3273817 Study Integrated Thesis

RWTH Aachen Institute of Ferrous Metallurgy

Study Integrated Thesis

Master Student Alireza Saeed-Akbari

Matr. –Nr. 268696

Subject:

Investigation on Material Characterization and

Mechanical Properties of Ultra High Strength

Boron Steel at High Temperatures

Supervisors: Univ. Prof. Dr.-Ing. W. Bleck M.Sc. Malek Naderi

December 2006

Acknowledgment I would like to deeply appreciate the cooperations and

considerations of all of the colleagues at the Institute of Ferrous

Metallurgy (IEHK), RWTH Aachen University. I should give a

special thanks to Mr. Malek Naderi for his kind and permanent

supports and helps during the project.

The workshops, metallography and warm deformation

laboratories' technicians are those who are mostly engaged with

the related experiments of the current work, for which I am

really grateful.

Abstract In new generation of car bodies, hot stamped boron steel parts are widely used in safe guards, bumpers, A and B pillars. By applying hot stamping technology, one can get a full martensitic microstructure in the final product. Thus, investigation on the behavior of boron steel at high temperatures would be the main objective. Current thesis presents a comprehensive set of laboratory works in terms of the mechanical simulation of the industrial hot stamping process. This was performed by means of the hot compression tests on a quenchable boron steel used for the hot stamping of the car bodies under the variety of thermomechanical conditions. The hardness and microstructural variations and dimensional changes through martensitic transformation were studied. Results of the present work show that the possibility for the formation of bainite and ferrite phases during the experiments at the temperatures far below the relevant phase regions in the CCT diagram must be taken into account. Furthermore, the lower isothermal compression temperatures and less amounts of the applied strain during the continuous forming and quenching experiments, increase the chance of having a full martensitic microstructure in the final product.

Table of Contents

1. Introduction 1

2. Literature Survey 3 2.1. Introduction, 3 2.2. Hot Stamping Process Background, 3

2.3. Metallurgical Fundamentals, 6 2.3.1. Effects of Boron and Carbon, 6 2.3.2. Characterization of Bainitic Microstructures, 10 2.3.2.1. Isothermally Formed Bainite, 11 2.3.2.2. Continuously Cooled Bainite, 12 2.3.3. Continuously Cooled Ferritic Microstructures, 17 2.3.3.1. Bainitic or Acicular Ferrite, 19 2.3.3.2. Granular Ferrite, 19 2.3.4. Martensite and Martensitic Transformation in Steels, 21 2.3.4.1. Chemical Composition Effect, 23 2.3.4.2. Cooling Rate Effect, 24 2.3.4.3. Austenization Temperature Effect, 25 2.3.4.4. Quenching Media Effect, 26 2.3.4.5. Lath Martensite, 29 2.3.4.6. Medium Carbon Martensite, 30 2.3.5. Thermomechanical Behavior of Boron Steels, 31

3. Experimental Procedure 35 3.1. Material Characterization, 35 3.1.1. Chemical Composition, 35 3.1.2. Microstructure, 35 3.1.3. CCT Diagram Design, 36 3.2. Isothermal and Non-isothermal Compression Tests, 37 3.3. Hardness and Metallography Tests, 41

4. Results and Discussion 41 4.1. High Temperature Isothermal Compression Tests, 41

4.1.1. The Hardness Values and Microstructural Evolutions, 41 4.1.1.1. Results, 41 4.1.1.2. Discussion, 47

4.1.2. The Deformation Data Analysis, 53 4.1.2.1. Results, 53 4.1.2.2. Discussion, 57

4.1.3. The Dilatation Data Analysis, 60 4.1.3.1. Results, 60 4.1.3.2. Discussion, 65

4.2. Simultaneous Deformation and Quenching Tests, 67 4.2.1. The Strain Magnitudes Effect, 67

4.2.1.1. Results, 67 4.2.1.2. Discussion, 74

4.2.2. The Strain Rate Effect, 77 4.2.2.1. Results, 77 4.2.2.2. Discussion, 80

4.2.3. The Austenization Soaking Time Effect, 84 4.2.3.1. Results, 84 4.2.3.2. Discussion, 88

4.2.4. The Initial Deformation Temperature Effect, 91 4.2.4.1. Results, 91 4.2.4.2. Discussion, 95

5. Conclusions 98 References i

Chapter 1 - INTRODUCTION 1

CHAPTER

ONE

INTRODUCTION Permanently increasing claims on passenger protection demand new solutions

of the automobile industry in consideration of steel light weight construction

[1]. The art or science of sheet metal stamping processes is challenged daily to

accommodate higher strength and thinner materials. Further, these materials

must be transformed into more complex shapes with fewer dies and increased

quality in the final part. High-strength and ultra high-strength steels have less

ductility, and hence less formability than lower strength steels. Thus, care must

be taken in part design and forming method selection. In addition, problems

such as springback1 are increased with yield strength and it must be accounted

for in the process design [2]. Hot stamping, i.e. simultaneous forming and

quenching, is a one-step manufacturing process for high and ultra high strength

steel profiles to be used as safety related structural components in car body

structures. In addition to the temperature-related improvement in the forming

properties of these steels, hot stamping makes it possible to further increase the

strength of materials during deformation. Within the process, thin-walled

profiles or blanks of hardenable steel are heated to the austenite region, placed

in a cooled forming tool where they are simultaneously hardened and formed to

the desired shape. Alternatively, the blanks are formed, clamped and

subsequently hardened by spray quenching and released. So the problem that is

solved by hot stamping is getting the advantages of the advanced high-strength

material's properties without the manufacturing limitations (such as springback)

in very high strength steels [3, 4].

Regarding the material selection to reduce weight and to improve the safety of

vehicles, car makers need very high resistance flat products with very good

1 Condition that occurs when a flat-rolled metal or alloy is cold-worked; upon release of the forming force, the material has a tendency to partially return to its original shape because of the elastic recovery of the material. This is called "springback" and influenced not only by the tensile and yield strengths, but also by thickness, bend radius and bend angle.

Chapter 1 - INTRODUCTION 2

formability and toughness for structural and impact resistant parts.

At present, these objectives can be fulfilled by the use of HSLA steels or

aluminum. If the shape is complex, the use of aluminum is a solution for

weight saving but its strength is limited and it is an expensive solution. The use

of HSLA steels is a satisfactory solution if the shape is not very complex.

These steels exhibit good strength, weldability and impact resistance. However,

an increase in strength decreases formability. To solve this problem, the

solution is to separate the required characteristics, i.e. good formability and a

very high strength. The way to realize this objective is in the use of new

structural sheet of boron steel with very good formability which permits

complex shapes to be obtained [4]1. To partly overcome and quantify the

mentioned challenges, modeling and simulation are gaining increasing

importance in the product development of structural components whose

manufacturing is based on the thermomechanical forming. However, finite

element simulations of coupled thermomechanical processes require accurate

and efficient simulation tools as well as relevant material data based on testing

of the mechanical and thermal properties under reliably simulated processing

conditions. The final mechanical properties depend on the microstructural

evolution [5].

The current work presents a comprehensive set of laboratory works in terms of

the mechanical simulation of the industrial hot stamping process by means of

hot compression tests on a 22MnB52 steel under the variety of

thermomechanical conditions. The mechanical properties and microstructural

data are given in two independent parallel sections regarding the high-

temperature isothermal compression and simultaneous deformation and cooling

processes (Chapter 4) to find the most optimized way of providing higher

strength and hardness values besides the most achievable industrial standards.

The experimental conditions, i.e. temperature and deformation parameters, are

chosen to be comparable with different industrial applications and limitations.

1 It should be noted that the addition of boron in the range of 0.0005-0.005%wt to certain steels increases the hardenability to great extent. 2 The experimental material is called BTR165 by the material provider.

Chapter 2 – LITERATURE SURVEY

3

CHAPTER

TWO

LITERATURE SURVEY

2.1. Introduction

Through the following pages, an overview regarding the fundamental concepts

and requirements of a successful hot stamping process and the microstructural

and mechanical characteristics of the ultra high strength boron steels are

presented. The role of boron on the thermomechanical behavior of steels and

the response of the experimental material to certain heat treating processes

followed by an appropriate high temperature deformation are described. The

current chapter is based on the definition of the hot stamping process, the

results and achievements of the previous investigations on different

optimization aspects of the hot deformation of boron steels, the role of alloying

elements, the relevant phase transformations and the final mechanical

properties of the studied materials.

2.2. Hot Stamping Process Background

The steels used in the automotive industry can be classified based on different

criteria. According to [6], the mentioned definitions are summarized as follow:

1. By metallurgical designation:

- Low-strength steels: interstitial-free (IF) and mild steels;

- Conventional high-strength steels: carbon-manganese (CMn),

bake hardenable (BH), interstitial free high-strength (IF-HS) and

high strength low-alloy steels (HSLA);

- Advanced-high-strength steels (AHSS): dual phase (DP),

transformation induced plasticity (TRIP), complex phase (CP)

and martensitic steels.


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2. Mechanical properties – tensile strength:

- Low strength steels, LSS: tensile strength <270MPa;

- High strength steels, HSS: tensile strength 270 – 700MPa;

- Ultra-high-strength steels, UHSS: tensile strength > 700MPa.

3. Mechanical properties – total elongation versus tensile strength.

Due to continuously higher demands from different organizations and severe

legislation on passive automotive safety and effort to reduce vehicle emissions,

the use of high- and ultra high strength components in both car body and

closures have increased drastically during the last two decades. The

components in the car body that are commonly made of ultra high strength

steels are shown in figure 2-1 [7]. The different components in accordance to

this figure are:

1. Door beam;

2. Bumper beam;

3. Cross and side members;

4. A-/B- pillar reinforcement;

5. Waist rail reinforcement.

Figure 2-1- Components in car body using ultra high strength steels [7].

The hot stamping process, as was introduced in chapter one, uses boron steel

blanks which are first austenitized at a temperature of ~900°C and then formed


5

and quenched between cold tools. The forming operation at elevated

temperatures allows complex geometries to be obtained due to the high

formability of the hot material. The quenching results in a material with a very

high yield and tensile strength, which falls into the category of martensitic ultra

high strength steels. Moreover, the hardened component shows a dimensional

accuracy comparable to that of mild steel products manufactured with

conventional forming methods. It must be noted that hot stamping of

quenchable steels is a non-isothermal sheet forming process. Due to this, the

calculation of the temperature evolution in the blank is crucial, because of its

significant impact on the material deformation as well as on its final

microstructure after phase transformation. This temperature evolution is

controlled by the heat exchange between the tools and the sheets [7, 8]. Finally,

hot stamping technology can be summarized by the following steps:

1. Punching of blanks;

2. Austenization in a furnace;

3. Forming and hardening;

4. In some cases surface treatment by blasting or pickling.

Figure 2-2 shows the mentioned process as described above.

Figure 2-2- Schematic description of the hot stamping process [7].


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2.3. Metallurgical Fundamentals

In this section as the major part of the current chapter, a brief introduction to

the effects of two important alloying elements – i.e. boron and carbon - on the

metallurgical characteristics of steels is given. Additionally, different phase

transformations in steels are evaluated based on their relation with the hot

deformation (simultaneous deformation and cooling process) of the ultra high

strength steels. The final part of current section is dedicated to an overview on

the effects of different deformation parameters on the phase transformations in

steels.

2.3.1. An Insight into the Effects of Boron and Carbon as Alloying Elements on the Hardenability of Steels

The simplest version of analyzing the effects of alloying elements on iron-

carbon alloys would require the consideration of a large number of ternary

alloy diagrams over a wide temperature range. However, Wever [9] pointed out

that iron binary equilibrium systems fall into four main categories (Figure 2-3):

open and closed γ-field (austenite) systems, and expanded and contracted

γ-field systems. This approach indicates that alloying elements can influence

the equilibrium diagram in two ways:

• by expanding the γ-field, and encouraging the formation of austenite

over wider compositional limits. These elements are called γ-stabilizers.

• by contracting the γ-field, and encouraging the formation of ferrite over

wider compositional limits. These elements are called α-stabilizers.

The form of the diagram depends to some degree on the electronic structure of

the alloying elements which is reflected in their relative positions in the

periodic classification [10]. Among different alloying elements in steels, the

role of boron and carbon is described here (due to their relation to the topic of

the current work).


7

Figure 2-3- Classification of iron alloy phase diagrams: a. open γ-field; b. expanded

γ-field; c. closed γ-field; d. contracted γ-field [10].

Regarding the expanding γ-field elements, carbon and nitrogen are the most

important elements in this group. Although the γ-phase field is expanded, its

range of existence is cut short by compound formation (Figure 2-3b). The


8

expansion of the γ-field by carbon, and nitrogen, underlies the whole of the

heat treatment of steels, by allowing formation of a homogeneous solid solution

(austenite) containing up to 2.0 wt % of carbon or 2.8 wt % of nitrogen.

Boron (together with the carbide forming elements tantalum, niobium and

zirconium) is the most important element under the contracted γ-field category.

The γ-loop is strongly contracted, but is accompanied by compound formation

(Figure 2-3d).

Although only binary systems have been considered so far, when carbon is

included to make ternary systems the same general principles usually apply.

For a fixed carbon content, as the alloying element is added, the γ -field is

either expanded or contracted depending on the particular solute [10].

The astonishingly large effect of a minute percentage of boron on hardenability

of steel has been of intriguing interest to metallurgists since the advent of

commercial boron steels. Compared to other elements commonly added to steel

to increase hardenability, commercial boron steels exhibit the following major

unique characteristics: loss in the hardenability effect of boron on austenitizing

at relatively high temperature and marked variation in the hardenability effect

of boron with carbon content of steel [11].

A number of hardenability mechanisms have been proposed to explain these

observations, but they are capable of explaining only part of the reported

evidence. Of the proposed mechanisms, four have survived to the present. All

assume that boron influences hardenability by retarding the nucleation of ferrite

and that it does not influence the thermodynamic properties of the bulk

austenite or ferrite phases. The first assumption is based on observations that

boron does not significantly change the growth rate of ferrite or the formation

rate of pearlite and martensite. The second assumption is based on the small

amount of boron present [12].

With the possibility of boron concentrations reaching significant levels (i.e. for

the austenite grain size of more than 30µm as described in [12]), a number of

mechanisms for retarding ferritic nucleation can be considered as follows:


9

- Reduction in austenite grain boundary energy: this occurs by the

diffusion of boron to austenite grain boundaries and lowering

their energy; hence, making them less favorable sites for ferrite

nucleation;

- Reduction in diffusivity: by decreasing the self-diffusivity of iron

in austenite grain boundaries;

- Reduction in number of sites: the first mechanism that of grain

boundary reduction, assumes that the austenite grain boundary

can be treated like a continuum with nucleation possible at any

site on the boundary. However, if the crystallographic nature of

the grain boundary is considered, one finds that even in high

angle boundaries, where one would expect nucleation to be most

rapid, there are regions of relatively high and low atom density.

