2.8.3 Hole expansion measurements (Stretch flangeability) . . . . . . . . . . . 21
3 Results 23
3.1 Microstructure investigations. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23
3.1.1 Recrystallization/austenitization kinetics . . . . . . . . . . . . . . . . . . 23
3.1.2 Dilatometric investigations . . . . . . . . . . . . . . . . . . . . . . . . . . 26
3.1.3 Annealing simulations . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
3.1.4 Quantitative analysis- Grain size measurements . . . . . . . . . . . . . . 443.2 Mechanical properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
3.2.1 Tensile testing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
3.2.2 Hardness measurements . . . . . . . . . . . . . . . . . . . . . . . . . . . 65
3.2.3 Hole expansion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 71
I
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4 Discussion 77
4.1 Grain refinement and microstructure investigations . . . . . . . . . . . . . . . . 77
4.2 Mechanical properties. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80
5 Summary 102
II
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Chapter 1
Introduction
The increasing automobile market demands for reduced fuel consumption as well as the need
to comply with the international environmental regulations regarding greenhouse gas emissions(GHG), resource reduction and recyclability have motivated and/or even forced the automotive
industry to produce more fuel-efficient vehicles by reducing their weight. In order to provide a
steel-based structural platform that fulfills the auto-makers requirements and takes advantage
of the new high strength steels, a new vehicle architecture based on novel design concepts has
been developed. The application of advanced high strength sheet steels exhibiting both high
strength and excellent formability offers the unique option of combining weight reduction (by
using thinner gauges of sheet material) with improved passive safety, optimized environmental
performance and manufacturing feasibility at affordable cost.
The most common steels used presently in the automobile industry are mild steels, whichare low-carbon steels characterized by a yield strength level of 140 MPa and an excellent deep
draw ability. Despite their forming and cost advantages over high strength steels, the ultimate
strength level of mild steels remains at relatively low levels, so that the crash performance is
mainly dependent on the sheet thickness. By consistent controlling of the alloy chemistry, con-
sidering the presence of interstitial carbon in ferrite, interstitial-free grades (IF) possessing an
ultra-low carbon content were produced. Traditional strengthening mechanisms such as solid
solution hardening (with phosphorous to be the most common hardening element), precipita-
tion hardening and grain refinement by carbides and/or nitrides were employed to increase the
strength of IF-steels, while maintaining their excellent formability. Micro-alloying with vana-
dium, niobium or titanium accompanied with fine carbide precipitation and grain refinement
leads to even higher strength levels and increases the ratio of yield to tensile strength. Bake-
hardening steels (BH) offer a combination of good formability during stamping and provide
an increased yield strength after the paint-baking process. To take advantage of the bake-
hardening effect a certain content of solute carbon and an appropriate aging heat treatment
are required, aiming at the supersaturation of carbon in the ferrite [1]. Due to their sharp
1
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upper yield point, BH-steels exhibit a high dent resistance, which makes them good candidate
materials for outer body panel applications. For the steel grades mentioned above, widely
designated as conventional steels, a reduced formability is an unavoidable consequence when
selecting steels with higher strength levels.
Figure 1.1: Strength-Formability relationship of thin sheet steels
To overcome this problem new types of high strength steels, the so-called Advanced High
Strength Steels (AHSS), were developed from the sheet steel suppliers in cooperation with the
automakers and design engineers. These grades exhibit higher rates of work hardening than
conventional steels as a result of their low yield strength to tensile strength ratio, show good
press formability and reach higher ultimate tensile strengths, therefore they have the potential
for significantly improved crash performance [2]. AHSS steels are multiphase steels consist
of hard islands of martensite, bainite and/or retained austenite dispersed in a ductile ferritic
matrix, in quantities and combinations sufficient to produce desired mechanical properties.
The multiphase AHSS family includes Dual-Phase (DP, ferritic-martensitic), TRansformation
Induced Plasticity (TRIP) and complex multiphase steels. Ferritic-bainitic steels, also known as
stretch-flangeable (SF) steels, are considered as a subgrade of the DP products. The mechanical
properties of conventional and AHSS thin sheet steels with respect to ductility and ultimate
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tensile strength are shown in Figure 1.1. As can be observed, a partial overlap between the
strength levels of different steel grades is possible.
According to the results of the ULSAB-AVC Program (Ultra Light Steel Auto Body-Advanced
Vehicle Concepts), an automotive body could be constructed by utilizing approximately 85 %
of AHSS, achieving a weight reduction of 25 % compared with a bench-marked average
base model and without any increase of the manufacturing costs. Figure 1.2shows that the
clear majority of autobody components is designed using dual-phase steels [3, 4]. Different
criteria such as formability, weldability, spring-back behavior and of course static and dynamic
properties play a significant role in the material selection, even though for some parts more
than one steel grades fulfill the mechanical property standards and hence are also applicable.
In particular, for a number of components that both DP and TRIP steels are viable candidates,
cost-effective DP grades were preferable.
Figure 1.2: ULSAB autobody structure steel grade distribution
Due to their special microstructural characteristics, ferritic-martensitic dual-phase (DP) steels
provide an attractive combination of strength and ductility and furthermore exhibit continuous
yielding behavior accompanied with a high work hardening rate. In order to meet the
different design requirements of individual components, various DP grades regarding tensile
strength and formability are produced industrially. This variation of mechanical properties
is mainly achieved by controlling the carbon content of the steel. Addition of other alloying
elements such as manganese, chromium, vanadium, molybdenum and nickel, individually or in
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combination, can also increase hardenability. Another method to increase the tensile strength
of dual-phase steels is to increase the martensite fraction by applying appropriate annealing
schedules, even though this procedure is followed by an expected loss in formability. In the
present study, grain refinement is proposed as an alternative way/solution to improve the me-
chanical properties of a low-alloyed dual-phase steel. To achieve this, severe plastic deformation
was applied to a pre-processed dual-phase steel by means of conventional cold-rolling, followed
by an appropriate final heat treatment with the aim to produce a homogeneous fine-grained
dual-phase ferritic-martensitic microstructure. The impact of grain refinement on the mechani-
cal properties of the dual-phase steel, regarding tensile properties, ultramicrohardness and hole
expansion behavior, is then investigated in this work.
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1.1 Low-alloyed dual-phase steels
Dual-phase (DP) steels were developed in the mid-seventies in order to satisfy the increasing
needs of automotive industry for new high strength steels which combine significant weight
reduction and improved crash performance, while keeping the manufacturing costs at affordablelevels. The high commercial potential of the newly developed alloy has motivated extensive
research in numerous laboratories, resulting in DP-grades having a wide range of chemical
compositions and being produced with various processing routes.
Dual-phase steels are characterized by a microstructure consisting of a fine dispersion of
hard martensite particles in a continuous, soft, ductile ferrite matrix. The term dual-phase,
firstly reported by Hayami and Furukawa [5] and thereafter adopted by the steel research
community, refers to the presence of essentially two phases, ferrite and martensite, in the
microstructure, although small amounts of bainite, pearlite and retained austenite may also be
present. Irrespective of the chemical composition of the alloy, the simplest way to obtain a dual-phase ferritic-martensitic steel is intercritical annealing of a ferritic-pearlitic microstructure in
the + two-phase field, followed by a sufficiently rapid cooling to enable the austenite to
martensite transformation.
Three basic approaches exist for the commercial production of dual-phase steels: (a) the
as-hot-rolled approach, where the dual-phase microstructure is developed during the con-
ventional hot-rolling cycle by careful control of chemistry and processing conditions [6
12], (b) the continuous annealing approach, where hot- or cold-rolled steel strip is un-
coiled and annealed intercritically to produce the desired microstructure [13, 14] and (c)the
batch-annealing, where hot- or cold-rolled material is annealed in the coiled condition
[1519].
Box- or batch-annealing was mainly considered for economical and practical reasons by steel-
makers where continuous facilities were not available. Dual-phase steels could be obtained by
means of batch-annealing in the intercritical region for approximately 3 h to ensure homogene-
ity, followed by very slow cooling. Due to the extremely low cooling rates, much higher alloying
contents were necessary to achieve the desired hardenability (i.e. 2.5 % Mn, 1.5 % Si, 1.0 % Cr).