If regions of low atom density are favored sites for nucleation of

ferrite, it is possible that boron poisons them either through filling

up the free volume by segregation there, or through precipitating

on them as borocarbides. If boron contaminates half the sites, the

ferrite "C" curve on TTT diagram would be shifted by a factor of

two. If it contaminated all the sites, the ferrite would have to form

elsewhere at a reduced rate. One attractive feature of this

mechanism is that there is no theoretical limit to the possible "C"

curve shift, and another is that in principle there can be sufficient

boron atoms present to saturate the sites. Because of these

features, one is not pressed to explain why small amounts of

boron can have a large effect.

- Nucleation of ferrite on borocarbides: Fe23(BC)6 precipitation is

a precursor to ferrite formation. The proponents of both the

reduction in grain boundary energy and the reduction of

diffusivity theories interpret this observation on the basis that

borocarbides precipitation draws boron out of the grain boundary

and removes the inhibition effect. Those who favor a site-


10

competition mechanism can interpret the above observation, also.

They suggest that borocarbides block ferrite nucleation only

when they are small but they encourage ferrite nucleation on their

own interfaces when they are sufficiently large [12].

The hardenability effect of boron element and the austenitizing temperature are

related in a pattern of a single peak curve, i.e. there is an increasing trend

followed by a drop after passing through a peak in the hardenability value as

the austenizing temperature increases. The quenching operation for the

structural boron steel should be performed at a lowest possible temperature

above Ac3, and the austenitizing quenching operation under 900°C can be

applied to all kinds of the boron steel, but the austenitizing case-hardening or

quench operation above 900°C is only applicable to the low alloy boron steel

with grains not likely to develop [13].

2.3.2. Characterization of Bainitic Microstructures The numerous terms created over the last 50 years to describe specific bainite

morphologies have led to some confusion, and it is suggested that the

commonly used terminologies do not adequately describe the full range of

bainitic microstructures which are observed [14]. Upper and lower bainite are

established terms describing microstructures which can easily be distinguished

using routine microscopy, and whose mechanisms of formation are well

understood. There are, however, a number of other descriptions of steel

microstructures which include the word 'bainite'. These additional descriptions

can be useful in communicating the form of the microstructure. But this must

be done with care, avoiding the natural tendency to imagine a particular

mechanism of transformation, simply because someone has chosen to coin the

terminology [15].

The morphological features of ferrous martensites have been rather well

characterized over the past decades. In comparison, the characterization of

bainitic microstructures and properties is much less complete. Bainite has


11

received relatively little attention, and a great deal of effort will be required to

understand the bainitic transformation more fully, particularly of bainite which

forms during continuous cooling [14].

In this part of the current chapter, the principles of the bainitic transformation

in addition to different bainite morphologies and its suggested definitions are

given. The target is to gradually bring readers' attention to the importance of

dividing the general concept of bainitic transformation into two major

categories: isothermally formed bainite and continuously cooled bainite. Later

on, the role of this second type of bainitic transformation – i.e. continuously

cooled bainite – will be discussed in chapter 4. It is then seen that having a

deep knowledge about the possible bainite morphologies can help to avoid any

misunderstandings of the final appeared phases in the microstructure of the

continuously cooled steels. As showing and describing all the possible bainite

morphologies are out of the discussions of the current report, the following

pages concentrate more on the definition and characteristics of the continuously

cooled bainite – i.e. granular bainite – in more details. For more information on

other bainitic transformation mechanisms and microstructures, please refer to

[14] and [15].

2.3.2.1. Isothermally Formed Bainite

Isothermal bainite is usually distinguished as 'upper' or 'lower' depending on

whether the carbides are distributed between individual ferrite regions or within

them, respectively. The carbides are usually cementite, although ε-carbide may

also be found in lower bainite. The difference between upper and lower bainite

is also based on whether the transformation temperature is above or below

approximately 350°C, although it has been shown that the distinction is not

universally applicable. Upper bainite comprises a lathlike morphology, and the

austenite/ferrite habit plane is thought to be near {111}γ/{110}α. Lower bainite

is generally reported to have a platelike morphology in isothermally

transformed steels, with an irrational habit plane somewhat further away from

{111}γ. The carbide in lower bainite generally consists of a single


12

crystallographic variant inclined to the apparent longitudinal axis of ferrite,

although multiple variants have also been reported (similar to those observed in

tempered martensite) [14].

2.3.2.2. Continuously Cooled Bainite

To the physical metallurgist, bainitic steels are recognized by the shape of their

CT diagram. In steels which are commercially important, a typical diagram

features the polygonal ferrite transformation (ferrite nose) shifted rightward to

regions of very slow cooling rate, thereby exposing a broad, flat bainite

transformation region. The advantage of having a CT diagram with a broad, flat

bainitic nose is that bainite with an almost constant transformation start

temperature can be produced over a wide range of cooling rates. Consequently,

bainite can be produced in heavy sections with little change in tensile

properties compared to thinner sections.

A typical CT diagram for a commercial bainitic steel is shown in figure 2-4 to

illustrate some of the important features of the bainitic transformation in

continuously cooled steels. In this figure, the bainitic transformation spans a

range of cooling rates from about 4°C/min to 600°C/min (measured between

800°C and 500°C). This range is typical of the rates experienced during

thermomechanical processing or heat treating in the commercial production of

steel components varying in thickness from 100 to 1000mm.


13

Figure 2-4- Continuous cooling-transformation diagram for a Ni-Cr-Mo steel.

Composition of steel (weight percent): 0.15C, 0.32Mn, 0.31Si, 2.72Ni, 0.41Mo [14].

Although the CT diagram in figure 2-4 indicates that bainitic microstructures

are generated over a wide range of cooling rates, the situation is complicated

because of the wide variations in microstructure which are actually observed.

For example, the light-optical microscope shows the appearance of an

'acicular' bainite microstructure (with some martensite) at a cooling rate of

461°C/min. At a much slower cooling rate of 3°C/min, a 'granular' bainitic

microstructure is produced.

In fact, the terms upper and lower bainite were originally used to describe

isothermal transformations in specific temperature regimes, but the terminology

is rather less meaningful (and even misleading) in describing the bainitic

transformation during continuous cooling were substantially different

microstructures can be obtained over a relatively constant range of

transformation temperatures.


14

Of all the unusual descriptions of bainitic microstructures, granular bainite is

probably the most useful and frequently used nomenclature. During the early

1950's, continuously cooled low-carbon steels were found to reveal

microstructures which consisted of 'coarse plates and those with an almost

entirely granular aspect', together with islands of retained austenite and

martensite (figure 2-5) [16].

Figure 2-5- Granular bainite in a Fe-0.15C-2.25Cr-0.5Mo wt% steel: Left picture,

light micrograph; Right picture, corresponding transmission electron micrograph [15].

Habraken and coworkers [17-20] called this morphology as 'granular bainite'

and the terminology became popular because many industrial heat-treatments

involve continuous cooling rather than isothermal transformation.

Granular bainite is supposed to occur only in steels which have been cooled

continuously; it cannot be produced by isothermal transformation.

Habrakan and Economopoulos [20] summarized their findings schematically in

the CT diagram which is presented here in figure 2-6.


15

Figure 2-6- Schematic representation of a CT diagram showing formation of granular

bainite (path I), upper bainite (path II), and lower bainite (path III) [14].

At relatively slow cooling rates, 'granular bainite' is formed (cooling path I).

At intermediate cooling rates (cooling path II), they reported the formation of

upper bainite. To form lower bainite, they suggested that an isothermal hold

just above the Ms temperature is required, as indicated by cooling path III in

figure 2-6.

One of the most complete studies on the nature of continuously cooled bainite

was carried out by Ohmori et al [21]. Their work examined various

microstructures which developed through both isothermal and continuous

cooling transformation in Ni-Cr-Mo steel. Using both replicas and thin foils,

they examined the fine morphological and crystallographic details in this alloy

and separated the various microstructures into three distinct classes which they

called bainite I, bainite II and bainite III (figure 2-7).


16

Figure 2-7- Schematic representation of the CT diagram of a Ni-Cr-Mo steel showing

three forms of bainite; Bainite I being a carbide-free form, bainite II being a form

similar to upper bainite, and bainite III being a form similar to lower bainite [14].

Bainite I consists of a carbide-free acicular ferrite with well-defined films of

retained austenite (and/or martensite) at the lath boundaries; bainite II is similar

to upper bainite, with cementite particles between the carbide free ferrite laths.

Bainite III is similar to the lower bainite, with cementite 'platelets' forming

within the laths. However, the acicular ferrite was found to be present in a lath

morphology, rather than in the plate morphology which is typically reported for

lower bainite.

The coarse plates referred to earlier in the current section (page 14) regarding

the granular bainite structure, do not really exist. They are in fact, sheaves of

bainitic ferrite with very thin regions of austenite between the sub-units

because of the low carbon concentration of the steels involved. Hence, on an

optical scale, they give an appearance of coarse plates (figure 2-5).

A characteristic (though not unique) feature of granular bainite is the lack of

carbides in the microstructure. The carbon that is partitioned from the bainitic


17

ferrite stabilizes the residual austenite, so that the final microstructure contains

both retained austenite and some high-carbon martensite. Consistent with

observations on conventional bainite, there is no redistribution of substitutional

solutes during the formation of granular bainite. The extent of transformation to

granular bainite is found to depend on the undercooling below the bainite-start

temperature. This is a reflection of the fact that the microstructure, like

conventional bainite, exhibits an incomplete reaction phenomenon.

The evidence therefore indicates that granular bainite is not different from

ordinary bainite in its mechanism of transformation. The peculiar morphology

is a consequence of two factors: continuous cooling transformation and a low

carbon concentration. The former permits extensive transformation to bainite

during gradual cooling to ambient temperature. The low carbon concentration

ensures that any films of austenite or regions of carbide that might exist

between sub-units are minimal, making the identification of individual platelets

within the sheaves rather difficult using light microscopy.

Finally, it is interesting that in an attempt to deduce a mechanism for the

formation of granular bainite, Habraken (1965) [18] proposed that the austenite

prior to transformation divides into regions which are rich in carbon, and those

which are relatively depleted. These depleted regions are then supposed to

transform into granular bainite [14, 15].

2.3.3. Continuously Cooled Ferritic Microstructures

In contrast to the equiaxed ferritic microstructures of conventionally hot-rolled

or cold-rolled-and-annealed steels, the ferritic microstructures formed by

decomposition of austenite and by virtue of alloying or rapid cooling, often

assume non-equiaxed morphologies. The temperature range in which the non-

equiaxed morphologies of ferrite form is intermediate to those at which

austenite transforms to equiaxed ferrite/pearlite and martensite. Therefore, this

range is the same as that in which bainitic microstructures form in medium-

carbon steels. However, the low-carbon steel ferritic microstructures formed at

intermediate temperatures differ in variety and form from classical bainitic


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microstructures. Figure 2-8 shows continuous-cooling-transformation (CCT)

diagram for an HSLA plate steel evaluated by Thompson et al. [23].

Different parts of the diagram are labeled by the letters PF, WF, AF and GF,

which stand for polygonal ferrite, Widmanstätten ferrite, acicular ferrite and

granular ferrite, respectively.

The general reaction of austenite to ferrite implies rejection of carbon into

retained austenite, according to the dynamic solubility limits of ferrite. At very

high cooling rates, even in very-low-carbon steels or irons with sufficient

hardenability, austenite may transform to martensite.

Figure 2-8- Continuous-cooling-transformation diagram of HSLA steel containing in

mass%, 0.06C, 1.45Mn, 1.25Cu, 0.97Ni, 0.72Cr, 0.42Mo [23].

In addition to the relatively well-characterized forms of ferrite which form from

austenite at high temperatures, types of ferrite which form from austenite at

intermediate temperatures are now commonly observed in continuously cooled

low-carbon steels [22].


19

Among different ferrite morphologies described by Kraus and Thompson [22],

'bainitic (or acicular) ferrite' and 'granular (or granular bainitic) ferrite' are

those related to the intermediate formation temperatures of ferrite. As the

mentioned temperature range is also applicable for the bainitic transformation,

its concept is briefly discussed here.

2.3.3.1. Bainitic or Acicular Ferrite

With increasing cooling rates, the austenite of low-carbon and ultra low-carbon

steels transforms to much finer ferrite crystals than described above. The most

commonly used terms for the resulting ferritic microstructures are bainitic

ferrite and acicular ferrite. The transformation temperatures for the formation

of these ferritic microstructures are clearly in the intermediate temperature

range as shown in the continuous cooling transformation diagram of figure 2-8.

Although the austenite decomposition is only to ferrite, coexisting with retained

austenite or M/A constituent, the microstructural arrangement of acicular

shaped ferrite crystals in groups of parallel laths is included in the Ohmori et al.

[21], bainite classification as BI bainite and in the Bramfitt and Speer bainite

classification [14] as B2, acicular ferrite with interlath austenite. Thus, the

literature describes the fine non-equiaxed ferritic intermediate temperature

austenite transformation product as both ferrite and bainite.

2.3.3.2. Granular Ferrite

Granular bainitic ferrite or granular ferrite, GF, has many similarities to bainitic

or acicular ferrite, but there appear to be morphological differences which merit

a separate category of austenite-to-ferrite transformation. Microstructures

consisting of granular bainite also form in the intermediate austenite

transformation range, as shown in CCT diagram of figure 2-8.

Although acicular and granular ferrites form over the same transformation

temperature range, the cooling rates which form granular ferrites appear to be

somewhat slower than those which form acicular ferrite [22].


20

Similar to acicular ferrite microstructures, the microstructure of granular ferrite

coexists with dispersed retained austenite or M/A particles in a featureless

matrix which may retain the prior austenite grain boundary structure. However,

in contrast to the acicular ferrite microstructures, the dispersed particles have a

granular or equiaxed morphology. TEM images show that the ferritic matrix

consists of fine ferrite crystals, containing high densities of dislocations,

separated by low-angle grain boundaries. As for acicular ferrite

microstructures, the low-angle boundaries explain the insensitivity of the

matrix ferrite crystals to etching for light microscopy. The ferrite crystals have

granular or equiaxed shapes which cause enclosed retained austenite or M/A

regions, by default, to have the granular or equiaxed shapes resolvable in light

micrographs.