On the other hand, dual-phase steel production in the hot strip mill demands precise control of
the transformation, because the transformation starts from single phase austenite. The
determination of an accurate CCT- diagram via dilatometry, where the influences of alloying
elements, heat treatment conditions and desired properties are integrated, is of great impor-
tance for the success of the process. The critical point is the correction of the CCT- diagram
to include the presence of strain in austenite (to simulate the real process conditions).
The use of continuous annealing lines (CAL) offers the advantages of high production rates,
better uniformity of the steel properties and, furthermore, the possibility of using lower alloyed
steels (having a lean chemistry). In continuous annealing lines three types of cooling are utilized
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[20]: (a) water-quenching, (b) gas-jet cooling and (c) air-cooling. The use of CAL equipped
with water quenching facilities makes possible an easy and economical production of dual-
phase high-strength steel. The basic heat cycle involves intercritical annealing and subsequent
water quenching to form the ferritic-martensitic microstructure. If necessary (according to
application), a tempering stage follows [13,2123].
From the above short review it becomes clear, that dual-phase steels can be produced by
cooling from the annealing temperature (intercritical or austenitic) by any cooling rate in the
range between batch-annealing and water-quenching. Since every steel producer has different
melting, rolling and cooling facilities, the choice of the alloying elements best suited to the
existing production capabilities is mandatory. Thus, a single widely accepted alloy composition
for each grade of the dual-phase steel family is out of consideration. There are many combi-
nations of alloying elements such as Mn, Si, Cr, Mo and V that can be added to low carbon
(0.1 wt. % C) iron to obtain the desired ferritic-martensitic microstructure. A basis with less
than 0.15 wt. % carbon (for weldability reasons) and 1.5 wt. % Mn is generally acceptable. To
achieve good ductility and toughness, an initial carbon content of about 0.1 wt. % C is ideal, so
that the carbon content and the morphology of martensite can be controlled. For strong and
tough martensite this value is approx. 0.4 wt. % C, depending on the total alloy composition.
For martensite carbon contents higher than 0.4 wt. % C twinned martensite may be formed.
Concerning the other alloying elements, manganese (Mn), chromium (Cr), molybdenum (Mo),
silicon (Si) and vanadium (V) are commonly used to increase the hardenability of austenite
[11,14,19,2428]. The role of Si and Mn is more complex, since they may also contribute to
solid solution strengthening of ferrite and hence to the strength level of the steel. The additionof elements such as Cr, Mo and V that promote carbide formation demands a careful process
control.
For each applied alloy chemistry there exists a critical value of the cooling rate which sets the
lower limit for the formation of a ferritic-martensitic microstructure. Applying cooling rates
lower than this value results in a ferritic-pearlitic microstructure while higher cooling rates are
capable of producing a microstructure consisting of martensite islands in a ferritic matrix. The
overcooling degree is responsible for the martensite fraction in the final product. Equivalently,
the critical cooling rate (which as previously shown is prompted/necessitated by the production
line capabilities) defines a minimum of alloying content for the dual-phase steels.Based on experimental results and accepting a similar alloying behavior of Cr and Mo as Mn
(mainly due to the hardenability effects), Tanaka et al. [14] expressed the influence of alloying
elements on the critical cooling rate in terms of a manganese equivalent, Mneq[%]:
logCR [K/s] = 1.73 (Mn)eq[%] + 3.95 (1.1)
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where
(Mn)eq= (Mn) + 1.3 (Cr) + 2.67(Mo)[%]. (1.2)
Though alloying elements may look versatile via equation eq. 1.1, their effects on microstructureand deformation behavior are not necessarily the same. It is clearly shown that the higher the
amount of alloying elements in the steel the lower is the critical cooling rate necessary to form
a dual-phase microstructure. This explains the reason why batch-annealing processes require
much higher alloying additions. Practically, the critical cooling rate is easily determined with
a series of simple cooling experiments for each individual alloy.
To understand the formation of the ferritic-martensitic microstructure and to be able to in-
terpret the products of the heat treatments it is essential to consider the phase transformations
taking place during heating, intercritical annealing and quenching. The formation of austen-
ite in low carbon C-Mn steels was studied by a number of investigators [ 2933]. In the case
of cold-rolled ferritic-pearlitic steels, the recrystallization of the cold-worked microstructure is
completed already before reaching the annealing temperature, even during the rapid heating
rates applied on most continuous annealing lines. By entering the intercritical two-phase re-
gion, austenite nucleates rapidly at pearlite or in the vicinity of cementite particles and grows
rapidly until the carbides are dissolved. A slower growth of austenite into ferrite is continued
at a rate initially controlled by carbon diffusion in austenite and finally by manganese diffusion
in austenite, until the system reaches the equilibrium state. Practically, due to the very short
annealing times, only carbon redistribution between the phases takes place, because substitu-
tional manganese diffuses much more slowly than interstitial carbon. This effect is referred to
as paraequilibrium.
Assuming that no manganese redistribution occurs, then a vertical section corresponding to
paraequilibrium conditions can be constructed for constant Mn content. Bearing that remark
in mind, Figure 1.3 demonstrates the production concept/scheme of a dual-phase steel by
intercritical annealing. According to the lever rule, for any given carbon content (C0) the
amount of austenite increases with increasing the intercritical temperature, becoming equal
to 100% at the Ac3 -temperature while, as a consequence, the carbon content of austenite
decreases, reaching its minimum value (C
=C0). In an analogous manner, for any given inter-critical annealing temperature the amount of austenite increases with increasing alloy carbon
content, becoming equal to 100 % at a carbon content corresponding to the/+ boundary
(C0=C). Additional alloying elements may cause some changes in the austenite formation
process and/or even widen or tighten the intercritical temperature field. Such alloying effects
can be roughly estimated by thermodynamical codes (e.g. ThermoCalc) but remain beyond
the scope of this work.
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Figure 1.3: Production of dual-phase steels by intercritical annealing. The equilibrium fractionsof austenite and ferrite as well as their carbon content at the annealing temperature can be easilyestimated by applying the lever rule.
The importance of intercritical annealing becomes apparent, since for a given alloy compo-
sition and for a pre-selected annealing temperature the maximum amount of austenite that
can be (ideally) transformed to martensite as well as its carbon content -i.e. its hardenability-
are already pre-determined. Therefore, high intercritical annealing temperatures result in high
austenite fractions of decreased hardenability while low annealing temperatures result in low
austenite fractions with increased hardenability.
Even though the products of the austenitic transformation are strongly dependent on the
intercritical annealing parameters, the cooling rate is the final decisive step for the production
of dual-phase steels. The combined influence of cooling rate and intercritical annealing on
the formed microstructures was studied individually by many investigators [26,28,3436]. In
each case, a very important parameter that should not be forgotten is the effect of alloying
elements on the stability of austenite, which can be qualitatively measured by means of the
martensite-start temperature (MS). For low carbon steels it was proposed by Andrews [37]
that:
MS [C] = 539 423C 30.4 Mn 17.7 Ni 12.1 Cr 7.5 Mo (1.3)
or similarly by Eldis [38] for dual-phase steel compositions (also in wt. %):
MS [C] = 531 391.2 C 43.3 Mn 21.8 Ni 16.2 Cr. (1.4)
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Quenching with high cooling rates from low intercritical temperatures ensures a complete
transformation of austenite into martensite because of the austenite stability; by applying
relatively low cooling rates the intercritically formed austenite partly transforms into ferrite,
enriching further the remaining austenite with carbon, increasing its hardenability and lower
even more the MS-temperature. In this case, isolated retained austenite particles may be
detected in the microstructure in the form of interlath films. However, the formation of
ferrite-carbide aggregates is usually unavoidable. Rapid quenching from high intercritical
temperatures produces even more complex effects: the decreased austenite hardenability
and the absence of time for the necessary diffusions during cooling, both reflected in a high
MS-temperature, result in the formation of autotempered martensite.