21

2.3.4. Martensite and Martensitic Transformation in Steels Perhaps the most important allotrope of iron is 'martensite', a chemically

metastable substance with about four to five times the strength of ferrite [24].

The name martensite is after the German scientist Martens [25]. A minimum of

0.4 wt% of carbon is needed in order to form martensite [24]. It was originally

described as the hard microconstituent found in quenched steels. Martensite

remains of greatest technological importance in steels where it can confer an

outstanding combination of strength (>3500 MPa) and toughness

(>200 MPa m1/2). Martensite can form at very low temperatures, where

diffusion, even of interstitial atoms, is not conceivable over the time period of

experiment. The highest temperature at which martensite forms is known as the

martensite-start, or Ms temperature. Although it is obvious that martensite can

form at low temperatures, this is not necessary to occur. Therefore, a low

transformation temperature is not sufficient evidence for diffusionless

transformation.

Martensite plates can grow at speeds of sound in the metal. In steel this can be

as high as 1100ms-1, which compares with the fastest recorded solidification

front velocity of about 80ms-1 in pure nickel. Such high speeds are inconsistent

with diffusion during transformation [25].

When the austenite is quenched to form martensite the carbon is 'frozen' in

place when the cell structure changes from FCC to BCC. The carbon atoms are

much too large to fit in the interstitial vacancies and thus distort the cell

structure into a Body Centered Tetragonal (BCT) structure [24]. The chemical

composition of martensite can be measured and shown to be identical to that of

the parent austenite. The totality of these observations demonstrate

convincingly that martensitic transformation are diffusionless [25].

The heat treatment process for most steels to get a full martensitic

microstructure involves heating the alloy until austenite forms, then quenching

the hot metal in water or oil, cooling it so rapidly that the transformation to

ferrite or pearlite does not have time to take place. The transformation into


22

martensite, by contrast, occurs almost immediately, due to the lower activation

energy.

Martensite has a lower density than austenite, so that the transformation

between them results in a change of volume. In this case, expansion occurs.

Internal stresses from this expansion generally take the form of compression on

the crystals of martensite and tension on the remaining ferrite, with a fair

amount of shear on both constituents. If quenching is done improperly, these

internal stresses can cause a part to shatter as it cools; at the very least, they

cause internal work hardening and other microscopic imperfections. It is

common for quench cracks to form when water quenched, although they may

not always be visible.

At this point, if its carbon content is high enough to produce a significant

concentration of martensite, the resulted product is extremely hard but very

brittle. Often, steel undergoes further heat treatment at a lower temperature to

destroy some of the martensite (by allowing enough time for cementite, etc., to

form) and help settle the internal stresses and defects. This softens the steel,

producing a more ductile and fracture-resistant metal. Because time is so

critical to the end result, this process is known as 'tempering', which forms

tempered steel [24].

Ideally, as mentioned before, the martensite reaction is a diffusionless shear

transformation, highly crystallographic in character, which leads to a

characteristic lath or lenticular microstructure.

The martensite reaction in steels is the best known of a large group of

transformations in alloys in which the transformation occurs by shear without

change in chemical composition. The generic name of martensitic

transformation describes all such reactions.

It should however be mentioned that there is a large number of transformations

which possess the geometric and crystallographic features of martensitic

transformations, but which also involve diffusion. Consequently, the broader

term of shear transformation is perhaps best used to describe the whole range of

possible transformations.


23

The martensite reaction in steels normally occurs athermally, i.e. during

cooling in a temperature range which can be precisely defined for a particular

steel. The reaction begins at a martensitic start temperature Ms which can vary

over a wide temperature range from as high as 500°C to well below room

temperature, depending on the concentration of γ-stabilizing alloying elements

in the steel [26].

2.3.4.1. Chemical Composition Effect

The martensite start temperature, Ms, is of vital importance for engineering

steels. Hence, great efforts have been made in predicting the Ms's of steels.

Obviously, chemical composition of a steel is a main factor in affecting its Ms

although the microstructure, (dislocation, vacancies, grain, twin, interphase

boundaries, and precipitates), external stress and plastic deformation, may

sometimes play an important role, too.

The Ms temperature of engineering pure iron is estimated as 540°C. C, Mn, Mo,

Cr and Si decrease the Ms while Mo increases the Ms. The analysis indicates

that most alloying elements have similar effects upon the Ms and A3

temperature.

The interactions between substitutional alloying elements can play an important

role in changing the Ms temperature. The Si-Mn interaction strongly increases

the Ms, while Si-Mo interaction significantly decreases the Ms. So far, there is

no proper physical explanation for this though supportive evidence has been

obtained from phenomenological result. Mn and Mo have the weakest

interaction. Si and Mo themselves have weak influence but their overall effect

depends further on the concentration of other alloying elements because of the

strong interactions found with other alloying elements [27].

Once the Ms is reached, further transformation takes place during cooling until

the reaction ceases at the Mf temperature. At this temperature all the austenite

should have transformed to martensite but frequently, in practice, a small

proportion of the austenite does not transform. Larger volume fractions of


24

austenite are retained in some highly alloyed steels, where the Mf temperature

is well below room temperature.

To obtain the martensitic reaction, it is usually necessary for the steel to be

rapidly cooled, so that the metastable austenite reaches Ms. The rate of cooling

must be sufficient to suppress the higher temperature diffusion-controlled

ferrite and pearlite reactions, as well as other intermediate reactions such as the

formation of bainite. The critical rate of cooling required is very sensitive to the

alloying elements present in the steel and, in general, will be lower as the total

alloy concentration is higher [26].

2.3.4.2. Cooling Rate Effect

In general, the martensitic transformation temperature is dependent on the

cooling rate when cooling rate is not high; above a critical cooling rate,

however, the starting temperature of the transformation is constant. Although

the constant starting temperature had been established many years ago, the

issue whether the Ms is constant and independent of the cooling rate was often

raised. In iron-base alloys, it is often observed that the transformation

temperature versus cooling rate curve show two plateaus when cooling rates

exceed a critical cooling rate (figure 2-9).

Figure 2-9- Relation between the transformation temperature of iron and the cooling

rate (0.006 – 0.039%C) [28].


25

In such a case, the plateau at the lower temperature is thought to be the Ms

temperature and the one at the higher temperature to be the A3 temperature (for

iron-base alloys), corresponding to the largest supercooling [28].

2.3.4.3. Austenization Temperature Effect

It has been reported that the higher the austenization temperature, the higher the

Ms temperature. Figure 2-10 shows an example, in which the broken line

indicates that the γ grain size increases as the austenization temperature

increases. Also, the longer the heating time, the higher the Ms temperature

(figure 2-11).

Figure 2-10 – Change of Ms temperature and austenite grain size with austenitizing

temperature (Fe – 0.33%C – 3.26%Ni – 0.85%Cr – 0.09%Mo; heating time 2 min for

800°C – 1000°C, 1 min for >1000°C) [28].


26

Figure 2-11- Change of Ms temperature with heating time of austenization (same

alloy as in figure 2-10; heating temperature 800°C) [28].

2.3.4.4. Quenching Media Effect

As to the interpretation of this fact, it must be noted that a lower quenching

temperature produces more frozen-in vacancies and hence more nucleation

sites. But it is uncertain how effective this phenomenon actually is. On the

other hand, a lower quenching medium temperature must produce a larger

thermal strain during quenching; hence it is expected to raise the Ms

temperature. This effect, however cannot be very large. A more likely cause of

raising the Ms temperature is the reduction of the energy needed for the

complementary shear during transformation, which originates in the

elimination of lattice imperfections due to heating to a higher temperature [28].

Each grain of austenite transforms by the sudden formation of thin plates or

laths of martensite of striking crystallographic character. The laths have a well-

defined habit plane and they normally occur on several variants of this plane

within each grain. The habit plane is not constant, but changes as the carbon

content is increased.

Martensite is a supersaturated solid solution of carbon in iron which has a

body-centred tetragonal structure, a distorted form of bcc iron. It is interesting

to note that carbon in interstitial solid solution expands the fcc iron lattice

uniformly, but with bcc iron the expansion is nonsymmetrical giving rise to

tetragonal distortion.


27

Analysis of the distortion produced by carbon atoms in the several types of site

available in the fcc and bcc lattices, has shown that in the fcc structure the

distortion is completely symmetrical, whereas in the bcc one, interstitial atoms

in z positions will give rise to much greater expansion of iron-iron atom

distances than in the x and y positions.

Martensitic planes in steel are frequently not parallel-sided; instead they are

often perpendicular as a result of constraints in the matrix, which oppose the

shape change resulting from the transformation. This is one of the reasons why

it is difficult to identify precisely habit planes in ferrous martensite.

Perhaps the most striking advances in the structure of ferrous martensites

occurred when thin foil electron microscopy was first used on this problem.

The two modes of plastic deformation are needed for the in-homogeneous

deformation part of the transformation, i.e. slip and twinning. All ferrous

martensites show very high dislocation densities of the order of 1011 to 1012

cm-2, which are similar to those of very heavily cold-worked alloys. Thus it is

usually impossible to analyze systematically the planes on which the

dislocations occur or determine their Burgers vectors.

The lower carbon (<0.5% C) martensites on the whole exhibit only

dislocations. At higher carbon levels very fine twins (5-10 nm wide) commonly

occur. In favorable circumstances the twins can be observed in the optical

microscope, but the electron microscope allows the precise identification of

twins by the use of the selected area electron diffraction technique. Thus the

twin shears can be analyzed precisely and have provided good evidence for the

correctness of the crystallographic theories discussed above. However,

twinning is not always fully developed and even within one plate some areas

are often untwined. The phenomenon is sensitive to composition.

The evidence suggests that deformation by dislocations and by twinning are

alternative methods by which the lattice invariant deformation occurs. From

general knowledge of the two deformation processes, the critical resolved shear

stress for twinning is always much higher than that for slip on the usual slip

plane. This applies to numerous alloys of different crystal structure.


28

Thus it might be expected that those factors, which raise the yield stress of the

austenite and martensite, will increase the likelihood of twinning. The

important variables are:

- carbon concentration;

- alloying element concentration;

- temperature of transformation;

- strain rate.

The yield stress of both austenite and martensite increases with carbon content,

so it would normally be expected that twinning would, therefore, be

encouraged. Likewise, an increase in the substitutional solute concentration

raises the strength and should also increase the incidence of twinning, even in

the absence of carbon, which would account for the twins observed in

martensite in high concentration binary alloys such as Fe-32%Ni.

A decrease in transformation temperature, i.e. reduction in Ms, should also help

the formation of twins and one would particularly expect this in alloys

transformed, for example, well below room temperature.

It should also be noted that carbon concentration and alloying element

concentration should assist by lowering Ms. As martensite forms over a range

of temperatures, it might be expected in some steels that the first formed plates

would be free of twins whereas the plates formed nearer to Ms would more

likely be twinned.

However, often plates have a mid-rib along which twinning occurs, the outer

regions of the plate being twin-free. This could possibly take place when the Ms

is below room temperature leading to twinned plates which might then grow

further on resting at room temperature.


29

2.3.4.5. Lath Martensite

The lath martensite structure is one of the most important structures in steels. It

is composed of fine substructures, i.e. "packets" which are a group of laths with

almost the same habit plane, and "blocks" which contain a group of laths with

almost the same orientation (figure 2-12). A prior austenite grain is divided by

several packets which are subdivided by blocks. It was recently shown that the

blocks are further subdivided by sub-blocks in low carbon steels [29].

Figure 2-12- OM images (3% nital etched) of lath martensite structures in the Fe-

0.2C-2Mn alloy: a) prior austenite grain size is 370µ m and b) 28µ m, respectively

[30].

This type of martensite is found in plain carbon and low alloy steels up to about

0.5 wt% carbon. The morphology is lath-like, where the laths are very long.

These are grouped together in packets with low angle boundaries between each

lath, although a minority of laths is separated by high angle boundaries. In plain

carbon steels practically no twin-related laths have been detected [26].

Since these packet and block boundaries are high angle boundaries, the

constituents are considered to be affective grains. Thus, the strength and

toughness of lath martensitic steels are strongly related to packet and block


30

sizes. It is known that both the block width and the packet size are proportional

to the prior austenite grain size. Usually, the packet size is taken as the

effective grain size for the strength and toughness of low carbon steels [30].

2.3.4.6. Medium Carbon Martensite

It is perhaps unfortunate that the term acicular is applied to this type of

martensite because its characteristic morphology is that of perpendicular plates,

a fact easily demonstrated by examination of plates intersecting two surfaces at

right angles (figure 2-13).

These plates first start to form in steels with about 0.5% carbon, and can be

concurrent with lath martensite in the range 0.5 %-l.0 % carbon. Unlike the

laths, the lenticular plates form in isolation rather than in packets, on planes

approximating to {225} and on several variants within one small region of a

grain, with the result that the structure is very complex [26].

Figure 2-13- Plate-like martensite microstructure.


31

2.3.5. An Overview of the Previous Investigations on the Thermomechanical Behavior of Boron Steels

Somani et al. [31] examined the effects of plastic deformation on dilatation

during the martensitic transformation in a B-bearing steel (i.e. Docol Bo 02)1.

Their results show that plastic deformation of austenite at high temperatures

enhances ferrite formation significantly and consequently, the dilatation

decreases markedly even at a cooling rate of 280°C/s.

It was found that, without plastic deformation, Ms and Mf were about 425°C

and 280°C, respectively. The change in diameter was about 0.53%

corresponding to a relative volume change of 3.2%. They mentioned that the

reason for the drastic decrease of dilatation and drop of the Ms value to 375°C

due to an increase in the prior plastic strain could be justified as a result of the

stabilization of austenite by means of plastic deformation and the presence of

retained austenite in this regard. There were, however, distinct differences in

the high temperature slopes of the dilatation curves. The slope in the deformed

specimens being smaller than that in non-deformed ones. This presumably

indicates that some ferrite formed at higher temperatures as strain-induced,

consequently, less martensite is present.

Microstructures examination also revealed that, at a cooling rate of 50°C/s, the

ferrite content was about 20~40%. Hardness measurements also confirmed that

the structure formed after severe plastic deformation was markedly softer,

about 295~375 HV10, compared to the martensite hardness of 490~500HV10.

However, martensite was still present in considerable amount, even though the

dilatation became very small. Therefore, they suggested that some other

factors, such as residual stresses due to prior plastic deformation may be an

additional reason for the decrease of dilatation.

Finally they found that, the severe plastic straining (strain 0.8~1.0) during

continuous cooling at 50°C/s results in a much lower final flow stress level

(800-950MPa at 300~200°C) than that obtained for martensitic structure in

isothermal tests (1650~1900MPa).

1 The chemical composition in wt% is: 0.22C- 0.29Si- 1.1Mn- 0.013P- 0.003S- 0.21Cr- 0.0034B- 0.05 Al- 0.0025N.


32

Another investigation by the mentioned authors [5] revealed that the Ms

temperature is lowered by about 25-70°C with increasing plastic strain from

0.16-0.39. As the reason for this, they proposed that, as a consequence of

ferrite formation, carbon becomes enriched in the remaining austenite, which

therefore transforms into martensite at a somewhat lower temperature.

It was also observed that ferrite with an ultra-fine grain size can be formed as

strain-induced by subjecting austenite to severe plastic straining at temperatures

slightly above Ar3. Hardness measurements also confirmed that the

microstructure formed after a high-temperature plastic deformation was

remarkably softer, 302-440 HV10, while the martensite had a hardness of 490-

510 HV10, which was justified due to the presence of ferrite in the

microstructure as described before.

In case of the hardness measurements, their image analysis data were in

contrast with their hardness values and made it rather impossible to determine

the martensite or bainite phases based on the optical microscopy images. They

believed that the distinction between the bainite and martensite phases might

require transmission electron microscopic examinations, which had, however

not been performed, because this matter was not very important in their

discussions.

To avoid the strain-induced phase transformation, it was suggested that the

consequences of the prior plastic deformation should be small enough or

disappear before the temperature reaches the ferritic regime level. This means

that forming should take place at a high temperature, >800°C, where the

driving force for the austenite decomposition is low, or the time should be long

enough for static recrystallisation to occur. Another, more realistic alternative

might be forming at low temperatures, such as <600°C, i.e. below the ferrite

regime. In that case, ferrite nucleation is not accelerated, although some

enhancement of bainite formation may occur. This may not be so detrimental,

however, due to the notably smaller strength difference between martensite and

bainite. Furthermore, in order to avoid straining to continue at the martensitic

stage, which would mean excessive forming loads, a major springback and high


33

residual stresses, forming should be finished above 420°C, which means that

the proper temperature range is quite narrow. Overall, it was proposed that,

minimization of the plastic strain, maximization of the cooling rate and/or

forming at 450-600°C may be suitable ways to avoid excessive ferrite

formation and to achieve the desired mechanical properties in formed and

quenched components [5].

In the last reviewed work here, Jun and coworkers [32] studied the effects of

thermomechanical processing on the microstructures and transformations of

low carbon HSLA steels with and without boron. Microstructures observed in

continuous cooled specimens were composed of pearlite, quasi-polygonal

ferrite, granular bainite, acicular ferrite, bainitic ferrite, lower bainite, and

martensite depending on cooling rate and transformation temperature. Fast

cooling rate depressed the formation of pearlite and quasi-polygonal ferrite,

which resulted in higher hardness. However, hot deformation slightly increased

transformation start temperature, and promoted the formation of pearlite and

quasi-polygonal ferrite. Hot deformation could also strongly promote the

acicular ferrite formation which was not formed in non-deformation condition.

Small boron addition effectively reduced the formation of pearlite and quasi-

polygonal ferrite and broadened the cooling rate region from bainitic ferrite and

martensite. Impurity boron segregates to grain boundaries and improves the

grain boundary cohesive strength. This causes the mentioned effective

suppression of pearlite and/or ferrite formation compared to other substitution

elements. Microhardness of granular bainite varied from 220 to 250 HV, which

resulted from high dislocation density and hard constituents. Transformation

mechanism of these bainite-like microstructures had both aspects of diffusional

and shear mechanisms. It was suggested that granular bainite forms because

carbon quickly diffuses away from the ferrite/austenite interface at relatively

slow cooling rates, preventing the formation of cementite. The increased

carbon content in the remaining austenite can stabilize austenite from further

transformation, and this entrapment of residual austenite leads to granular


34

bainite morphology. Shear mechanism for bainite-like transformation was

proposed to be more dominated as increasing cooling rates.

It was also found that the deformation causes the formation of acicular ferrite,

pearlite, and quasi-polygonal ferrite, otherwise prevents the martensite

compared to that of non-deformed condition. The corresponding transformation

curves of deformed CCT moves toward left side compared to those of non-

deformed CCT [32].

Chapter 3 – EXPERIMENTAL PROCEDURE

35

CHAPTER

THREE

EXPERIMENTAL PROCEDURE

In the current chapter, firstly the investigated material is introduced. Then the

complete set of the experimental conditions and parameters are discussed.

3.1. Material Characterization

3.1.1. Chemical Composition

The studied material is 22MnB5 steel. This is a hot rolled boron steel in form

of the plates produces by Benteler company, Germany. The manufacturer calls

the mentioned product as BTR165 steel. The chemical analysis of the

investigated steel is given in table 3-1.

Table 3-1- Chemical composition of the experimental alloy, wt-%.

C Si Mn P S Cr Ti B

0.24 0.27 1.14 0.015 0.001 0.17 0.036 0.003

3.1.2. Microstructure

This steel contains ferrite and pearlite phases (together with carbide) in

as-received condition. Figure 3-1 shows the microstructure of the BTR165 steel

in the rolling direction.


36

a b

Figure 3-1- Microstructure of the as-received BTR165 sheets in the rolling direction

a) 500X, and b) 1000X.

The image analysis data shows that the microstructure contains around 78%

ferrite besides 22%, combination of pearlite and carbide. Ferrite grain size was

measured to be comparable with 11 ASTM grain size standard.

3.1.3. CCT Diagram Design

The Continuous Cooling Transformation (CCT) diagram, figure 3-2, has been

produced by means of dilatometry tests, metallographic investigations and

hardness measurements. The circled numbers indicate the final hardness values

in the HV10 scale.


37

Cooling rate = 30°C/s

Figure 3-2- CCT diagram of BTR165 steel.

For the heating speed of 5°C/s, the eutectoid reaction temperature, Ac1, is

722°C and the start temperature of austenite to primary ferrite transformation,

Ac3, reaches 870°C. After austenization at 900°C for five minutes followed by

quenching the microstructure becomes fully martensitic (point M in figure 3-2).

Consequently, the steel is classified in the ultra high strength steels grade. The

martensite start point, Ms, lies at 410°C and the martensite finish point, Mf, at

230°C. It can be seen that a cooling rate greater than 30°C/s results in a

martensitic microstructure. At the lower cooling rates, bainite (zone B in figure

3-2) or even ferrite (zone F in figure 1), can be formed resulting in the lower

hardness and the lower strength levels.

3.2. Isothermal and Non-isothermal Compression Tests

In the hot stamping process, the material is subjected to a temperature history

of heating and subsequent high cooling rate to ensure the formation of

martensite. Since thermo-mechanical history of the material will affect its final

microstructure and properties, appropriate deformation and temperature

histories must be applied when carrying out the material testing. A

Baehr DIL 805 deformation dilatometer (figure 3-3) was used to create the


38

thermo-mechanical schedules. Such conditions were produced by several

isothermal compression and quenching tests as well as simultaneous forming

and quenching tests at temperature range between 600°C – 900°C. Different

strain rates between 0.1 s-1 – 10.0 s-1 for isothermal and 0.07 s-1 – 0.4 s-1 for

non-isothermal tests were applied. Due to the technical limitations in the Baehr

deformation dilatometer, higher strain rates and lower temperature limits could

not be applied in the simultaneous forming and quenching tests. In all of the

experiments, the samples were austenitized at 900°C for five minutes and

quenched to the compression temperature by 50°C/s. The above mentioned

processes are illustrated in figure 3-4. The yield and maximum stress, Ms and

Mf, were received from the dilatation tests.


39

Figure 3-3- Baehr DIL 805 deformation dilatometer, and the sample set up.


40

0 100 200 300 400 500 600 700 8000

200

400

600

800

1000

600°C650°C700°C750°C800°C

900°C

Isothermal deformation at constant temperatures for different strain rates:

strain rates = 0.1, 1.0 amd 10.0s-1

50°C/s

50°C/s

200°C/min

900°C, 5'Te

mpe

ratu

re (°

C)

Time (s)

0 100 200 300 400 500 600 700 8000

200

400

600

800

1000

50°C/s

Tem

pera

ture

(°C

)

Time (s)

Simultaneous forming and quenching strain rates = 0.07 - 1.0s-1

200°C/min

900°C, 5'

850°C

600°C

a b Figure 3-4- Schematic illustration of a) isothermal deformation, and

b) simultaneous forming and quenching experiments in the current work.

The experimental procedures were as follows: inserting the cylindrical

specimen (as shown in figure 3-5) in a vacuum chamber, resistance heating to

austenization temperature and performing the subsequent compression between

SiN2 anvils prior to controlled cooling. Molybdenum foils were used to prevent

the specimens to be pasted to the anvils, and the glass powder was utilized for

lubrication. The Pt/Pt-Rh10% thermocouple was welded to the specimen in

order to measure the temperature. The atmosphere was initially protected by

vacuum and then argon and helium shower were employed for a controlled

cooling.

5,0mm±0,1mm

10,0mm

±0,1mm

0,3mm±0,1mm4,0mm±0,1mm

0,3mm±0,1mm

Figure 3-5- Schematic illustration of the cylindrical specimen used during the

dilatation experiments.


41

3.3. Hardness and Metallography Tests

The deformed specimens in the dilatometry machine, were then cut off and sent

for the metallography and hardness (HV10 scale) measurements.

The microstructural images besides the image analysis data were extracted and

compared with the resulted hardness and mechanical data (dilatation tests).

Chapter 4 – RESULTS AND DISCUSSION

41

CHAPTER

FOUR

RESULTS AND DISCUSSION

4.1. High Temperature Isothermal Compression Tests

In this part of the chapter, the results of the high temperature isothermal

compression tests on the 22MnB5 steel (BTR165) are presented and analyzed.

The main topic has been divided into three sections: hardness and

microstructure, deformation, and dilatation data analysis.

4.1.1. The Hardness Values and Microstructural Variations

Firstly, the variations of the final hardness data (as quenched) by taking the

strain rates as the constant quantities were examined among different

isothermal deformation temperatures from 500°C to 900°C. As the second step,

the changes of the final hardness value were investigated based on the constant

deformation temperature by increasing the strain rates from 0.1 to 10.0s-1 at two

selected isothermal deformation temperatures, i.e. 750°C and 900°C. The

relevant microstructural images are presented for comparison.

4.1.1.1. Results

Figure 4-1, 4-2 and 4-3 show the variations of the final hardness values due to

the different deformation temperatures in three different constant strain rates.


42

500 600 700 800 900

320

340

360

380

400

420

440

460 451452442

409

368

HV

~

Deformation Temperature (°C)

Isothermal Deformation Temperature Increase

Strain Rate = 0.1s-1

Hardness Variations

329

Figure 4-1– Hardness variations by increasing the deformation temperature at the

strain rate of 0.1s-1.

500 600 700 800 900250

300

350

400

450

500

550

600

475

543

469

413

352

325

320

HV

~



Strain Rate = 1s-1

Hardness Variations

405




43

600 700 800 900250

300

350

400

450

500

402421

472

417

310

HV

~



Strain Rate = 10s-1

Hardness Variations



The diagrams demonstrate a continuous increasing trend at the strain rate of

0.1s-1 and the same trend with two decreasing sections at the lowest and the

highest deformation temperatures at the strain rate of 1.0s-1. At the strain rate of

10.0s-1, the increasing trend with some deviations from a sharp rise is observed.

A rather related effect, so-called “adiabatic heating”, at the strain rate of 10.0s-1

is presented in figure 4-4. As is seen, there is a sudden increase in the

temperature value (10-20 degrees), as the amount of exerted force is increased

by continuing the deformation process.

The variations of the hardness values regarding the increase of the strain rates

at the constant deformation temperatures of 750°C and 900°C are observed in

figures 4-5 and 4-6.


44

Figure 4-4 – A sample Force-time-Temperature (F-t-T) diagram of the 22MnB5

specimens regarding the isothermal deformation at 750°C and the strain rate of 10.0s-1

demonstrating the adiabatic heating effect.

400

405

410

415

420

425

430

435

440

445

450

1010.1

413

442

Strain Rate Increase (0.1~10s-1)Isothermal Deformation Temperature = 750°C

Hardness Variations

402

HV

~

Strain Rate (s-1)

Figure 4-5- Hardness variations by increasing the strain rate at the deformation

temperature of 750°C.


45

440

450

460

470

480

490

500

510

520

530

1010.1

451

475

Strain Rate Increase (0.1~10s-1)Isothermal Deformation Temperature = 900°C

Hardness Variations

521

HV

~

Strain Rate (s-1)

Figure 4-6- Hardness variations by increasing the strain rate at the deformation

temperature of 900°C.

a b

c

Figure 4-7- The final microstructure of the isothermally deformed samples at 750°C

by different strain rates; a) 0.1s-1, b) 1.0s-1 and c) 10.0s-1.


46

a b

c d

e f

Figure 4-8- The final microstructure of the isothermally deformed samples by the

strain rate of 0.1s-1 at different temperatures; a) 500°C, b) 650°C, c) 700°C, d) 750°C,

e) 800°C and f) 900°C.

Increasing the isothermal deformation temperature by the strain rate of 0.1s-1

results in the increase of martensite phase percentage in the final

microstructures which are shown in figure 4-8.


47

a b

Figure 4-9- The microstructure of the isothermally deformed samples by the strain

rate of 1.0s-1 at different temperatures; a) 550°C, b) 600°C.

Figure 4-9 gives the microstructure of the deformed samples at 550°C and

600°C by the strain rate of 1.0s-1. More ferrite could be observed in 600°C

sample than 550°C.

4.1.1.2. Discussion

Due to figure 4-1, a strain rate of 0.1s-1 leads into a residence time of five

seconds for the samples during the isothermal deformation before reaching the

final strain of 0.5 at the experimental temperature. Hence, this amount should

be added to the applied time for decreasing the temperature of the samples from

the austenization temperature (900°C) to the deformation temperature by the

cooling rate of 50°C/sec. Later, in section 4.1.2, it is seen that the mentioned

cooling rate prevents the samples from entering the isothermally formed ferrite

phase region during the experiments.

At 500°C, the microstructure mostly (more than 90%) contains the bainite

phase in addition to less evident areas of ferrite and martensite (figure 4-8a). It

could be assumed that during cooling, firstly the sample has entered the

continuously cooled ferrite phase region (see section 2.3.2.3). Consequently, by

further cooling to 500°C (from ferritic transformation temperature), and by

staying at this temperature and performing the deformation for five seconds,

the continuously cooled and isothermally formed bainitic transformations have


48

been started and accelerated (this complies with [15]). The final cooling below

the Ms temperature transforms the remained austenite into the martensite.