Dual-phase steels exhibit a number of superior mechanical properties, such as continuous
yielding behavior, low 0.2 % offset yield strength, high work hardening rate, high tensile strength
and remarkably high uniform and tensile elongations. The mechanical properties of dual-phase
steels arise from structural features, that is the fine dispersion of hard martensite particles
in a ductile ferritic matrix and all the related phenomena that accompany this coexistence.
The yielding and the work hardening behavior have been interpreted in terms of the high
dislocation densities and residual stresses arising in ferrite, as a consequence of the volume
expansion associated with the austenite to martensite transformation. The strength of dual-
phase steels was found to be dependent primarily upon the volume fraction and the carbon
content of martensite; solid solution strengthening of ferrite may also contribute to strength.
The excellent ductility reported for most of the dual-phase steels is the combined result ofmany factors. Among them are the volume fraction and the carbon content of martensite, the
ductility of martensite, topological parameters such as the martensite grain distribution in the
ferritic matrix, the alloy content of ferrite, the dislocation density in ferrite, the presence of
carbides and/or retained austenite. Additionally, lattice imaging from Koo and Thomas [24]
has revealed a good coherency at the ferrite/martensite interface, which prevents decohesive
interface failure during deformation and thus enables the full toughness of ferrite to be realized.
Tempering may be applied as part of the process in some continuous annealing lines, after
water-quenching from the intercritical temperature, to regulate the properties of the dual-phase
steel. It may also be an unavoidable side-effect of an operation, e.g. in a hot dip galvanizingline. Finally, tempering may be useful in some production practices such as bake hardening
(paint baking). The change in yield strength upon tempering is complex because of the relief
of residual stresses, carbon segregation to dislocations and the return of discontinuous yielding.
After tempering at low temperatures the yield strength increases but discontinuous yielding
returns only to the steels containing low volume fractions of martensite. When tempering at
high temperatures the yield strength decreases but discontinuous yielding appears in all steels.
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The tensile strength decreases, while post-uniform and uniform elongations increase due to the
change in hardness of martensite.
There has been a long discussion on the construction of a model correlating the mechanical
properties of dual-phase steels with their microstructural characteristics. Simple empirical rules
of mixtures [2628,3944] (usually linear regressions between the constituent phases proper-
ties, based on Mileikos theory of composites [45]) as well as more complicated/sophisticated
micromechanical models [4651] (introducing the importance of the effective in-situ prop-
erties and topological microstructural parameters) have been developed, most (if not all)
of them based on individual sets of experimental data. In some cases, there was a good
agreement between the predicted properties and the experimental results -even for martensite
fractions in the order of 80 %, in other cases some fair and well-established modifications have
to be made. Since each study refers to a specific alloy composition and a different kind of
heat treatment, a comparison between the obtained results contributes to the disagreements
reflected in literature over the past 25 years.
In order to meet the different design requirements of individual automobile-body components
for strength, crashworthiness, energy absorption, part complexity and dent resistance, a variety
of dual-phase grades exhibiting different strength and ductility levels is currently industrially
produced. Despite the numerous studies on the relationship between the mechanical properties
and the microstructural characteristics of dual-phase steels over the last decades, the chal-
lenge of increasing their formability at a constant strength level (or equivalently increasing the
strength while maintaining a high ductility) remains still unanswered.The statement of many researchers that for the improvement of properties of dual-phase
steels the ferrite should be fine-grained, free of ultrafine carbide precipitates and strength-
ened solely by alloying additions which have minimum effects on ductility is still not fully
explained/affirmed.
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1.2 Grain refinement
The formation of ultrafine-grained ferritic microstructures in low-alloyed and plain carbon steels
has been intensively investigated during the last decade, since grain refinement is expected to
have a beneficial effect on the mechanical properties of the steel. Though the term ultrafineis somewhat vague, it reflects the objective of the Japanese super metal project of producing
a strip having a ferrite grain size of 1 m throughout a minimum thickness of 1 mm [52].
According to the Hall-Petch relation (which is applicable to a variety of polycrystalline single-
and dual-phase metals), a decrease in grain size (d) results in an increase in yield strength (y):
y = 0+ kyd1/2. (1.5)
For a low-alloyed steel, for example, a decrease of the ferrite grain size from 5 m to 1 m wouldideally produce an increase of the yield strength by approximately 300 MPa. Additionally, it
has been reported that grain refinement improves the fatigue resistance of steels [ 53] and can
lead to superplastic behavior at high temperatures and appropriate strain rates [54].
The currently available techniques to obtain ultrafine grains are rapid solidification directly
from the melt, vapor deposition, cryogenic metal-forming, mechanical alloying and severe
plastic straining. Very small grains (with sizes in the nano-scale range) may be produced under
extreme conditions, often leading to impressive physical and mechanical properties. However,
due to the the limited production quantities and the very small grains individually formed
(powder-like), further consolidation and processing is required to produce a bulk material suit-able for structural use. Concerning the refinement methods, there exists a permanent conflict
between the achievable grain size in a material, the amount or the dimensions of the material
that can be processed in this way and -the most important- the cost of processing. In the
case of steel applications, an optimum compromise between these factors would be required [ 55].
Grain refinement in steels can be realized by microalloying. The addition of elements such as
Al, B, P, Sm or Ti in the microstructure can suppress the grain growth of ferrite, utilizing the
pinning effect of secondary phase particles and/or the dragging effect induced by solute atoms
[56]. However, the idea of controlling the grain size and hence the mechanical properties by
thermomechanical processing instead of the classical way of alloying is far more attractive, as
this would result in the production of steels with simpler chemistries and improved recyclability
and would lead to economic benefits as well.
To achieve steel grain refinement, substantial efforts involving severe plastic deformation
(SPD) by using conventional rolling equipment have been made. According to the dynamic
Strain Induced TRansformation method (SITR), proposed by [5761], grain refinement is pro-
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moted by the continuous intragranular nucleation of numerous small ferrite grains during hot-
or warm-rolling of austenite. The process involves a single-pass rolling of the steel strip at
a temperature just above the Ar3 (i.e. immediately above the temperature at which grain
boundary pro-eutectoid ferrite would begin to form) but below the Ae3, followed by air-cooling
or accelerated quenching depending on the desired microstructure. The required rolling reduc-
tion ranges between 35 and 40 %. Ultrafine ferrite (UFF) grains withdF
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(a) HPT (b) ECAP (c) ARB
Figure 1.4: Schematic illustration of the recently developed methods for severe plastic deformation(HPT and ECAP from [55], ARB from [66]).
nique for obtaining ultrafine-grained microstructures [6871]. According to the method, the
material is pressed through a die, where two channels form an L-shaped configuration with an
angle of 2 (Figure 1.4-b). The process imposes a severe strain on the sample by means of
shearing. Since there is no concomitant change in the cross-sectional dimensions of the samples,
repetitive pressings can produce very high effective strains (in the order of 10), achieving ho-
mogeneous grain refinement with grain sizes on the micro-meter scale. Moreover, the sample is
constrained so also less ductile materials can be processed. The factors influencing the method
are the pressing route by which the sample is rotated during the successive pressings, the dieangle which determines not only the strain per pass but also the geometry of deformation,
the die cross-section geometry, the speed and the temperature associated with pressing. Very
recent investigations revealed that ultrafine-grained ferritic-martensitic dual-phase steels can
be fabricated by ECAP, by applying an effective strain of 4 at 500 and subsequent intercrit-
ical annealing at 730 for 10 min. UFF grains with uniformly distributed blocky martensite
islands of 1 m were produced [72]. Despite the scientific interest and the partial success, the
application of the ECAP method is still restricted in terms of commercial capability while even
the future prospects for steel sheet production are questionable. The potential of up-scaling
ECAP is being currently investigated.Accumulative Roll Bonding (ARB) is also a newly developed technique to realize intense
plastic straining [66, 73, 74]. Following the illustration of Figure 1.4-c, one strip is neatly
placed on top of another strip and the two layers of material are joined together by warm
rolling. The rolled sample is then cut in two halves, which are again stacked together after
an appropriate surface preparation and are roll-bonded. The whole process can be repeated
without any change in the samples thickness and geometry, given that the reduction in thickness
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after each rolling pass is maintained/controlled to 50 %. This process can theoretically introduce
unlimited amounts of strain in the material. Effective strains of 8 have been reached for
aluminium and steels, resulting in grain sizes below 1 m. Critical factors for the success of the
method are surface preparation and cleaning, the deformation temperature and the amount of
induced strain. Although rolling at elevated temperatures is advantageous for joinability, too
high temperature would cause dynamic recrystallization and cancel the effect of accumulated
strain.
All the above proposed techniques are very attractive but seem to encounter several difficulties
in engineering applications. In each case, very high levels of strain or unrealistic inter-pass
times are required, complex processes and special equipment are involved and the production
capacities are still very low to cover the steel market demands. As a combined result of these
factors, the application of the novel refining methods in a continuous industrial process does
not seem possible in the near future.
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Chapter 2
Material production and experimental
procedure
2.1 Production of the material
The dual-phase steel investigated in this work was industrially produced by voestalpine Stahl
Linz. Its chemical composition is given in Table 2.1. Slabs with a thickness of 210 mm were
produced on a continuous casting machine. The slabs were reheated in a pusher-type furnace
to a temperature of 1250 and hot-rolled to a final thickness of 2.40 mm. The finishing
temperature was approximately 900 while the coiling temperature was about 600. Part
of the hot-rolled strip was milled to remove the surface scale and then cold-rolled to 1.00 mm
(cold-reduction of 58 %).
Table 2.1: Chemical composition of the investigated DP-steel.
DP steel C Si Mn Cr+Mo Fe
wt. % 0.1 0.15 1.5 0.8 Bal.
In the main bulk of this work strips from the hot-rolled ferritic-pearlitic material were used,
cut to specimens of 250 mm in length and 50 mm in width in order to satisfy the requirements
of the laboratory annealing and cold-rolling equipment. The industrially cold-rolled strip was
mainly used as reference material (denoted as R in the following) with regard to its response
to the same heat treatment schedules applied to the pre-processed and laboratory cold-rolled
material.
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2.2 Thermodynamical calculations
The phase diagram of the alloy as well as the critical temperatures where phase transformations
occur were initially calculated with the help of the program ThermoCalc, by using the database
TCFE3 and considering ferrite, austenite and cementite as the only phases present in the steel.The intercritical + two-phase temperature range was calculated between 710 and 815.
Additionally, the fraction of austenite at different intercritical annealing temperature steps was
determined, in a first attempt to set the maximum martensite fraction after cooling down from
the annealing temperature (Figures 2.1-a and 2.1-b). It should be taken into account that
all estimated temperatures and phase fractions represent thermodynamic equilibrium.
(a) (b)
Figure 2.1: Thermodynamic calculations: (a) concentration section through the phase diagram of
the investigated dual-phase steel (the dashed line indicates the carbon content) and (b) equilibrium
fraction of austenite during intercritical annealing.
2.3 Pre-processing
The first critical stage of the experimental procedure, described with the term pre-processing,
involves annealing of the as hot-rolled material in a Carbolite Three Zone Tube furnace (TZF)
with an adapted inert gas (Ar) supply to avoid prolonged oxidation. The specimens were
annealed at three different temperatures (760, 800 and 900) with a holding time of 5 min in
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an Ar atmosphere, followed by water-quenching to form a ferritic-martensitic (in the first two
cases) or a completely martensitic microstructure (Figure 2.2). The annealing temperatures
were selected by taking into account the thermodynamical calculations and the holding time
was set long enough to assure thermodynamic equilibrium.
(a) DP I (b) DP II (c) M
Figure 2.2: Micrographs of the pre-processed materials before laboratory cold-rolling, etched withLePera.
In the following chapters, the pre-processed materials-specimens will be denoted as DP I,
DPIIandM, according to the annealing temperature during pre-processing (Table 2.2).
Table 2.2: Denotation of pre-processed material.
Grade Pre-processing conditions Martensite fraction (%)DP I Tan= 760, tan= 5 min, WQ 23 %DPII Tan= 800, tan= 5 min, WQ 40 %M Tan= 900, tan= 5 min, WQ 100 %
2.4 Cold-Rolling trials
The ferritic-martensitic or completely martensitic microstructures produced during pre-
processing were cold-rolled in a Carl-Wezel laboratory mill with a roll diameter of 220 mm
at a roll peripheral speed of 16 mmin1 to a final thickness of 1.00 or 0.80 mm (cold reduction
of 58 % or 67 %, respectively). To avoid cracking of such work hardened microstructures, rolling
was performed in multiple passes.
The microstructures of the as hot-rolled ferritic-pearlitic starting material after industrial
cold-rolling as well as that of the pre-processed material after laboratory cold-rolling are shown
in Figure 2.3. There, the microstructures differ not only qualitatively but also quantitatively
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from each other. Micrograph (a) shows a rather coarse-grained ferritic-pearlitic microstructure,
while in micrographs (b) and (c) dual-phase ferritic-martensitic microstructures are shown,
possessing different martensite fractions due to the initial annealing conditions. The higher
the martensite fraction after pre-processing the finer is the obtained microstructure after
cold-rolling. This is also supported by the last micrograph (d), where a very fine martensitic
microstructure is shown.
(a) (b)
(c) (d)
Figure 2.3: (a) Industrially cold-rolled ferritic-pearlitic microstructure of the as hot-rolled startingmaterial (R), (b) Cold-rolled ferritic-martensitic microstructure, initially annealed at 760 (DP I),
(c) Cold-rolled ferritic-martensitic microstructure, initially annealed at 800 (DP II), (d) Cold-rolledcompletely martensitic microstructure (M), etched with Nital.
2.5 Dilatometric investigations
The influence of various annealing cycles on the microstructural evolution of the cold-rolled
material (with respect to grain refinement) was studied via dilatometry, with the samples cut
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in the rolling direction (10 x 4.0 x 1.0 or 0.80 mm, depending on the sheets thickness). The
dilatometric investigations were conducted on a Bahr dilatometer DIL 805 A/D. The selection
and the appropriate combination of the annealing parameters within a cycle (heating rate, an-
nealing temperature, holding time in the intercritical zone as well as the cooling rate) were made
with regard to the special features of the cold-rolled material, that is the ferritic-martensitic or
the pure martensitic microstructure and the deformation degree. The basic annealing schemes
are presented in Figure 2.4.
(a) (b)
Figure 2.4: (a) Intercritical annealing, (b) Repeated heating cycles oscillating between the+and
phase fields.
2.6 Annealing simulations and austenitization kinetics
All annealing simulations were conducted in the laboratory with the Multi-Purpose Annealing
Simulator (MULTIPAS) at voestalpine Stahl Linz on the cold-rolled pre-processed ferritic-
martensitic and martensitic as well as on the reference material. Special attention was paid
to the intercritical annealing heat treatments due to their applicability on an industrial scale.
Specimens with a thickness of 1.00 mm were preferred for this purpose, because they offer
the possibility of direct comparison with the standard industrially cold-rolled material (R,reference).
Although the phase transformation temperatures were calculated with ThermoCalc and the
recrystallization kinetics was investigated via dilatometry, additional annealing simulations were
conducted on the reference material to determine the austenitization kinetics of the alloy in
non-equilibrium conditions on a larger scale. Specimens were annealed in a temperature range
that covers both + two-phase andsingle-phase regions for different representative holding
times and then were water-quenched.
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2.7 Microstructure investigations
A light microscope (Olympus AX70) was used for the microstructural investigations of the
heat treated samples. To reveal the microstructure of the fine-grained materials the specimens
were conventionally prepared and then etched with LePeras etchant, which is a mixture of 1 %sodium metabisulfite in distilled water and 4 % picric acid in ethanol in a 1:1 volume ratio.