Therefore, the reported hardness value in the current sample is due to the effect

of the bainite phase instead of martensite phase in the microstructure.

Increasing the deformation temperature to 650°C results in the deformation

temperature to come out of the continuously cooled bainitic and ferritic

transformation regions; hence the sample is isothermally deformed in the

austenite phase region. By performing the isothermal deformation of austenite,

the nucleation sites of the continuously cooled bainite and ferrite phases are

increased. The jump of the ferrite phase percentage in the microstructure of

650°C sample in comparison with 500°C sample could be described based on

the probable more sensitivity of ferrite forming mechanisms to have more prior

nucleation sites produced by prior deformation (figure 4-8).

By the end of the deformation process and by further cooling, the austenite is

partially transformed to continuously cooled bainite and ferrite phases, while

more amount of it is transformed to martensite. This is because the residence

time is not enough for the isothermal bainitic transformation to occur. As a

result, the hardness value is increased by the increase of the martensite phase

percentage.

By increasing the deformation temperature, the ferritic transformation intensity

is gradually decreased; therefore besides the stability of the bainite phase

percentage, the amount of the final martensite is increased (figure 4-8). This is

because of the matter that the effect of prior deformation on generating the

nucleation sites of ferritic transformation is diminished as the deformation

temperature is increased. This fact can be described by the evaluation of the

700°C and 750°C samples. The 750°C sample mostly consists of bainite and

martensite phases and a very less amount of ferrite (figure 4-8 'c' and 'd'). The

light colored areas in the microstructure were determined as the bainite phase

due to the presence of the distributed carbides inside them.

By increasing the temperature to 800°C and 900°C, the final microstructure is

fully martensitic (figure 4-8 'e' and 'f'). This shows that the nose of the bainitic


49

and ferritic transformations in the CCT diagram has no coincidence with the

cooling line during the cooling process, although there is a deformation

residence time in the middle of the continuous cooling.

In figure 4-2, it takes seven seconds for the sample to reach the experimental

temperature of 550°C by the cooling rate of 50°C/sec. from the austenization

temperature of 900°C. It should be noted that the specimen enters the bainite

phase region during cooling before the deformation to be applied. The

combination of the time used for the cooling of the sample in the bainitic phase

region (which is around three seconds) and the residence time of the sample

during the deformation, i.e. 0.5 seconds (for the deformation up to the strain of

0.5 by the strain rate of 1.0s-1), leads into the bainitic transformation to be

developed to some extents. It means that the resulted bainite is produced mostly

by staying the sample at the bainitic phase region during cooling (continuously

cooled bainite) than the acceleration of the transformation by the isothermal

deformation (isothermally formed bainite). Looking into the microstructure,

figure 4-9a, more than 90% martensite is observed. This means that despite the

presence of two bainite forming mechanisms (isothermally formed and

continuously cooled), the bainitic transformation is not developed under current

experimental conditions. This shows that the lower deformation temperatures

may lead into the formation of more martensite by hindering the bainitic and

ferritic transformations.

Increasing the deformation temperature to 600°C, and coming out of the

bainitic transformation region, continuously cooled ferrite formation occurs

(due to the observed microstructural results in figure 4-9b), and by this, the

final hardness value of the sample is decreased. At the same time, the stay of

the sample in the bainitic phase region during cooling leads into the formation

of the continuously cooled bainite phase.

By further increase of the deformation temperature to 750°C, two factors have

a competitive effect on the hardness values. Firstly, the deformation of the

sample out of the continuously cooled bainitic and ferritic phase regions which

results in the presence of more austenite in the microstructure of the


50

experimental sample before the martensitic transformation temperature, which

means that the percentage of the martensite in the final structure and

consequently the hardness value is increased. Secondly, the formation

probability of the continuously cooled ferrite phase because of the isothermal

deformation, prior to the continuously cooled ferritic transformation

temperature, which results in the sample to enter a broad area of ferrite phase,

because of making more nucleation sites.

The dominant effect of the increase of the amount of martensite (as a harder

phase than bainite and ferrite) leads into the gradual increase of the final

hardness values of the samples.

By jump of the deformation temperature to the amounts more than 850°C to

900°C, the effect of the deformation on the stabilization of the austenite phase

prevents the sample to enter the ferritic and bainitic phase regions during the

cooling of the samples after the isothermal deformation. Therefore, in the 850

and 900°C samples, almost 100% martensite has been reported. The presence

of bainite phase in the 850°C sample could be justified due to the coincidence

of the cooling line and the bainitic transformation nose in the CCT diagram.

The decrease of the hardness value in the 900°C sample in comparison with

850°C (despite the presence of almost 100% martensite in both cases) could be

related to the difference of the martensitic structure in these two samples. This

is because in 850°C we have a 50°C decrease of the deformation temperature

from the austenization temperature (900°C), while in the case of deformation at

900°C, the deformation and the austenization temperatures are the same. It

could be imagined that the effect of the applied stress on the final martensitic

structure is diminished by a very high deformation temperature of 900°C in

comparison with 850°C. There are, however, other mechanisms such as

recovery and recrystallisation which may describe the same consequences.

In the case of isothermal deformation by the strain rate of 10.0s-1 (figure 4-3),

increasing the deformation temperature results in the increase of the hardness

value (because of an increase in the martensite content) by changing the place

where the sample enters the continuously cooled ferritic and bainitic phase


51

regions. This is the same as what is seen regarding the deformation by the

strain rates of 0.1 and 1.0s-1 (figures 4-1 and 4-2).

The related fluctuations in the increasing trend of the hardness diagram at the

strain rate of 10.0s-1 could be justified due to the presence of the temperature

increase (10-20°C) by means of the adiabatic heating1 (see figure 4-4) which

could change the location of the specimen in the phase regions and change the

phase percentage at higher deformation rates.

It takes three seconds for all the specimens to come from the austenization

temperature of 900°C to the deformation temperature of 750°C by the cooling

rate of 50°C/sec. Then there are 5, 0.5 and 0.05 - second periods of time for the

0.1, 1.0 and 10.0s-1 specimens respectively to be deformed isothermally at the

deformation temperature. Therefore, due to the mentioned residence times of

the samples, there is a more probability for the 0.1s-1 specimen to enter a

broader area of continuously cooled ferritic and bainitic phase regions than 1.0

and 10.0s-1 ones. The final microstructures show that the 0.1s-1 specimen is

heavily ferritic and bainitic and the remained parts have been transformed to

martensite (figure 4-7a). The reported hardness value (figure 4-5) is due to the

presence of more bainite beside the martensite in this specimen. In the 1.0 and

10.0s-1 specimens, the residence time at the deformation temperature is almost

the same (there is a 0.45 seconds difference between them). The more

important difference is a 10°C higher temperature in the case of 10.0s-1

specimen (because of the mentioned adiabatic heating) which in addition to a

shorter residence time (by 0.45 seconds) leads into the decrease of the bainite

and ferrite phases and the increase of the final martensite in the microstructure

of the isothermally deformed specimens.

The colorful role of the bainite in the hardness value of the 22MnB5 steel

results in the decrease of the hardness value by decreasing the bainite phase

percentage in the microstructure by increasing the strain rate from 0.1 to 10.0s-1

at the deformation temperature of 750°C.

1 Heat is generated by the plastic deformation at high strain rates. The heat generated in the material is either conducted and/or convected away to the surrounding or is used to increase the temperature of the material. When the heat generation rate is greater than the rate of heat loss, the temperature of the material is increased. Indeed, the heat dissipation is time dependant and therefore, in the low speed like in quasi-static processes, the heat can be transferred to the surrounding, dissipated and deformation occurs isothermally. During high-speed or fast enough processes, in the materials whose flow curves are temperature dependant, the flow stress is lowered simultaneously by the continuous rise of temperature due to adiabatic heating.


52

Regarding the deformed samples at 900°C (figure 4-6), all the specimens have

a rather full martensitic structure and the reported increase in the hardness data

could be justified due to the differences in the martensitic structure under

different deformation conditions and the presence of small amounts of bainite

phase which its structure and its arrangement beside a full martensitic matrix

may affect the final hardness values in different ways.

The hardness variation diagrams of the isothermally deformed specimens at

650°C and 700°C exhibit exactly the same behavior as for the 750°C specimen.

In addition, there is a same trend for the 800°C and 850°C specimens in

comparison with the 900°C specimen. Therefore, the relevant data is not

presented here, due to preventing further complications to occur. Based on

these trends, it would be possible to divide the whole experiments into two

high-deformation-temperature and low-deformation-temperature categories and

predict the response of the experimental material to the current heating and

deformation process regarding the temperature range at which the deformation

is done. At the lower deformation temperatures, the effect of any continuously

cooled bainite and ferrite phases has to be taken into account, while after the

high temperature isothermal deformation, the specimen does not enter the

mentioned transformation regions during cooling.


53

4.1.2. The Deformation Data Analysis

In this section, the whole resulted mechanical data from the dilatation

experiments, i.e. yield stresses, flow stresses, and work hardening rates are

presented and compared based on the variations of the experimental

temperatures and strain rates.

4.1.2.1. Results

Figure 4-10 shows the flow curves of the deformed samples at different

temperatures from 550°C to 900°C. It is seen that the stress level is

continuously decreased by increasing the deformation temperature.

0.0 0.1 0.2 0.3 0.4 0.5 0.6100

200

300

400

500

600

900°C850°C

800°C750°C700°C

650°C600°C

True

Stre

ss, σ

[MPa

]

True Plastic Strain, ε [-]

550°C

Figure 4-10- True stress – true plastic strain curves of the deformed samples based on

different deformation temperatures at the strain rate of 1.0s-1.

The flow curves of the deformed samples regarding different strain rates are

presented in figure 4-11. Coming from the highest to the lowest strain rates,

there is a drop in the value of the flow stress at 800°C.


54

0.0 0.1 0.2 0.3 0.4 0.5 0.60

50

100

150

200

250

300

350

0.1

1.0

True

Stre

ss, σ

[MPa

]


Strain Rate [s-1]10.0

Figure 4-11- True stress – true plastic strain curves of the deformed samples by

different strain rates (0.1, 1.0, and 10.0s-1) at 800°C.

Figures 4-12 gives the yield stress data due to the variations of the strain rate at

the constant selected deformation temperatures. The same data based on the

variations of the deformation temperature at the constant selected strain rates

are shown in figures 4-13 and 4-14. By increasing the strain rate, it is seen that

the yield stress value is increased, while by increasing the deformation

temperature, the mentioned value is decreased.


55

120

140

160

180

200

220

240

260

280

300

Deformation Temperature = 900°C

289

227

169

177

136

101

Yie

ld S

tress

(MPa

)

Strain Rate (s-1)0.1

Strain Rate Increase (0.1~10s-1

)

Isothermal Deformation Temperature = 750°C

Yield Stress Variations120

Deformation Temperature = 750°C

Figure 4-12– Yield stress variations by increasing the strain rate at the deformation

temperature of 750°C and 900°C.

400 500 600 700 800 9000

50

100

150

200

250

300

350

120

163169

194

218

Yie

ld S

tress

(MPa

)


Isothermal Deformation Temperature IncreaseStrain Rate = 0.1s-1

Yield Stress Variations

332

Figure 4-13–Yield stress variations by increasing the deformation temperature at the



56

400 500 600 700 800 9000

50

100

150

200

250

300

350

136156

166

227

232244

288

Yie

ld S

tress

(MPa

)


Isothermal Deformation Temperature IncreaseStrain Rate = 1s-1

Yield Stress Variations

300

Figure 4-14–Yield stress variations by increasing the deformation temperature at the


Work Hardening Rate (θ) Vs. StrainStrain Rate = 1.0 s Total Strain = 0.5

Deformation Start Temperature Variations

0

200

400

600

800

1000

1200

1400

1600

1800

0.00 0.10 0.20 0.30 0.40 0.50 0.60

True Strain, ε [-]

Wor

k H

arde

ning

Rat

e, θ

[Mpa

]

550°C

600°C

700°C

800°C750°C 850°C

900°C

650°C

-1

Figure 4-15– Work hardening rate variations versus true strain based on different

deformation temperatures; strain rate = 1.0s-1; total strain = 0.5.


57

The work hardening rate data based on the variations of the deformation

temperature and the variations of the strain rate are given in figures 4-15 and

4-16. The lowering effect of the deformation temperature increase on the work

hardening rate value is evident, while the mentioned value is not permanently

increased by increasing the strain rate.

Work Hardening Rate (θ) Vs. StrainDeformation Start Temperature = 800°C

Total Strain = 0.5Strain Rate Variations

0

200

400

600

800

1000

1200

0 0.1 0.2 0.3 0.4 0.5 0.6

True Strain, ε [-]

Wor

k H

arde

ning

Rat

e, θ

[Mpa

] Strain Rate = 1.0 s-1

Strain Rate = 10.0 s-1


Figure 4-16– Work hardening rate variations versus true strain based on different

strain rates; deformation temperature = 800°C; total strain = 0.5.

4.1.2.2. Discussion

In figure 4-12, the increase of the strain rate directly results in the increase of

the yield stress. This is because by increasing the strain rate from 0.1 to 10.0s-1,

a greater number of dislocations are produced at the beginning of the

deformation process. The sudden accumulation and hit of the mentioned

dislocations leads into a higher required stress level to enter the plastic state of

deformation. Based on the same note in section 4.1.1.2, the variations of the

yield stress by increasing the strain rate from 0.1 to 10.0s-1 at all the

deformation temperatures from 600°C to 900°C are the same. Therefore, here


58

only the 750°C and 900°C specimens were chosen as the group leaders of the

low- and high-deformation-temperature experiments.

Figures 4-13 and 4-14 show that rising the isothermal deformation temperature

leads into the decrease of the yield stress. This is because of more softness of

the structure and the activation of the temperature-dependent dislocation

motion mechanisms such as climb and cross slip.

By the evaluation of the work hardening rate diagrams, it is found that

increasing the isothermal deformation temperature leads into a drop in the work

hardening rate curve. The general observed trend of these diagrams can also be

found through the consideration of the flow curves behavior in figure 4-10. The

flow and work hardening rate curves clarify that the increase of the

deformation temperature ruins the work hardening phenomenon by the

activation of the dynamic recovery process which can also be found through

the smooth flow curves (figure 4-10) and the work hardening rate value which

reaches the zero state in the case of strains of more than 0.2-0.3.

Comparing the flow curves of the isothermally deformed samples at different

temperatures (figure 4-10), it is concluded that the increase of the deformation

temperature also evidently results in the drop of the stress level at the

beginning, during and the end of the plastic deformation. This is because the

dislocation motion as the most important factor for the deformation of the

metals becomes much easier as the temperature reaches the higher values.