This tint etching technique allows the distinction of phases by coloring, staining ferrite brown
and/or blue, bainite dark brown to black while martensite and retained austenite (hardly any
present in the materials of this work) remain white. Quantitative characteristics of the mi-
crostructure such as phase fractions and mean grain sizes of the constituents were determined
by line intercept measurements, considering ferrite as the dominating matrix phase and charac-
terizing martensite, bainite and/or retained austenite as a second phase. For severely deformed
microstructures after cold-rolling and before heat treatment or for annealed specimens where
LePeras etching agent could not produce the desirable effects, a 2 % Nital etchant (2 % nitricacid in ethanol) turned out to be an acceptable alternative.
In cases that a higher magnification analysis beyond the resolution capacity of light micro-
scope was required, e.g. for the investigation of ultrafine-grained microstructures or for the
identification of any third phase present (like bainite), Scanning Electron Microscopy (SEM)
was employed. The SEM observations were conducted on a LEO 1450 SEM and/or on a TOP-
CON SM-520 Field Emission Gun (FEG)-SEM on demand.
Microstructures of selected specimens representing critical heat treatments were further in-
vestigated by means of Transmission Electron Microscopy (TEM). TEM-specimens (thin foils)
were prepared by gentle mechanical thinning down to 80 m followed by electrolytic thinning
in 5 % perchloric acid in acetic acid. The thin foils were analyzed in a Philips CM20 STEM
transmission electron microscope, applying an accelerating voltage of 200 kV. For a first rough
overview of the microstructure, secondary electron (SE) images were taken from the thin foils.
The detailed analysis of the phases was carried out using bright and dark field techniques, while
the identification of the phases was performed using electron diffraction patterns.
2.8 Mechanical testing
2.8.1 Tensile testing
The mechanical properties of the steels were measured on a Roell-Korthaus RKM 250 testing
machine, according to European standard EN 10 002. All tensile specimens were machined
with their tensile axis parallel to the rolling direction and with a gage length of 80 mm. The
laboratory produced grades (DP I, DP II and M) were tensile tested in the as-annealed condition
while the industrially cold-rolled steel (R) was submitted to skin pass rolling of 0.5-1.0 % to
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improve roughness.
2.8.2 Ultramicrohardness and microhardness testing
The ultramicrohardness of the main constituents of the dual-phase microstructure (ferriteand martensite) was measured with a Kammrath & Weiss ultramicrohardness tester UMHT-3
equipped with a Vickers diamond square pyramid. The device was mounted in a LEO 1450
Scanning Electron Microscope (SEM). To identify the phases, the specimens were etched with
Nital etchant. The applied indentation load (15 mN) as well as other critical measuring param-
eters like indentation time and speed were selected to be the same for all measurements, so that
a comparison of the hardness between different phases and various heat treatments becomes
possible. Furthermore, to eliminate any possible errors derived from the optical measurement
of the indentation diagonals, all indentations and measurements in SEM were performed at the
same magnification (12000). Finally, the hardness was calculated from eq. 2.1:
HV=189 F
d2 , (2.1)
where d [m] stands for the mean length of the indentation diagonals and F [mN] for the
indentation load.
Additionally, the Vickers microhardness of the investigated dual-phase steels was determined
with a Micro-Duromat 4000 E microhardness tester from Reichert-Jung mounted in a light mi-
croscope (Reichert Metaplan 2, Leica AG). The indentation load was set to 150 p (approximately
1471 mN), producing Vickers impressions which are bigger by one order of magnitude compared
to ultramicrohardness measurements.
2.8.3 Hole expansion measurements (Stretch flangeability)
Hole flanging is a process widely applied in thin-sheet forming operations, which employs a
punch for producing structural parts with short necks that are subsequently used for assembly
with other components. During stretch flanging, the deformation mode at the edge of the hole
is a combination of bending and stretching which in some cases causes splitting failure and
therefore cannot be grasped by the conventional uniaxial tensile test.
The hole expansion behavior provides a way to measure the tendency of steels to split as a
hole is expanded under external forces and is characterized by the percentage increase in the
size of the hole at the moment that a crack forms [7579]. A schematic diagram of the hole
expansion equipment is shown in Figure 2.5. Selected dual-phase steel grades were cut to
125mm 125 mm square test pieces and before testing a 12 mm (0.15 mm) hole was punched
into the centre of each sample. The hole expansion test is conducted by expanding the initial
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punched hole using a 50 mm diameter punch; the punch driving is immediately stopped when
any crack (which extends all through the samples thickness) is observed at the edge of the hole.
The final diameter of the hole is measured by averaging two readings taken perpendicularly to
each other. The property is expressed as the ratio of the expanded hole size to the original hole
size, as defined by the following equation:
= df d0
d0 100, (2.2)
where [%] is the hole expansion ratio, df[mm] the average final hole diameter (after rupture),
and d0 [mm] is the initial hole diameter.
Figure 2.5: Schematic illustration of hole expansion testing equipment of voestalpine Stahl Linz
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Chapter 3
Results
3.1 Microstructure investigations
3.1.1 Recrystallization/austenitization kinetics
In order to study the austenitization/recrystallization kinetics regarding cementite dissolution
and austenite formation, specimens from the industrially cold-rolled material (R) were subjected
to heat treatments according to the annealing plan ofFigure 3.1which involves annealing not
only in the intercritical two-phase +ferritic-austenitic field but also in the pure austenitic
region. Five annealing temperatures covering a range of 100 in temperature intervals of 25
for holding times of 0 s, 10 s, and 100 s were applied, while the heating rate was set to 25 K/s
to reproduce the industrial conditions. After water-quenching the microstructure of the steelwas investigated.
Figure 3.1: Annealing plan to determine theaustenitization kinetics of the as cold-rolleddual-phase steel.
Figure 3.2: Influence of annealing tempera-ture and annealing holding time on the frac-tion of martensite.
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The fraction of martensite formed during water-quenching was determined by line intercept
measurements, providing the amount of the former austenite prior to the martensitic trans-
formation. Figure 3.2shows the influence of the annealing temperature for different holding
times on the martensite fraction. Although the martensite fraction increases dramatically in
the first 10 s of annealing, a further increase of the holding time has no significant effect except
for the annealing temperature of 775. Figure 3.3shows the microstructure of the steel (R)
as a function of annealing temperature for a holding time of 100 s. It is remarkable that for
annealing temperatures over 800 (even though the calculated + transformation
temperature is 815) the martensite fraction approaches the maximum value of 1.0.
(a)Tan= 750 (b)Tan= 775 (c)Tan= 800
(d)Tan= 825 (e)Tan= 850
Figure 3.3: Influence of the annealing temperature on the martensite fraction (former austenite
fraction before water-quenching) for a holding time of 100 s (LePera).
Figure 3.4 shows the microstructure of the steel after water-quenching from the annealingtemperature without a holding stage. For the lower annealing temperature (Figure 3.4-a), the
recrystallization is not completed and undissolved carbides are still detected in the microstruc-
ture. At higher annealing temperatures no carbides could be detected. It should be also noticed
that even after water-quenching from high annealing temperatures (825) a significant amount
of ferrite is present in the microstructure (Figure 3.4-d).
The zero holding time (0 s), which eliminates the possibility of equilibrium, provides valuable
data/information for the case that heat treatments schedules are applied involving flashing
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heating and cooling cycles (see Figure 2.4).
(a)Tan= 750 (b)Tan= 775 (c)Tan= 800
(d)Tan= 825 (e)Tan= 850
Figure 3.4: Influence of the annealing temperature on the martensite fraction (former austenitefraction before water-quenching) for a zero holding time (LePera).
The martensite fraction data are appended to the ThermoCalc diagram in order to compare
the experimental results with the numerical calculation. It was assumed that the martensite
fractions measured for a holding time of 100 s approach equilibrium status and additionally
that the amount of martensite represents the former austenite fraction during annealing. As
can be seen inFigure 3.5, the results are in good agreement with each other.
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Figure 3.5: Experimental data from martensite fraction measurements (square symbols) and equilib-rium fraction of austenite (line) as calculated with the program ThermoCalc, both given as a function
of annealing temperature.
3.1.2 Dilatometric investigations
The laboratory developed materials (DP I, DP II and M) were not submitted to the above water-
quenching process, since their recrystallization and austenitization kinetics are expected from
their design concept to be even faster. Nevertheless, systematic dilatometric investigations were
done in order to study the influence of the critical annealing parameters on the microstructureof these grades, aiming to clarify the recovery-recrystallization-grain growth mechanism and
finally to determine the optimum annealing conditions which lead to grain refinement.
Conventional annealing
Conventional heat treatment schedules following the annealing scheme of Figure 2.4-a were
applied to the laboratory cold-rolled microstructures. A relatively high heating rate of 25 K/s
from room to annealing temperature was applied, in order to simulate the industrial conditions
and additionally to minimize the time available for grain growth after recovery and recrystal-
lization. Three different annealing temperatures (Tan= 750, 800 and 840) were chosen
for the dilatometric investigations. According to the thermodynamical calculations, the first
two temperatures are located in the intercritical + while the third one in the austenite
phase field. However, the ferrite to austenite phase transformation as indicated by the dilata-
tion curves during the heating stage seems to be completed only at temperatures above 840
(see Figure 3.6). This important observation which represents the in-situ measurement of
the + transformation temperature does not actually contradict the results from the
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austenitization experiments, since in the measurement of the dilatation curve no holding time
is taken into account. This simply means that a temperature of 840 refers to intercritical
annealing conditions for a zero holding time and to austenitic annealing in case that a holding
time stage is applied. To avoid any possible misinterpretations in the discussion of the results,
Tan of 840 will be considered as austenitic annealing in the following pages.
Figure 3.6: Dilatation curve during heating. T indicates the temperature at which the ferrite to
austenite transformation is completed.
Holding times in the range between 5 s and 120 s were applied during intercritical annealing.
Recovery and recrystallization of the severely deformed microstructures were completed within
the first 30 s of holding stage for all materials investigated. The cold-rolled martensitic grade
(M) is already recrystallized at an even shorter holding time of 10 s. Microstructural observa-
tions revealed that holding times longer than 30 s are rather detrimental for the formation of an
ultrafine dual-phase microstructure, by leading to an undesirable grain growth (mainly of the
ferrite grains). Figure 3.7demonstrates the influence of holding time on the microstructure of
the grade M, intercritically annealed at 800 and quenched with 40 K/s to room temperature.
The specimens are treated with Nital etchant, which stains martensite brown or dark brown
and ferrite light brown to light cream/beige.Recrystallization is completed even for the shortest annealing time applied. The design and
the production history of the grade M as well as the high cooling rate after annealing eliminate
the possibility of presence of any carbide phases in the final microstructure. Holding time does
not seem to affect the fraction of martensite maintained in the steel after quenching but it is
proved to be a dominating parameter for the achievement of ultrafine-grained microstructures.
A holding time of 30 s provided the optimum solution with respect to grain refinement, as shown
in Figure 3.7-d. Further increase of the annealing time, for example by only 30 s, resulted to
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(a)tan= 5 s (b) tan= 10 s
(c)tan= 20 s (d) tan= 30 s
(e)tan= 50 s (f) tan= 60 s
Figure 3.7: Microstructures of the martensitic grade M after annealing at 800 for the subscribedholding times and quenching with a cooling rate of 40 K/s, etched with Nital.
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rapid grain coarsening of ferrite almost by an order of magnitude as shown in Figure 3.7-f.
Analogous observations regarding the time-dependence of grain refinement at certain annealing
conditions were also made for all laboratory grades investigated.
(a)Tan= 750 (b)Tan= 800 (c)Tan= 840
Figure 3.8: Microstructures of the martensitic grade M after annealing at different annealing tem-peratures for 30 s and quenching with a cooling rate of 40 K/s, etched with Nital.
Based on the experimental determination of an optimum holding time for grain refinement
(30 s in our case), the influence of annealing temperature on the microstructure was addition-
ally studied in the dilatometer before moving to a larger experimental scale. As shown in
Figure 3.8, annealing at higher temperatures leads to higher martensite fractions in the mi-
crostructure, consistent with the increased austenite fraction at these conditions. The difference
in martensite fraction between Figure 3.8-band Figure 3.8-c, referring to annealing temper-
atures of 800 and 840, respectively, is not clearly distinguishable, since for the holding time
applied the austenite fraction in both cases approximates the maximum. It should be noticed
that this temperature dependent increase of the martensite fraction is accompanied by a pro-
nounced alteration of the morphology and subsequently of the etching behavior of martensite
grains. Annealing at lower intercritical temperatures (e.g. at 750, Figure 3.8-a) produces
clear unstructured brown martensite, while higher annealing temperatures (Figures 3.8-b and
3.8-c) lead to the formation of dark-brown structured and generally coarser martensite (the
tints refer to the etching effect of Nital agent).
In all dilatometric investigations described above, a moderate cooling rate of 40 K/s wasapplied in order to select the most favorable conditions/parameters for grain refinement with
respect to annealing temperature and annealing time. Once this became clear, the influence
of cooling rate on the microstructural evolution was thoroughly examined. Greater attention
was paid to the grade M because of its starting microstructure peculiarity in comparison with
the other grades. In order to produce a wider possible range of dual-phase microstructures
depending on the cooling rate parameter, six cooling rates, that is 5, 10, 20, 40, 60 and 80 K/s,
were applied. Higher cooling rates were not achievable along the entire cooling stage, due
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to the restricted cooling capacity that the dilatometer -associated with the cooling gas used
(nitrogen)- could provide.
Figure 3.9 demonstrates the influence of cooling rate on the microstructure of the martensitic
grade M, after annealing at 840 for 30 s. By applying cooling rates below 10 K/s a rather
coarse multi-phase microstructure, containing ferrite, martensite and a third phase indicated
by the black tinted regions located on the ferrite grain boundaries (most probably bainite)
is formed. Grain refinement is achieved only at cooling rates higher than 40 K/s, which is
sufficient to avoid the presence of the third phase. Further increasing of the cooling rate results
also in ultrafine ferritic-martensitic microstructures with increased martensite fractions.
These observations are only qualitatively common to all grades (DP I, DP II and M). More
details about the impact of conventional heat treatments on the microstructure, with regard
to the individual annealing behavior of each laboratory grade investigated, will be described
and discussed in the section of larger scale and more comprehensive annealing simulations (see
3.1.3).
Non-conventional annealing
Non-conventional heat treatment schedules following the concept described in Figure 2.4-b
were applied to the laboratory cold-rolled material, particularly to the thinner grades (thick-
ness of 0.80 mm, corresponding to 67 % cold reduction), where recrystallization kinetics is ex-
pected to be extremely fast due to the higher cold deformation degree. The aim of these
heating-cooling cycles around the + transformation temperature was to provoke the
continuous formation of new grains, so to impede grain growth. By this repeated nucleation andpreventing the system from reaching an equilibrium state by rapidly changing the temperature,
an ultrafine-grained microstructure is finally obtained.
The selection of annealing parameters, i.e. the initial heating and final cooling rates, tem-
perature range and heating/cooling rates for the + / - cycling had to be done for each
laboratory grade separately, thereby taking into account its production history and its peculiar
microstructural characteristics. The special case in which only one cycle is applied actually
represents a limiting condition toward conventional annealing with 0 s holding time. It was
observed that more than four cycles have no further beneficial effect towards grain refinement;
on the contrary, this could even lead to grain coarsening.
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(a)CR= 5 K/s (b) CR= 10 K/s
(c)CR= 20 K/s (d) CR= 40 K/s
(e)CR= 60 K/s (f) CR= 80 K/s
Figure 3.9: Microstructures of the martensitic grade M after annealing at 840 for 30 s and quench-ing with different cooling rates, etched with Nital.