At 750 and 800°C, the work hardening rate curves do not follow the overall

decreasing trend by increasing the temperature and this is mostly because of the

data scattering than an understandable metallurgical event due to the

temperature increase.

The competitive effect of temperature and exerted strain on the work hardening

rate behavior could be observed comparing figures 4-11 and 4-16. A higher

work hardening rate is expected for a higher strain rate at higher temperatures,

while the work hardening itself faces a drop as the temperature increases.

Simultaneously, a higher strain rate produces more dislocations in a shorter

period of time and forces them to hit each other. This leads in a higher work


59

hardening rate at the beginning of the deformation and the disappearance of

this behavior by further deformation1.

As the deformation is continued, more dislocations hit each other and by this,

cause the work hardening rate to be deeply decreased for the strains more than

0.2 and finally reach a zero value for the strains more than 0.4. Considering the

deformed sample by the strain rate of 1.0s-1 (see figures 4-11 and 4-16), it is

found that the mentioned specimen is in a transition condition between the

effect of the higher temperature on more deformed samples (i.e. by higher

strain rates) to make their work hardening rate to be increased, and the effect of

a higher strain rate on decreasing the work hardening rate level as the

deformation is continued (i.e. because of the dislocation accumulation which

prevents the work hardening rate to be increased). Therefore, in spite of a logic

trend in the flow curves behavior (i.e. the higher the strain rate, the higher the

flow stress level), the highest work hardening rate is observed in the 1.0s-1

specimen, although all of the work hardening rate curves reach a zero value due

to the dynamic recovery phenomenon at the strains of more than 0.4.

The point to be considered is that the increasing effect of the dynamic recovery

is completely ruined when the dynamic recovery rate overcomes the rate of

dislocation generation by further deformation. At such moment, almost no

dislocation is remained in the microstructure and all of the newly born

dislocations will be swept away by the dynamic recovery process. Therefore,

no change in the amount of work hardening is observed and finally the work

hardening rate will be zero.

1 Due to [33] and during the initial stages of deformation, there is an increase in the flow stress as dislocations interact and multiply. However, as the dislocation density rises, so the driving force and hence the rate of recovery increases and during this period, a microstructure of low angle boundaries and subgrains develops. At a certain strain, the rates of work hardening and recovery reach a dynamic equilibrium, the dislocation density remains constant and a steady-state flow stress is obtained. During deformation at strain rates larger than ~1.0s-1 the heat generated by the work of deformation cannot all be removed from the specimen and the temperature of the specimen rises during the deformation. This may then cause a reduction in the flow stress as straining proceeds. In modeling the high temperature deformation behavior, it is very important that such effects are taken into account as mentioned in 4.1.1.2.


60

4.1.3. The Dilatation Data Analysis

In this section, the variations of the martensitic transformation start temperature

(Ms), the changes of the temperature range in which the martensitic

transformation occurs (i.e. Ms-Mf), and finally the dilatation curves of selected

experiments are presented and compared.

4.1.3.1. Results

The variations of the Ms and (Ms-Mf) data at the constant strain rates based on

different deformation temperatures and the variations of the microstructural

states are given in figures 4-17 to 4-22. Both Ms and (Ms-Mf) values almost

demonstrate an increasing trend due to the increase of the deformation

temperature.

400 500 600 700 800 900200

250

300

350

400

450

378

378

378378

378

Ms(

°C)


378


Ms Variations

Figure 4-17- Ms variations regarding the deformation temperature at the strain rate of

0.1s-1.


61

400 500 600 700 800 900200

250

300

350

400

450

408

414

390

393

367

385384

357


Ms Variations

Ms(

°C)



1.0s-1.

400 500 600 700 800 900200

250

300

350

400

450

393

407

380

398

363


Ms Variations

Ms(

°C)



10.0s-1.


62

Table 4-1- The variations of the phase percentage in the microstructure of the

22MnB5 isothermally deformed specimens regarding different deformation

temperatures and strain rates (M=martensite, B=bainite, F=ferrite). The phase

percentage in the blank fields has not been reported. Def.

Temp. /

Strain Rate

500°C

650°C

700°C

750°C

800°C

900°C

0.1s-1

100%B

60%M 27%B 13%F

66%M

34%F

76%M

24%F

96%M 4%B

100%M

1.0s-1

46%M 43%B 11%F

82%M

18%F

81%M

19%F

86%M 14%B

98%M 2% B/F

10.0s-1

30%M 63%B 7%F

96%M

4%F

400 500 600 700 800 9000

50

100

150

200

250

180

142

155

7975

111



Ms -Mf Variations

Ms-

Mf(°

C)


Figure 4-20- Ms-Mf variations regarding the deformation temperature at the strain rate

of 0.1s-1.


63

400 500 600 700 800 9000

50

100

150

200

250

212210

148

186

125112

128

118



Ms -Mf Variations

Ms-

Mf(°

C)



of 1.0s-1.

400 500 600 700 800 9000

50

100

150

200

250

193

208

143

157

121



Ms -Mf Variations

Ms-

Mf(°

C)



of 10.0s-1.

The dilatation curves of the experimental samples regarding three different

deformation temperatures are given in figures 4-23, 4-24 and 4-25 due to the

increase of the strain rates. As the strain rate is increased, the martensitic

transformation valleys in the dilatation curves are shifted to the right hand side


64

of the diagram. This behavior is seen in all three different experimental

temperatures.

0 100 200 300 400 500 600 700

-0.4

-0.3

-0.2

-0.1

0.0




Cha

nge

in L

engt

h (%

)

Temperature (°C)


Figure 4-23- Dilatation curves of the 22MnB5 steel by increasing the strain rate from 0.05 to 10.0s-1; austenization temperature = 900°C;

deformation temperature = 650°C.

0 100 200 300 400 500 600 700-0.5

-0.4

-0.3

-0.2

-0.1

0.0



Cha

nge

in L

engt

h (%

)

Temperature (°C)


Figure 4-24– Dilatation curves of the 22MnB5 steel by increasing the strain rate from 0.1 to 10.0s-1; austenization temperature = 900°C;

deformation temperature = 800°C.


65

0 100 200 300 400 500 600 700-0.5

-0.4

-0.3

-0.2

-0.1

0.0




Cha

nge

in L

engt

h (%

)

Temperature (°C)

Figure 4-25– Dilatation curves of the 22MnB5 steel by increasing the strain rate from

0.1 to 10.0s-1; austenization temperature = 900°C; deformation temperature = 900°C.

4.1.3.2. Discussion

By increasing the deformation temperature, Ms is decreased at all the strain

rates. This phenomenon is in contrast with the increasing effect of the

austenization time and temperature on the Ms value (see figures 2-10 and

2-11 in chapter 2). This could be justified due to the fact that, the austenite

matrix uses the higher temperature as an accelerating factor besides the exerted

strain (driving force for the new grains to be born) to make a finer austenite

grain structure after the isothermal deformation and before quenching. As

higher austenization time and temperature means higher austenite grain size (in

comparison with smaller grains due to the isothermal deformation at higher

temperatures), decreasing the grain size leads into the decrease of the Ms value

as described before.

The amount of martensite in the final microstructures is also increased at the

higher deformation temperatures (table 4-1). It could be justified regarding the

fact that, continuously cooled bainitic and ferritic transformations are hindered

with higher deformation temperatures as described before. As this is the only


66

way of consuming the austenitic matrix, more austenite remains till the

beginning of the martensitic transformation. Therefore, the final martensite

content is increased by isothermal deformation at higher temperatures.

Considering table 4-1 and figures 4-20 and 4-21, it is found that, despite the

value of Ms-Mf somehow gives an overview of the martensite phase percentage

in the final microstructure, but the occurrence of the bainite and ferrite phases

and their amounts may ruin the correlation between these two quantities. In

most cases (but not all cases), it is seen that the presence of bainite and ferrite

phases in the microstructure can decrease and increase the value of Ms-Mf

respectively.

It is seen through comparing figures 4-20, 4-21 and 4-22 that the overall trend

of three diagrams is increasing and the strain rate of 10.0s-1 shows a rather

higher level than two others. It means that by increasing the isothermal

deformation temperature, the amount of martensite and the value of Ms-Mf are

increased. Also the higher level of figure 4-22 (strain rate of 10.0s-1) shows that

the increase of the strain rate can increase the time required to produce a certain

amount of martensite.

In figures 4-23, 4-24 and 4-25, it can be found that by increasing the strain rate

at three deformation temperatures (650, 800 and 900°C), the martensitic

transformation regions of the dilatometry curves are shifted to the higher

temperatures. This fact beside what is seen in figures 4-20 to 4-22, is because

the residual stresses as a result of the deformation of the sample in the

austenitic phase region (the long range elastic stresses) leads into the increase

of the martensitic transformation start temperature (Ms) and the effect is

increased as the amount of deformation is increased [28].


67

4.2. Simultaneous Deformation and Quenching Tests

In the second part of the current chapter, the results of the simultaneous

deformation and quenching tests will be discussed in detail. This part has been

divided into four sections including: deformation duration effect, strain rate

effect, austenization soaking time effect, and deformation start temperature

effect.

4.2.1. The Strain Magnitudes Effect (Deformation Duration Effect)

The strain rate and the start temperature of the experiments were fixed at 0.1s-1

and 800°C respectively and the durations of the deformation tests were set to

be 1, 2, 3, 4 and 5 seconds (with respect to the total strain values of 0.1, 0.2,

0.3, 0.4 and 0.5). As the deformation is started at 800°C, the sample is cooled

down by the cooling rate of 50°C/sec. The cooling process is continued till the

sample reaches the room temperature. Therefore, the compression tests are

performed simultaneously during the cooling process.

4.2.1.1. Results

Figure 4-26 shows the various dilatation curves of the different deformation

periods from one to five seconds at the deformation start temperature of 800°C

and the strain rate of 0.1s-1.


68

100 200 300 400 500 600 700-0.8

-0.7

-0.6

-0.5

-0.4

-0.3

-0.2

-0.1

0.0

Deformation Time = 1 sec.




Cha

nge

in L

engt

h (%

)

Temperature (°C)


Figure 4-26- Dilatation curves of the 22MnB5 steel; deformation start

temperature = 800°C; strain rate = 0.1s-1; deformation duration = 1-5 seconds.

As can be seen in this figure, increasing the amount of the deformation

(basically shown by increasing the duration of the deformation process from

one to five seconds), slightly retards the picks and shifts the transformation

sections of the curves to the lower temperatures. Furthermore, the depths of the

valleys in these curves are substantially decreased by increasing the

deformation time from one to five seconds.

The Ms and Mf values were found, and the extracted Ms and (Ms-Mf) data were

compared in figures 4-27 and 4-28 respectively.


69

0.0 0.1 0.2 0.3 0.4 0.5 0.6320

330

340

350

360

370

380

390

400

378

364361

382

374

Strain Value IncreaseStrain Rate = 0.1s-1

Deformation Start Temperature = 800°CMs Variations

Ms(°

C)

Final Strain, ε

Figure 4-27- Ms variations regarding the final strain value.

0.0 0.1 0.2 0.3 0.4 0.5 0.60

20

40

60

80

100

120

140

160

180

100

105

87

159 136


Deformation Start Temperature = 800°CMs-Mf Variations

Ms-M

f (°C

)

Final Strain, ε

Figure 4-28- Ms-Mf variations regarding the final strain value.

It can be easily seen that the (Ms-Mf) values are decreased by increasing the

amount of the strain from 0.1 to 0.3 (figure 4-28) followed by an increase for

the strains above 0.3, while initially there is a decreasing behavior for the Ms

values (figure 4-27) followed by a rising trend due to the strains higher than

0.3.


70

Figure 4-29 shows how the amount of the martensite in the as deformed sample

varies with respect to the exerted strain. The graph shows a drop in the amount

of the martensite in the matrix which is followed by an increase for the strains

larger than 0.3.

0.1 0.2 0.3 0.4 0.5

0

20

40

60

80

100

45

58

~0

15

Am

ount

of M

arte

nsite

(%)

Final Strain, ε

98


Deformation Start Temperature = 800°CAmount of Martensite (%)

Figure 4-29- The variations in the amount of martensite in percentage regarding the

final strain value.

The effect of an increase in the amount of the applied strain on the flow curve

behavior of the 22MnB5 samples are shown in figure 4-30.


71

0.0 0.1 0.2 0.3 0.4 0.5 0.60

50

100

150

200

250

300

350

400

450

500

Final Strain = 0.4


Temperature at the End of

Deformation = 600°C

Final Strain = 0.5




Final Strain = 0.3




Final Strain = 0.2




True

Stre

ss, σ

(MPa

)

True Plastic Strain, ε [−]

Final Strain = 0.1




Figure 4-30– Flow curves variations by increasing the applied strain from 0.1 to 0.5;

strain rate = 0.1s-1; deformation start temperature = 800°C.

Going from the shortest to the longest deformation duration, there is a

continuous increasing trend for the flow curves behavior.

A sample time-temperature-Force (t-T-F) diagram of a specimen which was

deformed by the strain rate of 0.1s-1 and the start temperature of 800°C for two

seconds is shown in figure 4-31. As is seen, the temperature curve starts from

800°C and ends at 700°C (at the cooling rate of 50°C/sec. for two seconds),

while the specimen simultaneously starts to receive the load at 800°C and is

unloaded at 700°C.


72

0

1000

2000

3000

4000

5000

6000

7000

8000

571.5 572.0 572.5 573.0 573.5 574.0 574.5 575.0

Time (sec.)

Forc

e (N

)

680

700

720

740

760

780

800

820

Tem

pera

ture

(°C

)

Temperature

Force

Figure 4-31- A sample t-T-F diagram of the experimental material;

Strain rate = 0.1 s-1; Start temperature = 800°C, Duration of the test = 2 seconds.

Figure 4-32 shows the microstructures of the experimental specimens after the

deformation process regarding their different deformation times from one to

five seconds.


73

a

b

c

Figure 4-32- Microstructure of the 22MnB5 samples after the simultaneous

deformation and quenching process; deformation start temperature = 800°C,

strain rate = 0.1 s-1; a) deformation time = 1 sec., b) deformation time = 3 sec.,

c) deformation time = 5 sec.


74

4.2.1.2. Discussion

When austenite is plastically deformed, residual stresses and lattice defects are

introduced. The residual stresses are principally long-range elastic stresses that

raise the start temperature Ms and lower As temperature. This effect increases

with increasing the magnitude of deformation and reaches a saturation value.