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(a) (b)
(c) (d)
Figure 3.10: Ultrafine ferritic-martensitic microstructures produced from the following labora-
tory cold-rolled grades: (a) DPII (0.80 mm), (b) DP I (0.80 mm), (c) M (1.00mm) and (d) DP II
(1.00 mm). The applied heat treatments are given in the attached illustrations. Nital etchant.
Ultrafine homogeneous dual-phase steels produced by the thermal cycling technique are shown
in Figure 3.10. The specimens were etched with Nital: martensite appears dark brown to black
while ferrite remains light colored. Microstructures with a ferrite mean grain size between 2 and3 m and with martensite grains generally finer than 1 m could be obtained. Thinner grades,
corresponding to a higher cold deformation degree, react more sensitive than the thicker grades
to the dramatic changes of annealing conditions, in such a way that one or two flashing cycles
were enough to achieve grain refinement (Figure 3.10-a and 3.10-b). The small number of
heating-cooling cycles is simply translated into a shorter total holding time in the intercritical
or in the pure austenitic region and, finally, results in the formation of fine microstructures
possessing a relatively low martensite fraction. This agrees well with the recrystallization
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experiments described in 3.1.1, according to which for very short (or for zero) holding times
austenitization is not yet completed.
The in-situ variation of annealing parameters during heat treatment was successfully com-
bined with the rapid recovery and recrystallization kinetics of the laboratory cold-rolled pre-
processed grades, producing ultrafine dual-phase microstructures. Nevertheless, up-scaling the
steel production by this method faces serious difficulties, mainly due to the rapid temperature
changes necessary between the annealing segments and the complexity of controlling them with
the currently available industrial equipment.
3.1.3 Annealing simulations
Laboratory annealing simulations were conducted in order to reproduce the ultrafine mi-
crostructures on a larger scale and so being able to investigate mechanical properties. For
the final selection of annealing conditions the data collected from the dilatometric investiga-tions were taken into account.
The influence of the annealing temperature and of the cooling rate on the microstructure
was studied for all laboratory and industrial grades, since these parameters seem to have the
strongest impact on the microstructure evolution. To avoid grain coarsening, the heating rate
as well as the annealing holding time were held constant for all conventional heat treatment
schedules at 25 K/s and 30 s respectively. Two annealing temperatures were selected, the one
in the intercritical region (800, more or less standard in continuous annealing lines) and the
other in the austenitic region (840). Lower annealing temperatures (750) were excluded
from the simulations, because even though they provide with a variety of low martensite-
fraction steels they have a rather detrimental effect on grain refinement. For each annealing
temperature, a set of six cooling rates -the same as in the dilatometric investigations- was
applied, covering a wide range of microstructures regarding the formation of phases, the fraction
and the morphology of martensite and the mean grain size of ferrite.
Figure 3.11 shows the influence of the cooling rate on the microstructure of the laboratory
grade DP II after intercritical annealing at 800 for 30 s. The specimens are etched with
LePera. At cooling rates lower than 10 K/s a third dark/black colored phase (probably bainite)
is present in the microstructure, located on the grain boundaries and at the grain triple points
(Figures 3.11-a and 3.11-b). As the cooling rate increases, the third phase disappears while
the martensite fraction increases and, simultaneously, the ferrite mean grain size decreases
(Figures 3.11-cto3.11-f). Martensite grains remain fine for all cooling rates up to 40 K/s. In
the microstructures quenched with 60 K/s and 80 K/s also coarse martensite grains are found; in
the interior of these coarse grains a brown tinted substructure can be detected (Figures 3.11-e
and 3.11-f).
An analogous investigation for the martensitic grade M (by keeping the same annealing
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(a)CR= 5 K/s (b) CR= 10 K/s
(c)CR= 20 K/s (d) CR= 40 K/s
(e)CR= 60 K/s (f) CR= 80 K/s
Figure 3.11: Microstructures of the grade DP II annealed at 800 for 30s and quenched withdifferent cooling rates, etched with LePera.
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parameters constant) revealed a more pronounced influence of the cooling rate on the formed
microstructures, with respect to both grain refinement and martensite fraction. As shown in
Figure 3.12and in reference toFigure 3.11, a third phase could not be identified even for the
lowest cooling rates applied. Furthermore, coarse structured martensitic grains appear already
after cooling with 10 K/s, increasing in fraction/number with increasing cooling rate and finally
reaching grain sizes of the order of ferrite.
The cooling rate of 40 K/s proves to be a critical point in the direction of grain refine-
ment. Comparing the microstructure in Figure 3.12-cwith that inFigure 3.12-d, represent-
ing quenching with 20 K/s and 40 K/s respectively, the increase in martensite fraction as well
as the change in the martensite morphology - as indicated by its etching behavior - is distin-
guishable. This effect is followed by significant grain refinement of ferrite approximately by
an order of magnitude. Applying cooling rates higher than 40 K/s leads to higher martensite
fractions which are not accompanied with a proportional decrease in the grain size of ferrite.
The influence of the annealing temperature on the microstructure of the laboratory as well
as of the reference material is shown in Figure 3.13. For this purpose, a moderate cooling
rate of 20 K/s was applied in order to minimize the impact of the cooling rate on the formed
microstructures. At this condition, the industrial grade R reacts more sensitive to the annealing
temperature than the laboratory grades DP I and M. Increasing Tan from 800 to 840
results in a finer microstructure of R, possessing a higher martensite fraction (Figures 3.13-a
and 3.13-b). Nevertheless, small amounts of a dark-colored third phase are present in the
microstructure for both annealing temperatures.
On the other hand, the annealing temperature has no significant influence on the microstruc-tural characteristics of the grade DP I. Martensite fraction remains on the same level without
any morphological change while no third phase could be observed.
In the case of grade M, the higher annealing temperature has a major influence on the
martensite morphology by increasing the number of structured martensite grains, although
the total martensite fraction does not increase. A third phase could not be detected in the
microstructure.
A more representative comparison between all the grades investigated, regarding their an-
nealing behavior, is given in Figure 3.14. In order to emphasize the main differences, which
are of great importance for the interpretation of the mechanical properties, two extreme coolingrates of 5 K/s and 80 K/s were applied after annealing in the austenitic region (840) for 30s.
Comparing the microstructures of Figure 3.14 column-wise, it becomes obvious that the
cooling rate has a very strong influence on all materials with respect to grain refinement of
ferrite. This impact is more pronounced for the grades R and M than for grades DP I and DP II.
A cooling rate of 5 K/s is sufficient for a preliminary grain refinement of the dual-phase grades,
as shown in Figures 3.14-c and 3.14-e, partly explained from their already homogeneous
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(a)CR= 5 K/s (b) CR= 10 K/s
(c)CR= 20 K/s (d) CR= 40 K/s
(e)CR= 60 K/s (f) CR= 80 K/s
Figure 3.12: Microstructures of the grade M annealed at 800 for 30 s and quenched with differentcooling rates, etched with LePera.
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(a) R, Tan= 800 (b) R, Tan= 840
(c) DP I, Tan= 800 (d) DP I, Tan= 840
(e) M, Tan= 800 (f)M, Tan= 840
Figure 3.13: Microstructures of the grades R, DP I and M annealed at 800 (left column) and840 (right column) for 30 s after quenching with 20 K/s, etched with LePera.
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ferritic-martensitic starting microstructure. Irrespective of the grade, all microstructures in the
left column (referring to the lowest cooling rate) possess a small fraction of a third phase. For the
grade M this fraction is negligible (Figure 3.14-g). It should be underlined that the industrial
grade R preserves/retains this third phase (bainite and/or pearlite) even after quenching with
80K/s (Figure 3.14-b), a fact that differentiates this grade from the laboratory ones.
Additionally, the morphology of martensite is strongly dependent on the cooling rate. Low
cooling rates enhance the formation of white, clear, unstructured martensite grains. LePeras
etchant stains white also grains of retained austenite. However, magnetic volumetric measure-
ments showed that no retained austenite is present in the microstructures investigated. When
a dramatically different/higher cooling rate is applied (microstructures of the right column of
Figure 3.14), part of the martensite grains appear dark brown, which is generally true for the
coarser grains with a substructure in their interior.
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(a) R, CR= 5 K/s (b) R, CR= 80 K/s
(c) DP I, CR= 5 K/s (d) DP I, CR= 80 K/s
(e)DPII, CR= 5 K/s (f) DP II, CR= 80 K/s
(g)M, CR= 5 K/s (h) M, CR= 80 K/s
Figure 3.14: Microstructures of the grades R, DP I, DP II and M after annealing at 840 for 30s
and quenched with the lower and the higher cooling rates of 5 K/s (left column) and 80 K/s (right
column) respectively, etched with LePera.
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Further investigations were carried out by means of electron microscopy (involving SEM,
FEG-SEM and TEM) in order to identify all the phases present in the fine and ultrafine
microstructures. As the resolution scale increases, the microstructure characteristics become
location-dependent and no representative conclusions can be drawn about the overall fractions
and/or about the mean grain sizes of the phases. However, higher magnifications enable a more
detailed qualitative analysis.
To make the distinction between the phases possible, the creation of a topographic instead of
a color contrast is required, so that the grain boundaries between the phases are distinguishable.
For this reason all specimens for scanning electron microscopy were etched with Nital, which
preferentially etches ferrite, bainite and cementite and outlines their grain boundaries while
leaving martensite intact/undissolved.
Micrographs of the industrial grade R taken with a Field Emission Gun-Scanning Electron
Microscope at a magnification of 4000 are shown in Figure 3.15. The microstructures were
produced by annealing at 800 for 30 s and quenching with 60 K/s and were taken at different
locations of the same specimen. SEM micrographs reveal the presence of a third bainitic phase
in the microstructure of the reference grade R (Figure 3.15-b), supporting the observations
made in the light microscope. This third phase is mainly detected at the ferrite-martensite
grain boundaries and its fraction does not exceed 2 %.
(a) (b)
Figure 3.15: Micrographs of the industrial grade R annealed at 800 for 30 s and quenched with60 K/s. The symbols F, M and B marked on the grains stand for ferrite, martensite and bainite,respectively.
Quenching with cooling rates equal to or higher than 40 K/s (in the given example 60 K/s)
results in the formation of coarse martensite blocks, consisted of strong orientated laths even
within the same block (Figure 3.15-a). Since no grain boundaries can be identified between
the laths, these blocks are considered and evaluated as discrete martensite grains.
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Analogous difficulties, but this time regarding the identification/observation of grain bound-
aries between ferrite grains, are presented in Figure 3.15-b. There exist regions that resemble
martensite, which are present in the rim of ferrite grains without any distinct grain boundaries.
Light microscopy does not even allow the identification of these subgrains, irrespective of the
etchant used. To simplify the quantitative part of the evaluation and, moreover, to be on the
safe side in the grain size measurements, these grains are also considered as individual ferrite
grains.
The phenomenon described above becomes more pronounced and hence more critical for
the evaluation of the ultrafine microstructures produced from the laboratory grades, especially
after annealing at high temperatures and quenching with high cooling rates. Figure 3.16
shows FEG-SEM micrographs of the DP I grade, annealed in the austenitic region (840) and
quenched with 60 K/s. Although the existence of martensite subgrains partly surrounding the
coarser ferrite grains is evident, no grain boundaries between them could be detected. The
transition from ferrite to martensite is very smooth and the martensite subgrains appear
free of substructures, making the phase identification more difficult. The only indicative char-
acteristic of the martensite presence is a difference in contrast observed between the phases,
where martensite appears brighter due its higher concentration in carbon and in other alloying
elements (Cr, Mn, etc.). Ultramicrohardness measurements have confirmed these assumptions.
(a) (b)
Figure 3.16: Micrographs of the laboratory ferritic-martensitic grade DP I annealed at 840 for 30sand quenched with 60 K/s, (magn. 2000 and 8000 for (a) and (b) respectively).
Martensite grains exhibiting an unusual morphology are also shown in Figure 3.16. The
substructure in the interior of these grains, indicated by the white arrows (Figure 3.16-a), is
definitely different from the lath martensite structure described in Figure 3.15. Such complex
martensitic morphologies were already observed in the light microscope in the form of brown
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tinted substructures within the coarser white martensite grains (as stained with LePera). Ac-
cording to bibliographic references of related works [80, 81], the emergence of the martensite
substructure can be interpreted as a result of carbide precipitation during tempering - which
in the case investigated refers to autotempering during quenching, since no tempering stage is
applied after annealing.
(a) (b) (c)
Figure 3.17: TEM micrographs of grade DP I demonstrating: (a) ferrite and martensite grains,(b) martensite lath structure and (c) cementite precipitates within a grain of tempered martensite.
To clarify the substructure formation within martensite grains and to confirm the presence
of carbide precipitates, transmission electron microscopy (TEM) investigations were performed
on selected specimens, covering a representative range of grades and annealing conditions.
TEM micrographs of the ferritic-martensitic grade DP I are shown in Figure 3.17. The
investigated samples were annealed at 840 (- phase field) for 30 s and quenched with 5 K/s
and 80 K/s, represented by the images (a) and (b, c) respectively. For a direct comparison,
microstructures of the grade DP I subjected to the same heat treatment were earlier discussedin Figures 3.14-c and 3.14-d. The microstructure of the samples cooled down with 5 K/s
consists mainly of martensite and ferrite. Sometimes, bainite and/or cementite precipitates
could be found close to martensite grains. No tempered martensite could be detected, which is
in agreement with the light microscopy observations.
In the samples quenched with 80 K/s, martensite grains exhibiting a lath microstructure
could be observed (Figure 3.17-b). Additionally, significant amounts of martensite grains
with a substructure in the interior are found (Figure 3.17-c). The substructure is characterized
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by the presence of cementite precipitates, aligned along more than one habit plane variants.
These grains are identified as tempered (or to be more specific as autotempered) martensite
based on the observations of Bramfitt et al. [82], according to which the existence of multiple
carbide habit-variants is indicative of autotempered martensite [82]. In the following, the term
autotempered martensite will be used to describe such phase constituents.
Figure 3.18 provides a more detailed image of autotempered martensite formation, taken
from a sample of the martensitic grade M intercritically annealed at 800 and quenched with
80 K/s. As can been observed, autotempered martensite (marked on the micrograph) is located
in the middle of the grain and exhibits a different morphology from martensite located near
the martensite-ferrite interface. Higher resolution analysis confirms the presence of cementite
precipitates, revealing their multi-directional arrangement as well.
Figure 3.18: High resolution TEM micrographs of grade M, highlighting the formation of carbide
precipitates in autotempered martensite.
Additional TEM micrographs of the grade M are presented in Figure 3.19, corresponding
to dramatically different cooling rates after intercritical annealing. The microstructure of thesamples cooled down with 5 K/s (Figure 3.19-a) consists of fine martensite grains homoge-
neously dispersed in the ferritic matrix, without any tempered martensite present. Despite
the low cooling rate applied, no bainite could be detected. What should be also noticed, is
the presence of relatively coarse cementite precipitates located near martensite grains or in the
ferrite grain boundaries.
The samples quenched with 80 K/s (Figures 3.19-b and 3.19-c) show an ultrafine ferritic-
martensitic microstructure. The presence of autotempered martensite, existing in considerable
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fractions in this grade, was previously discussed in detail. Furthermore, negligible amounts
of retained austenite were identified, trapped in the form of thin plates between the laths of
martensite grains. However, the amount of retained austenite found is too small to be taken
into account in the quantitative analysis of the microstructure.
(a) (b) (c)
Figure 3.19: TEM micrographs of grade M intercritically annealed (Tan= 800) and cooled with
5 K/s and 80 K/s, for (a) and (b,c) respectively. The micrographs show: (a) coarse cementite precip-itates located around a martensite grain, (b) ultrafine ferrite and martensite grains and (c) thin filmsof retained austenite between the martensite laths.
3.