On the other hand lattice defects (i.e. short-range stresses) lower the martensite

finish temperature Mf and raise the Af temperature. In contrast, at temperatures

as high as 525°C, internal stresses are relieved; hence, the factors raising the Ms

will become lessened. However, there will still remain lattice defects that are

not annihilated by heating to 525°C. Some such defects can accelerate the

transformation below the Ms. This is why the amount of the martensite is

increased at a low degree of prior deformation. At higher amounts of

deformation, lattice defects which stabilize austenite are formed (carbon atoms

migrate to them) and lower the Ms and decrease the amount of martensite; that

is, the austenite is considerably stabilized. Finally, for more than 30%

deformation, only a further fluctuation of the carbon concentration occurs.

Consequently Ms is raised, the austenite becomes unstable and the amount of

martensite increases [28]. These facts are confirmed by the evaluation of

figures 4-27, 4-28 and 4-29.

Comparing the resulted data from figures 4-25 and 4-29, it is concluded that the

deep valley in the case of sample which was deformed for one second by the

strain rate of 0.1s-1 and the decrease of this depth by going to higher

deformation times (from one to three seconds), are due to the formation of the

higher amounts of martensite (98%) in the case of deformation for one second

and drop of this amount to almost 0% for the deformation up to three seconds.

The presence of the considerable martensite in the microstructure in spite of

less dilatation (i.e. dilatation curves of the deformed samples for four and five

seconds) has been discussed by Somani et al [31]. With respect to this work

and also to the microstructural evaluation of the current steel (figure 4-32), the

shallower depth of the valleys in the case of samples which were deformed for

four and five seconds can be justified regarding the presence of the other


75

effective parameters such as residual stresses due to prior plastic straining and

prior probable strain-induced ferrite formation which are able to decrease the

amount of dilatation as the deformation time is increased from three to five

seconds.

Considering the microstructural images, and later by the analysis of the flow

curves of the deformed specimens, also the possibility of the strain induced

and/or continuously cooled ferrite formation (instead of normal static austenite

to ferrite transformation at high temperatures), it is seen that besides the

observed granular bainite microstructure (which is evident by the presence of

carbide free regions in figure 4-32), almost significant amount of ferrite is

determined regarding the deformed samples for five seconds. This confirms the

latter justification for the decrease of dilatation as the strain induced

transformation or other disturbing mechanisms occur.

The extracted (Ms-Mf) values which were presented in figure 4-28 indirectly

show the amount of the remained austenite which can be transformed to

martensite by cooling down the sample below Ms. Also this confirms the data

which were discussed in figure 4-29. The less opportunity to produce

martensite, e.g. by the presence of the bainitic and ferritic transformations

(which leads into the less amounts of Ms-Mf in figure 4-28), the less martensite

is formed (figure 4-29).

The severe deformation of austenite prior to its transformation hinders the

growth of martensite, causing a reduction in the fraction of the transformation

in spite of an increased number of nucleation sites' density. In this regard, it has

been shown that, deformation during the thermomechanical processing of steels

also accelerates the rate of bainite reaction [15]. It was also demonstrated by

Samoni et al that the plastic deformation above Ms leads to strain-induced

ferrite formation [31]. These mean by doing the mechanical work on the

samples, we are changing the shape of the CCT diagram of the experimental

alloy in such a way that the areas of the bainitic and ferritic transformations are

shifted to the left hand side of the diagram. This leads into the fact that by

increasing the deformation time, the only way to reach a full martensitic


76

structure is by increasing the cooling rate to the amounts more than what were

examined here (50 °C/sec.) and probably out of the industrial possibility range

of cooling rate.

As is evident in figure 4-32, increasing the deformation time (or by other words

the amount of applied strain) results in an increase in the amounts of bainite

and strain induced ferrite phases in the microstructure, while a rather full

martensitic microstructure is achieved after the deformation process with the

lowest duration (the least amount of the exerted strain).

Figure 4-30 demonstrates that no phase transformation occurs during the

deformation for one to five seconds. This is revealed by a simple evaluation of

the slope of the flow curves that is not changed by further deformation. As the

final reported structures consist of bainite, ferrite and martensite, one can

conclude that whole transformations were taken place after the end of the

deformation process. However, the probability of ferrite formation prior to the

deformation must be taken into account.

Moreover, due to the simultaneous deformation and quenching nature of the

experiments, the most possible type of the produced bainite is 'granular'. This

fact was deeply introduced in chapter two. In each case, the observed ferrite

regions could be part of a carbide free granular bainite or an independent

continuously cooled or strain induced ferrite.

The difference between these types of ferritic microstructures is not clear at

least by means of optical microscopy during our experiments.

In figure 4-30, it is clear that, the bigger applied strains, the bigger work

hardening and more dislocation creation. Therefore, the flow curves were

monotonically increased.


77

4.2.2. The Strain Rate Effect

Taking the deformation start temperature as a constant value (800°C), the

variations of the strain rate and its consequent effects on the properties of the

experimental material are studied in this section.

4.2.2.1. Results

Figure 4-33 shows the flow curves of the deformed samples at 800°C regarding

their different strain rates.

0

100

200

300

400

500

600

0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35 0.40 0.45True Plastic Strain, ε [-]

True

Stre

ss,σ

(MPa

)

)s 0.4 = Rate(Strain 781°C = Temp.


)s 0.07 = Rate(Strain 693°C=Temp.

1-

1-

1-


)s 0.07 = Rate(Strain 622°C =Temp.

1-

1-

750°C = Temp. Final s 0.4= RateStrain -1




Figure 4-33- True stress – true plastic strain curves of the deformed samples by

different strain rates at 800°C; 0.07, 0.1, 0.2 and 0.4s-1.

As is seen in this figure, the value of the maximum stress increases by

decreasing the strain rate from 0.4 to 0.07s-1. In addition, the shape of the flow

curves in the samples which were deformed by higher strain rates, i.e. 0.2 and

0.4s-1 show not such a rising behavior as for the lower strain rates (i.e. 0.07 and

0.1s-1).


78

The variations of the maximum stress (the stress regarding the strain of 0.4)

and the hardness values have been demonstrated in figures 4-34 and 4-35.

There is a sharp drop in the amount of the maximum stress, while the as

quenched hardness value grows meaningfully when the strain rate increases.

0.05 0.10 0.15 0.20 0.25 0.30 0.35 0.40 0.45

300

350

400

450

500 Strain Rate Increase (0.07~0.4s-1)

Deformation Start Temperature = 800°CFinal Strain = 0.4

σ0.4 Variations

Final Temperature = 750°C



308

410

319

M

axim

um S

tress

, σ0.

4 (MPa

)

Strain Rate (s-1)

0.07

500


Figure 4-34- Maximum stress, σ0.4, variations by increasing the strain rate.

295

300

305

310

315

320

325

330

335

0.40.20.1


297

311



318

334

Final Temperature = 750°CStrain Rate Increase (0.07~0.4s-1)

Deformation Start Temperature = 800°CFinal Strain = 0.4

Hardness Variations

Strain Rate (s-1)

HV

~

0.07

Figure 4-35- Hardness variations by increasing the strain rate.


79

Figure 4-36, illustrates the microstructures of the deformed samples by the

strain rates of 0.07, 0.1, 0.2 and 0.4s-1. In all cases, considerable bainite and

martensite phases are observed.

a b

c d

Figure 4-36- Microstructures of the deformed samples by different strain rates;

a) 0.07s-1, b) 0.1s-1, c) 0.2s-1 and d) 0.4s-1. Start temperature = 800 °C,

final strain = 0.4.

Figure 4-37 shows the work hardening rate values of the samples deformed by

the strain rates of 0.07, 0.1, 0.2 and 0.4s-1 during the deformation process. After

the initial decreasing trend, all the graphs show a rather stable stage followed

by an increase at the end. The sample which was deformed by the strain rate of

0.07s-1 has the highest level among the others, while the 0.4 s-1 sample shows

the lowest level.


80

Work Hardening Rate (θ) Vs. Strain (ε)Deformation Start Temperature = 800°C

Total Strain = 0.4

0

500

1000

1500

2000

2500

3000

0 0.05 0.1 0.15 0.2 0.25 0.3 0.35 0.4 0.45


Wor

k H

arde

ning

Rat

e, θ

[MPa

]

0.07

0.1

0.4

0.2

Strain rate (s ) ~ 733°C

~ 586°C

~ 753°C 650°C

725°C

~763°C

777°C

~788°C

-1

Figure 4-37– Work hardening rate variations against true strain;

Start temperature = 800°C; Total strain = 0.4.

4.2.2.2. Discussion

The competitive role of the applied strain rate and the deformation temperature

affects the variations of the flow stress curves and the resulted work hardening

rate data presented in figures 4-33 and 4-37.

In the very initial stages of the flow curves in figure 4-33, the deformed

samples with the highest strain rates show the highest stress level while the

deformed samples with the lowest strain rates have the lowest levels. By

increasing the strain, this trend will be reversed. The mentioned behavior till

point A is normal, because a higher stress level is usually achieved with respect

to the higher strain rate.

The effect of the experimental temperature becomes more colorful as the

deformation continues. This is because by increasing the extent of the

deformation, the temperature distance among the deformed specimens by

different strain rates is increased. The higher the strain rate, the deformation is

finalized at a higher temperature. Due to this fact, the temperature level of the


81

deformed specimens by higher strain rates is higher than the lower strain rate

ones at a same strain. Therefore, from point A onward, the effect of the

temperature on the flow stress is going to be dominant. This is evident by

examining the flow curve behavior of the 0.07s-1 specimen which is

substantially increased as coming to the higher strains due to the much lower

temperature among the other tests at a same strain. As coming to point B, the

effect of the strain rate on the flow stress level is totally disappeared and a pure

temperature dependent order of the flow curves is observed. The lower the

temperature (i.e. at lower strain rates), the maximum stress is increased (figure

4-34).

Figure 4-36 shows that by increasing the strain rate from 0.07 to 0.4s-1 more

bainite and less martensite are observed in the final microstructures of the

deformed specimens. Almost no ferrite is evident in these pictures. This means

that, a simultaneous deformation and quenching has changed the shape of the

CCT diagrams of the specimens in a way that the nose of the bainitic

transformation is sharply shifted to the left hand side of the diagram. By this,

the higher strain rates normally mean that the higher amounts of bainite and

less martensite are formed.

An important point to be considered by the evaluation of the flow curves and

the microstructural images (figures 4-33 and 4-36) is that no transformation

occurs during the deformation by the strain rates of 0.1 to 0.4s-1, because there

is no change in the slope of the flow curves among these specimens. In

contrast, it is seen that at point B, the slope of the flow curve of the 0.07s-1

specimen is increased and this means that a new phase with a higher strength

than the austenite phase, i.e. bainite, is formed. A simple comparison of the

work hardening rate and the flow stress diagrams of this specimen around point

B results in defining the bainitic transformation start temperature to be around

600°C for it.

It could be imagined that the bainitic transformation is started somewhere

between point B and the strain of 0.3 (see and compare figure 4-33 and 4-37).

In the meanwhile, a pure temperature dependent stress level is transformed to


82

the combination of the lower temperature and the bainite phase effects on

increasing the flow stress level.

It is concluded that despite the role of temperature in determining the order of

the flow curves level at the strains of more than 0.2, the higher amounts of the

flow stress for the 0.07s-1 specimen can be emphasized by means of the bainitic

transformation effect.

Figure 4-35 compares the hardness values of the samples after quenching

regarding their different strain rates.

The criteria of the hardness values after quenching are as follow:

- The extent of the ferritic and bainitic transformations which consume

the austenite in the matrix and decreases the amount of the final

martensite in the structure;

- The extent of ferrite (soft phase) after quenching;

- The extent of bainite (rather hard phase) after quenching.

The increasing trend of the hardness values by increasing the strain rates in

spite of the decrease of the martensite content shows that the bainite phase has

a more powerful effect on changing the hardness values of the 22MnB5 steel in

these experiments than the martensite.

In figure 4-37, basically because of the constant cooling rate for all the

experiments (50°C/sec.) and the fact that the parameter which shows the time

effect is the strain rate, the deformed samples by the strain rates of 0.07 and

0.4s-1 show the lowest and the highest temperatures regarding an equal strain,

respectively.

Basically, it must be mentioned that the increase of the applied strain (i.e. by

continuing the deformation to the higher strains, or by a higher strain rate at a

same strain) increases the work hardening by the production of more

dislocations in the specimen. In contrast, the work hardening rate may be

decreased by means of further exerted strain. This is because as the material is

more deformed and work hardened, the changes of the work hardening value

are decreased. This is because the material is much stronger than before and in


83

spite of the increase of the work hardening, the work hardening rate (i.e. the

change in the work hardening value) is decreased.

Furthermore, the increase of the deformation temperature makes the work

hardening rate to be decreased by ruining the effect of the exerted strain due to

the softening effects of the higher temperatures.

The increasing trend of the work hardening rate from the strain of 0.1 to the

strain of almost 0.2 (in the case of 0.07s-1 sample) is due to the increase of

strain in the sample which is completely in the austenite phase. Moreover,

because of a low strain rate despite a high temperature, the increase of the

exerted strain has no decreasing effect on the work hardening rate; therefore by

increasing the strain in this area, the work hardening rate is increased.

Comparing the work hardening rates of the specimens till the strain of 0.3 in

figure 4-37, shows that the deformed samples by higher strain rates have a

lower work hardening rate due to a higher instant temperature at a same strain.

Also, the jump of the slope of the work hardening rate in the case of 0.07s-1

specimen after the strain of 0.3 is directly related to the occurrence of bainitic

transformation.


84

4.2.3. The Austenization Soaking Time Effect

In this section, the effects of three different austenization soaking times (i.e. 5,

10 and 15 minutes), on the consequent mechanical properties and

microstructural variations of the experimentally deformed 22MnB5 steel are

investigated.

4.2.3.1. Results

Figure 4-38 shows the flow curves of the deformed samples regarding their

different austenization soaking times. No meaningful deviation among the flow

curves is observed by increasing the soaking time from 5 to 15 minutes.

0.00 0.05 0.10 0.15 0.20 0.250

50

100

150

200

250

300

350

Austenization Time = 5 min.


True

Stre

ss, σ

[MPa

]

True Plastic Strain, ε [−]


Austenization Temperature = 900°CDeformation Start Temperature = 800°C


Final Strain = 0.2

Figure 4-38– Flow curves of the deformed samples regarding their different

austenization soaking times.

The dilatation results of the mentioned samples are demonstrated in figure

4-39. As can be easily seen, the locations of the valleys in terms of the

temperature have not been changed sharply in the case of austenization for 10

and 15 minutes, while it has been shifted to lower temperatures regarding the

shortest austenization soaking time.


85

100 200 300 400 500 600 700-0.5

-0.4

-0.3

-0.2

-0.1

0.0



Cha

nge

in L

engt

h (%

)

Temperature (°C)


Figure 4-39– Dilatation curves of the 22MnB5 steel by increasing the austenization

soaking time from 5 to 15 minutes; austenization temperature = 900°C; deformation

start temperature = 800°C; strain rate = 0.1s-1.

The amount of the phases in the as quenched microstructure of the samples and

the variations of the hardness values are shown in figure 4-40 which includes

the effect of the austenization soaking time. The amount of martensite is

increased while the amount of bainite and ferrite are decreased by increasing

the austenization soaking time from 5 to 15 minutes. The hardness values show

a drop followed by a rise during the same variations of the soaking time.


86

22

60

29

6

25

3

75

15

65

432

363.3385.7

0

20

40

60

80

100

120

140

160

180

200

0 5 10 15 20

Austenization Soaking Time (min.)

Phas

e Pe

rcen

t (%

)

0

50

100

150

200

250

300

350

400

450

500500

450

400

350

300

HV

~

martensite

bainiteferrite

Figure 4-40– Distribution of the microstructural phases and the hardness values

in terms of the austenization soaking time.

Figure 4-41 demonstrates the variations of the work hardening rate for the

above mentioned samples. As can be seen in this figure there is no evident

difference in the work hardening rate by changing the austenization soaking

time from 5 to 15 minutes.


87

Work Hardening Rate (θ) Vs. Strain (ε)Deformation Start Temperature = 800°CStrain Rate = 0.1 (s ),Total Strain = 0.2

Temperature at the end of deformation = 700°C

0

500

1000

1500

2000

2500

3000

3500

0 0.05 0.1 0.15 0.2 0.25


Wor

k H

arde

ning

Rat

e, θ

[MPa

]




-1

Figure 4-41– Work hardening rate variations of the deformed samples regarding

different austenization soaking times.

Figure 4-42 confirms the presence of the mentioned phases in figure 4-40 by

the demonstration of the microstructural images of the deformed samples

regarding the austenization soaking time.


88

a b

c

Figure 4-42– Microstructures of the deformed samples regarding different

austenization soaking times: a) 5 minutes, b) 10 minutes and c) 15 minutes.

4.2.3.2. Discussion

Basically, it can be imagined that increasing the austenization soaking time

may lead into the shift of the transformation curves to the right hand side of the

CCT diagram and retarding the nucleation and growth based transformations

such as bainitic and ferritic transformations. This means that further

austenization hinders the ferrite and bainite phases to occur. In addition,

considering the stability of the flow curves despite an increase in the

austenization soaking time (figure 4-38), and the fact that the deformation

process in all three cases has been finalized at 700°C, it can be concluded that

the mentioned temperature is higher than the ferritic transformation start

temperature. Therefore, all the specimens have been deformed based on the

same conditions in the austenitic region of the CCT diagram and no

transformation occurs during the deformation process. Moreover, it is found

that the probable change of the austenite grain size by increasing the


89

austenization soaking time has no colorful effect on the flow curve behavior of

the specimens.

Regarding the dilatation curves (figure 4-39), normally it is expected that

increasing the amount of produced martensite, results in increasing the depth of

the transformation valley in the dilatation diagram. As is seen in figure 4-39,

this trend is true due to the increase of the austenization soaking time from 5 to

10 minutes. Nevertheless, the decrease of the dilatation despite the increase of

the austenization soaking time (from 10 to 15 minutes) could be justified by

means of the other effective parameters – e.g. internal stresses – than the

occurrence of a more martensitic structure (in 15 minutes samples than 10

minutes ones) which normally increases the depth of the valleys in the

dilatation curves.

In addition, it is observed that decreasing the austenization soaking time to

5 minutes gradually shifts the martensitic transformation sections of the

dilatation diagrams to the left. It means that, shorter austenization time retards

the martensitic transformation, and a higher driving force (i.e. lower Ms) is

required to start the transformation.

This is because of the matter that, at higher austenization times, there is more

opportunity for the growth of the austenite grains. Therefore, the austenite to

martensite transformation becomes much easier. In other words, ferritic and

bainitic transformation are less probable to occur based on the same cooling

rate in comparison with the shorter austenization times. Hence, higher cooling

rate and/or lower Ms are required for the martensitic transformation to happen

after the austenization for a shorter period of time.

By the evaluation of the hardness data and the distribution of the different

phases in the microstructure (figures 4-40 and 4-42), it is found that one of the

most effective parameters for the hardness values to be defined is the amount of

bainite phase in the microstructure, although the amount of martensite has its

own natural effect on the mechanical properties. Due to this fact, despite an

increase in the amount of martensite (from 15 to 65%) and decrease of ferrite

(from 25 to 6%) by increasing the austenization soaking time from 5 to 10


90

minutes, the hardness value is decreased because of a decrease in the amount of

bainite phase (from 60 to 29%) in the final microstructure. As the amounts of

ferrite and bainite phases in the microstructures are almost stable (by increasing

the soaking time from 10 to 15 minutes), increasing the amount of martensite

from 65 to 75% leads into an increase in the hardness value.

There is no meaningful difference among the work hardening rate values of the

deformed samples regarding their different austenization soaking times as

shown in figure 4-41. This is because of the same reason as discussed in figure

4-38.


91

4.2.4. The Deformation Start Temperature Effect

The variations of the deformation start temperature and its effect on the

mechanical properties of the experimental material are studied here.

Furthermore, the microstructural changes and dilatation data are evaluated

based on their relation to the deformation temperature and conditions.

4.2.4.1. Results

Figure 4-43 gives the flow curves of the continuously cooled and deformed

specimens due to their different deformation start temperatures. As is evident,

increasing the deformation start temperature – i.e. from 700°C to 850°C -

continuously lowers the flow stress level during the deformation.

0

100

200

300

400

500

0 0.05 0.1 0.15 0.2 0.25 0.3


True

Stre

ss, σ

[MPa

]

Deformation Start Temperature = 700°C




Figure 4-43– Flow curves of the continuously deformed samples regarding different

deformation start temperatures.


92

0

1000

2000

3000

4000

5000

0 0.05 0.1 0.15 0.2 0.25 0.3

True Plastic Strain, ε

Wor

k H

arde

ning

Rat

e, θ

(MPa

)




Deformation Start Temperature =850°C

Work Hardening Rate (θ) Vs. Strain (ε)Strain Rate = 0.1 (s-1)

Total Strain = 0.3

Figure 4-44– Work hardening rate variations due to different deformation start

temperatures.

The resulted work hardening rate data from the flow curves of figure 4-43 are

presented in figure 4-44 against the increasing exerted strain. It is seen that at

lower amounts of deformation, the specimens which started to be deformed

from the lower temperatures have relatively higher work hardening rate level,

while the order is substantially decreased while coming to higher strains.

Increasing the deformation start temperature from 700°C to 750°C, the

hardness value is decreased (figure 4-45). Further rise of the deformation

temperature increases the hardness value as is shown in figure 4-45. The

mentioned increasing trend of the hardness data is almost linier against the

deformation start temperature.

Strain Rate = 0.1 (s-1)

(-)


93

650 700 750 800 850 900315

320

325

330

335

340

345

342

329.7

318.3

HV

~

Deformation Start Temperature (°C)

332.3

Figure 4-45– Hardness value variations at different deformation start temperatures.

Figure 4-46 demonstrates the variations of the dilatation curves of the deformed

specimens at different temperatures. There is a rather continuous shift for the

martensitic transformation region to the left hand side of the diagram by

decreasing the deformation temperature, while the depth of the transformation

valleys shows some fluctuations by coming to higher deformation start

temperatures.


94

100 200 300 400 500 600 700-0.5

-0.4

-0.3

-0.2

-0.1

0.0Deformation Start Temperature = 700°C



Cha

nge

in L

engt

h (%

)

Temperature (°C)


Figure 4-46– Dilatation curves of the continuously deformed specimens at different

deformation start temperatures.

a b

c d

Figure 4-47– Microstructure of the deformed samples at different deformation start

temperatures; a) 700°C, b) 750°C, c) 800°C and d) 850°C.


95

Microstructural variations of the specimens based on the deformation start

temperatures can be seen by the evaluation of figure 4-47. The amount of

bainite is decreased by increasing the deformation temperature to 750°C.

Coming to the higher temperatures, the fraction of bainite phase in the

microstructures is increased to great extent.

4.2.4.2. Discussion

The flow stress level in figure 4-43 has been totally affected by the deformation

temperature. The lower the deformation start temperature, the higher flow

stress levels are achieved. This is simply because the dislocation motion as the

reason for the formability of metals is intensified by increasing the

experimental temperature. Therefore, the materials can be deformed by lower

stresses at higher temperatures.

As mentioned in previous sections, no phase transformation occurs during the

current simultaneous deformation and quenching tests. This is found because

no change in the slope of the flow stress or work hardening rate curves is

observed. It can be concluded that, as the final structures in all the deformation

temperatures consist of ferrite, bainite and martensite, the start temperature of

the ferritic and bainitic transformations must be less than 550°C in the case of

sample which its deformation is started at 700°C. This is because it takes three

seconds for the strain of 0.3 to be achieved by the strain rate of 0.1s-1. As the

cooling rate is 50°C/sec., the temperature comes at 550°C after three seconds.

Therefore, for the bainitic and ferritic transformations to be appeared without

any effect on the flow curves behavior, the transformations start temperatures

must be less than 550°C. The mentioned calculated low temperature for a static

ferritic transformation shows that the reported ferrite must be categorized as the

strain induced ferrite or ferrite regions in the granular-carbide-free-bainite, and

not the result of a normal static transformation at high temperatures.

Considering figure 4-44, rather higher work hardening rate level in the case of

deformed specimens at lower temperatures in comparison with higher

temperature experiments can be distinguished at the strains of less than 0.2. At


96

more than this strain, the mentioned difference is vanished. This is because at

lower temperatures, the applied strain leads into the formation of the

dislocations which are not able to be ordered as fast as they are produced. At

lower strains, this may result in the decrease of the work hardening rate, but the

value of work hardening rate is higher for a lower deformation temperature due

to the higher efficiency of locking more dislocations at lower temperatures. At

higher strains, a kind of saturation in the work hardening rate value is reached.

It means that regardless of the deformation temperature, exerting more strains

on the sample does not change the work hardening rate in any way and a steady

state occurs. This is because for all the samples, there is a compromise between

the rate of the dislocation formation and the amount of ordered dislocations at

the deformation temperature. It means that at each temperature, the rate of

dislocation rearrangement is increased parallel – but not equal- to the rate of the

dislocation formation. Therefore, the specimen is work hardened continuously

by a constant work hardening rate.

Considering the hardness values in figure 4-45 and the microstructural images

in figure 4-47, it is found that the variations of the hardness value mostly

follow the amount of bainite in the final microstructure than ferrite or

martensite content. As can bee seen, the amount of bainite is decreased by

increasing the deformation temperature from 700 to 750°C, and then there is a

continuous increase in the amount of bainite as the deformation start

temperature increases from 750 to 850°C. The amount of bainite at 700 and

800°C specimens is almost the same. The same trend for the variation of the

hardness value is observed. Meanwhile the amount of martensite in the

microstructure is increased by increasing the deformation temperature from 700

to 750°C and after a decrease from 750 to 800°C, it is increased when coming

to 850°C. Except for the 750°C specimen, no meaningful variation for the

amount of ferrite in the final microstructure is observed. In the case of 750°C,

the amount of ferrite phase is increased.

Except for the deformation at 750°C which gives a different behavior, the

evaluation of the hardness and microstructural data shows that by increasing


97

the deformation start temperature, the bainitic transformation region is shifted

to the left hand side of the CCT diagram and that is more bainite is formed at a

same cooling rate by increasing the deformation temperature.

Figure 4-46 demonstrates that the martensitic transformation start temperature

is transferred to the lower temperatures as the simultaneous deformation start

temperature is decreased. It is found that the formation of martensite becomes

easier as the deformation takes place at higher temperatures.

This is justified due to the presence of less austenite in the microstructure of the

deformed samples at higher temperatures after finalizing the bainitic

transformation. Because of this, as more amounts of austenite in the

microstructure need a more powerful quenching medium (i.e. less quenching

temperature) to be transformed to martensite, the Ms temperature is decreased.

Although, due to the presence of different parameters (as mentioned in

previous sections), the depth of the valley does not directly show the amount of

martensite formed, but the decreasing behavior of Ms value by decreasing the

deformation temperature is also evident between 800 and 850°C specimens.

For this to be observed, one should consider the broad valley of the martensitic

transformation and its relevant Ms of more than 400°C in comparison with the

Ms value of less than 400°C for 800°C specimen.

Regarding the above mentioned justification and by going to the higher

temperatures in the case of dilatation curves of 800°C and 850°C specimens,

the bainitic transformation valleys can also be distinguished.

Chapter 5 - CONCLUSIONS 98

CHAPTER

Five

CONCLUSIONS

- The attention must be paid to the continuously cooled bainite and ferrite

formation which is observed during both isothermal compression and

simultaneous forming and quenching experiments. This is because even

in case of the isothermal compression, the specimens are under the

continuous cooling prior to and after the deformation.

- During the experiments, almost no phase transformation was observed

during the deformation. This was resulted from the consideration of the

slopes of the flow curves which do not show a sudden change due to a

phase transformation.

- The effect of the prior deformation on generating the nucleation sites of

ferritic transformation is diminished as the deformation temperature is

increased in isothermal compression tests.

- The lower isothermal compression temperature may lead into the

formation of more martensite by hindering the bainitic and ferritic

transformations.

- The colorful increasing effect of bainite phase percentage on the

hardness values beside the martensite content must be taken into account

when studying the 22MnB5 steel.

- During the isothermal compression tests, there is more opportunity for

the continuously cooled bainitic and ferritic transformations to occur at

lower temperatures than higher temperatures.

- The mutual effect of strain rate and deformation temperature on the

work hardening rate values must be considered.

- During the isothermal compression tests, in most cases, the presence of

bainite and ferrite phases can decrease and increase the value of Ms-Mf

respectively.

Chapter 5 - CONCLUSIONS 99

- When evaluating the dilatation curves, the effect of the internal stresses

and prior phase transformations must be taken into account.

- The severe plastic deformation of austenite prior to its transformation

hinders the growth of martensite causing a reduction in the fraction of

the transformation in spite of an increased number of nucleation sites'

density.

- Despite the change of the martensite content, increasing the

austenization time has no meaningful effect on the flow curves and the

resulted work hardening rate data regarding the simultaneous forming

and quenching experiments.

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